Transition Metal Oxide Thin Film based Chromogenics and Devices. A volume in Metal Oxides [1st Edition] 9780081017487, 9780081018996

Table of contents :
Content:
Front Matter,Copyright,Dedication,Author's Biography,Series Editor's Biography,Preface to the Series,AcknowledgmentsEntitled to full textChapter 1 - Introduction to Chromogenics, Pages 1-11
Chapter 2 - Introduction to Transition Metal Oxides and Thin Films, Pages 13-72
Chapter 3 - Electrochromic Thin Films and Devices, Pages 73-151
Chapter 4 - Thermochromic Thin Films and Devices, Pages 153-246
Chapter 5 - Photochromic Thin Films and Devices, Pages 247-288
Chapter 6 - Chromogenic Thin Film Photonic Crystals, Pages 289-329
Chapter 7 - Emerging Technologies∗, Pages 331-347
Index, Pages 349-357

Citation preview

METAL OXIDES SERIES

TRANSITION METAL OXIDE THIN FILMeBASED CHROMOGENICS AND DEVICES PANDURANG ASHRIT Series Editor GHENADII KOROTCENKOV

Elsevier Radarweg 29, PO Box 211, 1000 AE Amsterdam, Netherlands The Boulevard, Langford Lane, Kidlington, Oxford OX5 1GB, United Kingdom 50 Hampshire Street, 5th Floor, Cambridge, MA 02139, United States Copyright © 2017 Elsevier Ltd. All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, recording, or any information storage and retrieval system, without permission in writing from the publisher. Details on how to seek permission, further information about the Publisher’s permissions policies and our arrangements with organizations such as the Copyright Clearance Center and the Copyright Licensing Agency, can be found at our website: www.elsevier.com/ permissions. This book and the individual contributions contained in it are protected under copyright by the Publisher (other than as may be noted herein). Notices Knowledge and best practice in this field are constantly changing. As new research and experience broaden our understanding, changes in research methods, professional practices, or medical treatment may become necessary. Practitioners and researchers must always rely on their own experience and knowledge in evaluating and using any information, methods, compounds, or experiments described herein. In using such information or methods they should be mindful of their own safety and the safety of others, including parties for whom they have a professional responsibility. To the fullest extent of the law, neither the Publisher nor the authors, contributors, or editors, assume any liability for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions, or ideas contained in the material herein. Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library ISBN: 978-0-08-101899-6 (print) ISBN: 978-0-08-101748-7 (online) For information on all Elsevier publications visit our website at https://www.elsevier.com/books-and-journals

Publisher: Matthew Deans Acquisition Editor: Kayla Dos Santos Editorial Project Manager: Anna Valutkevich Production Project Manager: Debasish Ghosh Designer: Miles Hitchen Typeset by TNQ Books and Journals

DEDICATION To my beloved wife,

Anita Ashrit whose inner strength, relentless encouragement, and ability to keep smiling in all situations keep me in high spirits, and to our lovely daughter, the sunshine of our life,

Meeta Ashrit

Pandurang Ashrit, March 2017

AUTHOR’S BIOGRAPHY Professor Pandurang Ashrit is affiliated with the Department of Physics and Astronomy at the Université de Moncton, Canada. He is a named author on 75 documents, many of them dealing with thin films and chromogenic materials, and has patented a technique for chromogenically tuning photonic crystals. In 2013 he was granted an Innovation Award for Excellence in Applied Research by the government of New Brunswick.

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SERIES EDITOR’S BIOGRAPHY Ghenadii Korotcenkov Research Professor (2008epresent) School of Materials Science and Engineering, Gwangju Institute of Science and Technology, South Korea Ghenadii Korotcenkov received his PhD in Physics and Technology of Semiconductor Materials and Devices from the Technical University of Moldova in 1976 and his DrSci degree in Physics of Semiconductors and Dielectrics from the Academy of Science of Moldova in 1990 (Highest Qualification Committee of the USSR, Moscow). He has more than 40 years’ experience as a teacher and scientific researcher. For a long time he was the leader of a gas sensor group and manager of various national and international scientific and engineering projects carried out in the Laboratory of Micro- and Optoelectronics, Technical University of Moldova. His research had financial support from international foundations and programs such as CRDF, MRDA, ICTP, INTAS, INCO-COPERNICUS, COST, and NATO. From 2007 to 2008 he was an invited scientist at the Korea Institute of Energy Research (Daejeon). Since 2008 Dr. G. Korotcenkov has been a research professor at the School of Materials Science and Engineering at Gwangju Institute of Science and Technology in Korea. Specialists from the former Soviet Union know G. Korotcenkov’s research results in the study of Schottky barriers, metal oxide semiconductor structures, native oxides, and photoreceivers on the basis of IIIeV compounds such as InP, GaP, AlGaAs, and InGaAs. His current scientific interests, starting from 1995, include material sciences, focusing on metal oxide film deposition and characterization; surface science; and the design of thin film gas sensors and thermoelectric convertors. G. Korotcenkov is the author or editor of 35 books and special issues, including the 11-volume Chemical Sensors series published by Momentum Press (USA), 15-volume Chemical Sensors series published by Harbin Institute of Technology Press (China), 3-volume Porous Silicon: From Formation to Application published by CRC Press (USA), and 2-volume Handbook of Gas Sensor Materials published by Springer (USA). G. Korotcenkov is author and coauthor of more than 550 scientific publications including 20 review papers, 35 book chapters, more than 250 articles published in peer-reviewed scientific

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Series Editor’s Biography

journals, and more than 50 refereed papers published in international conference proceedings [h-factor ¼ 37 (Scopus) and h ¼ 44 (Google Scholar citation)]. He is a holder of 17 patents. He has presented more than 200 reports at national and international conferences, including 15 invited talks. G. Korotcenkov was coorganizer of several international conferences. His research activities are honored by an Award of the Supreme Council of Science and Advanced Technology of the Republic of Moldova (2004), Prize of the Presidents of the Ukrainian, Belarus, and Moldovan Academies of Sciences (2003), Senior Research Excellence Award of the Technical University of Moldova (2001, 2003, 2005), and National Youth Prize of the Republic of Moldova in the field of science and technology (1980), among others. G. Korotcenkov also received a fellowship from the International Research Exchange Board (USA, 1998), Brain Korea 21 Program (Korea, 2008e12), and Brainpool Program (Korea, 2015e17).

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

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Preface to the Series

field of metal oxide nanostructured morphologies (e.g., nanowires, nanorods, nanotubes, etc.) has become one of the most active research areas within the nanoscience community. The ability to create a variety of metal oxide-based composites and the ability to synthesize various multicomponent compounds significantly expand the range of properties that metal oxide-based materials can have, making metal oxides a truly versatile multifunctional material for widespread use. Small changes in their chemical composition and atomic structure can be accompanied by a spectacular variation in the properties and behavior of metal oxides. Even now, advances in synthesizing and characterizing techniques are revealing numerous new functions of metal oxides. Taking into account the importance of metal oxides for progress in microelectronics, optoelectronics, photonics, energy conversion, sensors, and catalysis, a large number of various books devoted to this class of materials have been published. However, one should note that some books from this list are too general, some books are collections of various original works without any generalizations, and others were published many years ago. But, during last decade, great progress has been made in the synthesis as well as the structural, physical, and chemical characterization and application of metal oxides in various devices, and a large number of papers have been published on metal oxides. In addition, until now many important topics related to metal oxides study and application have not been discussed. To remedy the situation in this area, we decided to generalize and systematize the results of research in this direction and to publish a series of books devoted to metal oxides. One should note that the proposed book series Metal Oxides is the first one devoted to consideration of metal oxides only. We believe that combining books on metal oxides in a series could help readers in searching for required information on the subject. In particular, we plan that the books from our series, which have a clear specialization by content, will provide an interdisciplinary discussion of various oxide materials with a wide range of topics, from material synthesis and deposition to characterization, processing, and then device fabrication and applications. This book series is prepared by a team of highly qualified experts, which guarantees it a high quality. I hope that our books will be useful and comfortable in use. I would also like to hope that readers will consider this Metal Oxides book series as an encyclopedia of metal oxides that enables one to understand the present

Preface to the Series

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status of metal oxides, to estimate the role of multifunctional metal oxides in the design of advanced devices, and then, based on observed knowledge, to formulate new goals for the further research. The intended audience of the present book series is scientists and researchers working or planning to work in the field of materials related to metal oxides, i.e., scientists and researchers whose activities are related to electronics, optoelectronics, energy, catalysis, sensors, electrical engineering, ceramics, biomedical designs, etc. I believe that this Metal Oxides book series will also be interesting for practicing engineers or project managers in industries and national laboratories who would like to design metal oxidebased devices, but don’t know how to do it, or how to select the optimal metal oxide for specific applications. With many references to the vast resource of published literature on the subject, this book series will serve as a significant and insightful source of valuable information, providing scientists and engineers with new insights for understanding and improving existing metal oxide-based devices and for designing new metal oxidebased materials with new and unexpected properties. I believe that this Metal Oxides book series would be very helpful for university students, postdocs, and professors. The structure of these books offers a basis for courses in the field of material sciences, chemical engineering, electronics, electrical engineering, optoelectronics, energy technologies, environmental control, and many others. Graduate students could also find the book series to be very useful in their research and in understanding the features of metal oxide synthesis and the study and application of this multifunctional material in various devices. We are sure that all of them will find information useful for their activity. Finally, I thank all the contributing authors and book editors who have been involved in the creation of these books. I am thankful that they agreed to participate in this project and for their efforts in the preparation of these books. Without their participation, this project would have not been possible. I also express my gratitude to Elsevier for giving us the opportunity to publish this series. I especially thank all the team at the editorial office at Elsevier for their patience during the development of this project and for encouraging us during the various stages of preparation. Ghenadii Korotcenkov

ACKNOWLEDGMENTS My sincere thanks to Elsevier International for this opportunity to write a book on transition metal oxide thin filmebased chromogenics, which is a rapidly growing field. Thanks are also due to Anna Valutkevich, Kayla Dos Santos, and Heather Cain at Elsevier for wonderfully coordinating the book project for its timely publication. I would also like to thank my research associates Drs. Gisia Beydaghyan and Bassel Abdel Samad at the Université de Moncton for their timely help and contribution of certain content of the book when I could not attend to it myself for personal reasons. I would like to express my great appreciation to Meeta Ashrit of the University of Ottawa and Francine Maillet of the Université de Moncton for their help in painstakingly organizing the references. I am deeply indebted to my teachers Professors M.A. Angadi (USA) and D.R. Balurgi (India) for introducing me to the field of thin films and for giving me an opportunity to do my higher studies in this field. On a very personal note, I would like to express my indebtedness to my mother (Indira Bai Ashrit), my aunt (Shanta Bai Ashrit), and my uncle (Bhimsenachar Ashrit) for their unconditional love during the early stage of my life. My profound gratitude to all the Ashrit brothers (Krishna, Gopal, Srinivas, and Shamsunder) and sisters-in-law (Radha, Tara, and Meera) for their support and sacrifices in making me what I am today. I am deeply indebted to Shamsunder and Meera for making our life wonderful in many ways. Thanks to my sisters (Radha Pappu and Sudha Gumaste) and brothers-in-law (Krishna Pappu and Mohan Gumaste) for their all-round support in our hours of need. Thanks are also due to my father-in-law, Mr. R.V. Kulkarni, whose persistent push to finish the book project kept me on my toes! I would also like to thank all my friends and colleagues in Moncton, Canada, for their support and encouragement during this project. My special thanks to my friend, colleague, and scientific collaborator, Professor A. Nait Ajjou, for his overall support over the years. Pandurang Ashrit, March 2017

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CHAPTER 1

Introduction to Chromogenics Chromogenics, originating from the Greek word chromo for color, in recent times refers to the field of study of materials that show a change in their optical properties (color) as a function of the ambient conditions they are subjected to. This word was used earlier in the field of photography to refer to materials that are sensitive to light, such as photographic paper covered with a silver halide emulsion [1]. In this case light was the trigger for the change in color of the paper (photochromism). However, the color change in the case of photographic paper was irreversible. More generally the word chromogenics, or chromogenic materials, is used now to refer to all materials that show a reversible change in their optical properties as a result of the application of various external stimuli such as light, electric field, heat, exposure to a gas, pressure, and more [2]. Chromogenic materials form an important class of smart materials showing a reversible change in various physical properties as a function of the applied stimulus. Following are some examples of such smart materials. Piezoelectric materials [3] are those in which a stress is induced upon the application of an electric field or in which, upon the application of a stress, a voltage is generated. Thermoelectric materials [4] are those that output a voltage when a temperature difference is applied across the material. Shape memory alloys [5] change their shape according to the temperature to which they are subjected. Chromogenic materials are specifically those smart materials that show a reversible optical property change under the application of the aforementioned stimuli. According to the stimulus applied to induce this color change, these chromogenic materials are termed photochromic (light induced), electrochromic (electric field induced), thermochromic (heat induced), gasochromic (gas exposure induced), and so on [6]. Chromogenic materials have become very important because of their optically dynamic and interactive nature. Their response to various external forces makes it easier to build a wide range of automated smart systems in which light control is needed. Let us consider these materials briefly.

1.1 ELECTROCHROMIC MATERIALS Electrochromic materials are those in which a reversible change in optical properties can be induced by the application of a small electric field. Transition Metal Oxide Thin FilmeBased Chromogenics and Devices ISBN 978-0-08-101747-0 http://dx.doi.org/10.1016/B978-0-08-101747-0.00001-5

© 2017 Elsevier Ltd. All rights reserved.

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Electrochromic coloration is the most widely studied and used form of chromogenics owing to the ease of control and reversibility of optical properties using an applied electric field. Liquid crystal displays are the most common forms of electrochromic devices known in the past few decades [7]. The change in their optical property from a transparent to an opaque state with the application of a small electric field originates from the change in the orientation of the liquid crystal molecules. This phenomenon has been used in a wide variety of electronic display devices from watches to televisions for quite some time. The disadvantages associated with the liquid crystal-based electrochromics are the relatively high power consumption required to maintain the on state and the switching of the devices only between the on and the off state in most cases. Similarly, electrodeposition or electrophoretic deposition systems are another popular example of electrochromic materials wherein, with the application of a small electric field, suspended metal particles held in a solution are deposited on the transparent electrode to form a film changing the optical properties of the device [8]. Hence, the presence of a liquid component is essential for the functioning of these devices. A wide range of organic [9e12] and inorganic materials [9,13] show electrochromic properties through other types of reversible chemical transformations. Among the inorganic compounds, many metal oxides exhibit an efficient form of electrochromic coloration. A number of molecular dyes and conducting polymers are also very efficient electrochromic materials. The electrochromic coloration in each of these types arises from various phenomena such as through a redox reaction or change in the oxidation state or formation of polarons, etc. Quite a number of applications based on electrochromic materials have been realized, such as smart windows, tintable rearview mirrors, electronic displays, etc. Some of these are listed in Table 1.1. Given in this table are also the stimuli that trigger the optical change along with the mechanism that leads to these changes.

1.2 PHOTOCHROMIC MATERIALS In photochromic materials a reversible change in optical properties occurs as a result of exposure of these materials to electromagnetic radiation. The reversibility of the photochromic coloration per se is not directly controllable but depends on the nature of coloration induced and the metastable state of the resulting product. To have a direct control of the discoloration other forms of activation such as an electric field or heat may be necessary.

Introduction to Chromogenics

Table 1.1 Examples of chromogenic materials, devices, and applications Chromogenic Activation Basis for the type mode optical change Applications

Polymer dispersed liquid crystals

Electric field

Suspended particle displays Microblinds

Electric field Electric field

Nanocrystals

Electric field

Reversible electrodeposition/ electrophoretic deposition

Electric field

Kerr effect

High electric field

Thermochromic

Heat

Gasochromic

Gas exposure

Photochromic

Light

Magnetochromic

Magnetic field

Change in orientation of liquid crystal molecules Alignment of the suspended particles Unrolling of the thin metal films to block the light Change in the structure of the composite material Deposition of suspended metal particles in the form of thin film on the transparent electrode Change in refractive index due to the alignment of the molecules Structural phase change at a critical temperature Metaleinsulator transition upon exposure to a gas Charge transfer between two electronic sites under the action of light Reorganization of magnetic particles

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References

Transparent displays

[17]

Smart windows Smart windows

[18]

Smart windows

[20]

Display devices

[8]

Tunable photonic devices

[21e23]

High-speed optical switches, smart windows Displays, tunable mirrors Photochromic glasses, memory storage applications Tunable color displays, tunable photonic crystals

[24]

[19]

[10]

[9]

[25]

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

The photochromic coloration is invariably activated by ultraviolet or visible radiation leading to conductivity and optical changes in the visible and infrared regions. Hence, many practical applications can be envisioned around the photochromic materials. Photochromic ophthalmic glasses that change their opacity in accordance with the ultraviolet light conditions are a popular example of this kind [14]. Metal halide-based devices change from a transparent state to a mirror-like reflective state upon exposure to a gas such as hydrogen following the application of a small electric field [15]. A wide range of organic and inorganic materials are known to exhibit photochromic coloration [16]. Under ultraviolet or blue light irradiation many organic molecules undergo transformation through the breaking of bonds. If a single bond (heterolytic) is broken in compounds such as spiropyrans, triarylmethane, and polymethine leading to coloration, this photochromism is short lived, as the system returns to its original state quickly. The hemolytic compounds lead to the formation of a pair of radicals under photo breakdown and coloration. The compounds showing this type of photochromism are pyrroles, hydrazines, disulfides, nitroso dimers, and more. These transformations are also not very stable and not interesting from the application point of view. Cisetrans isomerism photochromism exhibited by compounds such as urocanic acid, indigoid and cyanine dyes, polyenes, etc., is more stable and useful for applications. Some inorganic materials possessing large bandgaps exhibit a more metastable photochromism. Quite often the coloration involves the presence of defects such as F centers or impurities or dislocations. The photochromic coloration can be reverted only by other activations such as heat, electric field, etc. Quite a few inorganic metal oxides and semiconductors exhibit photochromism. It is their metastable coloration that makes these materials more useful from the application point of view.

1.3 THERMOCHROMIC MATERIALS These are materials that exhibit a reversible change in coloration under the influence of heat. Almost all materials undergo some reversible coloration change at a certain temperature. However, the term thermochromism is applied more specifically to materials that undergo an abrupt change (first order) in their optical properties at a critical temperature (Tc). Materials exhibiting such a thermochromic change undergo a reversible phase change. Hence, reversibility is an important fact to be considered for evaluating

Introduction to Chromogenics

5

thermochromic materials. Quite a few organic and inorganic materials are known to show a very efficient form of thermochromic change. A number of applications have been developed using these materials over a number of years in the areas of security, imagery, copiers, and temperature measurement. Among the organic materials, polydiacetylenes, poly(3-alkylthiophene), poly(p-xylyleneviologen dibromide), polyazomethines, and others exhibit thermochromism in various temperature ranges [6]. Silicon-based polymers such as silanes and siloxanes are known to exhibit thermochromic transformation, some showing a gradual change as a function of temperature, whereas others show an abrupt change. Among the inorganic compounds quite a number of transition metal (Ti, V, Nb, etc.) oxides and sulfides (Ag, Ni, Cr, etc.) show first-order thermochromic transformation over a wide range of temperatures. Among these inorganic compounds, it is vanadium dioxide (VO2) that exhibits a thermochromic phase transition at a temperature closest to that of room temperature, with a Tc value of around 68
C. Hence, this material has become the object of innumerable scientific research and industrial applications.

1.4 PIEZOCHROMIC MATERIALS These are materials showing reversible color change under the application of pressure. These materials are studied intensely because of their ability to detect stress and changes in pressure. Mechanochromic is another term used instead of piezochromic. There are three mechanisms of coloration of piezochromic materials: (1) As the pressure changes at the piezochromic surface (compression), it perturbs the electronic energy levels differently. In other words, electronic excitation from the ground state to higher levels takes place. (2) A firstorder phase transition occurs due to changes in the crystal structure when a pressure is applied on the surface. (3) The geometry of the molecules changes owing to applied pressure on the surface of the piezochromic material resulting in a color change. Piezochromic materials developed by Seeboth et al. [26] contain a mixture of polymers in cholesteryl derivatives. These materials respond to pressure and change their color reversibly from red to green. Some materials like dimethylglyoxime complexes exhibit a piezochromic response under a pressure of 63e150 kbar, showing a color change. Another piezochromic inorganic material is samarium monosulfide, which transforms from a semiconductor to its metal phase under a pressure of 6.5 bar [27].

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There are many applications for this type of material that can be used to detect a change in pressure on a surface. Some of these are piezochromic paint for tennis balls, volleyballs, basketballs, bathrooms, kitchen scales, offset printing ink, classic flexo ink, screen printing ink, and lamination. Some of the piezochromic materials are irreversible and the color at a defined pressure remains unchanged. Such materials are usually used to detect shock or to follow health-related parameters. An example of a piezochromic paint undergoing color change under pressure is shown in Fig. 1.1.

1.5 MAGNETOCHROMIC MATERIALS Materials in which a color change occurs instantly and reversibly from the application of an external magnetic field are termed as magnetochromic. Magnetochromism is also known as the modification of a material’s optical constants with an applied magnetic field. As the material is sensitive to an external magnetic field, the spin-induced electric polarization can be magnetically controlled and the spinecharge interaction is large. Candidates for magnetochromic materials with strong chargeespin coupling and that are magnetically controlled are spin-crossover complexes, colossalmagnetoresistive oxides, and low-dimensional molecular magnets [28]. Another type of material exhibiting magnetochromic properties are the multiferroic compounds, which are known to exhibit ferromagnetism, ferroelectricity, and ferroelasticity all in the same phase. MnWO4 is an

Figure 1.1 A piezochromic paint showing coloration under pressure. (Image taken from the OliKrom website.)

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7

example of such a multiferroic material in which spin-induced ferroelectricity and large magnetoelectric responses are observed [29]. Magnetochromic microspheres are another example, which are compatible with different types of dispersion media like water, alcohol, or even polymer solutions. Unlike other chromic materials, the change in color in this type of material is due to the magnetic fields acting on the orientation of the microspheres without any change in the intrinsic properties or structure [30]. It is actually the iron oxide (Fe3O4) contained in the microsphere that is responsible for the rotation (orientation array) due to the applied external magnetic field. By changing the microsphere orientation, the color switch is enabled. Magnetochromic materials are used in many applications such as display units like rewritable signage, papers, and posters. Another application that can be envisioned for these materials is paint and ink material for printing. In some cases they can be used for security reasons. An example of such a color change as a function of magnetic field is shown in Fig. 1.2.

1.6 CHEMICHROMIC MATERIALS Chemichromic materials are those that change their color in response to the chemical changes they are exposed to. There are many subgroups of chemichromic materials, such as: (1) gasochromic, those that usually show reversible coloration under exposure to an appropriate gas; (2) halochromic,

Figure 1.2 Gradual color change (from A to E progressively) as a function of applied external magnetic field on a magnetochromic material. (From M. Bichurin, V. Petrov, V. Leontyev, A. Saplev, Two-range magnetoelectric sensor, AIP Advances 7 (2017) 056317.)

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

which change color in response to pH changes; (3) ionochromic, those that show color change in the presence of ions in a medium; and (4) hygro/ hydrochromics or solvatochromic materials, which show a color change when they react with humidity or water. The best known examples of a gasochromic material are the WO3 [31] and V2O5 [32] thin films that undergo an efficient coloration upon exposure to hydrogen gas. For example, WO3 thin film changes from an initial transparent state to a dark blue state when exposed to hydrogen. It is used in double-pane windows in which the WO3 thin film is applied on one of the inside window surfaces and the hydrogen gas is captured in the gap between the two panes. An important application of WO3-based gasochromic devices is to detect hydrogen gas leakage to protect humans from a large exposure [31]. An example of this is shown in Fig. 1.3.

1.7 TRANSITION METAL OXIDE-BASED CHROMOGENIC MATERIALS In principle a variety of materials can exhibit chromogenic properties. However, in this book the principal focus is on the chromogenic materials and devices that are based on transition metal oxide (TMO) thin films. TMOs possess very unusual and highly interesting properties due to their partially filled d orbitals and their multiple oxidation states. It is possible to drive these oxides reversibly between their various metastable oxidation states by external activation. This activation triggers drastic changes in their optical, electrical, and other properties. Hence, TMOs form an important and highly convenient class of chromogenic materials. Prepared in thin film form, the TMOs provide an even more convenient means of controlling the various optical parameters such as transmittance, (A)

(B)

Figure 1.3 Hydrogen-detecting tape (A) before and (B) after leakage. (From the HySense Technology website.)

Introduction to Chromogenics

9

reflectance, absorption, and emissivity and the fundamental optical properties such as the refractive index (n) and extinction coefficient (k). These parameters are of incredible pertinence to chromogenic devices and their design. In addition, there has been a tremendous impetus to work with nanostructured thin films. The nanostructuring of thin films and the three-dimensional size effects imposed thereby lead to very interesting and new properties in the already well-known materials. An example of this is the nanostructured thin films of the very well-known material, titanium dioxide (TiO2). TiO2 in its bulk form is one of the most well-known and used materials in varied applications from paints to toothpaste to photocatalyzers and more. This is due to the abundant availability of this material as well as its nontoxicity and chemical stability. Even in its normal thin film form it is used extensively in applications such as dye-sensitized solar cells, bolometers, and others. These normal or so-called coarse grain thin films of TiO2 are not very well known for their electrochromic properties. However, it has been shown [33] that by preparing the same films in a nanostructured form, one can induce a fairly efficient electrochromic switching. Similarly, an enhancement in the electrochemical properties of nanostructured (NS) tin oxide (SnO2) thin films has been demonstrated [34]. The film properties of such NS thin films are so drastically different from those of the bulk or normal thin films that some consider it as a “new state” of the material [35]. There has been a tremendous emphasis placed on the study of the NS thin film form of the already wellknown materials to induce new properties in these materials. Hence, the main focus of this book is on the chromogenic properties of such NS thin films of TMOs. Both periodically NS and nonperiodically (randomly) NS thin films are addressed elaborately in view of their pertinent virtues related to their optical switching behavior.

1.8 SUMMARY Various forms of chromogenic materials are introduced in this chapter. Among these, the TMOs seem to provide a very convenient and versatile form of chromogenic coloration. These effective changes stem from the unusual atomic structure of the base transition metals of these oxides. Hence, these materials can be easily driven between their various metastable oxidation states. The reversible change in the optical properties in these materials can thus be triggered by heat (thermochromics), light (photochromics), electric field (electrochromics), exposure to gas (gas exposure),

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and more. Especially when prepared in thin film form, the TMOs provide a wide range of active control of various optical parameters with the application of various external stimuli. Hence, the study of TMO thin films has become extremely interesting and important from both research and technological points of view. The dynamic changes occurring in their physical and/or chemical structure leading to changes in optical and electrical properties are very interesting for studying the fundamental aspects. The interactive control they provide is very important for a number of industrial applications.

REFERENCES [1] http://notesonphotographs.eastmanhouse.org/index.php?title¼Chromogenic_Print. [2] C.M. Lampert, Chromogenic smart materials, Materials Today (March 2004) 28e35. [3] H.A. Sodano, D.J. Inman, G. Park, Comparision of piezoelectric energy harvesting devices for recharging batteries, Journal of Intelligent Material Systems Systems and Structures (JIMSS) 16 (10) (2005) 799e807. [4] http://academic.uprm.edu/pcaceres/Courses/Smart/SMD-8B.pdf. [5] M.H. Wu, L.M. Schetky, Industrial applications for shape memory alloys, in: Proceeding of the International Conference on Shape Memory and Super Elastic Technologies, Pacific Grove, California, 2000, pp. 171e182. [6] C.M. Lampert, C.G. Granqvist (Eds.), Large Area Chromogenics: Materials and Devices for Transmission Control, SPIE Institute Series, vol. IS4, 1990. Bellingham. [7] T.J. Sluckin, D.A. Dunmur, H. Stegemeyer, Crystals That Flow: Classic Papers from the History of Liquid Crystals, Taylor and Francis, 2004. [8] G. Oskam, P.M. Vereecken, X. Shao, J. Fransaer (Eds.), Semiconductors, Metal Oxides and Composites: Metallization and Electrodeposition of Thin Films and Nanostructures, vol. 25 (27), Electrochemical Society (ECS) Transactions, 2010. [9] P.M.S. Monk, R.J. Mortinmer, D.R. Rosseinsky, Electrochromism: Fundamentals and Applications, VCH Inc., Weinheim, 1995. [10] B. Scharifker, C. Wehrmann, Journal of Electrochemical Society 185 (1985) 93. [11] D.R. Rosseinsky, et al., Simple diffuse-reflectance monitoring of emerging surfaceattached species: bipyridilium, Prussian-blue and zinc-ferrocyanide voltammetry, Journal of Electroanalytical Chemistry 258 (1989) 233e239. [12] B. Grant, et al., Study of the electrochromism of methoxyfluorene compounds, Journal of Organic Chemistry 45 (1980) 702e705. [13] C.G. Granqvist, Handbook of Inorganic Electrochromic Materials, second ed., Elsevier, 2002. [14] R.C. Duncan Jr., D.L. Staebler, Inorganic photochromic materials, in: H.M. Smith (Ed.), Holographic Recording Materials, Topics in Applied Physics, 20, Springer, 1977, pp. 133e160 (Chapter 5). [15] W. Lohstroh, R. J.Westerwaal, B. Noheda, S. Enache, I.A.M.E. Giebels, B. Dam, R. Griessen, Self-organized layered hydrogenation in black Mg2NiHc switchable mirrors, Physical Review Letters 93 (2004) 197404e197408. [16] G.H. Brown, Photochromism, John Wiley and Sons Inc., 1971. [17] C.-W. Su, M.-Y. Chen, Polymer-dispersed liquid crystal applied in active-matrix transparent display, Journal of Display Technology 10 (8) (2014) 683e687. [18] http://www.smartglass.com/.

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[19] http://www.glasscanadamag.com/news/micro-blind-technology-2077. [20] A. Llordés, G. Garcia, J. Gazquez, D.J. Milliron, Tunable near- infrared and visible-light transmittance in nanocrystal-in-glass composites, Nature 500 (2013) 323e326. [21] P. Weinberger, John Kerr and his effects found in 1877 and 1878, Philosophical Magazine Letters (2008) 1e11 (iFirst). [22] M. Qasymeh, M. Cada, S.A. Ponomarenko, Quadratic electro-optic Kerr effect: applications to Photonic devices, IEEE Journal of Quantum Electronics 44 (2008) 740e746. [23] G. Korotcenkov, in: R.A. Potyrailo (Ed.), Integrated Analytical System Series Vol. 1, Springer, 2013. [24] P. Kiri, G. Hyettb, R. Binionsa, Solid state thermochromic materials, Advanced Materials Letters 1 (2) (2010) 85e106. [25] L. Zhuang, W. Zhang, Y. Zhao, H. Shen, H. Lin, J. Liang, Preparation and characterization of Fe3O4 particles with novel nanosheets morphology and magnetochromatic property by a modified solvothermal method, Scientific Reports 5 (9320) (2015) 1e6. [26] A. Seeboth, D. Lotzsch, R. Ruhmann, Piezochromic polymer materials displaying pressure changes in bar-ranges, American Journal of Materials Science 1 (2) (2008) 139e142. [27] P. Jin, Tanemura, Manufacturing Methods of Samarium Sulfide Thin Film, US Patent US6132568, October 17, 2000. [28] N. Kida, S. Kumakura, S. Ishiwata, Y. Taguchi, Y. Tokura, Gigantic terahertz magnetochromism via electromagnons in hexaferrite magnet, Ba2Mg2Fe12O22, Physical Review B 83 (2011) 064422-1-8. [29] S. Toyoda, N. Abe, T. Arima, S. Kimura, Large magnetochromism in multiferroic MnWO4, Physical Review B 91 (2015) 054417. [30] M. Bichurin, V. Petrov, V. Leontyev, A. Saplev, Two-range magnetoelectric sensor, AIP Advances 7 (2017) 056317. [31] G. Korotcenkov, Handbook of Gas Sensor Materials: Properties, Advantages and Shortcomings for Applications, Springer, New York, 2013. [32] P. Liu, S.-H. Lee, C.E. Tracy, J.A. Turner, J.R. Pitts, S.K. Deb, Electrochromic and chemochromic performance of mesoporous thin-film vanadium oxide, Solid State Ionics 165 (2003) 223e228. [33] A. Hagfeldt, N. Viachopolous, M. Graetzel, Fast electrochromic switching with nanocrystalline oxide semiconductor films, Journal of the Electrochemical Society 141, L82 (1994) 42e44. [34] D. Vasanth Raj, N. Ponpadian, D. Mangalraj, C. Viswanathan, Effect of annealing and electrochemical properties of sol-gel dip coated nanocrystalline V2O5 thin films, Materials Science in Semiconductor Processing 16 (2013) 256e262. [35] H. Gleiter, Nanostructured materials, Advanced Materials 4 (7/8) (1992) 474e481.

CHAPTER 2

Introduction to Transition Metal Oxides and Thin Films 2.1 BASICS OF TRANSITION METAL OXIDES Transition metal oxides (TMO) form one of the most fascinating classes of materials because of the unusual electronic structure of the base transition metal and its bonding with oxygen. The nature of TMOs is so intricate that the understanding of the correlation of their complex properties with their electronic and physical structure is said to promote an understanding of the behavior of all the other inorganic solids [1]. The partially filled d orbitals of the base transition element of TMOs are the basis for the wide range of oxides they can form as well as for their metastability in multiple oxidation states. It is these two properties that are of pertinence in their switching behavior between these states, which leads to drastic and reversible changes in their optical, electrical, and other properties. This switching can be brought about by providing small activation energies in the form of heat, light, electric field, pressure, and more, similar to the behavior of semiconductors. To understand the behavior of the TMOs it is very important to understand their crystal chemistry, which embodies both the chemical bonding that exists between the constituent species and the crystal structure of the TMOs. The aim of this chapter is to briefly outline the general structural and electronic properties of TMOs for a later discussion of the correlation that exists between these intimately related aspects and their optical, electrical, and switching properties. A detailed discussion of the electronic and structural properties of certain TMOs of more relevance to the chromogenic switching properties discussed in this book is also carried out. The effect of nanostructuring of TMO thin films on their switching behavior is also discussed briefly.

2.1.1 Transition Metal Oxide Electronic Structure It is the partially filled d orbitals of the transition metal and the electrone electron interactions in these orbitals that lead to most of the interesting and unique properties seen in TMOs. TMOs are formed by the charge exchange Transition Metal Oxide Thin FilmeBased Chromogenics and Devices ISBN 978-0-08-101747-0 http://dx.doi.org/10.1016/B978-0-08-101747-0.00002-7

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between the highly electronegative oxygen atoms and the less electronegative transition metal atoms. By losing different numbers of electrons in these orbitals to the oxygen atom, the transition metals can form various oxides. This exchange leads not only to a wide range of bonding but also to different structures and phases of the TMOs, especially when other compatible atoms are included in the structure and bonding. It is a formidable task to achieve a comprehensive theoretical model that can satisfactorily account for all the complex behaviors of the wide range of TMOs that can be formed in this way. Hence, only a brief framework of the underlying mechanism touching the most typical TMOs is discussed. The electronic properties of solids, in general, stem from the outermost or valence electrons. The two extreme theoretical considerations in dealing with the electronic properties of such solids are the following [2]: 1. The situation in which the electrons are localized with nearly no superposition of bands from the neighboring atoms. This situation is dealt with by both the ligand field theory and the band theory. A large energy (U) is required to transfer the electron in this band of one atom to an equivalent site of another atom. The electrons are strongly bound to the atom core, forming a narrow band with an energy spread W (W > U). A special situation arises, however, when W w U, in which the electrons are strongly correlated, as in the case of transition metals. The essential feature of such solids is that their behavior can be described only in terms of this electronic correlation. In TMOs we have the special case in which the localized and delocalized (itinerant) outer shell electrons can be present simultaneously. Hence, the electronic properties of TMOs are better described in terms of strongly correlated electrons than in terms of the behavior of a single electron. The chemical classification scheme is based on the progressive filling of electrons into the valence shells. Following this classification rule, the periodic table is formed of s, p, d, and f blocks [3]. Of these, it is, generally, the elements with partially filled d subshells that are referred to as the transition elements. In free space, the outermost shells of transition metal atoms are formed of incomplete d shells and filled s shells. These can be transition metals formed of 3d, 4d, and 5d subshells partially filled with electrons. The partially filled orbitals of these metals lead to a fairly good thermal and electrical conductivity as well as to multiple oxidation states. Hence, the electron transfer between these states

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leads to very interesting optical, electrical, and magnetic properties as well as to their interesting applications in various fields. The oxides of transition metals are formed by combining one or more of the transition metals with oxygen atoms. These oxides are endowed with interesting features such as: (1) facile electron exchange with oxygen ions under very moderate activation energies, (2) a wide variety of bonding and multiple structures, (3) unusually high carrier density in certain phases, and more. Essentially, the oxygen atom, being an electronegative element with four electrons in its outer 2p shell, needs two electrons to complete this shell. The less electronegative transition metals with inner incomplete d shells and outer complete s shells lose two electrons to form the various TMOs [4,5]. By this exchange and covalent bond formation various TMOs, from monoxides (e.g., MnO), to dioxides (e.g., VO2), to trioxides (e.g., WO3), to higher and complex oxides such as perovskite (e.g., LaMnO3), can be formed, in which the transition metal gives off its s electrons and the required number of d electrons. This covalent bond formation in TMOs leaves the transition element with different numbers of d electrons in its outer shell. The interesting properties in these oxides stem from these outer or frontier d electrons. For example, as per the scheme of filling the electronic orbits, the element tungsten (W), with an atomic number of 72, has the following electronic configuration: 1s22s22p63s23p64s23d104p65s24d105p66s2 4f145d4. Hence, this leads to the facile formation of the trioxide of tungsten (WO3) by sharing the six outer electrons with three oxygen atoms, shell 6 with two electrons in subshell s and the outermost shell (5) with four electrons in subshell d. The W atom becomes W6þ and each oxygen atom becomes O2. Similarly, vanadium dioxide (VO2), with vanadium’s atomic number of 23 and electronic configuration of 1s22s22p63s23p64s23d3, is formed by giving off two of its electrons from subshell 4s and two from the 3d subshell. The vanadium atom thus becomes V4þ and each oxygen atom becomes O2. Thus, in all the TMOs it is these outer d bands and the oxygen p bands that are the most significant in determining the electronic properties of TMOs. In addition to this oxide formation with a single or multiple transition element it is possible to have other alkaline or rare earth metal atoms that can also provide the needed electrons to the oxygen atoms. A wide range of transition metal-based oxides can thus be formed from the simple monoxides (MO) to the complex oxides of the form RxMmOn, where M represents the transition metal and R can be any other suitable atom that can be included. These TMOs can thus be created by adjusting the number of each atom in the compound (x, m, n) and having different numbers of electrons in the outer d bands of the transition metal.

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This variation in the number of electrons in the d band of the transition metal imparts a wide range of interesting electronic properties to the TMOs formed. The interelectron repulsion between the d electrons of the outer shell is the electron correlation not taken into account in the aforementioned simple band theory. The majority of TMOs not only are considered to be strongly correlated systems but also are treated as the prime examples of correlated systems. The approach of taking into consideration the electroneelectron correlation in the band structure calculation was first proposed by Mott [6]. His approach is especially pertinent to partially filled narrow bands such as the d bands in transition elements. In this case the energy degeneracy of the electrons in this band is removed because of electrostatic interactions. In Fig. 2.1 is shown a schematic of the density of states in such a case of many electron orbitals for 3d electrons (right side) in comparison with the one-electron consideration (left side) with a wide band. The Coulomb repulsion leads to the separation of states of 3dn and 3dnþ1 as shown in Fig. 2.1A. The 3dn levels are supposed to be completely filled and Ef is the Fermi level. If some energy is provided to the system to overcome this repulsion this would lead to the broadening of the d bands (Fig. 2.1B) and eventual overlap (Fig. 2.1C) of the two bands to form delocalized or itinerant bands. Mechanisms that lead to such broadening are electrostatic interactions with surrounding cations or anions. In situations such as the one shown in Fig. 2.1B with nearly overlapping bands, external experimental conditions such as heat and others can lead to the overcoming of the Coulomb repulsion, thus leading to Mott transitions of certain TMOs.

Figure 2.1 Energy density schematic of a transition metal atom with n 3d electrons. (A) 3d levels separated by Coulomb repulsion, (B) nearly overlapping 3d bands, and (C) band broadening due to overlap. (From P.F. Weller, Solid State Chemistry and Physics, vol. 1, Marcel Dekker Inc., 1973.)

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This also forms the basis for the classification of the TMOs as suggested by Zaaned, Sawatzky, and Allen [7]. With hardly any overlapping between the d orbitals of neighboring atoms in the TMOs, the d electrons are highly localized around the transition metal atom. This leads to a stronger Coulomb interaction between these electrons. The Coulomb interaction between the d electrons of the transition metal is treated in terms of the Hubbard value, U. The energy required to populate these d bands is given by UNd(Nd  1). The classification of the different TMOs is thus based on the strength of this correlation between the d electrons and the relative position between the frontier d bands and the lower oxygen p bands. With D (also called the charge transfer energy) representing the energy required to transfer an electron from the oxygen p band to the transition metal d band, the classification of TMOs can be carried out as follows. When D > U, the TMO is a Mott insulator and the energy regime in which this occurs is referred to as the MotteHubbard regime. However, when U > D, the TMO becomes a charge transfer insulator and the energy regime in which this occurs is referred to as the charge transfer regime. The energy schematic associated with this classification is shown in Fig. 2.2. The early transition metals such as vanadium (V) with fewer d electrons have oxygen p bands lower in energy and far from the Fermi level as shown in Fig. 2.2A. A higher energy (D) is required to transfer p electrons to the conduction band. The behavior of the TMOs in this lower energy regime is dominated by the low-energy excitations between the d bands. These early transition metal-based TMOs are Mott insulators. By contrast, in the late transition metals the number of d electrons increases but so does the positive charge at the ion core. This positive charge attracts electrons of the oxygen p bands thus lifting their energies to higher levels and bringing them closer to the Fermi level. The

Figure 2.2 Energy level schematics of a transition metal oxide of the type: (A) MotteHubbard insulator and (B) charge transfer insulator.

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Table 2.1 Transition metal series with partially filled d orbitals

oxygen p band is essentially placed between the upper and the lower Hubbard band as shown in Fig. 2.2B. These later transition metal-based TMOs are termed MotteHubbard insulators. Other TMOs in the middle, such as Mn with d5, show a combination of electronic behaviors of the two types seen. The changes occurring with increasing d electrons as one moves from the early transition metals to the late transition metals in the periodic table (Table 2.1) can be illustrated through the electronic structure of the monoxides, as an example. In cases in which there is an overlap between the wave functions from the neighboring atoms, consequently forming a band in place of individual atomic levels, the TMOs behave as conductors. This behavior is seen in early monoxides of the 3d row shown in Table 2.1, such as VO and titanium oxide (TiO), where owing to the large overlap these TMOs show a metallic behavior. As one moves to late transition elements the number of d electron increases as does the positive charge at the nucleus of the atom. As a consequence of this increasing nuclear charge there is less superposition of d orbitals of neighboring atoms leading to a less metallic behavior. Thus, for example, TiO with two electrons in the frontier 3d shell exhibits a metallic behavior, while nickel oxide (NiO) with eight electrons in its frontier 3d shell is an insulator. The Coulomb repulsion between the electrons is given in terms of the Hubbard energy, U, which is the energy needed to transfer an electron to a site already occupied. Hence, this can be written in terms of the energy needed to remove the electron from its initial site expending the ionization potential, I, and the energy needed to put it with another electron, EA, the electron affinity: U ¼ I  EA

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(A) Energy Upper Hubbard band

Fermi level

U

Lower Hubbard band

Bandwidth, W

(B)

Conducvity

Conductor Excited state

Energy Upper Hubbard band U

a Lower Hubbard band Ground state Insulator 1/a

Figure 2.3 (A) MotteHubbard splitting as a function of the bandwidth. (B) The variation of U as a function of the interatomic distance (a). Inset shows the transfer of an electron to a neighboring site upon excitation.

As shown in Figs. 2.1 and 2.2, this U represents the energy needed to transfer an electron between the already occupied d states, which are single atomic states when there is no interaction between the atoms, and develop into a band with increasing interaction and overlap. This energy gap between the d states is also called the MotteHubbard splitting, whose dependence on the bandwidth, W, is shown in Fig. 2.3A. The dependence of U on the interatomic spacing can be sketched as shown in Fig. 2.3B. Shown are the Hubbard bands in the ground state (lower Hubbard band), where each atom core has an electron, and the excited state (upper Hubbard band), where the

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electron is transferred to another atom with an electron by expending an energy U. A large a and no overlap between the atoms (W ¼ 0) corresponds to the insulator state, as a large energy U is required to transfer the electron to the excited state. With decreasing interatomic spacing and increasing atomic interaction (increasing W), the energy U required for transferring the electron to the excited state decreases. Eventually at a critical value of interatomic spacing when the two bands overlap, U becomes zero and the oxide undergoes an insulator-to-metal transition. Hence, the metal-to-insulator transition (MIT), called the Mott transition, seen in some TMOs is a combination of the electroneelectron repulsion and the overlapping of the atomic interaction between neighboring transition atoms.

2.1.2 Transition Metal Oxide Crystal Structure As mentioned earlier, the TMOs with their special electronic structure and chemical bonding crystallize in a wide range of structures. These are ReO3, spinel, rock salt, rutile, and cuprite and corundum-type structures [1]. The common feature observed in the TMO structure stems from the fact that the transition metal ion radius is smaller than the oxygen ion radius. This leads to the cubic close packing of anions and formation of octahedral and tetrahedral networks of the oxides with the smaller transition metal cations held in the holes. A summarized classification of the various TMOs in their bulk form according to the structure is given by Adela Muñoz-Pàez [8]. 2.1.2.1 Monoxides All the transition metals with 3d orbitals (first row of elements in Table 2.1) form monoxides, except scandium (Sc) and chromium (Cr). Five of the transition metals with 4d orbitals (Zr, Nb, Pd, Ag, and Cd of the second row in Table 2.1) form monoxide. Among the transition metals with 5d electrons (third row in Table 2.1), it is only mercury (Hg) that forms the monoxide. Most of these monoxides assume the rock salt (NaCl) crystal form as shown in Fig. 2.4A along with the octahedral unit. This structure can also be described in terms of the octahedral unit MO6 formed by corner-sharing oxygen ions and in which the transition metal cation (M) is held in the cavity of the octahedral as shown in Fig. 2.4A and B. While this octahedral structure is adopted by stoichiometric oxides, large deviations can be found in many of the TMOs, due to nonstoichiometry, deficiency or lack of certain atoms, and even atoms located out of their ideal position. Some such monoxides assume wurtzite (e.g., ZnO) and PtS (e.g., CuO) structures.

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Figure 2.4 (A) Rock salt structure with the octahedral unit: red circles (gray in print version) represent the oxygen ion and the open double circles represent the transition metal. (B) MO6 octahedra formed by corner-sharing oxygen ions.

Figure 2.5 (A) Rutile (TiO2 type) structure of the transition metal dioxides with the M4þ cation surrounded by hexacoordinated oxide anions in the octahedral unit. (B) Fluorite structure of the transition metal dioxides with the M4þ cations forming the facecentered cubic structure with the oxide anion held in the middle of each octant.

2.1.2.2 Dioxides The transition metals with 3d orbitals forming the dioxides are titanium, vanadium, chromium, and manganese. Almost all the transition metals with 4d and 5d orbitals form the dioxides except for cadmium. The crystal structure of transition metal dioxides is either fluorite- or rutile-like. These structures are shown in Fig. 2.5. The rutile form is made up of corner-sharing octahedral units of the type MO6 in a tetragonal cell as seen in Fig. 2.5A with the smaller M4þ cations surrounded by the hexagonally coordinated oxide anions. Many dioxides show a deviation from this structure with the

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Figure 2.6 Corundum (Al2O3 type) structure of the transition metal oxides (M2O3) with the M3þ cation held in the octahedral unit of hexacoordinated oxide anions.

alternate transition metal atoms being separated by long and short distances. The fluorite structure formed by the larger M4þ cations of zirconium and hafnium is made up of a face-centered cubic cell with the transition metal cations occupying the corners and the face center position as shown in Fig. 2.5B. The oxide anions are located at the eight 1/4 1/4 1/4 positions. 2.1.2.3 SesquioxidesdCorundum Type Most of the 3d transition metals except zinc form these oxides with a corundum-type structure (M2O3). The large M3þ cations are located in the tetrahedral units formed by the oxide anions as shown in Fig. 2.6 where the cations occupy most of the volume within the tetrahedral unit. 2.1.2.4 SpinelsdM3O4 The only transition metals showing this stoichiometry are manganese, iron, and cobalt with the 3d orbitals. Emanating from the multiple oxidation states of these oxides, the structure is formed of both tetrahedral and octahedral units as shown in Fig. 2.7. The adjacent octants of the unit cells are different from each other as shown in schematic (1) of this figure. The

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Figure 2.7 The spinel structure of the transition metal oxides (M3O4) formed by different octants in the unit cell (1), with the oxygen and metal bonding in these octants (2) and with the tetrahedral (3) and octahedral (4) units formed. (http:// encyclopedia2.thefreedictionary.com/Ferrite.)

Figure 2.8 Layered orthorhombic structure of vanadium pentoxide (V2O5) film with the up-and-down alternating VO5 square pyramidals. (http://www.chemtube3d.com/ solidstate/SS-V2O5.htm.)

M* cations and the oxide anion position and coordination are different in these adjacent octants and are shown in schematic (2) of the figure. The M* cations held in the tetrahedral and octahedral units are shown in schematics (3) and (4) of the figure. 2.1.2.5 M2O5 Only the transition metals vanadium, niobium, and tantalum are known to form these pentoxides. They form a layered structure, with a typical example shown in Fig. 2.8 for vanadium pentoxide (V2O5). The orthorhombic structure is formed of edge-sharing and corner-sharing VO5 square pyramids, which alternately face up and down. In addition to the five square pyramidal neighbors, the vanadium

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has a sixth distant neighbor, which forms a distorted VO6 octahedron. Nb2O5 and Ta2O5 exhibit similar layered and chain structures and multiple crystalline phases. 2.1.2.6 MO3 The only transition metals showing this stoichiometry are chromium, molybdenum, rhenium, and tungsten, most of which are of great pertinence to chromogenic properties. The simplest cubic structure of ReO3 is shown in Fig. 2.9A. The rhenium cations are considered to be located at the corners of the unit cell, while the oxide ions are located in the linking middle position on the edges between the Re cations. Hence, this structure forms a corner-sharing octahedron by the hexacoordinated O2 around the M6þ cation as shown in the figure. Molybdenum trioxide (MoO3) forms a layered and open structure

Figure 2.9 (A) The ReO3 structure formed by the corner-sharing ReO6 octahedral units in which the Re6þ cations are at the corners of the cube and the O2 anions on the edges. (B) MoO3 structure formed by the corner and two adjacent edges sharing MoO6 octahedral units. (http://www.geocities.jp/ohba_lab_ob_page/structure6.html.)

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in which the one MoO6 octahedral chain formed by the corner sharing also shares the adjacent edges with another chain forming a double layer as shown in Fig. 2.9B. WO3 is formed by the corner-sharing and slightly distorted WO6 octahedral units with a large open structure. 2.1.2.7 Perovskite Structure Perovskite is a structure formed of corner-sharing octahedral units and closely related to the ReO3 structure shown in Fig. 2.9A. The ReO3 structure as shown in this figure with the Re cations occupying the corner sites of the cube and oxide ions in the middle of the edges of the cube leads to a formidable hole in the middle of the cube. Perovskite structures are denoted by ABO3, with the large A ions being located in this hole with a coordination number of 12 by the oxide anions, as shown in Fig. 2.10A and

Figure 2.10 (A) The perovskite structure (ABO3) formed by the corner-sharing BO6 octahedral units with a large A ion in the center of the cell. (B) Line diagram of the perovskite transition. A, large central ion; B, metal cations; O, oxide ions. (http://www. solarchoice.net.au/blog/wp-content/uploads/Perovskite-crystal.png.)

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by the line diagram in Fig. 2.10B. These structures of the type AMO3, which form a tunnel structure, are very relevant to the electrochromic properties of the TMOs owing to the facile and reversible intercalation of foreign species. As seen in the case of WO3 earlier, the arrangement of the perovskite can be distorted in many cases. Nonstoichiometric perovskites of the form AxMO3 (e.g., tungsten bronzes, LixWO3) are very important for the electrochromic properties of TMOs. The intercalation of cations such as Hþ, Liþ, and Naþ along with the associated electrons in these atoms leads to the mixed valence states in the TMO network. The amount x of cation A that can be intercalated depends on the valence state of A and can vary over a wide range. It is also possible to have oxygen-deficient perovskites of the type AMO3x through the formation of other polyhedral chains to form the perovskite network and through the adoption of multiple oxidation states by the transition metal. These multiple oxidation states created in TMOs through nonstoichiometry and oxygen deficiency are very important and lead not only to interesting optical and electrical properties in these materials but also to their switching properties.

2.2 TYPICAL EXAMPLES OF TRANSITION METAL OXIDES SHOWING CHROMOGENIC PROPERTIES In this section a detailed discussion of the electronic and structural properties of the best-known TMOs showing chromogenic properties is given. These include WO3, the most promising and researched electrochromic material; the vanadium oxides (VO2 and V2O5) known for their electrochromic and thermochromic properties; and MoO3, showing very efficient electrochromic and photochromic properties. Some details of other similar TMOs such as NiO, chromium oxide (Cr2O3), niobium oxide (Nb2O5), and tantalum oxide (Ta2O5), known for their electrochromic properties, are also given. The mechanism of the chromogenic switching of these materials is also discussed briefly.

2.2.1 Tungsten Trioxide By far, WO3 has been the most studied and used TMO because of its efficient electrochromic and photochromic properties since the discovery of its electrochromic properties by S.K. Deb in 1969 [9]. The ReO3 and perovskite structures discussed earlier are very relevant to the properties of WO3 in its normal state and its electrochromically switched state, respectively. In its normal state WO3 has a ReO3 structure (Fig. 2.9) formed of

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corner-sharing WO6 octahedral units. Each tungsten ion with an oxidation state of W6þ is surrounded by six oxygen ions (O2) in this unit. As shown in Fig. 2.11A a network of the octahedral WO6 is formed in which between each two tungsten ions there is an oxygen ion in the linear chain. These octahedra are the building blocks of WO3 that lead to their various arrangements and thus to various structures. In perfect crystalline and stoichiometric form, the corner sharing at the oxygen sites leads to the regular network formation. This network of corner-sharing WO6 octahedra contains large empty spaces as shown in Fig. 2.11A, which play an important role in the WO3 electrochromism via intercalation of foreign ions into these spaces. At room temperature WO3 exhibits a monoclinic structure. Depending on the temperature and due to changes in the bond length and WO6 octahedra displacements, WO3 can have various structures from tetragonal to orthorhombic to monoclinic to triclinic. Substoichiometric forms of WO3 are known to display edge sharing of the WO6 octahedra as shown in Fig. 2.11B, referred to as the Magnéli phases [10,11]. These Magnéli phases in TMOs are variations derived from the initial perovskite crystal structure through edge sharing and structural distortion arriving from ordered oxygen deficiencies. The large space formed between the WO6 octahedral network in the WO3 perovskite-like structure (ABO3) in the absence of the third species ion A (Fig. 2.10) provides some degree of variation in the W position as well as for the WO6 octahedron orientation. These distortions are shown to lead to 11 different structures of WO3 [12]. Three of these structures are shown in Fig. 2.11C in comparison with the regular perovskite structure. The open space shown in Fig. 2.11 has been found to increase as the structure changes from perovskite to tetragonal to hexagonal to pyrochlore. While stoichiometric WO3 is highly transparent in the visible region of the spectrum and electrically insulating, oxygen-deficient WO3 exhibits slightly bluish to fairly brown coloration depending on the degree of deficiency. In this substoichiometric form, along with the optical absorption an electrical conductivity is present in the oxide. As discussed earlier, the electronic properties of WO3 emanate from the frontier orbitals of W and oxygen (O) as shown in Fig. 2.12 [13,14]. The valence bands correspond to the oxygen 2p bands and hence are completely filled as shown, while the empty conduction band is made up of the tungsten 5d orbitals. The Fermi level is positioned between these two bands as shown in the figure. The bandgap (D) between the valence and the conduction bands is reported to be 2.6e3 eV in polycrystalline WO3 thin films [15]. In comparison with the regular network, despite distortion of the WO6 octahedra and large

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(A) Open space

O-2 W+6

(B)

(C)

Figure 2.11 (A) WO3 structure formed by a network of corner-sharing WO6 octahedral units; each WO6 octahedron is also shown on the right. (B) WO3 network formed by edge-sharing WO6 octahedral units. (C) Tetragonal, hexagonal, and pyrochlore structures of WO3 in comparison with the regular perovskite structure. (From K. Bange, Colouration of tungsten oxide films: a model for optically active coatings, Solar Energy Materials and Solar Cells 58 (1999) 1e131.)

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29

WO3 W 5d

eg

4e

t2g

6e

Ef 12e

2p O3

6e

6e

2s

Figure 2.12 A simple band diagram of WO3. Bands arising from the empty tungsten 5d orbitals (top) and filled oxygen 2p orbitals (bottom) and Fermi level (Ef) are shown. (Adopted from L. Berggren, Optical Absorption and Electrical Conductivity in Lithium Intercalated Amorphous Tungsten Oxide Films (Ph.D. thesis), Uppsala University, 2004.)

variation in bond length within the octahedrons, of polycrystalline WO3, the amorphous phase is one in which there is no long-range ordering of the octahedral network [16]. Hence, the amorphous film can be considered as a random network of microcrystals of which the building blocks are still the WO6 octahedra. The bandgap values for amorphous WO3 are, in general, higher than those of crystalline WO3, being in the range of 3.2e3.5 eV [17]. In addition to the distortion, variation in bond length and angle, and oxygen deficiency. as well as the amorphous nature of WO3, the structure can be marred with various other defects such as interstitials, dangling bonds, and vacancies [18]. These high-density defects are very common in WO3 thin films during their preparation using various techniques and are important to the electrochromic performance of the films. The various deposition techniques commonly used for the preparation of the WO3 films are expected to lead to a number of species of the type WxOy where x and y vary over a wide range [11]. The electrochromic coloration in WO3 is expected to arrive through the double insertion of protons or alkali ions (Liþ, Naþ, Kþ, etc.) and electrons ensuing from a reversible chemical reaction, which can be represented through a simple transformation as follows: WO3 þ xMþ þ xe 4 Mx WO3 ðTransparentÞ

ðBlueÞ

(2.1)

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where M represents the type of ion inserted. Upon the intercalation of x number of ions and electrons into WO3, its optical properties change from transparent to blue in the visible region of the spectrum with the reversible formation of MxWO3, termed tungsten bronze. Owing to the perovskitelike open structure of WO3, ions can be intercalated into the material, bringing about an intense coloration. This process is fairly reversible until an optimum value of x is reached. Continued ion insertion beyond this point, though possible, can lead to irreversible reactions. For example, in the case of lithium ion insertion, the reversibility (intercalation and extraction) into WO3 can be maintained up to an x value of 0.7 [14]. The value of x that can be intercalated into the tungsten bronzes is not expected to be dependent on the size of the inserted ions [1]. Various models have been proposed to account for the role played by the inserted ions and electrons in determining the electronic, structural, and optical properties of the intercalated WO3 (MxWO3). These models range from the oxygen-vacancy-generated F centers [13] to intervalence charge transfer (IVCT) of inserted electrons [19] to polaron absorption [20] to oxygen extraction [21,22]. However, as seen earlier, the wide range of energy bands and structures made possible by the different arrangements of the octahedral units and the ensuing phases lead to a wide range of, and even inconsistent, experimental results. It is undoubtedly a daunting task to find a comprehensive theoretical model that can explain the mechanism underlying the electrochromic coloration in WO3. Most of the results pertaining to the electrochromic coloration of WO3 are discussed under two broad theoretical headings. 2.2.1.1 Crystalline Tungsten Trioxide The experimental results of polycrystalline WO3 are invariably discussed in the light of the Drude theory applicable to materials with a high free electron density [23]. As discussed earlier, in the case of crystalline WO3 with a longrange order and with sufficient overlapping of wave functions between neighboring sites with formation of bands, the inserted extrinsic electrons populate the bands and are not localized on any sites, becoming nearly free as in the case of metals or heavily doped semiconductors. Their behavior is characterized in terms of the plasma frequency, up, which depends on the free electron density, n, through the relation [24]: up ¼

4pne2 m

(2.2)

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31

connected to the dielectric constant ε(u) in the high (optical) frequency range by: u2p εðuÞ ¼ 1  2 (2.3) u With light incidence at frequencies (u) lower than the plasma frequency (up), the free electron cloud formed by the inserted electrons is capable of oscillating at the frequency of the imposed electromagnetic field, thus screening it from penetrating the material and in turn producing a reflectance at the same frequency. Hence, in crystalline WO3 the plasma cutoff frequency can be varied reversibly by controlling the density of electrons (ions), n, injected via the electrochromic phenomenon. Below the plasma frequency the light is reflected, and above this frequency, the light is transmitted. Hence, the optical behavior of crystalline WO3 is dominated by the density of electrons injected. The inserted ions that occupy the open spaces between the octahedral network (Fig. 2.10) essentially maintain the charge neutrality of the system. The electrochromic coloration in this case is through reflectance modulation. The variation of reflectance occurring through the insertion of increasing electron (ion) density is given in Fig. 2.13. Reflectance values calculated and corresponding to the dielectric constant of Eq. (2.3) are given for a crystalline WO3 of 0.2 mm thickness intercalated with various densities of electrons, ne (protons). As can be seen in this figure, with increasing ne, the reflectance increases and the plasma edge moves toward lower wavelengths. Hence, the reflectance modulation is more effective in the infrared region at moderate intercalation and

Figure 2.13 Calculated reflectance spectra as a function of the density of electrons (ne) inserted in a 0.2-mm-thick crystalline WO3 film. (From C.G. Granqvist, Handbook of Inorganic Electrochromic Materials, second ed., Elsevier, 2002.)

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becomes effective in the visible region only at high levels of intercalation. The film becomes more and more metallic with increasing free electron density. 2.2.1.2 Amorphous Tungsten Trioxide The situation in the case of amorphous WO3 in which a long-range order is lacking is much more complicated. The experimental results obtained in this case are analyzed in the light of the widely recognized and similar theories of IVCT [19] and polaron absorption [20]. Unlike in the case of crystalline WO3, owing to the lack of sufficient overlapping of the wave functions between neighboring sites, the inserted extrinsic electrons are localized on tungsten sites, changing the oxidation state of these sites from W6þ to W5þ. The coloration through electrochromic mechanisms is then widely accepted to be arriving through the transfer of the electron between these mixed oxidation sites upon the absorption of incident light with compatible energy, hn, as shown in Fig. 2.14. The ion associated with this electron, which essentially maintains the charge neutrality in the region, is shown to be occupying a nearby open space between the WO6 octahedra. The electron trapped on tungsten site A, transforming it to W5þ, is then e-

hν

W6+ W 5+

B A

M

+

W 6+ W 6+

Figure 2.14 Schematic of the trapping of an electron (e) on a W(A) site and its transfer to a neighboring W(B) site through the absorption of a photon (hn). The associated ion (Mþ) is in the nearby open space. (Adopted from T. Yoshimura, Oscillator strength of small-polaron absorption in WOx (x  3) electrochromic thin films, Journal of Applied Physics 57 (3) (1985) 911e919.)

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33

transferred to the neighboring tungsten site B through this photon absorption. These two steps can be represented as follows [25]: W6þ(A) þ W6þ(B) þ e / W5þ(A) þ W6þ(B)

(2.4)

W5þ(A) þ W6þ(B) þ hn / W6þ(A) þ W5þ(B)

(2.5)

According to the IVCT model, the coloration of amorphous WO3 is due to this absorption of light by the electrons for the intervalence transfer. Unless one is dealing with very nonstoichiometric tungsten oxide, IVCT is the most applicable model to account for the absorption modulation seen in amorphous WO3 films. A model closer to the IVCT is the polaron model [26] applicable to solids in which the variation in the bond lengths between oxygen and tungsten ions arriving as a result of the insertion of the extrinsic electrons and ions is taken into account. The occurrence of this variation has been experimentally shown in amorphous WO3 films [27]. Polarons are quasiparticles formed by the combination of electron (hole) and polarized lattice (phonon) pairs. The introduction of an extrinsic electron in the present context of amorphous WO3, for example, leads to the displacement of neighboring ions, which in turn change the frequency of the phonons (lattice vibrations). The electron in this case is expected to get self-trapped in the potential well created by distortion in the lattice due to its presence. Depending on the extent of the distortion this extrinsic electron creates in the lattice neighborhood, the terms small polaron (distortion of the order of lattice parameter) and large polaron (distortion extending over a large number of lattice points) are applied as shown in Fig. 2.15.

Figure 2.15 Schematic of a small polaron (left) and a large polaron (right) created by the trapping of an electron (e) on a W site. R, radius. (From O.F. Schimmer, G. Wittver, G. Baur, G. Brandt, Dependence of WO3 electrochromic absorption on crystallinity, Journal of Electrochemical Society 124 (5) (1977) 749e753.)

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E

4Ep q v A(W5+)

Ep

B(W6+)

Figure 2.16 Schematic of the lattice energy (E) as a function of the lattice coordinate around the A(W5þ) and B(W6þ) sites due to the trapping of the electron. Shown here also are the polaron energy (Ep) and the optical transition between the two sites (4Ep).

In the small polaron model the absorption of energy and transfer of the extrinsic electron localized on one site to another as shown in Fig. 2.14 is connected to the phonon created. The excess energy from this transfer is absorbed by the lattice for the creation of the phonon. Considering the one-dimensional picture of the lattice, the energy (E) of the parabolic form, Dq2, associated with it as one moves along the lattice coordinate (q) in the neighborhood of the above considered W(A,B) site is as shown in Fig. 2.16 [28]. Ep is the polaron energy due to the insertion of the extrinsic electron (ion), v represents the difference in the energy between the neighboring sites in the lattice. Then the vertical transition of the electron from the ground state of site A(W5þ) to the neighboring B(W6þ) site through the absorption of a photon is as shown in Fig. 2.16. The energy corresponding to this transition is 4Ep. Thus both the theories, IVCT and small polaron absorption, account to a large extent for the absorption modulation seen in amorphous WO3 through the electron transition between neighboring sites. However, the degree and energy range of their validity varies widely. In Fig. 2.17 is given a comparison of the agreement of the optical density spectra of these two theories with experimental data on amorphous WO3 films intercalated with electrons and Hþ [16].

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Figure 2.17 Optical density spectra for amorphous WO3 film intercalated with Hþ. Comparison of the experimental data with (top) intervalence charge transfer theory and (bottom) small polaron absorption theory.

It can be seen that while the IVCT theory is in good agreement with the experimental data over nearly the entire spectrum presented in Fig. 2.17, the polaron absorption theory is more suited to lower energy regions.

2.2.2 Vanadium Dioxide VO2 is the most-studied TMO, exhibiting a very efficient thermochromic switching near room temperature. This material is very well known to undergo a MIT at around 68 C, making it one of the most sought-after materials for chromogenic switching applications ever since the discovery

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Figure 2.18 The formation of various vanadium oxides and their phases. (From C.H. Griffiths, H.K. Eastwood, Influence of stoichiometry on the metal-semiconductor transition in vanadium dioxide, Journal of Applied Physics 45 (1974) 2201.)

of its thermochromic switching by Morin [29] in 1959. This temperature at which the transition occurs is referred to as the transition temperature, Tt. Although it is referred to popularly as MIT, the more appropriate term is semiconductor-to-metal transition because of the reasonably high to fairly high conductivities it exhibits in these, respective, states. This transition is accompanied by a change in the structure of the material. As mentioned before, the base transition metal can exhibit multiple oxidation states from 2þ to 5þ and thus can form a wide range of bonding with oxygen to form an unusually large number of oxides as shown in Fig. 2.18 [30]. VO2 with an oxidation state of 4þ is found to form in monoclinic or triclinic or tetragonal phases depending on the temperature. The structural changes and the accompanying phase change occurring with temperature can be better understood by commencing with the hightemperature rutile structure. At temperatures higher than the Tt of 68 C, the VO2 crystal exhibits a tetragonal structure of the rutile type. The smaller V4þ cations are surrounded by the hexagonally coordinated oxide (O2) anions as shown in Fig. 2.19A [31]. In this tetragonal structure the lattice constants are 4.5546 Å along the a axis and 2.8514 Å along the c axis [32]. Hence, the VeV separation is regular along this c axis. In each of these VO6 octahedra

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37

Figure 2.19 (A) VO2 tetragonal structure of the rutile type formed of corner-sharing and edge-sharing VO6 octahedra. (B) The VO6 octahedral chain formed along the c axis of the tetragonal structure by edge sharing of opposite sides of the units. (From V. Eyert, The metal-insulator transitions of VO2: a band theoretical approach, Annals of Physics 11 (2002) 9e61; P.A. Cox, Transition Metal Oxides e An Introduction to Their Electronic Structure and Properties, Oxford University Press, 1992.)

the hexagonally coordinated vanadium atom is equidistant from each of the corner oxide anions, thus bringing about a large charge screening and no metallic bonding between the vanadium atoms. The VO2 structure is made up of a mixture of corner-sharing and opposite edge-sharing octahedral units of the type VO6 in a tetragonal cell. The continuous stacking of these VO6 octahedra along the opposite edges of each unit forms a chain along the tetragonal c axis as shown in Fig. 2.19B [33] through the edge sharing. The neighboring octahedral chains thus formed are interlinked by the corner sharing of the octahedral. At temperatures below the Tt the pure VO2 crystals exhibit a monoclinic structure referred to as type M1 as shown in Fig. 2.20A. The lattice constants in this case for the monoclinic structure are 5.75 Å along axis a, 4.52 Å along axis b, and 5.38 Å along axis c with b ¼ 122.60 degrees. The rearrangement of the lattice leads to the doubling of the size compared to the high-temperature tetragonal structure. This structure can be considered as a distortion of the high-temperature VO2 structure discussed before because of two important changes. These changes between the monoclinic and the tetragonal structures can be seen more clearly in Fig. 2.20B [34]. The VeV distances become irregular along the rutile c axis owing to alternate short (2.65 Å) and long (3.12 Å) distances between them, thus leading to the pairing up of the nearby

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Figure 2.20 (A) VO2 monoclinic structure of M1 phase at temperature below the transition temperature (Tt). (B) Comparison of the tetragonal (high temperature) and monoclinic (low temperature) phases of VO2 with the localized electrons (monoclinic) and free electrons (tetragonal). (C) Stacking of the VO6 octahedra in the rutile (left) and monoclinic (right) phases. (From (A) V. Eyert, The metal-insulator transitions of VO2: a band theoretical approach, Annals of Physics 11 (2002) 9e61. (B) J.D. Budai, Metallization of vanadium dioxide driven by large phonon entropy, Nature 515 (7528) (2014) 535e539. (C) Y. Wu, et al., Decoupling the lattice distortion and charge doping effects on the phase transition behavior of VO2 by titanium (Ti4þ) doping, Scientific Reports 5 (9328) (2015) 1e8.)

cations along the monoclinic a axis and tilting of these bonds with respect to the rutile c axis. Shown in this figure are also the changes occurring to the d orbital electrons. In the monoclinic (low temperature) phase, owing to the VeV pairing up, the electrons are bound to the orbitals. At high temperature and in the ensuing tetragonal structure, the electrons are free. Thus VO2 transforms itself from a low-conducting (insulator-like) monoclinic phase at low temperature to a high-conducting (metallic) tetragonal phase when heated above the Tt. The figure also shows the associated lattice vibrations (phonons)

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Figure 2.21 (A) VO2 monoclinic structure of M2 phase at temperature below the transition temperature (Tt). (B) Planar schematic comparison of the M1, M2, and rutile (tetragonaldshown as dashed lines) phases of VO2. The red circles (dark gray in print version) correspond to the vanadium atoms. (From (A) V. Eyert, The metal-insulator transitions of VO2: a band theoretical approach, Annals of Physics 11 (2002) 9e61. (B) J.P. Pouget, H. Launois, J.P. O’Haenens, P. Merenda, T.M. Rice, Electron localization induced by uniaxial stress in pure VO2, Physical Review Letters 35 (13) (1975) 873e875.)

predominant at higher temperature, which are argued to be directly responsible for the stabilization of the metallic high-temperature phase and the resulting change in conductivity of VO2 [34]. This aspect of phase change is discussed further in the next section. A comparison of the changes occurring to the stacking of the edge-sharing and corner-sharing VO6 octahedra between the high-temperature (rutile) phase and the low-temperature (monoclinic) phase is shown in Fig. 2.20C [35]. VO2 has been found to crystallize in another closely related (M1-like) monoclinic structure, which can also be regarded as a distortion of the regular rutile structure. The existence of such a VO2 phase, referred to as M2, has been demonstrated through an external doping or through an applied strain along an axis [36,37]. The M2 phase structure of VO2 is shown in Fig. 2.21A along with a planar comparison of this phase with M1 and the regular tetragonal structure of the rutile type in Fig. 2.21B. Comparing the monoclinic structure of the M2 phase shown in Fig. 2.21A with that of M1 shown in Fig. 2.20A, the deviations between these two structures with respect to that of the VO2 rutile structure are quite evident. The lattice constants in the case of the M2 phase are shown to be 9.0664 Å along axis a, 5.7970 Å along axis b, and 4.5255 Å along axis c with b ¼ 91.88 degrees. In the M2 structure it is essential to designate the two vanadium and three oxygen atoms V1 and V2 and O1, O2, and O3, to discuss the variation in their

40

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distances compared to the M1 phase. The V1eV1 distance along the rutile c axis is short (2.538 Å) and the V2eV2 distance is long (2.933 Å). Similarly, the distances between the vanadium and the oxygen atoms are also different. The V1 atoms have a V1eO1 distance of 1.868 Å, V1eO2 distance of 1.852 Å, and V1eO3 distance of 2.089 Å. Similarly, the V2 atoms have distances of 1.931, 2.127, and 1.726 Å with respect to atoms O1, O2, and O3, respectively. These differences are also shown in Fig. 2.20B in a planar schematic of the monoclinic M1 and M2 as well as the rutile structures. As discussed earlier, while all the VeV chains align in a zigzag fashion in the M1 phase, only half of the vanadium atoms are aligned this way in the M2 phase. The nonzigzag vanadium atoms show a metallic bonding. However, both the monoclinic phases (M1, M2) exhibit a phase transition to the rutile-type tetragonal phase above the Tt. 2.2.2.1 Vanadium Dioxide Phase Transition The phenomenon of VO2’s MIT, or more precisely metal-to-semiconductor phase transition, as a function of temperature has been discussed and studied at length but continues to be enigmatic. Despite a vast amount of theoretical calculation, computer modeling, and experimentation [38e45], a clear-cut understanding of the underlying phenomenon is still lacking. The primary reason for this difficulty in finding a comprehensive model is the simultaneous occurrence of the structural and phase (metal to insulator) change at the Tt. Further, the three phases discussed earlier are structurally different but exhibit comparable change in properties (optical and electrical) with these transitions (M1erutile, M2erutile) occurring with the same activation energy [46]. The various experimental results pertaining to the transition are debated mainly in the light of two principal models, the Peierls insulator model and the Mott insulator model. According to the Peierls model one-dimensional solids show insulating behavior due to the periodic distortion of the lattice [1]. In the case of VO2 the monoclinic M1 and M2 phases undergo this distortion owing to the zigzag structure (dimerization) formed at low temperature (T < Tt) as discussed before. The resulting unit cell is doubled, forming a bandgap. At high temperature, however, because of the stabilization of the structure due to lattice vibrations, the VO2 crystals form regular chains between the V atoms leading to a high conductivity in this tetragonal rutile-like structure. These changes are shown in Fig. 2.22A. On the other hand, according to the MotteHubbard model the MIT is considered in the light of the electronic structure of VO2 as a function of the temperature. As discussed in the earlier section, VO2 is a prime example

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41

(A)

(B)

Monoclinic phase

Insulator state

Tetragonal (rule) phase

Metallic state

(C)

Figure 2.22 Schematic comparison of the (A) Peierls and (B) Mott transitions in VO2. (A) Monoclinic structure (left) to tetragonal (rutile; right) structure. (B) Insulator (left) to metallic (right) state. (C) Electronic structure of the monoclinic (insulator) and tetragonal (metal) state.

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

of strongly correlated systems. The electroneelectron interaction is described in terms of the Hubbard potential, U. There being a single electron in the outer orbital of each metal atom, to double the occupancy a large amount of activation energy is required in the form of heat or an externally applied electric field. It is also possible to bring this about by the overlapping of the wave functions creating the overlapping of the bands by decreasing the interatomic spacing. Both these effects increase the ratio between the bandgap (W) and U, eventually bringing about the band overlap as shown earlier in Figs. 2.1e2.3, and the transition from an insulator to a metallic state. This can also be seen in Fig. 2.22B. The electroneelectron interaction prohibits the double occupancy of a state by electrons of opposite spin. A large activation energy would be required to overcome the repulsive Coulomb potential between these electrons. This corresponds to the single-occupancy states as shown. However, at high temperature, with the thermal activation energy or high pressure that causes the overlapping of the bands the electrons can overcome the Coulomb repulsion to achieve the double occupancy. The electronic structure associated with the two states is also shown in Fig. 2.22C. As can be seen in this figure, in the monoclinic (insulator state) the dimerization and tilting of the V atoms causes the 3d band to split into a filled bonding dǁ state and an unfilled antibonding dǁ* state as well as the upward shift of the p* band. This creates a bandgap of 0.7 eV between the dǁ and the p* bands and of 2.5 eV between the dǁ and the dǁ* bands. This bandgap leads to the insulating behavior of the monoclinic VO2. However, in the case of the high-temperature tetragonal structure the crystal symmetry leads to the overlap and hybridization of the vanadium atom 3d levels and the oxygen atom 2p levels. The VO6 crystal field leads to the lifting of the 3d orbital degenerescence forming suborbitals: a doubly degenerate but less stable eg suborbital and a more stable triply degenerate t2g suborbital.

2.2.3 Vanadium Pentoxide V2O5 is one of the most studied TMOs for its use as an ion storage layer in batteries [47] and as a complementary counterelectrode layer in electrochromic devices [48]. It is also known to exhibit a thermochromic switching between semiconductor and metal state at 257 C [49]. The special layered structure of these materials is very conducive to ion intercalation in the open structures between the sheets leading to facile ion insertion and extraction as well as their transport. The high charge density capacity of V2O5 makes it an

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ideal candidate for ion storage in batteries. The ensuing optical changes occurring with this intercalation are highly beneficial for their use as ion storage layers with other TMOs such as WO3. As discussed in later chapters, unlike the cathodically coloring WO3, this higher oxide of vanadium exhibits an anodic coloration in the visible region of the spectrum, i.e., a discoloration in the visible region of the spectrum with the double insertion of ions and electrons. Hence, it is commonly used for ion storage in conjunction with WO3-based electrochromic devices. V2O5 has a d0 electronic configuration of the transition metal. Similar to the other TMOs discussed earlier, V2O5 exhibits an open structure due to the corner-sharing oxygen units. But unlike the octahedral VO6 units formed by the hexacoordinated oxygen surrounding the metal atoms, the building blocks in the case of V2O5 are the VO5 units.The vanadium atom is surrounded by five nearby oxygen atoms forming a pyramidal unit and another long-distance VeO bonding in the c direction as shown in Fig. 2.23A. It is this longer and weaker VeO bond that leads to the layered structure of V2O5 as shown in Fig. 2.23B. As seen here the planar layers are formed by the corner-sharing oxygen atoms as well as the edge sharing of the square pyramids. This leads to an arrangement of the square pyramids in an alternating upward and downward fashion to accommodate the two types of sharing. This structure also leads to a gap after every third row as shown in Fig. 2.23B and C. V2O5 is also known to form several polymorphs [50], namely, a-V2O5, d-V2O5, and g-V2O5, which differ in how the VO5 pyramids are arranged by various combinations of the tilts and edge and corner sharing. They all present an orthorhombic structure as shown in Fig. 2.23B. The lattice parameters along the three directions of the unit cell are a ¼ 11.5 Å, b ¼ 3.6 Å, and c ¼ 4.4 Å, approximately. The open and layered structure makes it possible to easily intercalate ions such as Liþ leading to the electrochromic coloration.

2.2.4 Molybdenum Trioxide MoO3 is another TMO exhibiting both efficient electrochromic and photochromic coloration and a d0 electronic configuration of the transition metal. These virtues make it a very attractive material for various applications such as gas sensors [47], electrochromic [48] and photochromic devices [49], catalysts [50], and lithium storage batteries [51]. Similar to most other TMOs discussed earlier, the structure is formed of a hexacoordinated Mo atom by six O atoms forming the octahedral units [52] as shown in Fig. 2.24A. MoO3 is known to

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(A)

V

(B)

(C)

Figure 2.23 (A) VO5 square pyramid formed by the pentacoordinated vanadium atom by four oxygen atoms forming the base of the pyramid and a vanadyl bond forming the apex; (B) cubic structure formed by the upright and inverted VO5 square pyramids; (C) layered structure of V2O5.

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(A) Oxygen Molybdenum

(B)

a

b

c

Figure 2.24 (A) MoO3 octahedral unit formed by the hexacoordinated molybdenum atom by six oxygen atoms. (B) Crystal structure of the three polymorphs of MoO3, a-MoO3 (a), b-MoO3 (b), and h-MoO3 (c).

exhibit three polymorphic structures, namely, a-MoO3 showing an orthorhombic structure, b-MoO3 showing a monoclinic structure, and a hexagonal h-MoO3 structure [53e55]. These structures differ from one another in the way the basic octahedral units bond together. The crystal structures of the three polymorphs are as shown in Fig. 2.24B. In the case of the orthorhombic a-MoO3 the structure is made up of a double chain of edge-sharing MoO6 octahedral units as shown in Fig. 2.24B (a) along the ab plane. These double chains form a zigzag structure as viewed along the c axis. The double chains are held by strong covalent forces along the a and c directions through the edge

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sharing of the octahedral units. Along the b direction the double chains are formed by weak van der Waals forces, thus giving a-MoO3 a layered structure. As seen in Fig. 2.24B (b), the crystal structure of the metastable b-MoO3 showing a monoclinic unit cell closely resembles the ReO3 structure similar to that of WO3 discussed earlier. The MoO6 octahedral units are held together by the corner-sharing oxygen atoms in the c direction and by the edge-sharing octahedral units in the a direction. The other metastable form of MoO3, h-MoO3 showing the hexagonal units cell, is formed by a zigzag network of MoO6 octahedral units, which are formed by the corner-sharing oxygen atoms in the c direction [Fig. 2.24B (c)]. This kind of connectivity between the adjacent corner oxygen atoms linking the MoO6 octahedral units leads to a hexagonal symmetry and tunnels with open space (w1.5 Å) into which alkali metal ions can be inserted. This open structure makes the h-MoO3 a sought-after material for lithium storage batteries with a high charge capacity [56].

2.3 INTRODUCTION TO THIN FILMS AND NANOSTRUCTURE In view of the rich physics and chemistry exhibited by the TMOs and their dynamic behavior, the preparation of these materials in thin film or lowdimension form would enhance even more the scope of their study and exploitation of their dynamic properties for applications. The preparation of TMOs and their devices in thin film form would also facilitate their study under the application of various external stimuli. The field of thin films, devoted to the study of solid materials with one of their dimensions reduced to a few nanometers to a few hundred nanometers, has evolved into an important and independent branch of materials science and solid-state physics in the past few decades. This is due to the fact that the thin film state is drastically different from the bulk state of materials, presenting extremely interesting and controllable properties. Thin films form a part of the so-called advanced materials obtained by a careful and innovative modification of hitherto well-known materials. Advanced materials have played a pivotal role in the technological evolution throughout the history of humanity. The earliest examples of this are the alloys in which two or more metals with wellknown but fixed properties were mixed in various proportions to artificially generate a new range of properties that did not exist earlier. Similarly, throughout history humankind has invented a great number of advanced materials to address the demands created by technological advancement. Thin

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films are one such and fairly recent approach to creating advanced materials by reducing the dimensions of already well-known materials to the nanometer scale. Thin films have become an even more important aspect in modern-day society because of their high effectiveness, providing the required properties in extremely minute quantities. Hence, they have been touted to be the “green” solution to the increasing demand on natural resources by the exploding population and the increasing demand for human comfort worldwide through technology [57]. Though the knowledge and applications of thin films date back to at least a couple of millennia, with such applications as gold and silver foils [58], it was not until recently that a serious scientific study of the various aspects of thin films was undertaken and their industrial application potential unleashed. This has spawned important developments in the fields of thin film optics, microelectronics, solar cells, semiconductors, thin film sensors, and much more. The thin film technology started flourishing in the middle of the 20th century with the advent of high-speed vacuum technology. High-quality thin films could be grown in a controlled and reproducible way by physical vapor deposition (PVD) methods in chambers with very low pressures. The simplest of the PVD methods is thermal evaporation, in which the bulk material to be deposited is evaporated or sublimated in a vacuum chamber and the vapor is condensed on a suitable substrate in the form of a thin film. Nowadays a wide variety of physical and chemical methods are available to deposit thin films. Among the other physical methods are e-beam evaporation, pulsed laser deposition (PLD), sputtering, and a large number of variations of these basic methods. Among the chemical methods are chemical vapor deposition (CVD), spin coating or dip coating from a precursor solution, electrodeposition, electroless deposition, and the variations of these methods [59]. Most of these physical and chemical methods have been adopted with variation for use in research laboratories and/or industry to suit the particular need in terms of quality, production rate, costeffectiveness, upscalability, and more. Generally, when we speak of the various properties of materials, we allude to the bulk properties of these materials as they present themselves in nature. This three-dimensional (3-D) state of solids is defined by the macroscopic extension of the material in space along the three directions (x, y, z). In Fig. 2.25A, for example, is shown a bulk cube of a material. The geometrically normalized properties in this bulk state are fixed as long as the three dimensions are macroscopic. However, upon the reduction of one of these dimensions to the order of an electron mean free path in metals (few tenths of nanometers)

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Figure 2.25 Schematic of transition from a bulk (3-D) material to a switchable nanostructured thin film indicating the scope of research and development in the field of thin films.

or a light wavelength (few hundred nanometers), drastic changes appear in the electrical, optical, and other properties of these materials because of the severe restrictions imposed on the movement of the electrons and/or the way light interacts at the boundaries of the reduced material. This is the thin film or the 2-D state, in which the third dimension is negligible compared to the other two. One such “normal” or so-called “coarse grain” thin film with reduced dimension along the z axis supported on a substrate is shown in Fig. 2.25B. In this “normal” film it is assumed that the internal film structure and density

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of the material remain the same as in the bulk and that only one of the dimensions is reduced. This is the so-called 1-D size effect. This thin film state of the material is defined by several important aspects that make it very different than the bulk. It has an enormous surface-to-volume ratio, as a small quantity of material is spread over a large surface. Hence, the thin film state is important for all surface-sensitive properties. The thin film state in reality, however, is also marred by low density, high porosity, high density of defects, and high stress. The thin film properties are extremely sensitive to the film preparation conditions. Almost all the parameters involved in the deposition of a thin film can change the properties in a significant way. It is a very daunting challenge in the field of thin films to control all these parameters precisely and to deposit thin films with highly reproducible properties. Nevertheless, if one succeeds in this control, one can almost induce the required property into these films. Thus, an enormous volume of research to this day in thin films work is dedicated to establishing the correlation between the deposition parameters used and the properties imparted to the resulting films. These parameters include the deposition methodology, film thickness, nature of the substrate, deposition rate, pressure during deposition, substrate temperature, and more. The initial thin film work carried out in the second part of the 20th century was concentrated on getting high-quality films with a density approaching that of the bulk material. A great amount of research and technological effort was invested in studying and applying the thickness effect until the virtues of porous and granular structures began to unravel. In addition to this 1-D size effect, the interior of the film can be made up of grains of nanometer dimensions, thus inducing a 3-D size effect. Depending on the size of these grains (few nanometers to a few hundred nanometers), very interesting optical and electrical property changes can be observed. The presence of such grains, which imparts a random nanostructure to the film, adds exponentially to the internal surface as there is an enormous boundary around each of the grains. This has important consequences for surface-related properties such as gas sensing, light interaction, and more. Further, the nanostructure thus imparted can also be periodic whereby the grains within the film can be periodically arranged. In addition to providing a large internal surface, such periodic structures lead to even more interesting optical, photonic, and electrical properties, especially in the light wavelength region [60,61]. A schematic of such a randomly and periodically nanostructured film is shown in Fig. 2.25C. In addition to these passive changes imposed by the size effect, if these periodically and randomly nanostructured thin films are made up of optically

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(A) I0

I0

R

R Skin depth (I = I0e-α d) A

Bulk (3-D) state

(B)

I0

I0

R

R RD

A

A

TD

T

T Thin film (2-D) state

Figure 2.26 Schematic of light interaction with (A) a bulk metal and (B) a continuous thin film (left) and a nanostructured or discontinuous thin film (right). A, absorbed light; I0, incident light; R, reflected light; T, transmitted light.

switchable chromogenic materials (Fig. 2.25D) as discussed earlier, one can induce dynamic and reversible optical changes. These changes can be brought about by the influence of various external forces such as electric field, heat, light, pressure, and more. In addition, depending on the scale of the nanostructure these dynamic and static changes can be brought about in various wavelength regions of the electromagnetic spectrum. This will enable the user to control the various optical parameters of interest in an interactive way in different wavelength regions, thus increasing tremendously the scope of research and technological applications of smart systems. Given in Fig. 2.26 is an illustration of the optical effects brought about by the nanostructuring of the thin films. Light incident (I0) on a block of a metal is mostly reflected (R) and the small fraction of the light that penetrates the metal is quickly absorbed (A) within a short distance with an exponential attenuation [62]. This distance is referred to as the penetration or skin depth (d), which, along with the absorption or attenuation coefficient (a) of the material, determines the exponential decay of the light intensity in the metal as shown in Fig. 2.26A. The light interaction with a polished bulk metal surface is, hence, mostly carried out in terms of the specular reflected (R) light. However, by working with a very thin film, the same metal can be rendered partially transparent with film thicknesses in the skin-depth range. As shown in Fig. 2.26B, left, at these film thicknesses or the 2-D state of the metal, the optical behavior of the film is richly characterized in terms of

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the specular reflectance (R), specular transmittance (T), and absorption (A) of the film, assuming smooth interfaces and only the 1-D size effect (continuous film). In this thin film form all three optical parameters can be significantly altered by varying the film thickness of the metal. Also, if the film is nanostructured with a significant surface roughness on the impinging light wavelength scale, in addition to the three specular optical parameters a significant diffuse reflectance (RD) and diffuse transmittance (TD) can be expected as shown in Fig. 2.26B, right. Again a significant variation in all the optical parameters can be brought about by varying the nanostructure imparted. An example of the wide range of variation in the optical properties that can be achieved in copper thin films by varying the deposition parameters and subsequently the nanostructure is shown in Fig. 2.27. The results obtained in our laboratory have shown that a small variation in the film thickness from 20 to 2.5 nm, combined with variation in the substrate temperature, induces a large variation in the transmittance to cover nearly the entire optical range shown here. The thickest copper film of 20 nm deposited at room temperature cuts most of the light going through the film in the nearinfrared region and has a small transmittance in the visible region of the spectrum. With decreasing film thickness the transmittance rises quite steadily in both spectral regions. For a 2.5-nm-thick film of copper deposited at a substrate temperature of 300 C, the transmittance is nearly 90%. Hence, one can easily tailor the optical behavior as per need by choosing the right combination of deposition parameters. An enormous amount of research in thin films is, thus, dedicated to this aspect of establishing the correlation between the deposition parameters used and the ensuing film properties. The mapping out Copper thin film

Transmittance (%)

100

300ºC

2.5 nm

80 60

2.5 nm

40

10 nm

20 0

7.5 nm

5 nm

20 nm 300

850

1400

1950

2500

Wavelength (nm) Figure 2.27 Variation in the specular transmittance of a copper thin film with the deposition parameters.

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of the film properties versus deposition parameters is also a daunting challenge because of the very sensitive dependence of the properties on deposition parameters as well as the large number of parameters involved in the preparation of the thin films. However, if such a precise control of all parameters related to the preparation of the film is achieved one can easily induce the required properties into these films. Added to these factors and the aforementioned static properties of the films, if the nanostructured thin films are made up of optically switchable or chromogenic materials, all these optical parameters can be drastically and reversibly switched between their metastable states. In addition, as already pointed out earlier, this nanostructure induced in the films can be random or periodic, giving further provision for variation in photonic properties. Hence, the combination of thin film nanostructuring and applying it to chromogenic materials provides an enormous scope for research and development not only from the point of view of understanding these materials and their behavior but also for their important applications in various fields.

2.4 THIN FILM NANOSTRUCTURING TECHNIQUES A great number of physical and chemical techniques are available these days for the preparation of thin films. In most of these techniques the nanostructuring can be achieved by using a combination of appropriate deposition conditions. Among the PVD methods are thermal evaporation, e-beam evaporation, PLD, sputtering, and a large number of variations of these basic methods. Similarly, among the chemical methods are CVD, spin coating or dip coating from a precursor solution, electrodeposition, electroless deposition, and the variations of these methods. In the following section some of the nanostructuring techniques associated with these physical and chemical methods are elaborated.

2.4.1 High-Pressure Evaporation/Sublimation and Condensation The technique of thermal evaporation is the simplest of the PVD methods. The material to be deposited in thin film form is evaporated or sublimated normally under high-vacuum conditions in an actively pumped-down vacuum chamber. Generally, the high vacuum used for evaporation is compatible with the mean free path of the evaporated species (atoms, molecules, particles) that is necessary for their rectilinear trajectory between the source and the substrate without getting scattered. For this “line of sight” deposition the mean

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free path needs to be in the range of the deposition chamber dimensions. In the most common cases the parameters that control the film properties and structure are the film thickness, nature of the substrate, substrate temperature, chamber pressure, and rate of deposition. As seen in Fig. 2.27, each of these parameters can be adjusted to impart a nanostructure to the film. For example, if the evaporation is carried out at room temperature, the particle flux arriving with a high kinetic energy upon impinging the substrate surface loses its energy quickly to the surface and sticks to the surface, creating the nuclei for the further growth of a film. A large number of nucleation sites are formed throughout the substrate surface owing to the limited mobility of the species. This condition is conducive to the early formation of a continuous film. The nucleation process is quite probably a random process with isotropic distribution of the nucleation sites [63]. However, if the same evaporation/sublimation is carried out under the same conditions but at a higher substrate temperature, the particles arriving at the substrate surface will still have some thermal activation energy to move around the substrate to find a thermodynamically suitable site. This movement of the particles leads to relatively lesser nucleation sites and eventually to a discontinuous or granular film growth. Hence, at higher substrate temperature a continuous film begins to form only at higher film thicknesses. The rate of deposition, on the other hand, has the reverse effect of a discontinuous or granular film formation at lower deposition rates. This is due to the high flux of incoming species that at normal substrate temperature tend to form a large number of nucleation sites throughout the substrate surface and hence a continuous film early on [63]. Hence, a combination of a higher substrate temperature and a lower rate of deposition is thus expected to give a higher probability of a granular or a nanostructured film at lower thicknesses. This approach to thin film deposition under high-vacuum conditions relies predominantly on the thermodynamics between the impinging species on the substrate and the substrate conditions. It is these conditions that dictate the early stage nucleation and the subsequent growth of the film on the substrate. An efficient form of nanoparticle formation in large quantities is the one in which the nucleation is initiated in the gas phase, i.e., before the evaporated/sublimated species arrive at the substrate surface [64]. Such an approach is expected to give nanoparticles whose dimensions can be precisely controlled through the control of the deposition conditions. Although work on the fabrication of such ultrafine particles applied to noble metal films dates back to the early part of the 20th century [65,66], its evolution and application to various oxides and other compounds continue to this day [67e72].

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If the deposition chamber is maintained at a relatively higher pressure using an inert gas such as argon the mean free path of the evaporated/sublimated species is drastically reduced. According to the kinetic theory of gases, the relationship between the mean free path (l) and the chamber pressure (p) for a certain gas is given by: kT l ¼ pffiffiffi (2.6) 2ppD2 with a direct dependence on the absolute temperature (T) of the gas and an inverse dependence on the square of the diameter (D) of the gas molecules [73]. However, if two different types of molecules are involved, the mean free path is given by: 4kT l ¼ sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi   M1 1þ ppðD1 þ D2 Þ2 M2

(2.7)

where M1 and M2 are the molecular masses of the background (vacuum chamber) gas and the impinging evaporated/sublimated species, respectively, and D1 and D2 their corresponding diameters [74]. However, this relation cannot be applied directly to the case of evaporated/sublimated species because of the large variation and dependence of the dimensions (M2 and D2) on other deposition conditions. However, the order of mean free path in the vacuum chamber can be estimated in terms of the chamber gas at different pressures. Given in Table 2.2 are the mean free Table 2.2 Mean free path ( ) versus chamber pressure (p) for argon gas at room temperature Pressure, p (Pascal) Mean free path, (m)

1  102 1  101 1  100 1  101 1  102 1  103 1  104 1  105 1  106 1  107 1  108 1  109 1  1010

7.04  105 7.04  104 7.04  103 7.04  102 7.04  101 7.04  100 7.04  101 7.04  102 7.04  103 7.04  104 7.04  105 7.04  106 7.04  107

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path values for argon gas at room temperature in the pressure range commonly used in thin film deposition. When the mean free path is reduced to less than the dimensions of the deposition chamber by increasing the gas pressure with a neutral gas such as argon, the evaporated/sublimated species undergo an increased collision with the argon molecules and are deviated from their rectilinear path between the source and the substrate. These species undergo a high number of collisions very close to the source and quickly lose their energy. They are then carried around by the background carrier gas under a convective flow entering into further collisions before condensing on the substrate or any surface inside the vacuum chamber. In these situations there is an increased probability of not only the species colliding with the argon gas molecules but also the evaporated particles colliding with one another. These interparticle collisions are expected to lead to an in-flight coalescence growth and aggregation before the particles condense on a surface. The nanoparticles thus formed can be condensed on a surface to form a film with a substantial thickness. Such growth of fine metal particles has shown the dependence of the particle size on evaporation rates, carrier (inert) gas pressure, and sourceesubstrate distance [66]. Thus the degree of nanostructuring (grain size) can be expected to be controlled via these deposition parameters. The condensation on the substrate in the form of a nanostructured thin film can be rendered more efficient by cooling the substrate with a cold finger below the ambient temperature in the chamber. An efficient cooling of the substrate establishes a thermal gradient toward the substrate for the convectively drifting particles [75]. This approach has been applied to the formation of nanostructured thin films of electrochromic WO3 [71]. WO3 thin films deposited under high- and low-pressure conditions have been studied. The nanostructured films deposited by the high-pressure sublimation and condensation method exhibit very interesting optical and other properties compared to the films deposited under ultrahigh-vacuum conditions. A better nanostructure variation and control has been reported [72] using the same technique but with a slight variation. In this second technique the sourceesubstrate distance was varied systematically in the presence of relatively high pressure (102 Pa) argon gas. Similar to the earlier technique, the sublimated particles undergo a high collision rate with the argon gas molecules and between themselves before arriving at the cooled substrate. At shorter distances the number of collisions is relatively less, providing for a smaller grain growth. With increasing sourceesubstrate distance and with a higher number of collisions one can

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Figure 2.28 Variation of the roughness (Ra) and the grain size (Gs) with sourcee substrate distance in tungsten trioxide (WO3) thin films.

expect films made up of bigger grains on the substrate. Experimental results show a systematic increase in grain size, surface roughness, and porosity of the films with increasing sourceesubstrate distance. Hence, using the technique of high-pressure sublimation and condensation, along with the chamber pressure and/or sourceesubstrate distance one can impart the required nanostructure to the films deposited. Examples of the results obtained in our laboratory are shown in Figs. 2.28 and 2.29. The variations

Figure 2.29 Variation of the roughness (Ra) and the grain size (Gs) with chamber pressure in tungsten trioxide thin films.

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Figure 2.30 Scanning electron microscopy image of a WO3 thin film reactively deposited on an Si/SiO2 substrate. (From A. Ponzoni, E. Comini, M. Ferroni, G. Sberveglieri, Ultrasensitive and highly selective gas sensors using three-dimensional tungsten oxide nanowire networks, Thin Solid Films 490 (2005) 81e85.)

in grain size and surface roughness with increasing sourceesubstrate distance are shown in Fig. 2.28. Variations in the same parameters with increasing chamber pressure are shown in Fig. 2.29. Both deposition conditions are quite effective in terms of nanostructure control. In this approach there is also the possibility of the background carrier gas being used to induce reactive deposition of the film. Ponzoni et al. have used this approach for the deposition of nanostructured WO3 thin films for gas sensing studies [76]. High-pressure sublimation of tungsten wire was carried out with oxygen gas to reactively form the oxide with a very rich nanostructure as shown in Fig. 2.30.

2.4.2 Oblique/Glancing Angle Deposition Another method that has become increasingly important and popular for the deposition of nanostructured thin films is the so-called glancing angle deposition (GLAD) [77,78]. This method originated from its predecessor, oblique angle deposition (OAD), used very extensively for a very long time [79e81]. The films produced by holding the substrates at an angle with respect to the incoming particle flux are known to have a columnar structure along the incidence direction of the flux. This is due to the shadowing effect on certain parts of the substrate by the nucleation sites already formed. As shown in Fig. 2.31A, the species arriving after the setting in of the nucleation and under low-mobility conditions stick to the surface of the nucleation sites and contribute to the rapid growth of inclined columns. Such columnar films are highly porous and are very interesting

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Figure 2.31 Schematics of the (A) oblique angle deposition and (B) glancing angle deposition methods.

anisotropic films with high birefringence. Generally, under similar conditions of film deposition, including the material deposited, rate of deposition, chamber pressure, substrate temperature, and background pressure, there is an empirical relation between the particle flux angle (a) and the column growth angle (b) given by [82]: tanðbÞwð1=2ÞtanðaÞ

(2.8)

This tangent rule is commonly applicable at smaller particle flux incidence angles (a < 50 ). A much improved relation for this dependence proposed by Tait et al. [83], which is applicable for the whole range of particle flux angles (a) under the aforementioned fixed film deposition conditions, is given by: b ¼ a sin½1  cosðaÞ=2

(2.9)

OAD has been applied in innumerable thin film depositions to induce a columnar nanostructure allowing the control of film porosity and consequentially many other properties of the films [84e88]. In some cases highly

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porous films with densities as low as 15% of the bulk material were obtained along with a zigzag nanostructure of the columns by a systematic manipulation of the substrate with respect to the particle flux [89]. To add a new dimension to this nanostructure control via OAD, Azzam proposed the rotation of the substrate while making the oblique deposition [90], as shown in Fig. 2.31B. His proposal was aimed at the realization of helicoidal bianisotropic media nanostructures by continuously rotating the substrate during the growth of the film through OAD. To overcome the average effect of cluster formation through the continuous rotation, especially at high speeds, it was suggested by A. Lakhtakia et al. [91] to make a stepwise rotation of the substrate to better control the pitch of the chiral nanostructures, providing an even better control of the size of the helical columns comparable to the visible light wavelengths and related applications. This method of inducing a controlled nanostructure by combining the oblique deposition and a versatile rotation of the substrate was put to practical use by Robbie et al. [92] for the fabrication of the first magnesium fluoride (MgF2) and calcium fluoride (CaF) thin films with helical columns with pitches comparable to the wavelengths of visible light. This versatile technique has come to be known popularly as GLAD and the thin films deposited by this method are referred to as being “nanosculpted.” In addition to the columnar structure control by the shadowing effect during the film growth and the column shape control by substrate rotation, other parameters related to film deposition also play a significant role. The morphology of the columns can be controlled by varying the obliqueness of the angle, the deposition rate, and the rotational speed [93,94]. Since this initial work, the GLAD method has been used in innumerable works to obtain a wide variety of nanosculpted films induced with very special optical, electrical, and other properties [95e100]. A few examples of the nanostructures obtained by the GLAD method are shown in Fig. 2.32. The images correspond to the progressive degree of nanosculpting achieved with the increased degree of substrate manipulation added to OAD. In Fig. 2.32A is shown the microstructure of a film deposited at an extreme oblique angle of 85 degrees without the manipulation of the substrate during the film deposition. The film deposited under such conditions is made up of columns highly inclined in the direction of the particle flux. However, if the substrate is manipulated with respect to the particle flux incidence during the film deposition a zigzag nanostructure can be obtained as shown in Fig. 2.32B. Here the substrate was rotated to change the particle flux incidence by 180 degrees back and forth while holding the

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(A)

(B)

(C)

Figure 2.32 Glancing angle deposition method with (A) oblique angle deposition at 85 degrees, combined with (B) 180 degrees back and forth manipulation of the substrate during film deposition and (C) constant rate rotation of the substrate during film deposition. (From K. Robbie, M.J. Brett, Sculptured thin films and glancing angle deposition: growth mechanics and applications, Journal of Vacuum Science and Technology A: Vacuum, Surfaces, and Films 15 (3) (1997) 1460e1465.)

substrate at a constant angle of 85 degrees with respect to the normal to the substrate. Combining the same oblique incidence with a constant rotation instead of two back and forth positions a helical structure such as the ones shown in Fig. 2.32C can be obtained. In addition to the dynamic manipulation of the substrate during film deposition, a myriad of other nanostructures including multilayer composite films can be created by controlling the various parameters related to the deposition itself, such as the rate of deposition and mobility of the adatoms. These examples are shown in Figs. 2.33 and 2.34. Because of its versatility the GLAD method is also proposed for the fabrication of 3-D photonic crystals with large photonic bandgaps (PBGs) [101]. An important aspect of the growth of photonic crystals is the periodicity of the grown films. In general, the GLAD method, which relies on the random formation of nucleation sites, self-shadowing, and subsequent

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(B)

(C)

(D)

(E)

(F)

61

Figure 2.33 Glancing angle deposition method applied to nanosculpt six different films: (A) MgF2, (B) MgF2, (C) Cr, (D) CaF2, (E) Ti, and (F) Pt. (From K. Robbie, M.J. Brett, Sculptured thin films and glancing angle deposition: growth mechanics and applications, Journal of Vacuum Science and Technology A: Vacuum, Surfaces, and Films 15 (3) (1997) 1460e1465.)

growth of the nanostructured film, does not ensure a periodicity of the structure. However, if a periodically patterned substrate is used as a template for this growth the conditions for initiation of nucleation and self-shadowing will be met and 3-D photonic crystals can be grown. Kennedy et al. [102] have clearly demonstrated this possibility by growing tetragonal square photonic crystals by using prepatterned silicon wafers as substrates.

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Figure 2.34 A multilayer composite film with varying nanostructure obtained by the glancing angle deposition method: from bottom to top, layer 1, TiO2 helix, righthanded, three turns; layer 2, SiO2 zigzag; layer 3, SiO2 vertical posts; and layer 4, SiO2 helix, left-handed, three turns. (From A.C. van Popta, J.C. Sit, M.J. Brett, Trends in Nanotechnology Conference (TNT2004), September 13e17, Segovia-Spain (Poster), 2004.)

Hence, the GLAD method is an extremely viable approach to depositing nonperiodic and periodic nanostructures. Its versatility also stems from the fact that almost any PVD method can be used to produce the particle influx toward the substrate. By precise manipulation of the substrate the thin film can be nanosculpted as per need. Hence, nanostructured thin films of almost any material can be prepared by using the GLAD method. The only limitation associated with the GLAD method is that it is not easily upscalable to large area depositions. However, for research and applications involving small areas this method is the best option for depositing thin films with periodic and nonperiodic nanostructures.

2.4.3 Other Methods The nanostructuring of thin films has also been achieved by various other physical and chemical methods. In view of the focus here on the nanostructuring of TMO thin films and their chromogenic properties, only the methods pertaining to this subject are discussed briefly.

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2.4.3.1 SoleGel Method In this wet-chemical method, whose name originates as a combined short form of solution and gel, a molecular precursor solution is used to obtain the gel through the ensuing chemical process. The precursor is generally formed of the pertinent transition metal alkoxide of the form M(OR)n or MO(OR)n in water or an organic solvent. M and R here represent, respectively, the metal and the alkyl group. Under suitable conditions this precursor undergoes hydrolysis and polycondensation reactions to form colloids suspended in a continuous polymer network, the gel [103]. Following the preparation of the gel, the required thin films can be obtained by either the dip-coating or the spin-coating method. In the former case, the substrate to be coated is dipped in the prepared gel, dried, and eventually thermally or chemically treated to remove the organic solvents, leaving behind a thin film of the metal oxide, as shown in Fig. 2.35A. In the spin-coating method a small quantity of the gel is dropped on the substrate and spun at a high speed to spread and dry the solution uniformly and efficiently on the substrate, forming a uniform film, as shown in Fig. 2.35B. This film is further thermally treated to remove the remaining organic solvent, leaving behind the required thin film. In both approaches the film thickness and nanostructure control can be achieved by the various parameters involved in the process. Structure-directing surfactants and block copolymers can also be used in the preparation of the precursor solution to control the nanostructure to induce various morphologies in the films [104e107]. In the dip-coating method the important parameters are the molar concentration of the solution, relative humidity, duration of immersion, and withdrawal speeds, as well as the postdeposition treatment [108]. These solegel-based chemical methods have the greatest advantage in terms of their simplicity and upscalability, and thus are used for industrial fabrication of thin films. In the spin-coating method the parameters of importance are molar concentration of the solution, volume of solution placed on the substrate, spinning speed, and duration, as well as the postdeposition treatment [109]. The spin-coating method, however, is more advantageous for research laboratories because of the material wastage involved. It is generally suitable for uniform application of thin films over small areas. The solegel route has been applied extensively for the preparation of various TMO thin films using both the spin-coating and the dip-coating method [110e112].

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Figure 2.35 Schematics of (A) the dip-coating method and (B) the spin-coating method of thin film deposition.

2.4.3.2 Chemical Vapor Deposition CVD is another important chemical process for producing thin films of high quality in which the precursor is in a volatile form possessing sufficient reactivity. These gaseous-phase precursors react and/or decompose to produce the required solid film on the substrate held in a reaction chamber. Other gaseous by-products may also be formed from these reactions, which are removed by a continuous carrier gas flow in the system. The carrier gas occupies the main volume of the reaction chamber, similar to the organic solvent in the solegel method. The CVD process is divided into two types, low-pressure CVD (LPCVD) and atmospheric-pressure CVD. The process itself and the reactivity of the precursor compounds in the LPCVD process

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can be further enhanced by the presence of a plasma or laser or by high temperatures in the range of 200e2000 C [113]. Nanostructured WO3 and MoO3 thin films have been deposited by various forms of CVD to enhance their electrochromic and/or ion storage capacity [114e117].

2.5 SUMMARY Among the inorganic compounds, TMOs are an extremely interesting class of materials owing to the special orbital structure of the base transition metals. The partially filled d orbitals of these metals lead to a wide range of bonding with oxygen, forming TMOs with a large range of properties from ionic to metallic. The various chemical bonding structures of TMOs also lead to a range of physical structures and mixed frameworks such as hexagonal, octahedral, and tetrahedral. The wide range of possibilities and the multiple oxidation states make the TMOs highly attractive from a chromogenic properties point of view, in that a dynamic and reversible change in optical and electrical properties can be induced through external forces such as light, heat, electric field, gas exposure, and more. Hence, the most important aspect of the TMOs that is relevant for chromogenic properties is the possibility of switching them between their multiple states and structures through external activation in various forms. This reversible switching brings about a distinct change in their optical, electrical, and magnetic properties. Because these changes are associated with both physical and chemical changes in the TMOs, these are fascinating materials both from an industrial application point of view and for scientific research to understand the complex and reversible changes occurring in these materials. The chemical and physical structure of the various TMOs and their relevance to chromogenic properties are established in this chapter. Some of the most well-known TMOs exhibiting chromogenic properties are discussed in more detail. WO3 is the most versatile and studied TMO for its electrochromic properties. VO2 has been studied exhaustively for its highly efficient thermochromic properties. MoO3 is another TMO well known for its photochromic properties. These naturally manifesting fascinating switching properties of the bulk TMOs can be made even more interesting and attractive by working with the low-dimension form of these materials. Prepared in thin film form, one gains control over the switching of many more optically relevant properties such as transmittance, reflectance, and absorptance, as well as the fundamental optical constants (refractive index). Such a versatile control does not exist in the bulk materials. In addition, the

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switching properties depend sensitively on the thin film preparation conditions, which in turn control the nanostructure of the TMO thin films. The nanostructure of these TMO films can, thus, be varied via a strict control of the various deposition parameters. Hence, the important work in the field of TMO-related chromogenics on the fundamental side is to understand, establish, and model the exact correlation that exists between the switching properties and the nanostructure imparted. On the applied side, this knowledge can be used to tailor the chromogenic switching properties of the TMOs as per need of the applications. A wide range of physical and chemical techniques of TMO thin film nanostructuring is available these days. In addition to the optical switching properties, one can envision the preparation of periodically nanostructured TMO thin films. This possibility would open up new vistas for TMO-based chromogenics by providing a precise and dynamic control over photonic properties such as PBG, which is relevant for manipulation of light propagation in these materials. Research work on such periodic structures, called photonic crystals, based on TMOs and their chromogenic properties is at present in its infancy. A wide range of challenges and opportunities exists in basic and applied research on TMO-based chromogenics, especially in thin film and lowdimension (nanostructure) form. On the theoretical side, though many models exist and are adequate to explain the basic switching behavior of TMOs in bulk form, their application to nanostructured TMOs and their chromogenic behavior becomes a bit askew. A comprehensive model that is applicable to all structures is yet to be established in each of the chromogenic cases. On the experimental side, nanostructure being extremely sensitively dependent on the physical and chemical deposition conditions, a major challenge is to fabricate nanostructured TMO thin films in a reproducible way. However, if one can precisely control all the relevant parameters and establish a precise correlation between thin film properties and nanostructure, plenty of opportunity exists to tailor the TMO thin films and their chromogenic properties. The work with the most potential in the near future with regard to TMOs seems to be in the area of fabrication of TMO-based photonic crystals and their devices. This work is very important from the fundamental point of view as well as for industrial application in the photonics sector. The wide range of TMOs and their chromogenic properties (photochromic, electrochromic, thermochromic, gasochromic, etc.) make for plenty of opportunities for basic and applied research.

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[50] Guojun Shi, et al., Thin-Film b-MoO3 supported on a-Fe2O3 as a shell-core catalyst for the selective oxidation of methanol to formaldehyde, ChemCatChem 4 (2012) 760e765. [51] M.C. Rao, K. Ravindranadh, A. Kasturi, M.S. Shekhawat, Structural stoichiometry and phase transitions of MoO3 thin films for solid state batteries, Research Journal of Recent Sciences 2 (4) (2013) 67e73. [52] N.A. Chernova, M. Roppolo, A.C. Dillon, M.S. Whittingham, Layered vanadium and molybdenum oxides: batteries and electrochromics, Journal of Materials Chemistry 19 (2009) 2526e2552. [53] Large area chromogenics: materials and devices for transmission control, in: C.M. Lampert, C.G. Granqvist (Eds.), SPIE Institute Series, vol. IS4, 1990. Bellingham. [54] A. Cavalleri, T. Dekorsy, H.H.W. Chong, J.C. Kieffer, R.W. Schoenlein, Evidence for a structurally-driven insulator-to-metal transition in VO2: a view from the ultrafast timescale, Physical Review B: Covering Condensed Matter and Materials Physics 70 (16) (2004) 161102(R). [55] D. Di Yao, et al., Electrodeposited a and b phase MoO3 films and investigation of their gasochromic properties, Crystal Growth & Design 12 (2012) 1865e1870. [56] J. Zhou, et al., Synthesis of hexagonal MoO3 nanorods and a study of their electrochemical performance as anode materials for lithium-ion batteries, Journal of Materials Chemistry A 3 (2015) 7463e7468. [57] G.B. Smith, C.G. Granqvist, Green technology: solutions for sustainability and energy in the built environment, CRC Press, Boca Raton, FL, U.S.A., 2010. [58] J.E. Greene, Applied Physics Reviews 2 (2014) 011101, 1e11. [59] M. Ohring, The Materials Science of Thin Films, Academic Press, 1992. [60] J.D. Joannopoulos, S.C. Johnson, N. Winn, R.D. Meade, Photonic Crystals: Molding the Flow of Light, second ed., Princeton University Press, 2008. [61] E. Yablonovitch, Photonic band-gap crystals, Journal of Physics: Condensed Matter 5 (16) (1993) 2443e2460. [62] E. Hecht, Optics, second ed., Addison Wesley, 1990, pp. 108e113 (Chapter 4). [63] L. Ecketova, Physics of Thin Films, second ed., Plenum Press, 1986, pp. 112e116 (Chapter 4). [64] C.G. Granqvist, L.B. Kish, W.H. Marlow (Eds.), Gas Phase Nanoparticle Synthesis, Kluwer, Dordrecht, The Netherlands, 2004. [65] A.H. Pfund, Inert gas evaporation, Physical Review 35 (1930) 1434. [66] H.C. Burger, P.H. van Cittart, Zeitschrift für Physik 66 (1930) 210e217. [67] C.G. Granqvist, R.A. Buhrman, Ultrafine metal particles, Journal of Applied Physics 47 (1976) 2200e2220. [68] D.M. Mattox, G.J. Kominiak, Deposition of semiconductor films with high solar absorptivity, Journal of Vacuum Science and Technology 12 (1975) 182. [69] S. Iwama, K. Hayakawa, T. Arizumi, Ultrafine powders of TiN and AIN produced by a reactive gas evaporation technique with electron beam heating, Journal of Crystal Growth 56 (2) (1982) 265e269. [70] Y. Kimura, C. Kaito, Production of refractoryemetal ultrafine particles by gas evaporation method and their surface oxide layer, Surface and Coatings Technology 202 (17) (2008) 4159e4162. [71] P.V. Ashrit, Dry lithiation study of nanocrystalline, polycrystalline and amorphous tungsten trioxide thin-films, Thin Solid Films 385 (2001) 81e88. [72] B. Abdel Samad, J. Thibodeau, P.V. Ashrit, Preparation of nanostructured tungsten trioxide thin films by high pressure sublimation and condensation, Applied Surface Science 350 (2015) 94e99. [73] www.pfeiffer-vacuum.com.

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[74] R.W. Carlson, in: G.L. Weissler, R.W. Carlson (Eds.), Methods of Experimental Physics (Vacuum Physics and Technology), vol. 14, Academic Press, 1979, pp. 6e9. [75] J.S. Mea, S. Gauvin, P.V. Ashrit, Design of a physical vapor transport cell for time controlled deposition of nucleation phase organic thin films, Review of Scientific Instruments 78 (2007), 043902/1e8. [76] A. Ponzoni, E. Comini, M. Ferroni, G. Sberveglieri, Ultrasensitive and highly selective gas sensors using three-dimensional tungsten oxide nanowire networks, Thin Solid Films 490 (2005) 81e85. [77] M.M. Hawkeye, M.T. Taschuk, M.J. Brett, Glancing Angle Deposition of Thin Films: Engineering the Nanoscale, Wiley, 2014. [78] M.M. Hawkeye, M.J. Brett, Glancing angle deposition: fabrication, properties, and applications of micro-and nanostructured thin films, Journal of Vacuum Science and Technology, A 25 (5) (2007) 1317e1335. [79] Tomoyoshi, Y. Taga, Thin film retardation plate by oblique deposition, Applied Optics 28 (1989) 2466e2482. [80] D.K. Pandya, A.C. Rastogi, K.L. Chopra, Obliquely deposited amorphous Ge films. I. Growth and structure, Journal of Applied Physics 46 (1975) 2966e2975. [81] K. Hara, M. Kamiya, T. Hashimoto, K. Okamoto, H. Fujiwara, Columnar structure of obliquely deposited iron films prepared at low substrate temperatures, Thin Solid Films 158 (1988) 239e244. [82] J.M. Nieuwenhuizen, H.B. Haunstra, Microfractography, Philips Technical Review 27 (1966) 87. [83] R.N. Tait, T. Smy, M.J. Brett, Modelling and characterization of columnar growth in evaporated films, Thin Solid Films 226 (1993) 196e201. [84] S.J. Bolchot, H. Wilman, Quantitative relations between the orientation axis tilt, the angle of incidence of the vapour and the surface form in vacuum-condensed films, Thin Solid Films 69 (2) (1980) 191e199. [85] N.G. Nakhodkin, A.F. Bardamid, A.I. Novoselskaya, Effects of the angle of deposition on short-range order in amorphous germanium, Thin Solid Films 112 (1984) 267e277. [86] R.T. Kivaisi, Optical properties of obliquely evaporated aluminium films, Thin Solid Films 97 (2) (1982) 153e163. [87] S.P. Svensson, T.G. Andersson, Film thickness distribution at oblique evaporation, Journal of Vacuum Science & Technology 20 (1982) 245e247. [88] N.G. Nakhodkin, A.I. Shaldervan, Effect of vapour incidence angles on profiles and properties of condensed films, Thin Solid Films 10 (1) (1972) 109e122. [89] K. Robbie, et al., Fabrication of thin films with highly porous microstructures, Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films 13 (3) (1995) 1032e1035. [90] R.M.A. Azzam, Chiral thin solid films: method of deposition and applications, Applied Physics Letters 61 (26) (1992) 3118e3120. [91] A. Lakhtakia, W.S. Weighlhofer, Proceedings of the Royal Society of London, A 448 (1995) 419e437. [92] K. Robbie, M.J. Brett, A. Lakhtakia, First thin film realization of a helicoidal bianisotropic medium, Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films 13 (6) (1995) 2991e2993. [93] K. Robbie, M.J. Brett, Sculptured thin films and glancing angle deposition: growth mechanics and applications, Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films 15 (3) (1997) 1460e1465. [94] R. Messier, et al., Engineered sculptured nematic thin films, Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films 15 (4) (1997) 2148e2152.

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[95] D.-X. Ye, et al., Manipulating the column tilt angles of nanocolumnar films by glancing-angle deposition, Nanotechnology 13 (5) (2002) 615e618. [96] K.D. Harris, K.L. Westra, M.J. Brett, Fabrication of perforated thin films with helical and chevron pore shapes, Electrochemical and Solid State Letters 4 (6) (2001) C39eC42. [97] M. Malac, R.F. Egerton, Observations of the microscopic growth mechanism of pillars and helices formed by glancing-angle thin-film deposition, Journal of Vacuum Science & Technology A: Vacuum, Surfaces, and Films 19 (1) (2001) 158e166. [98] T. Smy, et al., Three-dimensional simulation of film microstructure produced by glancing angle deposition, Journal of Vacuum Science & Technology, A 18 (5) (2000) 2507e2512. [99] Y.P. Zhao, D.-X. Ye, G.-C. Wnag, T.-M. Lu, Proceedings of the Society of PhotoOptical Instrumentation Engineers (SPIE) 5219 (2003) 59e73. [100] A.C. van Popta, J.C. Sit, M.J. Brett, Trends in Nanotechnology Conference (TNT2004), September 13e17, Segovia-Spain (Poster), (2004). [101] O. Toader, S. John, Square spiral photonic crystals: robust architecture for microfabrication of materials with large three-dimensional photonic band gaps, Physical Review, E 66 (1) (2002), 016610/1e18. [102] S.R. Kennedy, M.J. Brett, O. Toader, S. John, Fabrication of tetragonal square spiral photonic crystals, Nano Letters 2 (1) (2002) 59e62. [103] C.J. Brinker, G.W. Scherer, Sol-Gel Science: The Physics and Chemistry of Sol-Gel Processing, Academic Press, 1990. [104] C.T. Kresge, et al., Ordered mesoporous molecular sieves synthesized by a liquidcrystal template mechanism, Nature 359 (1992) 710e712. [105] M. Muthukumar, C.K. Ober, E.L. Thomas, Competing interactions and levels of ordering in self-organizing polymeric materials, Science 277 (5330) (1997) 1225e1232. [106] S. Badilescu, P.V. Ashrit, Study of sol-gel prepared nanostructured WO3 thin films composites for electrochromic applications, Solid State Ionics 158 (1e2) (2002) 187e197. [107] Y. Djaoued, et al., Low temperature sol-gel preparation of nanocrystalline TiO2 thin films, Journal of Sol-Gel Science & Technology 24 (3) (2002) 247e254. [108] J.P. Chatelon, C. Terrier, J.A. Roger, Influence of elaboration parameters on the properties of tin oxide films obtained by the sol-gel process, Journal of Sol-Gel Science and Technology 10 (1997) 55e65. [109] B.W. Shivaraj, H.N. Narasimha Murthy, M. Krishna, S.C. Sharma, Investigation of influence of spin coating parameters ont eh morphology of ZnO thin films by taguchi method, International Journal of Thin Films Science and Technology 2 (2) (2013) 143e154. [110] C.O. Avellaneda, K. Dahmouche, L.O.S. Bulhões, A. Pawlicka, Characterization of an all sol-gel electrochromic device WO3/ormolyte/CeO2-TiO2, Journal of Sol-Gel Science and Technology 19 (1) (2000) 447e451. [111] B. Orel, A. Surea, U.O. Krasovec, Acta Chimika Slovakia 45 (4) (1998) 487e506. [112] Y.X. Guo, et al., Facile preparation of vanadium oxide thin films on sapphire (0001) by sol-gel method, Journal of Sol-Gel Science and Technology 70 (1) (2014) 40e46. [113] https://www.diva-portal.org/smash/get/diva2:714031/FULLTEXT01.pdf. [114] Y.-S. Lin, T.-H. Tsai, S.-W. Tien, Atmospheric pressure plasma jet-synthesized electrochromic organomolybdenum oxide thin films for flexible electrochromic devices, Thin Solid Films 529 (2013) 248e252. [115] L. Meda, R.C. Breitkopt, T.E. Haas, R.U. Kirss, Investigation of electrochromic properties of nanocrystalline tungsten oxide thin film, Thin Solid Films 402 (1e2) (2002) 126e130.

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[116] C.M. White, et al., Flexible electrochromic devices based on crystalline WO3 nanostructures produced with hot-wire chemical vapour deposition, Thin Solid Films 517 (12) (2009) 3596e3599. [117] A.C. Dillon, et al., Metal oxide nano-particles for improved electrochromic and lithium-ion battery technologies, Thin Solid Films 516 (5) (2008) 794e797.

FURTHER READING [1] S.K. Deb, Opportunities and challenges in science and technology of WO3 for electrochromic and related applications, Solar Energy Materials and Solar Cells 92 (2008) 245e258. [2] J.B. Goodenough, The two components of the crystallographic transition in VO2, Journal of Solid State Chemistry 3 (4) (1971) 490e500.

CHAPTER 3

Electrochromic Thin Films and Devices 3.1 BASICS OF THE ELECTROCHROMIC MECHANISM Electrochromic (EC) materials, as discussed earlier, are those in which a reversible coloration can be induced via the application of a small electric field. As discussed in Chapter 1, various materials are known to exhibit the EC phenomenon emanating from a wide variety of physical and consequent optical changes occurring with the application of an external electric field. In the transition metal oxides (TMOs), which are of pertinence here, the EC coloration can be conveniently described as occurring through the double injection of electrons and ions into the material bringing about a reversible chemical change in the composition. The ensuing reaction with the application of the electric field can be represented by: Mx Oy þ zIþ þ ze 4Iz Mx Oy

(3.1)

where M represents the transition metal, O the oxygen, I the ion inserted, e the electron. With the application of the electric field, z number of ions (such as Hþ, Liþ, Naþ, Kþ, Agþ) and electrons are injected into the TMO to form a metastable compound. The significant difference in the optical properties between the initial TMO (MxOy) and the resulting compound (IzMxOy) determines the efficacy of the EC phenomenon. Depending on the direction of the coloration in the visible region of the spectrum, the TMOs are classified as either cathodic or anodic materials. In cathodic TMOs such as tungsten trioxide (WO3), molybdenum oxide (MoO3), titanium oxide (TiO2), etc., the direction of the optical change is from a fairly transparent state to a colored state as represented below for WO3 [1]: WO3 þ zIþ þ ze 4 Iz WO3 ðTransparentÞ

ðDeep blueÞ

(3.2)

In anodic TMOs such as vanadium pentoxide (V2O5), chromium oxide (Cr2O3), nickel oxides (NiOx), etc., on the other hand, the optical change

Transition Metal Oxide Thin FilmeBased Chromogenics and Devices ISBN 978-0-08-101747-0 http://dx.doi.org/10.1016/B978-0-08-101747-0.00003-9

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upon EC coloration is from an initially visibly colored state to a colorless state as shown below for V2O5: V2 O5 þ zIþ þ ze 4 Iz V2 O5 ðGreenish yellowÞ

ðTransparentÞ

(3.3)

Some of the important aspects related to the EC phenomenon are the following: 1. wavelength region in which the optical switching occurs 2. coloration efficiency (CE) of the given material at a certain wavelength of the spectrum 3. charge capacity of the TMO for the particular species inserted and extracted 4. mixed conductivity (ion and electron) of the TMO for the particular species inserted and extracted 5. speed of EC coloration of the TMO As discussed at length in Chapter 1, the optical properties and optical switching occurring through the chromogenic phenomenon can be enhanced and better controlled in various modes by working with the thin film form [one-dimensional (1-D) effect] of the TMOs. In addition, the nanostructuring (3-D effect) of the TMOs not only is expected to enhance even more the range of the optical properties measured and controlled, but also has a significant impact on the switching performance of the TMO thin films (speed and CE). This is due to the fact that the EC switching is stringently dependent on the electrical transport of the inserted/extracted species in the films as can be clearly seen from the aforementioned reversible reactions. The enormous internal surface, grain boundary scattering, and porosity of the films significantly alter both the optical and the electrical properties and as a consequence the EC switching properties of the TMO thin films. Thus all the aforementioned important aspects related to the EC phenomenon are dependent on the film’s nanostructure. Hence, an enormous impetus is currently placed on the nanostructuring of the various TMO thin films to improve their EC efficiency. One of the best examples to illustrate the importance of nanostructuring is the work on titanium dioxide (TiO2) thin films by Hagfeldt et al. [2]. According to this work, conventional TiO2 (anatase) films deposited as normal coarse-grain films hardly show any EC activity or coloration [3]. However, the nanostructured TiO2 films show an excellent EC efficiency. This can be attributed to the high porosity and the large internal surface available in

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the nanostructured films, which greatly facilitate the transport of the inserted/extracted species as well as providing a large internal surface for the EC interaction. Another similar example is the indium tin oxide (ITO) thin films. Generally, ITO is a very well-known transparent conductor used in various industrial applications. When optimized for high transparency and electronic conductivity, ITO films are not known for their EC activity. However, when prepared in nanostructured form these films are shown to exhibit very efficient EC coloration as well as a high-speed optical and electrical response [4]. The results on the performance of a number of nanostructured TMO thin films are discussed in the following section. A wide range of TMO thin film nanostructuring techniques are also available now. Some of these are mentioned in Chapter 2. The various techniques essential for the study of EC behavior of the different TMO thin films and their devices are described in the following sections.

3.1.1 Electrochromic Device Configuration As can be surmised from the EC reactions given earlier, for the double insertion of ions and electrons to occur in TMO thin films upon the application of an electric field, other components are essential. A complete EC device is nothing but an optically active battery in reverse. In a battery the voltage stored is supplied to power an external circuit, whereas voltage needs to be applied by an external source to an EC device to insert or extract ions into or out of the base EC layer of a TMO thin film. With this insertion into or extraction out of the TMO layer occurs the reversible reaction cited earlier accompanied by a change in coloration. For stable operation the EC device generally needs to be a five-layer structure as shown in Fig. 3.1. Shown in this figure are the essential components for the functionality of an EC device. The TMO thin film forms the base of the EC device around which all the other components are designed. The double insertion into this layer, shown in Fig. 3.1, as the base EC layer, brings about the EC switching of the device. The ions required for the EC mechanism are stored in the ion storage (IS) layer, also referred to as the counterelectrode. The EC and counterelectrode layers are isolated by an electronically insulating but ion-conducting layer. These three components are flanked by the transparent conducting (TC1, TC2) layers on the outside, essential for the application of the electric field to the system. Upon application of the electric field in the mode shown in Fig. 3.1, the ions (Iþ) stored in the IS layer are transported across the ion conductor (IC)

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

Transparent conductor (TC2) I+

I+

I+

Ion storage layer

I+

I+

I+

e-

e-

Ion conducting layer Base EC layer e-

e- e Transparent conductor (TC1) e

Figure 3.1 Schematic of the essential components of a thin film electrochromic (EC) device.

layer and inserted into the initially transparent EC layer. Electrons (e) are inserted at the same time into this EC layer from the adjacent TC1 layer. This double insertion brings about the coloration of the EC layer. When the polarity is reversed the inserted species are transported back to the initial layer and the EC layer returns to its initial state. Hence, the EC coloration mechanism hinges on both the optical and the electrical properties of the component layers. Following are the essential requirements for each of the layers of the EC device shown in Fig. 3.1. 1. EC layer: This layer is the basis of the EC operation of the system. The design of the EC device begins with this layer and has to take into account the optical and electrical traits required as well as the thickness of the film. On the optical side, it is desirable to have the maximum CE in the spectral region of interest, i.e., display the maximum optical change for minimum charge (Q) injection. Hence it is desirable to have a TMO film with high transmittance in its normal state that is capable of showing low transmittance in the colored state via high reflectance or absorption in the spectral region of interest. Because this EC coloration depends stringently on the amount of charge (ions and electrons) injection, it is desirable to have a TMO film with a high charge capacity to deliver the optical change required. On the electrical side, because the EC mechanism emanates from the double

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injection into this layer, the film is expected to have a high ionic and electronic conductivity. High-efficiency cathodically coloring TMOs such as WO3 and MoO3 are very suitable for this layer. Depending on their structure, these films are capable of meeting the aforementioned requirements in a very efficient way. An enormous amount of research and development work has been carried out in the past few decades on the use of WO3 films as EC layers. It is by far the most studied TMO for this use because of its optical and electrical merits. 2. Counterelectrode layer: The counterelectrode layer plays an almost equally important role, being the IS layer, supplying the amount of ions required for the coloration of the EC layer. This layer has to be designed to accommodate at least the minimum charge (ions) required for the desired degree of coloration of the EC layer. In terms of the optical properties, this layer is expected to have a high transmittance in the spectral region of interest in the normal (charged) state and little to no transmittance change in the discharged state. Because it is the IS layer, it is expected to have a high charge density capacity, which will suffice for an efficient coloration of the EC layer. Because both charge species (ions and electrons) are expected to flow in and out of this layer during the EC operation, this layer, similar to the EC layer, is expected to have a high mixed (electronic and ionic) conductivity. Anodically coloring TMOs such as V2O5 and Cr2O3 films are well suited to use as ion-storing counterelectrodes in conjunction with the cathodically coloring TMOs. In such devices the coloration can be complementary to both the EC and the counterelectrode layer coloring, one under ion insertion and the other under ion extraction, respectively, according to Eqs. (3.2) and (3.3). The combined optical effect can render the CE of the system very high. 3. IC layer: This is the layer that separates the EC and the counterelectrode layer. Only the ions are expected to be transported across this layer under the influence of the applied electric field. The optical requirement for this film is that it be highly transparent in the spectral region of interest in all the modes of operation of the system, i.e., optically invariant. On the electrical properties side, the stringent requirements are that the layer be highly electronically resistive to minimize leak currents and be a very good IC. The choice of materials for the ion-conducting layer depends on the type of ion (Hþ, Liþ, etc.) chosen for the EC operation. Well-known transparent lithium and protonic conductors used in batteries are generally good for EC devices [5].

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4. TC layers: The TC layers on either side of the EC stack are used to apply the electric field to the system to transport the charge species to the pertinent layers. The optical requirement for these layers is that they be highly transparent in the spectral region of interest and optically invariant under all modes of operation of the system. For a uniform application of the electrical field for the EC operation and uniform coloration/discoloration, these layers are expected to have a high electronic conductivity (low resistivity). Industrially well-known transparent conductors such as ITO, aluminum-doped zinc oxide, indium-doped cadmium oxide, etc., fulfill these optical and electric requirements very satisfactorily. All the layers need to be robust mechanically for thousands of cycles of EC coloration and discoloration (bleaching) [1].

3.2 CHARACTERIZATION OF ELECTROCHROMIC MATERIALS AND DEVICES The operation of an EC device is a complex phenomenon involving the back-and-forth transport of charge species between layers along with the ensuing optical changes. Hence, the design and operation of EC materials and their devices involve various precise measurements of the electrical, optical, and kinetic parameters.

3.2.1 Double-Insertion Methods As mentioned before the design of each EC device begins with the base EC layer for the test of its CE. Hence, the fundamental characterization relates to the insertion and extraction of the charges. The EC layer is studied as a single layer and its optical/electrical properties are examined under double insertion/extraction. Two of the most commonly used methods for the testing of TMO thin films under double insertion are the electrochemical method and the dry method. These methods are described next in detail. 3.2.1.1 Electrochemical Method The TMO under test is deposited on a substrate precoated with a TC layer as shown in Fig. 3.2A. This sample is then inserted into an electrochemical cell (Fig. 3.2B) containing the appropriate electrolyte solution as a reservoir of the ions required for the EC mechanism. The sample under study forms the working electrode, which is juxtaposed to a counterelectrode and a

Electrochromic Thin Films and Devices

(A)

Substrate

79

(B)

TC

TMO

Figure 3.2 (A) Schematic of a transition metal oxide (TMO) thin film deposited on a substrate coated with a transparent conducting layer (TC). (B) Schematic of a threeelectrode electrochemical cell used for the electrochromic characterization of a TMO film. C, counterelectrode; I, current; R, reference electrode; U, constant potential; W, working electrode.

reference (R) electrode. Counterelectrodes such as platinum (Pt), lithium (Li), or carbon (C) and R electrodes such as calomel, lithium, or Ag/AgCl are most commonly used. For lithium ion (Liþ) insertion, for example, lithium perchlorate (LiClO4) or LiAsF6 salts dissolved in propylene carbonate (PC) are used as the electrolyte with the solution containing Liþ and  ClO 4 /AsF6 as the source of ions [1,6]. Under the application of an electric field with a negative polarity on the working electrode the Liþ ions migrate toward this electrode and are inserted into the TMO thin film to bring  about the EC coloration. The ClO 4 /AsF6 ions migrate toward the counterelectrode and are deposited over the surface. Upon reversal of the polarity the ion species are extracted from the respective electrodes and returned to the solution bringing about the discoloration of the TMO film. For the insertion of other types of ions, such as protons (Hþ) or sodium, for example, electrolytes such as sulfuric acid (H2SO4) or sodium chloride (NaCl) in methanesulfonyl chloride/aluminum chloride are used, respectively [7,8]. The use of the electrochemical method of insertion of ions also permits an in-depth study of the electrochemical behavior of the TMO thin film of interest. A wide variety of tests can be carried out to determine the suitability of TMO thin films for their EC behavior. The most common

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

Chronoamperometry

Cyclic voltammetry

V

Chro

V

I

t 0

0 I

0

t

t

V

I +Q

Q -Q

0

t

0

V

0

t

Figure 3.3 Schematic of the various time-dependence techniques used for the electrochemical kinetics of a transition metal oxide thin film.

tests are chronoamperometry (CA), cyclic voltammetry (CV), and chronopotentiometry (CP) (also called the galvanostatic measurements) [6,9e11], as shown in Fig. 3.3. In CA, shown in Fig. 3.3, left, a constant potential is applied between the working (W) electrode and the R electrode, and the current related to the insertion of ions is measured between the W electrode and the counterelectrode. From the time (t) variation in the current (I) curve, one can deduce the charge (Q) inserted or extracted into the TMO thin film. In the CV technique shown in Fig. 3.3, middle, the potential is ramped up and down between two extreme values of V as shown, at different scan rates (mV/s). Along with this cyclic variation in the potential, the current related to the electrode oxidation/reaction with insertion and extraction of the ions can be followed on the positive and negative arms of the potential. The charge (Q) inserted and extracted can be deduced once again from the I dependence on t. The scan rate appropriate for closed loop of the IeV curve has to be chosen, in which case the amount of charge inserted will be nearly equal to the charge extracted. Suboptimal scan rates lead to open IeV cycles, indicating an imbalance between charge inserted and charge extracted, which leads to irreversible kinetics and charge accumulation in

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the TMO electrodes. In galvanostatic (chronopotentiometric) measurements the small current applied to the TMO electrode is controlled and the time variation in the voltage is measured. The TMO film under study can be tested over a large number of CA, CV, and CP cycles to test the durability of the sample for ion insertion and extraction. A powerful technique for the transport measurements of the TMO thin films is impedance spectroscopy [9]. With this technique, treated in detail by many authors [6,9,10], many parameters related to the transport of the inserted/extracted species can be discerned, such as ionic diffusion and ionic and electronic conductivity. The experimental measurements of the electrical behavior of the TMO under the application of an alternating current signal over a wide range of frequencies are fitted to an equivalent circuit to find the different electrical component behaviors within the sample. 3.2.1.2 Dry Lithiation Method This is a lesser known and used method compared to the electrochemical method discussed earlier. The electrochemical method is used both for the characterization and for the insertion of the desired species. The TMO under study is necessarily brought into contact with liquid electrolytes. For further characterization of this sample or for the fabrication of EC devices, the sample surface has to be cleaned in an appropriate manner to get rid of the electrolyte. This cleaning can be cumbersome and may entail some undesirable effects on the film. For such cases the method of dry lithiation can be very advantageous. In the dry lithiation method the TMO film under study is exposed to lithium atom vapors under vacuum. It has been found that the EC effect, similar to that of the separate insertion of ions and electrons [Reaction (3.1)], can be induced in the TMO of interest. This reaction can be represented as: Mx Oy þ zLio 4Liz Mx Oy o

(3.4)

where Li represents a lithium atom. The coloration with this dry method of insertion proceeds similar to that of the electrochemical method [12]. The source of lithium atoms in the dry method can be either from lithium compounds such as lithium niobate (LiNbO3), lithium nitride (Li3N), or lithium iodide (LiI) or from the lithium foils [12e14]. Lithium being a highly reactive material, especially in the presence of humidity and air, it is essential to take precautionary measures in the handling of this material and to carry out the insertion procedure under vacuum. In terms of the reactive nature of lithium

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

and the aforementioned lithium compounds, LiNbO3 has a distinct advantage in terms of the handling of the material in air before the deposition procedure. Lithium foils, for example, can be e-beam evaporated under vacuum to give off lithium atoms, which diffuse and intercalate into the TMO film being tested to induce EC coloration. However, they can pose severe limitations in terms of handling and storage. Li3N and LiI are highly hygroscopic and can be hazardous, causing skin, lung, and eye irritation if not handled properly. LiNbO3, on the other hand, is a very well-known material used in a wide range of applications, including military ones, because of its very effective electrooptical and nonlinear optical properties. It is an extremely stable material with no known hazardous effects [15]. Lithium niobate (LiNbO3) powder can be thermally evaporated under controlled conditions to give off lithium atoms. When heated to around 815 C under high vacuum it decomposes and liberates lithium atoms. These lithium atoms arriving at the exposed surface of the TMO film of interest diffuse and intercalate into it, bringing about a very efficient coloration similar to the electrochemical method. As in any thin film deposition, the amount of lithium inserted can be measured by placing a quartz crystal thickness monitor close to the sample. To know the charge (Q) inserted into the film the effective mass thickness thus measured can be calibrated using the electrochemical method. For purely calibration purposes, the TMO thin film under study can be deposited on a transparent conductor similar to the one shown in Fig. 3.2A. This TMO film on the TC substrate can then be intercalated with lithium atoms in a vacuum system to different effective mass thicknesses. The Li-intercalated film can then be inserted into an electrochemical cell (Fig. 3.2B) and with the application of a small electric field the inserted lithium can be extracted. The amount of charge (Q) extracted can be calculated from the extraction current (I) dependence on time (t). Hence, for each type of TMO film a calibration curve can be established between the effective mass thickness deposited and the charge extracted. One such calibration curve is shown in Fig. 3.4. As can be seen from this example, a fairly linear dependence is seen in the region studied, giving a charge (Q) of 6.25 mC/cm2 for every 10 nm of effective mass thickness intercalated. The optical density change (DOD) associated with this insertion is shown in Fig. 3.4B. However, precautions should be taken not to overcharge the system, which could drive the TMO into irreversible kinetics with the formation of higher oxides in the system [16]. There are several distinct advantages of the dry lithiation approach. In the EC study of a TMO of interest, there is no need to bring the film into

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Figure 3.4 (A) Calibration curve of the dry lithiation method; Li effective mass thickness versus charge extracted in an electrochemical method. (B) Optical density change (DOD) as a function of the Li mass thickness. (From P.V. Ashrit, G. Bader, F.E. Giroouard, V.-V. Truong, Proceedings of the SPIE 1401 (1990) 119e129.)

contact with a liquid for the ion insertion. There is also no need for a transparent conductor on which to deposit the TMO. The film can be deposited on any substrate such as glass, which greatly facilitates the electrical and optical characterization of the individual films [17]. The greatest

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

advantage is in the fabrication of EC devices, which becomes cumbersome if contact with a liquid is involved. However, the main disadvantage of this method is that the inserted ions cannot be taken out for kinetic studies. Hence, the method can be conveniently used for the CE study of TMO thin films and for device fabrication.

3.2.2 Optical Characterization One of the most important applications of the optical switching of TMOs is in the smart windows [18] proposed for energy management in buildings, automobiles, and airplanes. Windows have become an important and integral part of modern architecture. Transparent windows also occupy a large surface area in automobiles. Hence, it is imperative to take into consideration this aspect in interior space energy management. Further, the light comfort level is also important. The temporal dynamic nature of the energy and/or visible light available on earth and the seasonal variations in climate call for dynamically adjustable windows as per the need of the season and the time of the day. Hence, TMOs are expected to play an important role in this regard and in energy management because of their dynamic optical behavior through the EC phenomenon. Relentless research and development work is hence focused on the development of TMO-based EC smart windows. Although EC-based smart windows have become a reality, even on a commercial scale [19], improvements in their switching efficiency (coloration and speed of switching), cost-effectiveness, and durability are being sought persistently by working with nanostructured TMOs to better tailor the EC optical switching. From the energy and visible light application point of view, the spectral region of most interest in smart windows is the solar spectrum, which spans from around 0.3 to 2 mm, comprising ultraviolet (0.3e0.4 mm), visible (0.4e0.7 mm), and near-infrared (0.7e2 mm) radiation. This and the other spectral parameters of interest related to smart windows are shown in Fig. 3.5 [1]. The blackbody radiation spectrum derived from Planck’s radiation law [20] is given in Fig. 3.5A: ul dl ¼ du ¼ 8phc

dl 
bhc l5 e l  1

(3.5)

It can be seen that as the temperature rises, the radiation emitted by the blackbody increases, and the peak of this radiation (lmax) moves to lower

Electrochromic Thin Films and Devices

85

(A)

(B)

(C)

(D)

Figure 3.5 Spectral parameters of importance in light and energy management. (A) Blackbody radiation spectrum around room temperature, (B) extraterrestrial solar radiation spectrum, (C) atmospheric absorption due to various gases at sea level, and (D) photopic efficiency of the human eye. (Adapted from Large area chromogenics: materials and devices for transmission control, in: C.M. Lampert, C.G. Granqvist (Eds.), SPIE Institute Series, vol. IS4, Bellingham, 1990.)

wavelengths. The extraterrestrial solar radiation reaching the top of the atmospheric envelope is given in Fig. 3.5B. This can be approximated to the radiation emitted by a blackbody at the temperature of 6000K. However, for the consideration of energy applications the optical properties of the atmospheric envelope should be considered. The atmospheric absorption bands are shown in Fig. 3.5C. It can be seen that a nearly transparent window exists between 0.3 and 2 mm, though marred by some strong

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

absorption bands in the higher wavelength region. This region lets most of the solar radiation reach the surface of the earth. Another important window in the atmosphere from an energy consideration point of view is in the region between 8 and 13 mm, also shown in Fig. 3.5C. Comparing this window with the spectral region of emittance of the blackbody shown in Fig. 3.5A, it can be seen that a significant part of the radiation emitted by a blackbody at around room temperature can escape to space, thus contributing to radiative cooling. Hence, the terrestrial temperature is a balance between the solar radiation gain occurring mostly between 0.3 and 2 mm and the radiative loss occurring between 8 and 13 mm. Finally, given in Fig. 3.5D is the photopic efficiency of the human eye, which is the visible region of the solar spectrum, with the most sensitivity around 0.55 mm. Hence, in the design of energy management and visible-light-related devices such as the smart windows, these considerations are extremely important. The optical characterization of the EC switching of TMO thin films is often carried out in terms of the integrated optical transmittance or reflectance of the film under consideration. The integrated optical parameters are derived from the area under the spectral curve before and after the optical switching to establish the optical efficiency of the material being tested. The main parameters of interest are the integrated solar transmittance (Tsolar) and integrated photopic transmittance (Tphotopic), which are calculated from the optical spectra of the film as follows: R 2 mm Tsolar ¼

T ðlÞ4ðlÞdl

0:3 mm R 2 mm 0:3 mm

4ðlÞdl

(3.6)

and R 0:7 mm Tphotopic ¼

T ðlÞjðlÞdl

0:4 mm R 0:7 mm 0:4 mm

jðlÞdl

(3.7)

where T(l) is the spectral transmittance of the TMO being tested in the respective regions of the spectrum, and 4(l) and j(l) are, respectively, the solar irradiance spectrum and the luminous efficiency of the photopic vision of the human eye. To calculate these optical parameters of the thin film of interest, the T(l) spectrum is multiplied by the solar j(l) spectrum available at different air mass (AM) values. Similarly, the T(l) spectrum is multiplied by the photopic j(l) to calculate the photopic performance. For these

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87

calculations related to solar spectrum, it is important to consider the sun’s elevation with respect to the earth’s surface, which changes with the time of the day. This is taken into consideration through the AM ratio between the solar radiation path length at noon (shortest path length) and the path length at an angle q. Hence, the AM ratio is expressed as 1/cosq. Generally these values are calculated for an AM value of 1.5. These integrated values of the TMOs are calculated before and after the EC switching to establish the CE of the films. To calculate the EC CE, the charge (Q) required to bring about a certain degree of coloration needs to be considered. Hence, by definition, CE is calculated at a certain wavelength as: CE ðlÞ ¼ DOD ðlÞ=Q (3.8) where the optical density change, DOD, is given at a wavelength as:   Tn ðlÞ DOD ðlÞ ¼ log (3.9) Tc ðlÞ where Tn and Tc are the normal-state and colored-state transmittances, respectively. The overall CEs of the TMOs of interest for energy applications can thus be calculated also from the integrated transmittance values in the normal (Tn) and colored (Tc) states in the spectral regions of interest. The same considerations can be applied to materials switching in reflectance mode by working with the integrated reflectance values. The EC optical switching can also be carried out at a more fundamental level by measuring the change in the optical constants, the refractive index (n) and the extinction coefficient (k), of the films before and after coloration [21]. Optical ellipsometry is a very powerful tool for such an in-depth characterization [22]. In most cases the optical constants derived can also be used in determining the porosity of the nanostructured films by optical means [23,24]. The film porosity (P) can be calculated from a simple equation involving the bulk refractive index of the material, nB; the thin film refractive index, nF; and the refractive index of air, nair: nF ¼ ð1  PÞnB þ Pnair

(3.10)

3.2.3 Other Characterization Techniques Other techniques generically applicable to thin films are used for the in-depth characterization of the EC properties of thin films. These include

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

X-ray photoelectron spectroscopy, X-ray diffraction (XRD), atomic force microscopy (AFM), micro-Raman studies, Fourier transform infrared, transmission electron microscopy and scanning electron microscopy (SEM), electrical resistivity, and temperature coefficient of resistance measurements. The EC phenomenon of the TMO thin films being a complex mechanism involving charge transport, the switching of optical and electrical properties, composition, and structural change, it is desirable to apply a large number of techniques to gain an in-depth understanding of the process.

3.3 ELECTROCHROMIC PROPERTIES OF NANOSTRUCTURED TRANSITION METAL OXIDE THIN FILMS In the following sections are given the results of the EC performance of the various TMO thin films. As emphasized earlier, the nanostructuring of these films has a profound effect on their EC performance because of the enormous internal surface created in these films and the porosity, which strongly influences the optical, electrical, and charge transport properties. Further, the nanoparticles that form the TMO film can be themselves amorphous or crystalline, further influencing these properties. An enormous amount of work has been carried out to elucidate the importance of nanostructured films. In this section the results relating to nonperiodically nanostructured TMO thin films are discussed. The discussion of the periodically (orderly) structured thin films or the so-called photonic crystals is deferred until Chapter 6.

3.3.1 Nanostructured Tungsten Trioxide Thin Films WO3 thin films are the most studied among the TMOs for their EC properties. A colossal amount of research and development work has been published in the past few decades following the report of the EC phenomenon by S.K. Deb in 1969 [25]. It has been well established that WO3 is the most versatile and suitable TMO exhibiting the EC phenomenon. The most important factors that have led to this immense study and use of WO3 in EC-related work are the following: 1. WO3 thin films exhibit EC behavior in both polycrystalline and amorphous forms, though the phenomena underlying the two types of coloration are totally different. 2. They exhibit an efficient coloration in the solar spectral region, making them the most sought-after materials for smart windows envisioned for light and energy management applications.

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3. WO3 films, depending on their micro- and nanostructure, exhibit a very high CE. 4. They possess a high ionic charge density that is adequate for an efficient EC switching. 5. They are known to exhibit a fast EC response depending on their micro- and nanostructure and phase. 6. WO3 films exhibit a very high physical and chemical stability. 7. WO3 thin films can be prepared by a wide range of techniques, which makes the tailoring of their nanostructure and EC performance very easy. The EC coloration can be conveniently represented by the reversible reaction shown earlier (Eq. 3.2) under the double insertion of ions and electrons. Assuming the widely accepted intervalence charge transfer [26], the EC coloration can be seen as follows. The ReO3-like network of the WO6 octahedral units shown in Fig. 2.14 makes for a large open space into which ions can be intercalated. Under the double insertion of ions and electrons in an amorphous WO3 film as already discussed in Chapter 2, the electrons are localized on a W6þ site and are transferred to neighboring W sites by absorption of a photon as shown in Fig. 3.6A. Hence, the EC coloration in amorphous WO3 films is via absorption modulation as shown in Fig. 3.6B, where the intensity of the transmitted light (IT) is predominantly altered through absorption of the EC-switched WO3 film. On the other hand, in the case of the crystalline WO3 film, owing to the long-range order and the overlapping wave functions of the neighboring W sites, the inserted electron is nearly free to move around in the crystal. Hence, the optical behavior of the WO3 film is predominantly due to these free electrons, similar to a metal. The externally inserted electrons contribute to the free electron plasma (up) as discussed earlier (Eq. 3.2). In this case the intensity of the transmitted light (IT) is predominantly dependent on the change in reflectance of the WO3 film as shown in Fig. 3.6C. There is a distinct difference in the optical behavior between the WO3 films in these two phases. In amorphous WO3 films the absorption modulation generally manifests in the lower wavelength region (visible and near infrared), while the reflectance modulation of crystalline WO3 films is at higher wavelengths (near infrared) for the density of lithium generally inserted [27]. Further, in the absorption-based coloration in amorphous WO3 films, with increasing double insertion the degree of coloration changes in a fixed wavelength region, i.e., coloration deepens around the visible and near-infrared region. In

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

(A)

(B)

(C)

I

I

ΔIR

ΔIA

ΔIT

ΔIT

Figure 3.6 (A) Schematic of the corner-sharing WO6 octahedral units, the interstitial location of the inserted ion ( ), and the intervalence transfer of the inserted electron between W sites by photon absorption. (B) Optical modulation with dominant change in film absorption (DIA). (C) Optical modulation with dominant change in film reflectance (DIR). DIT, change in intensity of transmitted light.

crystalline WO3 films, on the other hand, with increasing double insertion, the reflection front moves to lower wavelengths, i.e., from the near-infrared to the visible region. Hence, the latter type of films exhibit a wavelength selectivity with increased ioneelectron insertion. In addition to this effect of phase (crystalline vs. amorphous) of the WO3 film many other factors strongly influence the optical, EC, and electrical behavior of these films. As noted in Chapter 2 (Fig. 2.11), the WO3 structure can be formed of a range of networks of either cornersharing or edge-sharing WO6 octahedrals or can have tetragonal, hexagonal, or pyrochlore structures. The film substoichiometry plays an important role. Each of these factors can lead to a distinct combination of optical and electrical properties in normal and EC-switched states. Thermally evaporated WO3 thin films deposited under varying conditions, for example,

Electrochromic Thin Films and Devices

91

have shown distinctly different properties due to their changing nanostructure [28]. WO3 thin films deposited under high vacuum (5  106 torr) and low substrate temperature (200 C) were amorphous with a fairly smooth surface and exhibited a bluish tint due to slight nonstoichiometry. When inserted with lithium atoms by a dry method these films showed an intense blue coloration based in absorption modulation where the absorption peak was centered around the wavelength of 900 nm. The same film heated in air at high temperature (400 C) was transformed into a fairly stoichiometric film with a polycrystalline state exhibiting high transmittance and strong interference peaks. The film also exhibited a grainy surface and a conglomerate with a good intercrystalline contact. Such a polycrystalline WO3 film when intercalated with lithium atoms showed a very effective EC coloration over a wide range of the spectral regions studied. This broadband EC coloration was attributed to the mixed crystalline phase of the film leading to absorption modulation at lower wavelengths and reflectance modulation at higher wavelengths. Such a film seems to be made up of WO3 nanocrystals held in the amorphous WO3 matrix. In the same work [27], the WO3 films deposited at a low temperature (4e5 C) under a relatively high pressure (5  104 torr) and later heat treated in air at a high temperature (400 C) were nanostructured with large grains (400e500 nm). They exhibited a slight diffuse reflectance, appearing faintly cloudy. Such films exhibited a very high transmittance (90%) with a nearly total suppression of interference peaks. These films when intercalated with lithium atoms showed a very effective coloration with an absorption peak toward the higher wavelengths (w1300 nm). From the absence of any EC change in reflectance in the spectral region concerned one can surmise that the coloration is entirely based in absorption modulation stemming from the amorphous nature of the nanoparticles forming the film. The optical results pertaining to the three types of films studied and the CE in the spectral region studied are shown in Fig. 3.7. From these spectra it can be seen that the film nanostructure plays an extremely important role not only in inducing different optical properties in the film but also in the EC switching properties. From this work it is seen that polycrystalline film exhibits the most efficient and broadband EC coloration, while the nanostructured film exhibits high transmittance and a selective coloration. The integrated solar and photopic transmittance values calculated from these curves are given in Table 3.1. These results indicate the distinct possibility of tailoring the TMO films’, in general, and WO3 thin films’, in particular, optical and EC switching

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Figure 3.7 Transmittance spectra for WO3 amorphous (AM), polycrystalline (PC), and nanostructured (NS) film in the as-deposited state and in the lithium-inserted state (curves A, B, C, D, E, and F correspond to 0, 5, 7.5, 10, 15, and 20 nm of lithium, respectively) and their coloration efficiency at the insertion of 20 nm of lithium. (From P.V. Ashrit, Thin Solid Films 385 (2001) 81e88.)

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Table 3.1 Integrated solar and photopic transmittance values for various WO3 films Tsolar% Tsolar% Tphotopic% Tphotopic% Sample (normal) (colored) (normal) (colored)

Amorphous WO3 Polycrystalline WO3 Nanostructured WO3

69 75 87

23 22 36

72 75 91

22 30 59

Tphotopic, integrated photopic transmittance; Tsolar, integrated solar transmittance.

properties via their nanostructure control. The nanostructured WO3 films in the aforementioned example exhibit the highest transmittance and the coloration is more effective in the solar spectral region than in the visible. Such films are suitable for smart windows for which only the heat (infrared) modulation of the solar wavelengths is desired, as the films maintain a high degree of visible transmittance even in the colored state. Crystalline nanoparticles (nanorods and nanospheroids) of WO3 have been deposited by hot-wire chemical vapor deposition to enhance the CE, durability, and fast proton insertion kinetics [29]. A significant improvement in EC properties as well as cyclic durability under proton injection was demonstrated. Such films show a high potential for application in EC display devices. WOx films made up of ultrafine nanocrystals with dimensions around 6 nm have also been deposited by reactive gas deposition and studied by infrared spectroscopy under lithium insertion [30]. A scanning electron micrograph and the very efficient EC coloration induced under electrochemical insertion of lithium into this film are shown in Fig. 3.8. The electrodeposition method of thin film preparation is another interesting possibility, which, along with the virtues of cost-effectiveness and upscalability, gives a facile approach to the nanostructuring of TMO thin films [31,32]. Nanostructured or so-called mesoporous thin films of WO3 have been prepared by pulsed (5-ns duration) electrodeposition using peroxytungstic acid [33]. A good control of the WO3 film morphology with a nanoporous surface and extremely fine nanoparticles has been demonstrated by varying the pH value of the precursor. Coarse-grained WO3 films (referred to as “transparent”) electrodeposited with a high pH value (1.92) precursor and a continuous current were compared with mesoporous WO3 films electrodeposited with a lower pH value (0.8) solution and a pulsed current. A significant change in the morphology of the two types of films is shown in Fig. 3.9A. A significant improvement in the proton charge injection and extraction can be achieved by working with mesoporous films, as also seen through the CV cycles shown in

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Figure 3.8 (A) Scanning electron micrograph of a fine-grained nanocrystalline WO3 film deposited by reactive gas deposition. (B) Transmittance spectra of a WO3 film deposited on an indium tin oxide substrate in as-deposited, lithium-intercalated, and bleached states. (From J.L. Solis, A. Hoel, V. Lantto, C.G. Granqvist, Journal of Applied Physics 89 (2001) 2727e2732.)

Fig. 3.9B. Such mesoporous films are deemed important for EC performance and devices. The mesoporous WO3 films deposited under these conditions can yield films with crystals of a few nanometers in size. Electrodeposited WO3 films with large aggregates (micrometers) using the peroxytungstic acid precursor and a continuous current were also achieved with good results [34]. The effect of aging of the precursor and the formation of grape-like conglomerates of WO3 have been demonstrated. As shown in Fig. 3.10A, the time evolution of the WO3 crystals toward highly structured nanoaggregates with a large internal surface is quite evident. The high porosity and the enormous surface created are very

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(B)

(a)

95

(b)

(a) (b)

Figure 3.9 (A) Comparison of scanning electron micrographs of an electrodeposited mesoporous WO3 film (pH 0.8) (a) and a coarse-grained WO3 film (pH 1.92, “transparent”) (b). (B) Comparison of the cyclic voltammetry (CV) cycles of the two films shown in (A). (From B. Yang, H. Li, M. Blackford, V. Luca, Current Applied Physics 6 (2006) 436e439.)

advantageous for the intercalation and deintercalation of ions and the EC mechanism. An even more significant improvement in the insertion and extraction of lithium ions into these WO3 films deposited with an optimally aged solution can be seen through the CV cycles shown in Fig. 3.10B. A good EC CE is reported in devices based on these conglomerate WO3 films. However, such WO3 films made up of micrometer-sized conglomerates may lead to large light scattering from the film surface.

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(A) (a)

(b)

(c)

(d)

(B)

Figure 3.10 (A) Scanning electron micrographs of a WO3 film electrodeposited with an optimally aged precursor solution. (B) Comparison of the cyclic voltammetry cycles of two electrodeposited WO3 films showing the effect of precursor solution aging.

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Several works have elucidated the advantage of working with hydrated tungsten trioxide films even for lithium intercalation [35e37]. Unlike the anhydrous WO3 films made up of a network of edge- and/or cornersharing WO6 octahedral units, the oxide dihydrate (WO3e2H2O), for example, is expected to be made up of corner-sharing distorted WO5(H2O) octahedral units presenting a layered structure. The remaining water is located between these layers. The presence of this water is expected to facilitate the transport of inserted lithium ions in these nanosheets [38]. Such ultrathin nanosheets of tungsten oxide dihydrate (WO3e2H2O) have been deposited by a novel lamellar inorganiceorganic hybrid approach based on tungstic acid and exfoliation with excellent results for application in the fabrication of flexible EC devices [39]. A schematic of this procedure of the formation of the nanosheets is shown in Fig. 3.11A. Large-area tungsten oxide dihydrate sheets with a thickness of around 1.4 nm are fabricated. The films have been shown to be more flexible compared to bulk WO3e2H2O, also with a high lithium diffusion coefficient leading to very fast EC response and coloration contrast. A schematic of the nanosheet WO3e2H2O-based flexible EC device (ECD) is shown in Fig. 3.11B. The EC performance under lithium ion insertion of such a device is shown in comparison with its bulk counterpart-based device in Fig. 3.11C. Transmittance spectra of the device in colored and bleached states with the application of 3 V is shown in Fig. 3.11D. The CE recorded for the nanosheets ECD is threefold larger than that of the counterpart ECD. The solegel method of thin film deposition is also a very attractive one for obtaining mesoporous thin films with many advantages [40]. In addition to the distinct advantage of the upscalability and cost-effectiveness, this method offers the advantage of facile control of the composition and nanostructure of the thin films by doping or mixing other particles such as polystyrene beads in the precursor solutions [41]. Hence, both the physical and the chemical aspects of the resulting films can be controlled. Structure-directing agents and templatebased techniques are often used to achieve highly porous films with controlled structure by the solegel and electrochemical methods [42,43]. A particular example of this is the use of polyethylene glycol (PEG) as the structure-directing agent for the deposition of nanostructured WO3 thin films [42]. Mesoporous WO3 films were obtained from a colloidal tungstic acid stabilized by PEG. The nanostructure evolution was studied as

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Figure 3.11 (A) Schematic of the fabrication of nanosheets of WO3e2H2O by the lamellar inorganiceorganic hybrid method. (B) Comparison of the lithium intercalation kinetics between WO3e2H2O nanosheets and the bulk counterpart. (C) A flexible electrochromic device (ECD) based on WO3e2H2O nanosheets. (D) Transmittance spectra of the WO3e2H2O nanosheet-based ECD in bleached and colored states. ITO, indium tin oxide; PET, poly(ethylene terephthalate). (From L. Liang, et al., Scientific Reports 3 (2013) 1e8, Article ID: 1936.)

a function of annealing time. This colloidal solution approach was shown to yield WO3 films with a rich nanostructure both with and without the PEG additive. However, with the addition of this structure-directing agent, the nanostructure was remarkably better controlled, as can be seen from the

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(A)

(B) (a)

(b)

400°C-10 h

400°C-30 min

(c)

(d)

500°C-30 min

450°C-30 min

(e)

550°C-30 min

(f)

600°C-30 min

Figure 3.12 (A) Scanning electron microscopy (SEM) image of a WO3 film deposited from a tungstic acid colloidal solution without the structure-directing agent and annealed at 550 C. (B) SEM images of a WO3 film deposited from tungstic acid colloidal solution with the structure-directing agent (polyethylene glycol) and annealed at different temperatures and durations. (From C. Santato, M. Odziemkowski, M. Ulmann, J. Augustynski, Journal of American Chemical Society 123 (2001) 10639e10649.)

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comparison of the SEM images shown in Fig. 3.12A and B. Under different annealing conditions the WO3 nanoparticles forming the films are shown to undergo a systematic transformation from a needle-like structure to a plate-like structure, while their dimensions change from a few nanometers to around 100 nm. The film deposited without PEG, too, has an extremely rich nanostructure with particle dimensions of a few hundred nanometers. The EC activity of the highly transparent nanocrystalline WO3 films was verified by a large number (>10,000) of CV cycles under the insertion of protons and lithium ions. These films are reported to show a very efficient EC performance and stable operation with some samples showing the highest CE of 40 mC/cm2. A similar effect of annealing temperature on the nanocrystalline structure of WO3 thin films is also seen in films obtained by a surfactant-assisted electrodeposition method [43]. In electrodeposited WO3 thin films obtained from a solution of peroxytungstic acid containing PEG (400) and cetyl trimethyl ammonium bromide, the role of the surfactants and annealing is elucidated. Although a good CE is exhibited by the lowtemperature-annealed WO3 films, the CE drops sharply with increasing temperature. An electrodeposition method was also used by Dinh et al. [44] to deposit mixed nanostructured Ti/W oxides to enhance the durability and large area application of these EC films. WO3 nanoparticles from a peroxytungstic acid were electrodeposited into a porous film of TiO2. A change in the nanostructure of the mixed oxide film yielded a further fine-grained morphology as shown in Fig. 3.13A and B. A fairly high efficiency of lithium ion insertion/extraction, cyclic operation, and transmittance change are shown in Fig. 3.13CeE, respectively. A very good coloration/bleaching response time of the devices based on these mixed oxide films as well as a high CE is reported. Nanorod-like [45] and plate-like [46] WO3 films have also been deposited by a hydrothermal process without using surfactants or structure-directing agents. Although these films exhibit a high EC CE and cyclic durability the coloration/bleaching response is too slow. As discussed in Chapter 1, many powerful physical vapor deposition (PVD) techniques are used for the controlled nanostructuring of thin films. The most interesting and recent technique is the glancing angle deposition (GLAD) method whose virtues are described in Chapter 1. This method has been used for the nanostructuring of WO3 films for their EC properties [47,48]. A systematic variation of the substrate tilt angle with respect to the particle flux coming from the source along with the rotation of the substrate is expected to lead to a “nanosculpting” of the deposited film. The high

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(B)

(C)

(D)

(E)

Figure 3.13 (A) Scanning electron microscopy (SEM) image of the nanostructured TiO2 film compared with (B) an SEM image of the mixed WO3/TiO2 nanostructure. (C) Cyclic voltammetry cycles of the mixed oxide film. (D) Coloration/bleaching cycles of the mixed oxide film. (E) Transmittance change of the mixed oxide film with the application of various potentials between the standard calomel electrode (SCE) and the working electrode: 0.5 V (curve 1), 0.4 V (curve 2), 0.3 V (curve 3), 0.2 V (curve 4), 0.1 V (curve 5), 0 V (curve 6), 0.1 V (curve 7), 0.2 V (curve 8) and 0.3 V (curve 9). (From N.N. Dinh, D.H. Ninh, T.T. Thao, V.-V. Truong, Journal of Nanomaterials 2012 (2012) 7, Article ID: 781236.)

degree of controlled porosity is interesting from the point of view of tailoring of the optical properties as well as the ion transport-dependent EC properties. An example of this work is shown in Fig. 3.14. The evolution and growth of the nanostructure as a function of the substrate tilt angle and

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with film thickness at a particular angle (85 degrees) is shown in Fig. 3.14A. Hence, with the combination of these two parameters, the nanostructure can be tailored over a wide range. The porosity of the WO3 films studied in this work also varies over a wide range (11%e68%) as a function of the tilt angle. Hence, the effective mass thickness available for the lithium ion injection has to be considered for any EC mechanism analysis. In Fig. 3.14B are given the transmittance change, the CE, and the correlation between the CE and the effective mass thickness. Owing to the high porosity induced, the glancing angle-deposited WO3 thin films are highly transparent, with a high suppression of interference. These films exhibit a broadband high CE in the middle of the solar spectral region, with CE values reaching 60 mC/cm2 at low quantities of Li insertion. Such films are capable of showing a very effective switching between a very highly transparent normal state and a very highly colored state, especially if a very large equivalent mass thickness of WO3 film is employed, as can be seen in Fig. 3.14B (CE vs. equivalent mass thickness). Such effective EC coloration with a small amount of charge inserted is very beneficial for the stable operation and charge transport kinetics of the devices based on these films. In addition to this GLAD method in which substrate manipulation is the control parameter for the film nanostructure control, other PVD

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Figure 3.14 (A) Evolution of the atomic force microscope images of WO3 films deposited by the glancing angle deposition method as a function substrate tilt angle, (a) 0 degrees, (b) 40 degrees, (c) 70 degrees, (d) 80 degrees, and (e) 85 degrees, and as a function of film thickness, (f) 22 nm, (g) 56 nm, (h) 85 nm, (i) 197 nm, and (j) 445 nm (all images taken at 0.25  1 mm). (B) Transmittance spectra of two 60-degree-tilt glancing angle-deposited WO3 thin films of similar thickness (a, b) inserted with different quantities of lithium and coloration efficiency (CE) dependence on wavelength (c) and equivalent mass thickness of the film (d). (From G. Beydaghyan, G. Bader, P.V. Ashrit, Thin Solid Films 516 (2008) 1646e1650.)

methods have been used very effectively by manipulating the deposition conditions as variable parameters [28,49e51]. Two of the parameters of distinct importance are the chamber pressure and sourceesubstrate distance. Under high vacuum conditions, the mean free path of the evaporated

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particles or species is generally very large, being on the order of the vacuum chamber dimensions. Thus the evaporated particles emanating from the source travel in nearly rectilinear paths in all directions and condense on any surface that is at a relatively cooler temperature, including the substrate on which the film deposition is desired. This has been found to give smooth and normal, or the so-called “coarse-grained,” WO3 films [28]. However, if the chamber pressure is increased, the mean free path of the evaporated/ sublimated particles decreases, inducing an in-flight formation of the grains due to a higher probability of interparticle collision and coalescing. This factor can be used very effectively for the growth of nanostructured WO3 films, as already discussed earlier, showing very interesting EC properties [28]. A similar approach was taken by Baloukas et al. [49] for the fabrication of WO3 EC effect-based interference filters. Instead of the regular thin film interference stack made of two different high- and low-refractive-index materials [52], the same material was used for depositing porous and dense WO3 thin films to create interference or a Bragg mirror stack [53]. This approach of using deposition conditions to control the porosity (and the refractive index) of the films when applied to switchable TMOs such as WO3 and MoO3 is very advantageous. It is possible not only to create a photonic bandgap (PBG) in the spectral region of interest but also to spectrally tune the PBG. WO3 thin films have been radio-frequency sputter deposited under varying chamber pressures and their EC properties verified under proton insertion/extraction [49]. In Fig. 3.15A are shown the CV cycles of such films. As can be seen from this figure, the ion insertion/ extraction cycles depend very sensitively on the chamber pressure used and the nanostructure induced into the WO3 films. The lowest pressure (5 mtorr) sputtered films show very meager EC activity, while the best and optimum results are shown for a film deposited at 20 mtorr. The transmittance change with time in a 27-layer interference filter is shown in Fig. 3.15B and the coloration/bleaching cycles of the filter are shown in Fig. 3.15C. A very good coloration is seen in the visible region of the spectrum. The film deposited at 20 mtorr pressure showing the best EC activity has a CE of 48 cm2/C. In addition to the chamber pressure the sourceesubstrate distance is also expected to play an important role in inducing a controlled nanostructure in WO3 films. At a given chamber pressure and mean free path of the evaporated particles the time of flight of these particles also seems to be important. The longer path between the source and the substrate under the conditions of shorter mean free path provides for a higher number of

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Figure 3.15 (A) Cyclic voltammetry cycles of WO3 films sputter deposited at various pressures under proton injection/extraction. (B) Transmittance change under proton injection of a 27-layer porous/dense WO3 interference filter as a function of coloration time. (C) Coloration/bleaching cycles of the same filter at two different wavelengths. (From B. Baloukas, J.-M. Lamarre, L. Martinu, Solar Energy Materials and Solar Cells 95 (2011) 807e815.)

interparticle collisions and a higher probability of growth. Hence, the sourceesubstrate distance can be a control parameter in inducing a particular nanostructure in the films. WO3 thin films were deposited by thermal sublimation and condensation by varying the sourceesubstrate distance in an argon-backfilled chamber at a pressure of 3  104 torr and condensed on a substrate [24]. The rich nanostructure induced in the films as a function of the sourceesubstrate distance is given in Fig. 3.16A. The WO3 film porosity, surface roughness, and grain size varied within ranges of around 50%e60%, 2.5e5 nm, and 10e35 nm, respectively, for a change in

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(A)

(B)

(a)

(b)

(c)

(d)

(e)

(f)

Figure 3.16 (A) Atomic force microscopy images of 1  1 mm dimension of high-pressure sublimated and condensed WO3 films at various sourceesubstrate distances: (a) 13 cm, (b) 15 cm, (c) 16 cm, (d) 18 cm, (e) 20 cm, and (f) 21 cm. (B) Transmittance and coloration efficiency (CE) of the WO3 films as functions of wavelength. (From B. Abdel Samad, J. Thibbodeau, P.V. Ashrit, Applied Surface Science 350 (2015) 94e99.)

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sourceesubstrate distance of 13e21 cm. The films were tested for their EC behavior under dry lithium atom insertion. The transmittance changes and the associated coloration efficiencies are shown in Fig. 3.16B.These films exhibited an excellent EC coloration with the optimum samples exhibiting CEs in excess of 100 cm2/C. The samples deposited at the sourceesubstrate distances of 18 and 20 cm exhibited the most solar spectral change. Note: An important thing has to be emphasized regarding the utilization of the parameter CE in its current form. As stated in Eq. (3.8), CE is obtained from the optical density change (DOD) at a certain charge (Q) injection. Although it is a convenient method of characterizing the EC efficiency of a film, it can mislead at very small values of Q inserted. In the initial stages of coloration, for a very small insertion of charge there is a drastic change in the optical properties. This manifests as a high CE value due to the ratio between DOD and Q. However, this change in coloration may not be sufficient for practical purposes. For example, in the aforementioned works, a CE of 80 cm2/C is found in porous sputter-deposited WO3 films for the insertion of 0.7 mC of charge, while the change in transmittance is only 6.8% [49]. Similarly, very high values of CE on the order of 150 cm2/C are found in some nanostructured WO3 films with the insertion of very small amounts of lithium [24]. Although these CE values are correct from a calculation point of view, they cannot be used neither for comparison between different films nor for the practical design of ECDs based on these values. One has to take into consideration the overall DT or DOD induced by the inserted charge. This consideration has to be applied whether one is dealing with a chemical insertion method or a dry lithiation method or even protonic insertion or lithium insertion. Another important aspect in evaluating the EC performance of the WO3 and other TMO thin films is the mixed conductivity. As is evident from Eqs. (3.2) and (3.3), the films deposited need to have good ionic and electronic conductivities to be temporally efficient in the EC coloration and bleaching. Hence, in addition to the optical properties, one has to optimize the electrical properties as well as the charge insertion kinetics of the film. Attempts are being made relentlessly to improve the EC performance of TMO thin films by inducing a nanostructure in them. However, the simultaneous study of the optical and electrical properties along with the other parameters connected with transport of the inserted charge is not systematic and is rare. Increasing the nanostructure of TMO films may be favorable to ionic conductivity but will decrease the electronic conductivity because of the large grain boundary scattering imposed on electronic transport. Normal (dense or

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coarse grained) films, on the other hand, have a better electronic conductivity but a limited ionic conductivity due to the limited internal surface available for the ion transport. Hence, it is extremely important to optimize both the ionic and the electronic conductivity in these films by judiciously tailoring the nanostructure. Almost all the work related to EC studies of TMO thin films through their nanostructuring tends to concentrate heavily on improving the optical performance (CE) rather than the equally important electrical properties [54]. Even in some of the works in which the optical and electrochemical kinetics are carried out in detail the aspect of simultaneously optimizing the two conductivities is very patchy [55]. An attempt to achieve this optimization of mixed conductivity through nanostructuring has been made [56]. WO3 thin films have been deposited under varying chamber pressures to systematically change their nanostructure and their ionic and electronic conductivities as well as their lithium diffusion coefficient. The nanostructure induced in these WO3 films through this approach is shown in Fig. 3.17A. The WO3 films deposited under a high vacuum (1  105 torr) present a very smooth surface with closely packed tiny grains. As the chamber pressure increases, the film nanostructure evolves systematically with increasingly larger aggregates leading to large internal surface and porosity. The surface roughness and the grain size also increased with pressure. As expected, the diffusion coefficient of lithium ions also increased with increasing chamber pressure. The electrochemical activity in the WO3 films deposited at different pressures was verified through the closed-loop CV cycles for the insertion/extraction of lithium ions. The most important result of this work, the mixed conductivity, is shown in Fig. 3.17B. It is seen that with increasing nanostructure the ionic conductivity increases by 2 orders of magnitude, while the electronic conductivity decreases by the three orders of magnitude within the scope of the work carried out. This is a significant result as it indicates the possibility of optimizing the mixed conductivity of WO3 and other TMO thin films through their nanostructuring. The optimum mixed conductivity was found for nanostructured WO3 films deposited at a pressure of 5.7  104 torr. The EC performance of such a mixed conductivity optimized film is expected to be very efficient.

3.3.2 Nanostructured Vanadium Pentoxide Films V2O5 films are very well known for their anodic type of EC coloration, i.e., these initially yellowish films become transparent under the double

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Electrochromic Thin Films and Devices

(A)

(b)

(a)

(c)

(d)

(e)

(B)

1.40E-09

1.20E-06

Ionic Conduc vity (Ω-1cm-1)

1.20E-09

1.00E-06

1.00E-09

8.00E-07

8.00E-10 6.00E-07

ionic

6.00E-10

electronic

4.00E-07

4.00E-10

2.00E-07

2.00E-10

Electronic Conduc vity (Ω-1cm-1)

Ce = -1E-06P + 1E-09

Ci = 0.4494P2 + 0.0006P + 2E-08 0.00E+00 1.00E-05

0.00E+00 1.90E-04

3.70E-04

5.50E-04

7.30E-04

9.10E-04

Working pressure (torr)

Figure 3.17 (A) Atomic force microscopy images of a WO3 film deposited at (a) 1  105 torr, (b) 5  105 torr, (c) 1  104 torr, (d) 5  104 torr, and (e) 6  104 torr. (B) Ionic and electronic conductivities of WO3 films deposited at different pressures. (From B. Abdel Samad, P.V. Ashrit 2017 (to be published work).)

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insertion of ions and electrons [1]. The ensuing reversible reaction under the double insertion can be represented by Eq. (3.3). The cathodic EC coloration of V2O5 films is also verified to take place under the insertion of lithium atoms [57], similar to that of WO3 films, and is used for the fabrication of all-solid ECDs [58]. This reaction can be represented by: V2 O5 þ zLio 4 Liz V2 O5

(3.11)

The combination of a cathodically coloring WO3 film as the base W electrode and anodically coloring V2O5 film as the counterelectrode (for IS) is deemed very advantageous. In addition to acting as IS layers, the yellowish coloration of V2O5 films when ions are extracted out of it to color the WO3 is very complementary to the operation of an EC system based on these two layers. Both the layers color and discolor in the visible region of the spectrum, making the CE of the overall device very effective. In addition, as in other TMO thin films, the nanostructure plays an important role in the EC properties of these films. Hence, a lot of research work has been focused on the nanostructuring of V2O5 thin films to enhance their optical, electrical, and EC properties. V2O5 thin films are also sought after for their use in thin film lithium batteries because of their high charge storage capacity. As seen in Chapter 2, V2O5 exhibits a doublelayered structure with a lot of open space, which is conducive to facile intercalation of ionic species and to their transport. It has been verified that the increasing nanostructure of V2O5 leads to increasing electrochemical activity as well as to the evolution toward a nanostructure with the number of lithium intercalation/deintercalation cycles [59]. The double-layer structure of V2O5 films, as seen earlier in Chapter 2, emanates from the alternately up- and down-facing edge-sharing VO5 pyramidal structures, with every third row changing because of the corner sharing. These double-chain sheets are held together with weak VeO bonds in the c direction. A high density of lithium can be intercalated reversibly in these layered structures. However, the films display phase change depending on the amount of charge inserted. The generally initially a-V2O5 with an orthorhombic structure tends toward ε-LixV2O5 with increasing lithium (x) insertion in the range 0.35e0.7. A further increase in lithium to one atom per unit forms the d-LiV2O5 phase [60]. These phase transformations are remarkably reversible. Phase irreversibility sets in for higher values of x for lithium leading to the formation of the g-LixV2O5 phase. Even higher amounts of lithium intercalation are reported with the

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formation of the u-Li3V2O5 phase with a tetragonal structure. It is to be noted that even with these irreversible phases of g-LixV2O5 and u-Li3V2O5 the lithium itself can be reversibly cycled over a large limit, albeit with severe limitations in the discharge rates [61,62]. The amount of lithium that can be intercalated in these irreversible phases (g, u) corresponds to very high charge capacities that are desirable for lithium battery applications. However, in applications related to electrochromics a much smaller lithium charge that can be easily and reversibly accommodated in the reversible phases of a, ε, and d may suffice for the efficient operation of an EC system in terms of coloration and kinetics. The effect of deposition conditions on the EC performance of ammonium-doped nanostructured V2O5 thin films prepared by a novel chemical bath deposition has been studied [63]. In Fig. 3.18A are shown the SEM micrographs of the resulting (NH4)0.3V2O5e1.25H2O films. As can be seen from these images, the 200- and 1220-nm-thick films exhibit a very rich and varied morphology that is conducive to lithium ion intercalation depending on the conditions of preparation. Deposition time and temperature seem to play an important role in determining the morphology and the thickness of the films. Though at different thicknesses, the films are made up of ribbon-like strands whose width, length, and density (porosity) change according to the conditions of deposition. In the thinner films the ribbons are longer (500 nm), with pore diameters of the same dimensions. In the thicker films the length of the ribbons is reduced (250 nm) and the packing density of them increases, thus decreasing the porosity. The lithium intercalation behavior of these films was verified by CV cycles and the optical changes were induced reversibly. In Fig. 3.18B and C are shown the CV cycles and the transmittance changes occurring in some of these typical films. These films show fairly reproducible CV cycles, with the thinner film exhibiting two anodic and two cathodic peaks, while the thicker film exhibits a single anodic peak and two cathodic peaks. These differences illustrate the difference between the amorphous (thicker) film and the crystalline (thinner) film. Depending on the deposition conditions, most of the films studied here have a very high transmittance (>80%) in the normal state in the higher visible to near-infrared region. With the insertion of lithium the optical band edge moves to lower wavelengths giving the typical visible change from yellowish to clear. In the higher wavelength region studied there is an about 20% reduction in the transmittance after lithiation. These changes are shown for a (NH4)0.3V2O5e1.25H2O film prepared under two different

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Figure 3.18 (A) Scanning electron microscopy micrographs of (NH4)0.3V2O5e1.25H2O films deposited under different conditions: (a, b) at 50 C for 30 min, (c) at 85 C for 3 min, and (d) at 85 C for 40 min. Scale bars: (a, c) 0.5 mm, (b) 0.1 mm, and (d) 10 mm. (B) Cyclic voltammetry cycles for (NH4)0.3V2O5e1.25H2O films deposited at 85 C for different times: (a) 3 min (thin) and (b) 20 min (thick). (C) Transmittance spectra of (NH4)0.3V2O5e1.25H2O films in the as-deposited and lithium-intercalated state: (a) film deposited at 50 C for 40 min with aging and (b) film deposited at 85 C for 7 min with color/bleach cycles. (From M. Najdoski, V. Koleva, A. Samet, Journal of Physical Chemistry C 118 (2014) 9636e9646.)

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conditions in Fig. 3.18C. It can be seen that, in general, the films exhibit a good stability with respect to time and color/bleach cycles. Lithium intercalation studies were carried out in electrophoretically deposited V2O5 thin films yielding an orthorhombic a phase with high porosity [64]. Li intercalation using an electrolyte of LiClO4 in PC was studied in two applied potential ranges, narrower (0.4 to 1.1 V) and wider (0.4 to 1.6 V), leading to lower (25 mA/cm2) and higher (50 mA/cm2) current densities as well as to lower (300 mAh/g) charge capacities, respectively. Very interesting observations were made in the two ranges studied. The SEM image of the as-deposited V2O5 film is shown in Fig. 3.19A. The film retained this morphology with the insertion and extraction of lithium using the narrower voltage range. However, the film morphology changed drastically when intercalated with lithium using the wider voltage range, as shown in Fig. 3.19B. The film developed cracks along the grain boundaries. However, in both regimes the intercalation and deintercalation cycles were reproducible over a large number of cycles, albeit with some changes, as shown in the form of CV cycles in Fig. 3.19C. The CV cycles during the narrower voltage range nearly exhibit three peaks in both the cathodic and the anodic cycles corresponding to lithium intercalation and deintercalation. With increased lithium cycling the kinetics seem to stabilize with the same distinct peaks and an increased amount of charge inserted and extracted as indicated by a slight expansion of the CV cycle [Fig. 3.19C (a)]. In comparison to this, the intercalation over a higher range of voltage, as shown in the CV cycles [Fig. 3.9C (b)], seems to lead to some irreversible processes. The CV peaks in the 1st cycle and the 10th differ significantly and the overall area of the cycle expands, indicating a higher amount of charge inserted and extracted. The transmittance changes associated with the lithium ion insertion and extraction of these films is shown in Fig. 3.19D. The as-deposited film exhibits a fairly high transmittance (>60%) in the higher wavelength range studied, with a strong absorption in the lower visible wavelengths. This is the typical spectrum of a-V2O5, rendering it with the deep yellowish color. However, with the insertion of lithium the film exhibits a very intense coloration over the entire spectral region studied, though the absorption band edge seems to move to lower wavelengths. It is also to be noted that the authors report this coloration at a very high current density of 100 mA/

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Figure 3.19 Scanning electron microscopy images of an electrophoretically deposited V2O5 film (A) in the as-deposited state and (B) after 50 cycles of lithium insertion/ extraction. (C) Cyclic voltammetry cycles of the film in the lower voltage range (0.4 to 1.1 V) (a) and higher voltage range (0.4 to 1.6 V) (b). (D) Transmittance spectra of the film in the as-deposited and after the lithium-inserted and lithium-extracted cycles. (From Y. Wang, Guozhong, 51 (2006) 4865e4872.)

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cm2, which is much higher than the current density they report even in the higher voltage range of their study. The interesting feature is that the films seem to be stable under repeated coloration and bleaching cycles at this high current density. From an ECD point of view the significance of this work is that one can stay in the lower voltage regime with very reproducible and stable electrochemical kinetics and optical and structural properties. The amount of lithium inserted and extracted in this voltage regime is sufficiently large to induce a significant anodic EC coloration desired for the use of V2O5 films as counterelectrodes. The optical and structural properties of sputter-deposited V2O5 thin films were verified under dry lithiation [65]. Films were prepared with varying amounts of oxygen in the chamber. In Fig. 3.20A and B is shown the optical behavior of some of these films under the insertion of lithium atoms. In the as-deposited state the film typically exhibits a high transmittance (>70%) in the higher wavelength region studied. A very sharp transmittance cutoff region exists because of the absorption band located below 500 nm. With increasing amounts of lithium atoms inserted the absorption edge moves systematically toward lower wavelengths bringing about the anodic discoloration to the film, rendering it clear in the visible region of the spectrum. It can also be seen for both the V2O5 samples shown here that at higher wavelengths there is a small decrease in the transmittance, also referred to as cathodic coloration. Such opposite colorations occurring in two different spectral regions in V2O5 have been studied in detail by Wu et al. [66]. The cathodic coloration occurring in the higher wavelength region is attributed to the small polaron absorption in the V2O5 films and the anodic coloration is due to the blue shift of the absorption edge, both occurring under double insertion of lithium and electrons. A similar effect can be assumed in the dry lithiated V2O5 films. For EC applications, especially for the use of V2O5 films as counterelectrodes (IS), it is advantageous to have a dominant anodic behavior with a small or no cathodic component. The films shown in Fig. 3.20A and B are very interesting from this point of view. The predominant optical change is in the lower wavelengths with only a slight decrease in transmittance at higher wavelengths. In Fig. 3.20B are shown the changes occurring in the nanostructure of one such film with lithium insertion. The as-deposited film shows a very fine structure with small grains. With increasing lithium insertion the structure gradually evolves, with the grains getting

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Figure 3.20 (A, B) Typical transmittance spectra of two of the sputter-deposited V2O5 films as functions of lithium atom insertion. (C) Atomic force microscopy images of the evolution of the nanostructure of the films with the insertion lithium atoms: (a) 0 nm, (b) 5 nm, (c) 10 nm, (d) 15 nm, (e) 20 nm, and (f) 30 nm. Raman spectra of films deposited at two different substrate temperatures, (D) 25 C and (E) 250 C. (From R. Balu, P.V. Ashrit 2008 (unpublished work).)

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larger in most parts up to an insertion of 20 nm (12.5 mC/cm2). At the highest level of lithium insertion studied (19.5 mC/cm2) a drastic modification of the surface occurs with the formation of large islands. This is perhaps the saturation level, as can also be seen from the transmittance changes in both the lower and the higher spectral regions of Fig. 3.20A, indicating the probable onset of irreversible processes. The changes occurring in these films with lithiation were followed by Raman spectroscopy. Fig. 3.20D and E shows the Raman profile of two such films deposited at room temperature and at 250 C. The V2O5 films are mainly characterized by two peaks, one at 935 cm1 and the other at 1012 cm1. These peaks and their changes with intercalation are more prominent in the high-temperature-deposited film. The peak at 935 cm1 is attributed to the VO4þ]O stretching band and the peak at 1012 cm1 is attributed to the VO5þ]O stretching band. The appearance of both peaks can be attributed to the presence of a mixed oxide phase comprising both V2O5 and VO2. The strong peak seen at 147 cm1 in the room-temperaturedeposited film indicates the presence of a layered structure in the V2O5. Various other peaks correspond to various V2O5 stretching. With the insertion of lithium the peak at 147 cm1 starts increasing initially and then decreases. This effect can be attributed to the effect of lithium occupying the space between the layers, thus increasing the bond length between the layers. At the highest insertion of 20 nm (12.5 mC/cm2) there seems to be a structural damage and amorphization of the oxide [67]. With increasing equivalent double insertion of lithium and electrons, the electrons reduce V5þ to V4þ sites, thus creating more V4þ]O bonds as indicated by the increase in the intensity of the peak at 935 cm1. Once again, at the highest amount of intercalation (20 nm) the peak corresponding to 935 cm1 is lost, indicating the collapse of the structure.

3.3.3 Nanostructured Molybdenum Trioxide Films MoO3 films are very well known to exhibit an efficient EC and photochromic coloration [27,68]. With a crystal structure very similar to that of WO3, perovskite-like, MoO3 exhibits a reversible intercalation behavior similar to WO3: MoO3 þ zIþ þ ze 4 IzMoO3 (3.12) ðTransparentÞ

ðDarkÞ

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The coloration under lithium atom insertion too has been verified to proceed in a similar fashion [69]: MoO3 þ zLio 4 Liz MoO3 ðTransparentÞ

ðdarkÞ

(3.13)

Hence the EC behavior of MoO3 films is very similar to that of WO3, being equally efficient. However, the absorption band in the case of MoO3, in general, is located at a lower wavelength region than in the case of WO3, which gives it a darker appearance than the deep blue coloration seen in WO3 films. The CE exhibited by these films is parallel to that of WO3 thin films. Despite these great virtues the work related to the EC properties is nowhere comparable to that dedicated to the WO3 films. There is a lot of work on MoO3 films connected to lithium batteries [70e72]. This is due to the fact that MoO3, similar to V2O5, presents a layered structure especially in the phase. This is an orthorhombic phase formed by the corner-sharing MoO6 octahedra. The sheets thus formed are held together by weak van der Waals forces forming a layered structure. This open structure between the layers is very conducive to IS and transport, making MoO3 a sought-after material for high ion density storage as counterelectrodes in lithium batteries. The combined merits of this material for high-density ion intercalation and the efficient EC effect make it another ideal material for ECDs. New traits and improved EC performance have been achieved in MoO3 films by working with the nanostructured form. Sputter-deposited MoO3 films were studied for their EC properties by the dry lithiation method [69]. AFM images of the films are shown in Fig. 3.21 in the as-deposited state, and the evolution of the nanostructure with increasing insertion of lithium atoms is also shown. The average grain size of 48 nm found in the as-deposited film inflates gradually to 50 nm (2.5 nm Li), 85 nm (5 nm Li), and 180 nm (10 nm Li). Further lithium insertion beyond 10 nm, as seen in these images, seems to distort the surface structure. This can be clearly seen in the case of insertion of 15 and 20 nm of lithium, where the inflated grains collapse into separate particles. The transmittance changes occurring with the insertion of lithium are shown in Fig. 3.21G. In the as-deposited state the MoO3 film exhibits a very high transmittance on the order of 80% throughout the spectral region studied. With the insertion of lithium a fairly effective EC coloration ensues, with the peak of modulation centered in the visible region. As also

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Figure 3.21 Atomic force microscopy images of a sputter-deposited MoO3 film with the insertion of various amounts of lithium atoms: (A) 0 nm, (B) 2.5 nm, (C) 5 nm, (D) 10 nm, (E) 15 nm, and (F) 20 nm. (G) Transmission spectra of the MoO3 film with the insertion of these amounts of lithium atoms. (From A. Taj, P.V. Ashrit, Journal of Materials Science 39 (2004) 3541e3544.)

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seen in the AFM images, the optical modulation seems to saturate beyond the insertion of 15 nm of lithium. Peak CE for this MoO3 sample for an insertion of 20 nm (12.5 mC/cm2) was around 40 C/cm2 and was situated around a wavelength of 750 nm. The dry lithiation method was also used to study nanostructured MoO3 thin films deposited by the GLAD method [73]. This method is very well suited to tailoring the MoO3 film nanostructure to adjust the optical and EC properties. The films were prepared by systematically adjusting the substrate tilt parameter from 0 to 85 degrees at a constant rotation speed of 20 rpm in all cases. In Fig. 3.22A are shown the AFM images of the samples

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(C)

Figure 3.22 (A) Atomic force microscopy images of the evolution of grain size as a function of the tilt angle of the glancing angle-deposited MoO3 films. Images on the left side are for as-deposited films and those on the right side are for annealed films with tilt angles of 0 degrees (a and e), 60 degrees (b and f), 80 degrees (c and g), and 85 degrees (d and h). Absorption spectra of (B) as-deposited and (C) annealed MoO3 films. From the top, the substrate tilt angles during deposition are: (a) 0, (b) 20, (c) 40, and (d) 60 degrees. (From G. Beydaghyan, M. Boudreau, P.V. Ashrit, Journal of Materials Research 26 (1) (2011) 56e61.)

deposited at room temperature and at different inclinations before and after annealing at 400 C. At lower tilt angles, i.e., between 0 and 60 degrees, the films show a very fine-grained structure (5e50 nm) as seen in Fig. 3.22A (a) and (b). With increasing tilt angle beyond this, a fairly big change occurs in the grain size due to the increasing shadowing effect at large angles as seen in Fig. 3.22A (c) and (d). After the films were annealed at 400 C, a dramatic change occurred in each of them. The aggregates seen before annealing seem to coalesce to form a layered structure with domains on the order of 1e5 mm as seen in Fig. 3.22A (e) to (h). The possible formation of a predominantly a-MoO3 phase after annealing was verified by XRD studies. The film porosity as calculated from the refractive indices varied dramatically with increasing tilt angle. The porosity (p) in a 0-degree tilt was around 6%e12%, varying all the way up to around 80% at the highest

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tilt studied, 85 degrees. The equivalent mass thickness, which was obtained by multiplying the physical thickness (t) by (1  p), was found to vary equally drastically from 240 to 31 nm with tilt angle varying between 0 and 85 degrees. This wide range of variation in the film nanostructure attainable by the GLAD method makes it possible to tailor the optical and EC properties. However, one has to pay attention to the optimization of EC properties in such porous films. As seen in this case of MoO3 thin films, while tuning the porosity continuously attention has to be paid to the actual material (equivalent mass thickness) available in the film to induce EC coloration. Also as mentioned earlier, the mixed ionic and electronic conductivity too has to be optimized [56]. Hence, a judicious choice of the nanostructure to optimize the ionic conductivity (through porous nanostructure), electronic conductivity (through a good intergranular contact), and mass thickness (sites) has to be made to induce a good coloration. The EC response of the MoO3 films deposited by the GLAD method at various tilt angles is shown in Fig. 3.22B and C as a function of the quantity of lithium inserted. The as-deposited MoO3 film, which was initially very transparent (w80%), exhibits a very deep coloration centered around 800 nm. Absorption of up to a maximum of 50% is seen in films with 0-degree tilt with the maximum lithiation of 12.5 nm. There is a slight diminution in the optical absorption with lithiation as the tilt angle increases. This is due to the decrease in the effective mass thickness available at higher tilt angles. The behavior of annealed MoO3 films (Fig. 3.22C) is quite different, especially at higher wavelengths. It is seen that the optical absorption band is more or less spread out through most of the spectral region studied. The absorption is very effective in films deposited at 0 degrees with peak absorption around 50% in the vicinity of 1000-nm wavelength and decreasing only slightly at other wavelengths. However, with increasing tilt angle it is seen that the absorption, though more or less uniform throughout the spectral region studied, decreases quite swiftly. This work demonstrates the wide range of possibilities that exist for tailoring the nanostructure of MoO3 films, in particular, and TMO films, in general, for their optical, electrical, and EC properties. The as-deposited MoO3 films of this work showing efficient EC coloration in the lower spectral region are suitable for applications such as display devices, while the annealed films showing a wideband absorption are deemed suitable for smart windows for solar energy management. The tailoring of MoO3 film nanostructure to enhance EC properties was also demonstrated by Patil et al. [74]. MoO2 films were initially

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Figure 3.23 (A) Scanning electron microscopy image of an electrodeposited and airannealed MoO3 thin film. (B) Transmission spectra of the air-annealed films in the electrochemically colored and bleached states. DOD, change in optical density. (From R.S. Patil, M.D. Uplane, P.S. Patil, International Journal of Electrochemical Science 3 (2008) 259e265.)

electrodeposited fluorine-doped tin oxide (FTO)-coated glass using an ammonium molybdate precursor solution and subsequently thermally oxidized in air to achieve orthorhombic a-MoO3 films. The amorphous nature of the samples before and the crystalline nature as well as the a-MoO3 phase after the annealing were confirmed by XRD studies. In Fig. 3.23A is shown an SEM image of the annealed MoO3 film. As can be seen from this image the film develops a rodlike structure in which these elongated particles are well coalesced into one another and yet having a good porosity that is necessary for the ion intercalation. The EC coloration behavior through transmittance change is shown in Fig. 3.23B. A fairly effective coloration typical of MoO3 under proton injection is seen through most of the visible region of the spectrum and into the near-infrared region. A CE of 34 cm2/C is reported at the wavelength of 630 nm. The electrochemical kinetics are not reported in this work. Such films with high porosity for facile ion transport and good connectivity of the structure for electron transport may be conducive to kinetics associated with the coloration. The evolution of electrodeposited MoO3 film nanostructure and the consequent evolution of the films’ EC behavior were studied in detail by McEvoy et al. [75]. MoO3 was electrodeposited on ITO-coated glass using a precursor solution of aqueous peroxo-polymolybdate and sintered to

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various temperatures from 25 C to 450 C. The films sintered at lower temperatures (350 C) were crystalline, exhibiting the stable a-MoO3 phase. At intermediate temperatures the films had a mixed phase of a-MoO3 and b-MoO3. The structural evolution from a finegrained smooth surface to the development of microcrystals with crystalline planes is evident as shown in Fig. 3.24A. The EC performance of these films was examined under the insertion of lithium ions. In Fig. 3.24B and C are shown, respectively, the absorption spectra and the absorption change as a function of the insertion/extraction of lithium. The authors report that the mixed-phase crystalline films of MoO3 sintered at higher temperatures (>250 C) exhibit a higher coloration compared to the mixed-valence hydrated amorphous films of MoO3, which were sintered at low temperature (100 C). However, it was found that, while in films prepared at low temperature the insertion/extraction of lithium was facile and reversible, the films prepared with high-temperature sintering did not exhibit completely reversible cycles. In another work reported by the same group the authors studied in detail the inhomogeneity of coloration of polycrystalline MoO3 using a new approach that combines time-lapse transmission optical microscopy with step potential electrochemical insertion of lithium into the MoO3 films [76]. The images obtained by this technique are shown in Fig. 3.24D. These time-lapse images taken at the wavelength of 633 nm were followed as a function of lithium insertion into these mixed-phase MoO3 films. As can be seen from these images the inhomogeneous coloration of the different domains is quite evident. Raman microprobe spectroscopy experiments have shown that the regions showing the least coloration consisted of a nanocrystalline monoclinic b-MoO3 phase and the domains showing the maximum coloration were attributed to the presence of a predominantly orthorhombic a-MoO3 phase. Other regions with intermediate coloration were ascribed to the mixed phases. This technique of chronoabsorptometric imaging is unique and can be a powerful tool in terms of understanding the diffusion behavior of TMO thin films for EC studies. Unlike the other electrochemical and impedance spectroscopy techniques used for such characterization, which give an averaged out behavior of the lithium diffusion, the chronoamperometric imaging is expected to yield an in-depth insight into the charge transport phenomenon in such disordered TMOs.

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Figure 3.24 A) Atomic force microscopy images (10  10 mm) of electrodeposited MoO3 films sintered at various temperatures. (B) Absorption spectral response (a) and electrochemical response (b) for the MoO3 film sintered at 100 C. (C) Absorption spectral response (a) and electrochemical response (b) for the MoO3 film sintered at 250 C. (D) Chronoabsorptometric imaging of MoO3 without lithium (a) and with lithium (b). (From T.M. McEvoy, Keith J. Stevenson, J.T. Hupp, X. Dang, Langmuir 19 (2003) 4316e4326; T.M. McEvoy, K.J. Stevenson, Journal of American Chemical Society 125 (2003) 8438e8439.)

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Yung-Sen et al. [77] have deposited highly efficient EC thin films of organo-molybdenum oxides (MoOxCy) with the aim of fabricating flexible ECDs. These films were deposited on poly(ethylene terephthalate) (PET) substrates coated with ITO thin films using molybdenum carbonyl [Mo(CO)6] powder and a plasma torch. Nanoporous thin films of around 300 nm thickness made up of MoOxCy nanocrystals were formed with high mechanical flexibility on the PET/ITO substrates. A field-emission SEM image of the film surface is shown in Fig. 3.25A. The EC behavior of this film was studied under lithium intercalation. The highly porous surface of the MoOxCy film seems to be very favorable to the insertion/extraction of lithium with very good CE in the visible region of the spectrum as shown in Fig. 3.25B. As can be seen in this figure the film changes from a uniform (neutral) and high-transmittance (>80%) state to a highly uniform and

Figure 3.25 (A) Field-emission scanning electron microscopy image of a MoOxCy film deposited on a poly(ethylene terephthalate)/indium tin oxide substrate. (B) Transmittance spectra of the same film in the lithium-intercalated (colored; Tc) and deintercalated (bleached; Tb) states. (C) Cyclic voltammogram of the same film under various cycles of insertion and extraction. (FromY.-S. Lin, T.-H. Tsai, S.-W. Tien, Thin Solid Films 520 (2013) 248e252.)

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colored state. This optical change corresponds to a CE of around 37.5 cm2/ C. The most interesting point about these films is their cyclic stability. In Fig. 3.25C are shown the cyclic voltammograms of the films, which indicate their reversible operation against lithium insertion/extraction even after 200 cycles of coloration and bleaching. Such films with high CE on flexible substrates are very interesting from the application point of view.

3.4 A FEW EXAMPLES OF ELECTROCHROMIC DEVICES BASED ON NANOSTRUCTURED THIN FILMS As discussed in earlier sections, nanostructured TMO thin films play a vital role in ECDs. A few important examples of ECDs comprising primarily nanostructured TMO thin films as the base EC layer are discussed here. As previously emphasized in Section 3.1.1, for a stable and efficient EC operation a five-component system, as shown in Fig. 3.1, is deemed essential. The EC layer, which forms the basis of the operation of the entire ECD, consists essentially of various TMO thin films. The most essential optical switching of the device occurs in this layer and all the other layers are designed around this TMO-based EC layer to provide the necessities of the mechanism as per Eq. (3.1). The basis of this EC coloration being the double injection and extraction of the species (ions and electrons) into this TMO layer, logically the design of the ECD must start by setting the specifications of the TMO-based EC layer. These specifications include the minimum film thickness and CE intended for the particular TMO used. One has to define equally the wavelength region of interest where this coloration needs to occur. Hence, the TMO needs to be highly transparent in its normal (bleached) state and show the maximum coloration in the wavelength region of interest under the insertion of charge. In addition to this optical switching, the TMO should have a high charge capacity as well as a high mixed (ionic and electronic) conductivity. A complete characterization of the TMO thin film of interest as an individual layer needs to be carried out and all the essential parameters required for EC operation need to be optimized prior to attempting a device fabrication. This prior optimization study should establish the ionic charge, Q (Hþ, Liþ, etc.), needed to electrochromically color the EC layer to a certain specified degree [DOD (l)]. Based on this information, an appropriate IS layer needs to be designed and optimized. This could be another TMO with a much weaker CE in the wavelength region of interest. It is desirable to have a TMO-based IS

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that is highly transparent in the wavelength region of interest and does not undergo much coloration (low CE) under the double insertion. In addition, similar to the EC layer, the IS layer needs to have a high mixed conductivity as well as a high charge capacity. As discussed earlier in Section 3.1.1, the use of weakly or complementarily (anodically) coloring vanadium pentoxide (V2O5) as the IS layer in conjunction with the strongly coloring tungsten oxide (WO3) as the EC layer, for example, is a common practice [1]. A similar example is the use of an anodically coloring nickel oxide (NiO) layer in combination with a cathodically coloring WO3 layer [78]. The IS layer should be studied and optimized as an individual layer. The minimum film thickness necessary to accommodate the aforementioned specific charge Q should be determined by the electrochemical insertion/ extraction of the species of interest. Following this, the appropriate IC layer for the type of charge species þ (H , Liþ, etc..) needs to be identified. The requirements for the IC layer are: (1) a high transmittance in the spectral region of interest, (2) a high ionic conductivity for the type of ions used in the device, and (3) a high electrical resistivity. The IC layer is inserted between the EC and the IS layers to isolate the two layers electronically while allowing the transport of ions. A wide range of solid, liquid, and gel ICs, both organic and inorganic, are available in the literature [1]. The most appropriate ion-conducting layer needs to be studied and integrated in conjunction with the particular ECeIS combination chosen. This three-layer combination is then flanked on either side by a TC layer, which is used to apply the necessary electric field to activate the EC coloration and bleaching. The TC layers need to be highly transparent in the spectral region of interest and possess a high electronic conductivity. This five-component system is either deposited on a single substrate (monolithic ECD) or held between two substrates (sandwich ECD) as shown in Fig. 3.26. Each type of configuration presents its own challenges and advantages. In the monolithic configuration (Fig. 3.26A), in which all five layers need to be deposited on a single substrate, the challenge is to optimize each of the layers in the EC system for its specific functionality as an integral part. Severe limitations apply to the conditions one can use for the optimization of a particular layer owing to the presence of other already deposited layers in the stack. Hence, one has to be prudent as to what conditions can be used as the five-layer stack is built up. The advantage of the monolithic system is that once these limitations are overcome and conditions for optimization of each layer are fixed, the fabrication and operation of such a system as an integral unit is

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Figure 3.26 (A) Schematic of a monolithic all-solid thin film electrochromic device (ECD). (B) Schematic of a “symmetric” sandwich ECD. (C) Schematic of an “asymmetric” sandwich ECD (half cell). EC, electrochromic layer; IC, ion conductor; IS, ion storage layer; TC, transparent conductor.

quite simplistic and highly reproducible. The sandwich structures (Fig. 3.26B), on the other hand, pose relatively less stringent limitations to the optimization of individual layers, as part 1 and part 2 are prepared separately and put together with a spacer and a sealant inserted with the appropriate ion-conducting material. Such a sandwich structure involves an EC layer on one substrate and an IS layer on the other substrate. In Fig. 3.26C is shown a device in which the intermediate component between part 1 and part 2 acts both as the IS and the IC. Quite a number of ECD configurations, as discussed later, based on these essential layers are proposed for the fabrication of inorganic ECDs. Some of the five essential components of the ECD can be in liquid or gel or solid form and in some cases a single component can play multiple roles. Further, the substrates

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used to support the entire thin film system can be rigid or flexible. All these factors lead to a wide range of ECD configurations. ECDs of the form TC/ EC/IC/IS/TC, in which the acronyms represent the various layers shown in Fig. 3.1 and discussed earlier (Fig. 3.26B), are referred to as “symmetric” cells. Simplified ECDs of the form TC/EC/ICeIS/TC (Fig. 3.26C), in which the IC layer also plays the role of IS, are referred to as “asymmetric” or “half” cells. It is more common to use symmetric cells for device fabrication, while the “half” cells are used for the electrochemical testing of the TMOs and for their optimization [78]. In some device configurations, instead of providing external sources for the ECD operation, the ECDs are integrated directly with photovoltaic (PV) devices [79] or the two systems are put in tandem [80] to make them self-powered systems when exposed to light. A few examples of ECDs of each type are discussed next.

3.4.1 Electrochromic Devices With Liquid Electrolytes The use of liquid electrolytes as the source of ions (as well as IS) needed for the EC mechanism and also as the IC is very common in most device fabrication work. It is the simplest and most convenient configuration although marred by sealing problems. ECDs based on complementary electrodes of nanocrystalline titanium dioxide (TiO2) and tungsten trioxide (WO3) thin films were fabricated and tested by Bonhôte et al. [81]. The electrodes in this symmetric configuration consisted of a 0.2-mm-thick nanostructured WO3 film deposited on F-doped SnO2 transparent conductors on one side and a 5-mm-thick nanocrystalline TiO2 film derivatized by a monolayer of a phosphonated triarylamine. A schematic of the derivatized TiO2 electrode is shown in Fig. 3.27A. It consists of a nanocrystalline TiO2 film deposited on a transparent conductor on which a compact monolayer of a phosphonated triarylamine is adsorbed. The average TiO2 particle size was 20 nm. The derivatized molecular monolayers on the TiO2 surfaces on the conducting substrates were shown to become electroactive owing to electron exchange with their conductive substrates. A lateral exchange of charge between the monolayer molecules was shown to take place [82]. A schematic of the symmetric device is shown in Fig. 3.27B. As can be seen in this figure, the device was formed by spacing the two glass plates with the transparent conductors (F:SnO2) and their respective electrodes (WO3 and TiO2) 100 mm apart and filling this space with lithium(I) bis(trifluoromethanesulfonyl)imide solution in methoxyethoxypropionitrile as the electrolyte. The molecular monolayer used here was transparent in its

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Figure 3.27 (A) Schematic of the nanostructured TiO2 film derivatized with a molecular monolayer of phosphonated triarylamine. (B) Schematic of the operation of the EC device comprising the nanostructured WO3 film and the derivatized nanostructured TiO2 complementary electrodes. (Adapted from P. Bonhôte, E. Gogniat, M. Grätzel, P.V. Ashrit, Thin Solid Films 350 (1999) 269e275.)

reduced state and became highly absorbing in the higher visible wavelengths, thus becoming blue when oxidized. The application of a small positive potential to the TiO2 electrode resulted in the oxidation and coloration of the monolayer by lateral charge percolation with the conducting substrate. The simultaneous cathodic coloration of the nanostructured WO3 occurred as the lithium ions diffused, bringing about the redox reaction of this electrode. The complementary and simultaneous reactions occurring in both the electrodes lead to an intense coloration of the device in the visible and near-infrared region of the spectrum as shown in Fig. 3.28A. In addition these devices showed an extremely rapid response as shown in Fig. 3.28B.

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Figure 3.28 (A) Transmittance spectra of a nanocrystalline TiO2 and WO3 electrochromic device (ECD) in the colored state with the application of 1.5 V (lower curve) and the bleached state at 0 V (upper curve). (B) Electrochromic response of the ECD following a voltage step from between 0 and 1.5 V to 0 V. (C) Coloration and bleaching cycles of the ECD submitted to 14,400 cycles of 5 s at 0 V (dashed lines) and 5 s at 1.5 V (solid lines). (Adapted from P. Bonhôte, E. Gogniat, M. Grätzel and P.V. Ashrit, Thin Solid Films 350 (1999) 269e275.)

Under cyclic colorationebleaching operation the films were also found to be extremely stable, showing a slight variation after 14,400 cycles as shown in Fig. 3.28C. Such devices combining not only the nanostructure approach but also molecular electrochromism are extremely advantageous for fabrication and applications. In addition to using only the dye-sensitized nanostructured TiO2 in the capacity of complementary layers against the cathodically coloring WO3 thin films, the PV nature of these electrodes is also exploited for powering the EC mechanism. The so-called photoelectrochromic (PEC) devices as shown in Fig. 3.29A are self-powered under illumination [79]. The photovoltage produced by the dye-sensitized TiO2 electrode under illumination drives the EC mechanism in the WO3 film through the injection of electrons and the charge-compensating Li ions from the liquid

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Figure 3.29 (A) Schematic of a photoelectrochromic (PEC) device based on dyesensitized TiO2 and WO3 thin film. (B) Transmittance spectrum of the PEC device in the bleached and colored states. (From D. Benson, R. Crandall, S.K. Deb, J.L. Stone, StandAlone Photovoltaic Powered Electrochromic Windows, (US Patent 5384653), January 24, 1995).)

electrolyte. A fairly efficient EC coloration as shown in Fig. 3.29B was demonstrated. Such integrated self-powering PVeEC systems are very advantageous for large-area application because of their simplicity in fabrication and economic operation.

3.4.2 Solid-State Electrochromic Devices Quite a few ECDs involving only solid-state components have been proposed, which include both monolithic and sandwich configurations. The sandwich-type solid-state ECDs based on polymer electrolytes act as good ICs. A wide range of polymeric proton and lithium ICs are available [1] for this purpose. An example is the so-called photovoltachromic cell (PVCC), which is equivalent to the aforementioned PEC with liquid electrolytes instead of solid polymer electrolytes [80]. Using a hybrid of PEG and titanium (Ti) the authors demonstrated the fabrication and successful EC operation of a solvent-free device. One electrode consisted of a dye-sensitized TiO2 nanoparticle film and the other electrode was patterned WO3/Pt. A schematic of the device is shown in Fig. 3.30A. Both the electrodes were deposited on FTO as transparent electrodes. The hybrid PEG/Ti supplemented with LiI, I2, and 1,2-dimethyl-3propylimidazolium iodide (DMPII) formed the solvent-free polymer between the two electrodes, forming a sandwich configuration. The coloration and bleaching cycles of this device containing an optimal amount of DMPII are shown in Fig. 3.30B. The tunability of the color of

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Figure 3.30 (A) Schematic of a photovoltachromic cell (PVCC) based on dye-sensitized TiO2 nanoparticle film and patterned WO3/Pt film connected to an external load resistance R. (B) Coloration and bleaching cycles of the PVCC recorded at 788-nm wavelength in the light-on and light-off states. FTO, fluorine-doped tin oxide; PEG, polyethylene glycol. (From M.-C. Yang, H.-W. Chao, J.-J. Wu, Nanoscale 6 (2014) 9541e9544.)

the PVCC with varying degree of addition of DMPII as well as its longterm coloration and bleaching stability was demonstrated. One of the earliest monolithic all-solid ECDs was fabricated and tested by Ashrit et al. [57]. As mentioned earlier, the preparation of such monolithic systems is a formidable challenge as each of the layers in the device has to be optimized for its optical, electrical, and/or EC functionality as an integral layer of the device. This causes severe restrictions on the thin film preparation as the conditions of preparation become restrictive with the addition of each layer. In this work, the authors report the successful fabrication of an ECD with the material configuration glass/ITO/LixV2O5/ LiBO2/WO3/Au. Each of the layers under study was initially deposited on a glass substrate as an individual layer. The optically active layers of V2O5 and WO3 films were intercalated with various amounts of lithium for optimization work. The dry lithiation method was the most convenient for this work [12]. The transmittance spectra of WO3 and V2O5 thin films of 200 nm thickness are given in Fig. 3.31. The thermally evaporated amorphous WO3 films were deemed the most suitable for this work, showing an excellent cathodic coloration under lithium insertion of up to 25 mC/cm2 with the formation of LixWO3, as shown in Fig. 3.31A. The reversibility of this charge was verified by electrochemical insertion and extraction of a similar sample deposited on an ITO-coated glass substrate. This established the amorphous and stable WO3 film deposited at 200 C with a reversible charge capacity of 25 mC/cm2 (40 nm effective mass thickness) of lithium

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Figure 3.31 (A) Transmittance spectra of a 200-nm-thick WO3 thin film with the insertion of various effective mass thicknesses of lithium: (curve a) 0 nm, (curve b) 10 nm, (curve c) 20 nm, (curve d) 30 nm, and (curve e) 40 nm. (B) Transmittance spectra of 200-nm-thick V2O5 film with the insertion of various effective mass thicknesses of lithium: (curve a) 0 nm, (curve b) 10 nm, (curve c) 20 nm, and (curve d) 50 nm. (From P.V. Ashrit, et al., Solid State Ionics 59 (1993) 47e57.)

as the base layer around which the other layers and the ECD were tailored. Similarly, the V2O5 film was sputter deposited on a glass substrate to a thickness of 200 nm and intercalated with up to 50 nm effective mass thickness (31 mC/cm2) of lithium. Only a weak EC anodic coloration was displayed by these films, as shown in Fig. 3.31B, with the formation of LixV2O5. The initially bright

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yellow film became more and more transparent with the insertion of increasing amounts of lithium, due to the well-known shift in the absorption edge from the visible to the ultraviolet region. As can be seen in this figure, even at the insertion of 31 mC/cm2 the film exhibits a fairly high transmittance in the solar spectral region. Hence, a 200-nm-thick sputter-deposited V2O5 film intercalated with 50 nm effective mass thickness of lithium was deemed suitable for lithium IS in the device. Based on these results and the suitability of lithium borate (LiBO2) thin films in the capacity of efficient lithium ICs [12], an all-solid monolithic system with the configuration glass/ITO/LixV2O5/LiBO2/WO3/Au was put together with a thin film of gold (15 nm) as the outer transparent conductor. For the fabrication of the device, V2O5 thin films were first sputter deposited onto ITO-coated glass substrates. They were then intercalated with 40 nm effective mass thickness (25 mC/cm2) of lithium. Subsequently thin films of LiBO2, WO3, and Au were deposited to form the device. The results of the EC operation this system are shown in Fig. 3.32. As shown in Fig. 3.32A a highly efficient EC switching was demonstrated by such a device although the overall transmittance in the normal and completely bleached states of the system was very low owing to the outer gold thin film. The switching operation demonstrated the feasibility of this EC system and of the various components used therein. A slow but excellent reversibility of the EC operation of the device is evident from the coloration and bleaching cycles shown in Fig. 3.32B and C. The device was optically and electrically stable and reversible over more than 1000 cycles of coloration and bleaching. The system also exhibited an excellent optical memory, retaining its extreme colored and bleached states for well over 8 h with no significant change. Hence, this configuration is extremely interesting and carries a high potential for ECD fabrication. Based on this initial work, Ashrit et al. [58] developed two other all-solid monolithic and more transparent EC systems. The materials configuration of their first system was glass/ITO/LixV2O5/LiBO2/WO3/ ITO, wherein the outer gold electrode of their earlier work [57] was replaced by a more transparent ITO film. The deposition of this outer transparent conductor with a high degree of solar and visible transmittance as well as a high electronic conductivity was the most challenging aspect in such monolithic systems. This layer needed to be deposited under severe limitations of deposition conditions without altering the already optimized properties of the underlying layers. Hence, a high substrate temperature or a high degree of postdeposition thermal treatment, which is generally

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Figure 3.32 (A) Transmittance spectra of an all-solid monolithic electrochromic (EC) device with different applied voltages: (curve a) 2 V, (curve b) 1 V, (curve c) 0 V, (curve d) 1 V, and (curve e) 2 V. (B) Coloration and bleaching cycles of the EC system measured at a wavelength of 815 nm with an applied voltage of 3 V. (C) Cyclic voltammetry (a) and transmittance change recorded at a wavelength of 632 nm (b) of the all-solid device after several hundred cycles. (From P.V. Ashrit, et al., Solid State Ionics 59 (1993) 47e57.)

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necessary to obtain good quality ITO films, was deemed impossible. The challenges associated with the deposition and optimization of the outer ITO electrode were dealt with by sputter depositing the film at room temperature and simultaneously bombarding the film during and after the deposition with high-energy argon atoms using a fast atom beam (FAB) gun [83]. These room temperature atom-assisted deposited ITO films had a resistivity of 3  104 U cm and solar and visible transmittances of over 80%. The EC switching properties of this first system with the configuration glass/ITO/LixV2O5/LiBO2/WO3/ITO are shown in Fig. 3.33. As can be seen from Fig. 3.33A an excellent optical switching is exhibited by this device with the application of 2 and 3 V. The integrated solar and photopic transmittance values in the bleached and colored states are as follows: Tsolar (bleached) ¼ 58%, Tphotopic (bleached) ¼ 65% Tsolar (colored, 2 V) ¼ 14%, Tphotopic (colored, 2 V) ¼ 18% Tsolar (colored, 3 V) ¼ 9%, Tphotopic (colored, 3 V) ¼ 13% This optical performance of the ECD is deemed fairly adequate in view of its application as a smart window in the light of the criteria established for this application [84], which are: Tsolar (bleached) ¼ 50%e70%, Tphotopic (bleached) ¼ 50%e70% Tsolar (colored) < 10%e20%, Tphotopic (colored) < 10%e20% The coloration and bleaching cycles of this monolithic ECD were very reproducible, though very slow, requiring 6 min for each cycle at room temperature. To verify the stability of its EC operation, the device was cycled under varying conditions of temperature and humidity. These results, recorded at a wavelength of 801 nm, are shown in Fig. 3.33B and C. As can be seen in Fig. 3.33B, the cyclic performance of the device actually improves with increasing temperature in terms of both its optical switching and its speed of switching. The switching speed at the high temperature of 60 C tested was five times lower than that at room temperature. This cyclic operation at high temperature was quite reversible without significant damage to the device. In a similar study the switching of the device under varying humidity conditions was examined with the exposure of the device for 24e48 h to two levels of relative humidity, 25% and 76%. As shown in Fig. 3.33C, relative humidity does not seem to affect the performance of the device significantly. Hence, the outer transparent electrode of ITO seems to act to some extent as a protective layer. The significant variation in the operation of the EC system as a function of temperature seems to originate

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Figure 3.33 (A) Transmittance spectra of an all-solid monolithic electrochromic device (ECD) in the bleached state (curve a) and in the colored state with the application of 2 V (curve b) and 3 V (curve c). (B) Coloration and bleaching cycles of the ECD recorded at 801 nm at different temperatures: (a) 25 C, (b) 40 C, and (c) 60 C. (C) Coloration and bleaching cycles of the ECD recorded at 850 nm at different relative humidity levels: (a) 25% and (b) 70%.

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Figure 3.34 (A) Transmittance spectra of an all-solid monolithic electrochromic device (ECD) with an MgF2 protective coating in the as-fabricated state before (curve a) and after (curve b) the application of the coating and in the colored state (curve c). (B) Coloration and bleaching cycles of the ECD with an MgF2 protective coating measured at a wavelength of 801 nm with an applied voltage of 3 V. (From P.V. Ashrit, F.E. Girouard, V.-V. Truong, Solid State Ionics 89 (1996) 65e73 .)

from all of the solid ionic layers employed in the system, but principally from the ion-conducting layer of LiBO2. The same authors [58] further improved the all-solid monolithic system with the materials configuration glass/ITO/LixV2O5/LiBO2eLiF/WO3/ ITO/MgF2. This improvement was undertaken in view of the observation of the earlier discussed system’s time-dependent degradation 3 months after its fabrication. Addition of an outer MgF2 layer as a protective coating was undertaken. To achieve a high-quality MgF2 thin film on the ECD, room temperature sputter deposition of MgF2 with simultaneous bombardment by fast argon atoms by the FAB gun was carried out. The improved configuration also employed a LiBO2eLiF composite film as IC because of its superior optical and electrical performance. Shown in Fig. 3.34 is the performance of this improved all-solid monolithic ECD. In Fig. 3.34A are shown transmittance spectra of the device in the as-fabricated state before and after the application of the MgF2 protective overcoating as well as in the colored state of the device with the application of 3 V. An excellent EC modulation as seen in this figure can be achieved in the visible and near-infrared regions of the spectra. The integrated transmittance values are as follows: Tsolar (bleached) ¼ 51.6%, Tphotopic (bleached) ¼ 59.6% Tsolar (colored) ¼ 13.4%, Tphotopic (colored) ¼ 21.2%

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Figure 3.35 (A) Schematic of an all-solid monolithic electrochromic device (ECD) for emittance control. (B) Infrared reflectance modulation behavior of the ECD under the application of 2 V. a-, amorphous; c-, crystalline.

Once again it can be seen that these optical values are close to the optical criteria established for the smart window application of these devices. The coloration and bleaching cycles of this device are shown in Fig. 3.34B. These cycles are quite rapid in comparison with the earlier shown (Fig. 3.33B) cycles of the unprotected ECD operating at room temperature (25 C). Many all-solid monolithic ECDs with reflectance modulation capability are also built for emittance control in the far-infrared (2e40 mm) spectral region. For example, Franke et al. [85] have used amorphous WO3 (a-WO3) and crystalline WO3 (c-WO3) in tandem in the capacity of IS and EC layers, respectively, in their device as shown in Fig. 3.35A. A thin film of tantalum pentoxide (Ta2O5) with a high infrared (IR) transmittance was employed as the ion-conducting layer. The optical modulation under lithium ion injection in a-WO3 is predominantly at lower wavelengths (visible and near-IR), being via absorption modulation (intervalence transfer of electrons) with no effect on its high IR transmittance. Hence, it can be employed as an IS layer with no consequence in the far-IR region. The EC optical modulation in c-WO3, on the other hand, being primarily via reflectance, that occurs in these films manifests in the IR region. Hence, the c-WO3 film can be employed as the base EC layer with efficient reflectance modulation under the application of an electric field, as shown in Fig. 3.35B. The electrodes employed in this case are an aluminum film as the inner electrode with high IR reflectance and an aluminum grid film as

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the outer electrode with a very high IR transmittance. The IR reflectance modulation essentially happens in the c-WO3 layer with the double injection and extraction of lithium ions and electrons under the application of an electric field. As can be seen in Fig. 3.35B an effective emittance modulation occurs through most of the spectral region studied. A 20% emissivity change is reported by the authors. The coloration and bleaching cycles of this all-solid monolithic system, however, were quite slow, being around 30 min for each cycle.

3.4.3 Electrochromic Devices With Gel Electrolytes A wide variety of sandwich-structure ECDs incorporating gel electrolytes are proposed. Such systems are the most convenient to fabricate and operate. The gel electrolytes, also amenable to semisolid EC systems, present many advantages over the liquid electrolyte-based EC systems as well as the all-solid monolithic systems. As highlighted earlier, the all-solid monolithic systems are marred by quite a few challenges regarding the deposition and optimization of subsequent layers on the same substrate. The coloration and bleaching cycles of all-solid systems are invariably slow owing to the inherent solid-state low ionic conductivity of the solid IC used. Liquid electrolyte-based ECDs also present quite a few challenges concerning (1) leaking of the electrolyte, (2) proper sealing of the device, (3) chemical stability of the liquid and the layers in the presence of the liquid, (4) upscalability, and more. Hence, devices based on gel electrolytes have attracted a lot of attention. They are easy to fabricate because of their sandwich structure and the proposed electrolytes present a very high ionic conductivity. An example of this is an ECD proposed by Garino et al. [86], exhibiting a very fast coloration and bleaching response. This system, as shown in Fig. 3.36A, incorporates an electrochemically deposited V2O5 and a sputter-deposited WO3 film as anode and cathode, respectively. Each of these layers was prepared on separate ITO-coated substrates. The sandwich-structure ECD was fabricated by putting these substrates together with a gel polymer electrolyte based on Bisphenol A ethoxylate dimethacrylate and PEG methyl ether methacrylate (PEGMA). In addition to using PEGMA as a reactive diluent the polymer was mixed with bis(trifluoromethane) sulfonamide lithium salt as a source of Liþ and ethylene carbonateediethylcarbonate as an organic plasticizer. This combination gel polymer also acted as the spacer between the two electrodes as well as the glue holding the two electrodes together. The optical performance of the

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Figure 3.36 (A) Schematic of a quasi-solid-state polymeric gel-based electrochromic device (ECD). (B) Transmittance spectra of the ECD with applied potential recorded at a wavelength of 700 nm. (C) Coloration and bleaching cycles of the ECD with the application of 2 V. ITO, indium tin oxide; PET, poly(ethylene terephthalate).

ECD is shown in Fig. 3.36B. A fairly good optical modulation is seen as a function of increasing applied potential. The bulk of the optical modulation manifests in the near-IR region. Such ECDs may be interesting from the thermal management of solar energy point of view as they maintain a relatively high visible transmittance and show a good optical modulation in the near-IR region. The optimized gel polymer used exhibits an ionic conductivity comparable to that of liquid electrolytes and is expected to be electrochemically stable. The authors report a visible coloration response of a few seconds. In Fig. 3.36C are shown the coloration and bleaching cycles of the ECD recorded at a wavelength of 700 nm. As can be seen from this

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figure very fast coloration (2 s) and even faster bleaching (1 s) occur with the application of 2 V with respect to the V2O5 electrode. The ECDs were reported to have a good optical memory and a good reversibility after 1000 cycles of coloration and bleaching. A similar approach to fabricating all-solid polymer electrolytes was taken by Guofa et al. [87], by using a gel turned solid electrolyte containing a mixture of polymethylmethacrylate (PMMA), polycarbonate, and gel turned solid electrolyte infused with PC and lithium perchlorate (LiClO4). Inkjet-printed nanoparticulate electrodes of nickel oxide (NiO) and WO3 were used as complementary layers in the sandwich EC structures. A comparative study of the devices with and without the NiO layer was reported. In Fig. 3.37A are given the SEM images of the nanoparticulate thin films. A porous but continuous structure that is very important for the double injection/extraction into the active electrodes is seen in these films. A schematic of the ECD formed by sandwiching the inkjet-printed WO3 on one ITO-coated substrate and V2O5 on the other using the polymer electrolyte is shown in Fig. 3.37B. The structure was assembled with the electrolyte in the gel form and eventually transformed into solid electrolyte by air drying at room temperature. The optical switching performance of the ECD with and without the NiO electrode is shown Fig. 3.37C. The inset photos in this figure show the degree of coloration in the two cases. The device without the NiO film shows a small optical modulation in the higher wavelength region. The inclusion of NiO film in the device reduces the transmittance of the device throughout the spectral region studied but leads to a significant optical modulation. As can be seen in the photos, visibly the coloration of the device is quite intense compared to the device without the NiO film. The variation in the optical density (DOD) as a function of the inserted charge is shown in Fig. 3.37D. In the ECD without the IS layer of NiO, the DOD saturates after an insertion of around 1 mC/ cm2, while the device with the NiO film shows a linear increase in DOD up to an insertion of 10 mC/cm2 studied, corresponding to a CE of 136.7 cm2/C. Hence, a symmetric ECD containing the complementary electrodes of nanoparticulate NiO and WO3 is superior in all respects. PMMA electrolyte-based ECDs have also been successfully fabricated for use as dynamically tintable visors for helmets [88]. In these sandwich structures the electrode films were deposited on two polyester substrates. The substrates consisted of ITO-coated polyester (CP Films, type OC 80/ 7) on which reactively sputter-deposited WO3 film was the base EC layer

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Figure 3.37 (A) Scanning electron microscopy images of inkjet-printed NiO (left) and WO3 (right) films. (B) Schematic of a polymethylmethacrylate-based electrochromic (EC) system with WO3 as the EC layer and NiO as the ions storage layer. (C) Transmittance spectra of the EC device with and without the NiO film in the bleached and colored states. (D) Comparison of the variation in optical density (DOD) with inserted charge density in an EC device with and without the NiO film. ITO, indium tin oxide. (Adapted from G. Cai, et al., Nanoscale 8 (2016) 348e357.)

and NieV oxide was the counterelectrode. The blue WO3 films were sputter deposited in the presence of a gas mixture of argon/oxygen/ hydrogen in the ratio of 4:2:13 [89]. The charge insertion in the initially transparent NieV oxide film was carried out by a novel method in which the films were exposed to ozone originating from UV irradiation in the presence of oxygen [90]. The two polyester sheets, each containing the pertinent electrodes, were laminated together using the PMMA electrolyte

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formulated in a special way [91] by roll pressing at a temperature of 80 C and edge sealed with epoxy. The coloration and bleaching of such devices is found to be around 100 s per cycle and the extent of coloration is reported to depend on the wavelength. The device seems to exhibit more intense coloration in the higher visible and near-IR wavelengths with a DT of around 50% as seen in this figure. The reversibility of the coloration and bleaching cycles is excellent for this device. No significant change in the device’s switching behavior is seen, even after several thousand cycles. In addition, the device is reported to exhibit an excellent optical memory, retaining its coloration in open circuit for several hours. In terms of its virtues related to switching stability, CE, optical memory, and facile and economic fabrication this device seems to be quite apt for many applications. The company ChromoGenics in Sweden is fabricating visors for motorcycle helmets based on these flexible foil-based ECDs [92,93]. Other rigid or glass-based EC technologies [94,95] that are commercially successful are expensive and have relatively limited scope of applications and have not made inroads into all types of window applications. The foil-based approach for ECDs is thus touted as the best path to large-scale commercialization of EC technology.

3.5 SUMMARY TMO thin films exhibit one of the most convenient forms of EC coloration, i.e., reversible coloration under the application of a small electric field. Unlike in the case of devices based on liquid crystals (LCDs), which undergo an oneoff switching, the EC coloration in TMO thin films can be gradual and can be made wavelength selective. Further, unlike LCDs, large voltages are not required to maintain the colored state of the TMO-based ECDs, as the latter display a good optical memory. The reversible and stable coloration of the TMOs is rooted in their metastability in multiple oxidation states. Hence, by a small electric field activation these materials can be switched between their oxidation states with large reversible changes in their optical and electrical properties. Tungsten trioxide (WO3) is by far the most studied TMO owing to its number of virtues that are important from an application point of view. These include its coloration in both amorphous and polycrystalline states exhibiting, respectively, absorption and reflectance modulation. The coloration in both cases takes place in the solar spectral region, making it the most sought-after material for energy-

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related applications. Molybdenum oxide (MoO3) films show nearly the same qualities as WO3 films, but somehow the work carried out on this material with regard to energy-related devices and application is not on par with that of WO3 films. Hence, a tremendous opportunity exists to explore this material from the energy application point of view. While both MoO3 and WO3 exhibit cathodic coloration, i.e., transformation from transparent to colored state under the application of an electric field, TMOs such as vanadium pentoxide (V2O5) and nickel oxide (NiO) exhibit anodic coloration, i.e., transformation from an initial colored state to a transparent state under the application of an electric field. In both types of coloration, the color changes under the application of an electric field ensue because of a double (ions and electrons) injection into or extraction from these materials. The EC coloration depends on the transport of these species. In the thin form one can affect the transport mechanism by tailoring the nanostructure of the TMO thin films. An enormous amount of research and development work is being carried out worldwide to improve the EC coloration of the various TMO thin films by nanostructure control. Most of the work on nanostructured TMO films as of this writing is concentrated on improving the optical performance of the films, and excellent results have been achieved. However, the coloration mechanism is invariably slow in such reported works. The EC mechanism being inherently rooted in both the optical and the electrical (transport) properties, it is essential to study simultaneously the effect of nanostructuring on both of these aspects. The effect of nanostructuring on the transport of both the electrons and the ions needs to be studied carefully and optimized for both. Lots of opportunities exist for each of the TMO thin films for such an in-depth study of the correlation between the nanostructure effect and the CE as well as the ion and electronic conductivity. The technique of chronoabsorptometry used by McEvoy et al. [76], in which the TMO coloration can be mapped out by time-lapse optical microscopy with step potential electrochemical insertion, can be very powerful for studying the EC phenomenon. The method of dry lithiation for studying individual films and for device fabrication can be very convenient in EC-related work [12e14]. On the theoretical side, of the many models suggested, three are most commonly used. These are the intervalence transfer of electrons and the large and small polaron models commonly applicable to amorphous films exhibiting absorption modulation. The Drude model, applicable to metals, and its variation, applicable to heavily doped semiconductors, are commonly used

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for explaining the dynamic optical behavior of polycrystalline films. It seems that under general conditions of TMO thin film preparation the films are invariably made up of crystals embedded in a host amorphous medium. Hence, a good approach would be to treat the films as a TMO composite amorphousepolycrystalline phase wherein the double insertion leads to a simultaneous evolution of both absorption and reflectance modulation manifesting in different spectral regions. This makes for a wideband EC coloration in the visible and near-IR region. A comprehensive model that not only takes into account the combined optical (absorptionereflectance) modulation but also the respective components associated with each type is highly desirable to account for all the optical changes occurring. This needs to be handled with the effective medium theories that are available. A very good opportunity exists for the development of such a comprehensive model. In experimental work such an evolution with heat treatment has been seen in the case of WO3 films. An initially a-WO3 film exhibiting purely absorption modulation gradually exhibits a combination of absorptionereflectance modulation with heat treatment. A precise accounting of the optical changes occurring in nanostructured films in various spectral regions (visible to near-IR) can be even more complicated. Various configurations can be envisioned for ECD fabrication, from monolithic to sandwich structures encompassing liquid, gel, or solid components in conjunction with the suitable TMO layer. A combination of a cathodically coloring EC layer such as WO3 and an anodically coloring counterelectrode layer is the most common and convenient form of fabrication of devices. Despite the large volume of work carried out in this field, a lot of challenges need to be overcome to make EC applications a reality on a large scale. This also provides plenty of opportunity for further improvement of existing materials and their study from both scientific and industrial research points of view.

REFERENCES [1] Large area chromogenics: materials and devices for transmission control, in: C.M. Lampert, C.G. Granqvist (Eds.), SPIE Institute Series, vol. IS4, 1990. Bellingham. [2] A. Hagfeldt, N. Viachopolous, M. Graetzel, Journal of Electrochemical Society 141, L82 (1994) 42e44. [3] W.J. Macklin, R.J. Neat, Solid State Ionics 694 (1992) 53e56. [4] R. Nakajima, Y. Yamada, T. Komastsu, K. Murashiro, T. Saji, K. Hoshino, RCS Advances 2 (2012) 4377e4381.

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[5] V.-V. Truong, F.E. Girouard, P.V. Ashrit, Inorganic Ion Conductor, Part 8, pp. 386e401 in Ref. [1]. [6] C. Ho, I.D. Raistrick, R.A. Huggins, Journal of Electrochemical Society 127 (2) (1980) 343e350. [7] Yoon-Chae, et al., Journal of Electrochemical Society 152 (2) (2005) H201eH204. [8] L. Su, J. Winninck, P. Kohl, Journal of Power Sources 101 (2001) 226e230. [9] M. Balkanski, Microionics e Solid State Integrable Batteries, Elsevier Science Publishers, B.V., 1991. [10] R.A. Huggins, Ionics 8 (2002) 300e313. [11] A.J. Bard, L.R. Faulkner, Electrochemical Methods e Fundamentals and Applications, second ed., Wiley, 2001. [12] P.V. Ashrit, G. Bader, F.E. Giroouard, V.-V. Truong, Proceedings of the SPIE 1401 (1990) 119e129. [13] T. Yoshimura, M. Watanabe, Y. Koike, M. Tanaka, Japanese Journal of Applied Physics 22 (1983) 157. [14] Y. Young, Z. Jiayu, G. Peifu, T. Jinfa, Solar Energy Materials & Solar Cells 46 (1997) 349e355. [15] http://www.lambdaphoto.co.uk/pdfs/Inrad_datasheet_LNB.pdf. [16] P.V. Ashrit, G. Bader, F.E. Girouard, V.-V. Truong, Proceedings of the SPIE 1149 (1989) 7. [17] B. Abdel Samad, P.V. Ashrit, Solid State Ionics 283 (2015) 52e55. [18] B.A. Korgel, Nature 500 (2013) 278e279. [19] http://sageglass.com/portfolio/. [20] F. Reif, Fundamentals of Statistical and Thermal Physics, McGraw Hill, 1965. [21] K.V. Madhuri, G. Bader, P.V. Ashrit, International Journal of Applied Engineering Research 9 (2014) 2061e2073. [22] G. Bader, P.V. Ashrit, V.-V. Truong, Applied Optics 37 (1998) 1146e1151. [23] D.J. Taylor, P.F. Freig, S.L. Hietala, Thin Solid Films 332 (1998) 257e261. [24] B. Abdel Samad, J. Thibbodeau, P.V. Ashrit, Applied Surface Science 350 (2015) 94e99. [25] S.K. Deb, Applied Optics 8 (1969) 192e195. [26] B.W. Faughnan, R.S. Crandall, P.M. Heyman, RCA Review 36 (1975) 177e197. [27] C.G. Granqvist, Handbook of Inorganic Electrochromic Materials, second ed., Elsevier, 2002. [28] P.V. Ashrit, Thin Solid Films 385 (2001) 81e88. [29] Se-Hee, et al., Advanced Materials 18 (2006) 763e766. [30] J.L. Solis, A. Hoel, V. Lantto, C.G. Granqvist, Journal of Applied Physics 89 (2001) 2727e2732. [31] E.A. Meulenkamp, Journal of Electrochemistry 144 (1997) 1664. [32] P.K. Shen, A.C.C. Tseung, Journal of Materials Chemistry 2 (1992) 1141. [33] B. Yang, H. Li, M. Blackford, V. Luca, Current Applied Physics 6 (2006) 436e439. [34] G. Leftheriotis, P. Yianoulis, Solid State Ionics 179 (2008) 2192e2197. [35] A. Bessiere, et al., Journal of Applied Physics 91 (2002) 1589e1594. [36] Z. Xie, et al., Journal of Materials Chemistry 22 (2012) 19904e19910. [37] X.L. Wei, P.K. Shen, Electrochemical Communications 8 (2006) 293e298. [38] P. Judeinstein, J. Livage, Journal de chimie physique 90 (1993) 1137e1147. [39] L. Liang, et al., Scientific Reports 3 (2013) 1e8, 1936. [40] A.E. Aliev, H.W. Shin, Solid State Ionics 154-155 (2002) 425e431. [41] S. Badilescu, P.V. Ashrit, Solid State Ionics 158 (2003) 187e197. [42] C. Santato, M. Odziemkowski, M. Ulmann, J. Augustynski, Journal of American Chemical Society 123 (2001) 10639e10649.

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[43] M. Deepa, et al., Acta Materialia 55 (2007) 6095e6107. [44] N.N. Dinh, D.H. Ninh, T.T. Thao, V.-V. Truong, Journal of Nanomaterials 2012 (2012) 7, 781236. [45] J. Wang, E. Khoo, P.S. Lee, J. Ma, Journal of Physics Chemistry C112 (2008) 14306e14312. [46] Z. Jiao, et al., Journal of Physics 43 (2010) 285501. [47] G. Beydaghyan, G. Bader, P.V. Ashrit, Thin Solid Films 516 (2008) 1646e1650. [48] G. Beydagyhan, J.-L.M. Renaud, G. Bader, P.V. Ashrit, Journal of Materials Research 23 (1) (2008) 274e280. [49] B. Baloukas, J.-M. Lamarre, L. Martinu, Solar Energy Materials and Solar Cells 95 (2011) 807e815. [50] L.F. Reyes, et al., Journal of the European Ceramic Society 24 (2004) 1415e1419. [51] T. Takahashi, et al., Thin Solid Films 515 (2007) 6567e6571. [52] H.A. Macleod, Thin Film Optical Filters, Institute of Physics Publishing, London, 2001. [53] G. Beydaghyan, et al., Applied Physics Letters 97 (2010) 163104. [54] H. Wei, et al., Energy Environment Focus 2 (1) (2013) 112. [55] R. Mukherjee, P.P. Sahay, Journal of Materials Science: Materials Electronics 26 (2015) 6293e6305. [56] B. Abdel Samad, P.V. Ashrit, 2017 (to be published work). [57] P.V. Ashrit, et al., Solid State Ionics 59 (1993) 47e57. [58] P.V. Ashrit, F.E. Girouard, V.-V. Truong, Solid State Ionics 89 (1996) 65e73. [59] S.T. Lutta, H. Dong, P.Y. Zavaliz, M.S. Whittingham, Materials Research Bulletin 40 (2005) 383e393. [60] N.A. Chernova, M. Roppollo, A.C. Dillon, M.S. Whittingham, Journal of Materials Chemistry 19 (2009) 2526e2552. [61] C. Delmas, et al., Solid State Ionics 69 (1994) 257. [62] C. Leger, S. Bach, P. Soudan, J.-P. Pereira-Ramos, Journal of Electrochemical Society 152 (2005) A236. [63] M. Najdoski, V. Koleva, A. Samet, Journal of Physical Chemistry C 118 (2014) 9636e9646. [64] Y. Wang, G. Cao, Ekectrochimica Acta 51 (2006) 4865e4872. [65] R. Balu, P.V. Ashrit 2008 (unpublished work). [66] G. Wu, et al., Thin Solid Films 485 (2005) 284e289. [67] R. Baddour-Hadjean, V. Golabkan, J.-P. Pereira-Ramos, A. Montoux, D. Lincot, Journal of Raman Spectroscopy 33 (2002) 631. [68] T. He, et al., Langmuir 17 (2001) 8024. [69] A. Taj, P.V. Ashrit, Journal of Materials Science 39 (2004) 3541e3544. [70] A.C. Dillon, et al., Thin Solid Films 516 (2008) 794e797. [71] C. Julien, A. Khelfa, J.P. Guesdon, A. Gorenstein, Applied Physics A59 (1994) 173e178. [72] M. Winter, J.O. Besenhard, M.E. Sphar, P. Novak, Advanced Materials 10 (1998) 725e763. [73] G. Beydaghyan, M. Boudreau, P.V. Ashrit, Journal of Materials Research 26 (1) (2011) 56e61. [74] R.S. Patil, M.D. Uplane, P.S. Patil, International Journal of Electrochemical Science 3 (2008) 259e265. [75] T.M. McEvoy, K.J. Stevenson, J.T. Hupp, X. Dang, Langmuir 19 (2003) 4316e4326. [76] T.M. McEvoy, K.J. Stevenson, Journal of American Chemical Society 125 (2003) 8438e8439. [77] Y.-S. Lin, T.-H. Tsai, S.-W. Tien, Thin Solid Films 520 (2013) 248e252.

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[78] S. Papaefthimiou, G. Leftheriotis, P. Yianoulis, Thin Solid Films 343-344 (1999) 183e186. [79] D. Benson, R. Crandall, S.K. Deb, J.L. Stone, Stand-Alone Photovoltaic Powered Electrochromic Windows (US Patent 5384653), January 24, 1995. [80] M.-C. Yang, H.-W. Chao, J.-J. Wu, Nanoscale 6 (2014) 9541e9544. [81] P. Bonhôte, E. Gogniat, M. Grätzel, P.V. Ashrit, Thin Solid Films 350 (1999) 269e275. [82] P. Bonhôte, et al., Journal of Physical Chemistry B 102 (1998) 1498e1507. [83] V. Thi Phat Minh, et al., in: Proc. Window Innovations’ 95, June 5 and 6, 1995, pp. 463e470. Toronto. [84] S.E. Selkowitz and C.M. Lampert, Application of Large-Area Chromogenics to Architectural Glazings, Part 2, pp. 22e45 in Ref. [1]. [85] E.B. Franke, et al., Applied Physics Letters 77 (7) (2000) 930e932. [86] N. Garino, et al., International Journal of Electrochemistry (2013) 1e10, 138753. [87] G. Cai, et al., Nanoscale 8 (2016) 348e357. [88] A. Azens, G. Gustavsson, R. Karmhag, C.G. Granqvist, Solid State Ionics 165 (2003) 1e5. [89] A.P. Giri, R. Messier, Materials Society Symposium Proceedings 24 (1984) 221. [90] A. Azens, L. Kullman, C.G. Granqvist, Solar Energy Materials & Solar Cells 76 (2003) 147e153. [91] J. Stevens, (U.S. Patent 528843). [92] www.chromogenics.se. [93] C.G. Granqvist, Solar Energy Materials & Solar Cells 92 (2008) 203e208. [94] https://www.sageglass.com/. [95] http://www.saint-gobain-sekurit.com/.

FURTHER READING [1] G.A. Niklasson, C.G. Granqvist, Journal of Materials Chemistry 17 (2007) 127e156.

CHAPTER 4

Thermochromic Thin Films and Devices 4.1 BASICS OF THERMOCHROMIC MECHANISM As discussed in Chapter 1, thermochromic materials are those that exhibit a reversible change in optical properties in certain spectral regions as a function of temperature. Quite a few materials are capable of undergoing such a reversible change. Similar to the electrochromic materials, the thermochromic materials are also very important from the point of view of applications because of their interactive nature dependent on the ambient temperature that they are subjected to. They are also important from the fundamental research point of view because of the reversible physical and chemical changes occurring in them with temperature. Almost all the materials undergo some form of thermochromism in a certain range of temperatures in the sense that they undergo some continuous and reversible optical change when heated. This is the phenomenon of incandescence occurring because of the excitation of the electrons in the atoms that absorb the thermal energy and reradiate in the form of light. However, we are interested here more precisely in the form of thermochromism that is of the first order, i.e., in which an abrupt and discontinuous phase change occurs in certain materials at a specific critical temperature called the transition temperature (Tt). There are a wide variety of organic and inorganic materials exhibiting this type of abrupt thermochromism [1]. In view of their pertinence to this book, only thermochromic properties of the inorganic transition metal oxide (TMO) thin films are discussed here. Among the TMOs exhibiting thermochromism, such as titanium trioxide (Ti2O3) [2] and vanadium oxides (V2O3, V2O5, VO, and VO2) [3], it is the latter oxides that have attracted enormous research and application interest owing to the efficient form of thermochromic switching they exhibit at different critical temperatures originating from their transition from a semiconductor to a metallic state. Among the vanadium oxide series it is vanadium dioxide (VO2) that has attracted an exorbitant amount of interest because of the proximity of its Tt to room temperature and the possibility of varying this Transition Metal Oxide Thin FilmeBased Chromogenics and Devices ISBN 978-0-08-101747-0 http://dx.doi.org/10.1016/B978-0-08-101747-0.00004-0

© 2017 Elsevier Ltd. All rights reserved.

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Figure 4.1 Phase transition temperatures of different vanadium oxides. (From C.H. Griffiths, H.K. Eastwood, Journal of Applied Physics 45 (1974) 2201e2206.)

temperature using suitable dopants [4]. By preparing VO2 in thin film and nanostructured forms as well as doping these with the right dopants, a wide variety of optical and electrical switching at various temperatures can be induced in this material [5e7]. This wide range of possibilities has opened up enormously the scope for research and application of this material. Vanadium being a transition metal exhibits multiple oxidation states from þII to þIV, thus forming a wide range of oxides as shown earlier in Chapter 2 (Fig. 2.18). It is quite clear from this figure that VO2 forms in a very narrow and strict range of parameters. In Fig. 4.1 are shown the Tt’s of these various vanadium oxides with respect to room and liquid nitrogen temperatures [8].

4.1.1 Phase Transition The phase transition between the insulator and the metallic states at the Tt is the most important parameter that characterizes VO2. This phase transition is followed by abrupt changes in both optical and electrical properties of this material. Hence an enormous amount of experimental and theoretical research is devoted to the study of the phase transition ever since the discovery of this effect by Morin in 1952 [3]. VO2 was found to show a resistivity decrease of three orders of magnitude at the Tt and a positive temperature coefficient of resistivity above this Tt, indicating a metallic

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behavior. Subsequently a structural distortion was also found to exist in this material when this Tt was crossed [9,10]. Then the existence of a narrower bandgap was proposed by Adler et al. [11], along with the discovery of the existence of a bandgap and the variation of Tt by applying a stress along the c axis of the VO2 crystal. Following this, it was found that the metal-toinsulator transition (MIT) is a first order transition and that the Tt depends linearly on the applied hydrostatic pressure [12]. The historical developments related to vanadium oxides and, in particular, VO2 structure and phase transition are given in Table 4.1 [13]. The basic mechanism underlying the thermochromic switch occurring in certain TMOs was initially named an MIT because of a sudden and high order of change of resistance observed at a certain critical Tt [14]. However, this was later renamed more appropriately as a metal-to-semiconductor transition [15] following the magnitude of the resistance change that occurs in these materials. For single crystals of vanadium dioxide (VO2), for example, the Tt value is 68 C. As already discussed in Chapter 2, the crystal structure of pure VO2 changes from a room temperature monoclinic phase Table 4.1 Historical development of the phase transition in vanadium oxides Year Important milestone References

1952 1955 1956 1959 1960

1965 1967 1969 1970 1971 1972 1975

Observation of an MIT in V2O3 Proposition of tetragonal rutile-type structure in VO2 Discovery of a distortion in the monoclinic structure of VO2 crystal Observation of MIT in VO and VO2 Outlining of the important of VeV interaction in VO2 structure and proposition of semiconductor-to-metal transition Presence of a covalent VeV bond; suggestion of a one-electron model for insulator phase Introduction of a bandgap due to crystallographic distortions Role of phonons in stabilization of the metallic state Role of strong electronic interaction in MIT Proposition of VO2 band structure above and below the transition temperature Discovery of the M2 metastable state of VO2 VO2 (M2) metastable state stabilization

MIT, metal-to-insulator transition.

Menth et al. [16] Magneli and Andersson [17] Andersson [10] Morin [3] Goodenough [9]

Umeda et al. [18] Adler et al. [11] Berglund et al. [12] Rice et al. [19] Goodenough [6] Marezio et al. [20] Pouget et al. [21]

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M1 to a high-temperature (>Tt) tetragonal phase. Despite the vast amount of work on this material to explain clearly the underlying phenomenon of this transition, a comprehensive theoretical model seems to be obscure so far. This is due to the fact that both the structural and the phase (metal to semiconductor) changes occur simultaneously. Added to this, there exists another monoclinic insulator phase referred to as the M2 in VO2, which is obtained by a slight distortion to the M1 phase of pure VO2 as shown earlier in Fig. 2.21. These three phases, the monoclinic M1, monoclinic M2, and tetragonal phase, are structurally different but exhibit comparable changes in properties (optical and electrical) with these transitions occurring with the same activation energy [7]. The various experimental results pertaining to the transition are debated mainly in the light of two principal models, the Peierls insulator model and the Mott insulator model. According to the Peierls model discussed in Chapter 2, one-dimensional solids show insulating behavior due to the periodic distortion of the lattice. In the case of VO2 the monoclinic M1 and M2 phases undergo this distortion owing to the zigzag structure (dimerization) formed at low temperature (T < Tt). The resulting unit cell is doubled, forming a bandgap. At high temperature, however, the VO2 crystals form regular chains between the V atoms leading to a high conductivity in this tetragonal rutile-like structure as shown in Fig. 2.22A. According to the MotteHubbard model as discussed in Chapter 2, the MIT is considered in the light of the electronic structure of VO2 as a function of the temperature. Mott initially considered the discontinuous phase transition in terms of a simple cubic single-electron atom lattice model with a lattice constant a and examined the variation of the free electron density with the variation of a [22]. This variation of free electron density with a to bring about the insulator-to-metal transition can arrive in three different ways as shown in Fig. 4.2 [23]. In the first case a gradual increase in free electron density occurs as the lattice constant a is decreased gradually. One cannot locate the point of transformation between the metallic and the insulator phase. A second situation is where a threshold value of a is needed before the onset of free electron density (N) increase. The change from this point onward of N is again continuous with no specific point of transformation. It is the third situation showing the first order (discontinuous) change in N that is referred to as the Mott transition and is pertinent to the MIT in VO2. As per this Mott model, at large values of a, the solid is essentially an insulator as the electrons were localized on the atom site in the absence of

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Figure 4.2 Variation of the free electron density (N) as a function of the inverse lattice parameter (1/a) of a crystalline solid. (Line 1) A gradual change from insulator to conductor. (Line 2) Onset of a gradual change at a threshold value of the lattice constant. (Line 3) Sudden change from insulator to metal at a critical value. (From N.F. Mott, Metal-insulator Transitions, Taylor & Francis Ltd., London, Harper & Row Publishers Inc.,/Barnes & Noble Import Division, New York, 1974.)

large activation energy (U ) required to overcome the Coulomb repulsion from neighboring site electrons. However, when a is sufficiently small to bring about an overlap of neighboring electron states the solid exhibits a metallic behavior. This simple model in which electroneelectron interaction is not considered does not clearly account for the systematic effect of temperature in bringing about the transition. The applicability of the Mott transition to VO2 can also be argued because the simple model does not explicitly require a crystalline phase, i.e., phase transition can occur even in amorphous VO2. Experimentally this is a very rare occurrence as seen in many works related to VO2 thin films prepared by various techniques. However, nearly amorphous VO2 thin films prepared under some special conditions show a transition to a metallic state [24]. Hubbard considered this model in terms of the exchange interaction between the single electrons of the solid [25,26]. As discussed in Chapter 2, the TMOs are characterized by a narrow d band of bandwidth W. Hubbard’s arduous mathematical treatment can be conceptually simplified by considering two extreme regimes: (1) in the situation in which the activation energy U needed to transfer an electron from one atomic site to an equivalent site of another atom is much higher than the d bandwidth,

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W (W > U ), the solid behaves as a conductor with a high free electron density. The most interesting situation occurs when the two energies are of the same order, W w U, in which the electrons become strongly correlated and phase transition can occur under the right conditions. This is the essential feature of TMOs such as VO2 in which localized and delocalized electrons can coexist. Hence, the thermochromic behavior of this material is better described in terms of strongly correlated electrons than in terms of single electron behavior. As commented by Jorgenson et al. [27], in the treatment of experimental results in the light of the various approaches, one has to choose the theory that is more appropriate for the experimental situation used. This is especially important in the discussion of the thermochromic behavior of VO2 thin films as a wide range of deposition parameters exist that impart different micro- and nanostructures to these films. A distinct difference in the treatment of VO2 thin films as far as the phase transition is concerned is that it can be viewed as a percolation phenomenon as confirmed experimentally [28,29]. In the vicinity of the Tt, VO2 is expected to be made up of metallic and insulator (semiconductor) components and hence its behavior can be theoretically treated using the Bruggeman effective medium theory [30]. Here, one arrives at the macroscopic behavior of the composite medium considered as a simple mixture in terms of the properties of the microscopic components forming the medium. The Bruggeman effective medium approximation (EMA) can be given as [31]: εm  εeff ε  εeff ð1  f Þ þf ¼0 (4.1) εm þ 2εeff ε þ 2εeff where f refers to the filling factor of one of the components with its dielectric constant ε in the matrix of a second component with its dielectric constant εm filling the rest of the medium. One arrives at the effective dielectric constant, εeff, using Eq. (4.1). In the percolation treatment pertaining to the phase transition of VO2 the transition is referred to as the percolation transition. This phase transition is shown in Fig. 4.3A in terms of the resistance change occurring in VO2 in the vicinity of the Tt. As per the percolation treatment the filling factor of the metallic phase VO2, initially small or nonexistent, increases with temperature around Tt. The

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Figure 4.3 (A) Change in resistance of VO2 as a function of temperature around the transition temperature: coexistence of metallic and dielectric VO2 phases. (B) Schematic of the development of the VO2 metallic phase as a function of temperature and the percolation. (From C.O.F. Ba, et al., Journal of Vacuum Science & Technology A, 34 (3) (2016) 031505e1-8.)

metallic component in the VO2 system increases slowly initially and increases rapidly upon reaching the percolation threshold. Below the percolation threshold the VO2 metallic inclusions are formed as individual entities without connectivity between them. Upon reaching the percolation threshold this connectivity is established as reflected by the sudden dip in the resistance shown in Fig. 4.3A. The schematic of this change occurring with temperature is shown in Fig. 4.3B. A direct observation of this transition has been made using scanning near-field infrared microscopy where the gradual development of the metallic VO2 phase with temperature in the vicinity of the Tt takes place [29]. This is shown in Fig. 4.4 in which the initial development of metallic puddles on a nanoscale as an extremely sensitive function of temperature

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341K

342.4K

342.6K

342.8K

343K

343.6K

Figure 4.4 Scanning near-field infrared micrographs of VO2 in the close vicinity of the transition temperature showing the percolation in progress (image area 4 mm  4 mm in all cases). (From M.M. Kazilbash, et al., Science 318 (2007) 1750e1753.)

and the sudden extension of this to nearly the entire sample are quite evident. The filling factor that corresponds to the percolation threshold is around 0.5 [32,33]. Budai et al. [34] took an important and new thermodynamic approach to the MIT of VO2 following the very early proposition by Hearne [35] based on electronephonon interaction. According to this innovative

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work [35], the dielectric state of VO2 is due to the d band splitting and uncrossing brought about by the distortion. At high temperature a strong electronephonon interaction is expected to result in phonon softening raising the vibrational entropy of VO2 (rutile). This increased entropy was deduced to be the driving force for the phase transition. Based on their quantitative experimental and theoretical work, Budai et al. [34] concluded that the MIT in VO2 is indeed driven by the large phonon entropy in the hightemperature rutile state. As stated earlier, in the two most important approaches taken to explain the MIT in VO2, the electron-lattice instability-driven Peierls model and the electroneelectron correlation-based MotteHubbard model, the aspect of lattice vibrations (phonons) is not taken into account. Also not considered are the energy and entropy changes occurring with the transition. By calculating the phonon dispersion in the high-temperature rutile state and the low-temperature monoclinic state of VO2, it is demonstrated that the MIT is predominantly a phenomenon driven by strongly anharmonic phonons rather than by electrons. It is this large vibrational entropy of the rutile state above the Tt that stabilizes it. The authors examined experimentally the temperature dependence of the phonon density of states (PDOS). As shown in Fig. 4.5A (a), the PODS spectrum changes above and below the Tt. PDOS is shifted to a lower energy in the rutile state (>Tt) compared to the monoclinic state (Tt) drive VO2 toward a tetragonal (rutile) phase with delocalized electrons. In the

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Figure 4.5 (A) Experimental and theoretical phonon density of states spectra showing the phonon stiffening in the M1 (298K) and phonon softening in the rutile (381K) phase. (B) Changes in the crystal structure of VO2 in the high (rutile)- and low (M1)temperature states. (From J.D. Budai, et al., Nature 515 (2014) 535e539.)

low-temperature (300 nm) show an efficient TC switching [39] similar to the one shown in Fig. 4.7A. The transmittance in the near- to high-infrared region drops to near zero and the reflectance in this spectral region increases correspondingly. The change in optical properties in the visible region of the spectrum is nearly negligible. The TC switching contrast depends quite strongly on the VO2 deposition conditions as shown in Fig. 4.7B in which the film thickness is shown to have a predominant effect on the optical and optical switching properties. As seen in this figure, with increasing film thickness the near-infrared transmittance of the sputter-deposited VO2 films increases while the transmittance in the visible region of the spectrum decreases. While this decrease in the visible region due to increasing absorption is a common observation in all the films, the increase in infrared (IR) transmittance with increasing film thickness is quite surprising. It can also be seen from this figure that the optical contrast following the TC switching increases with increasing film thickness. The other deposition conditions such as the substrate temperature also affect to some extent the optical and switching behavior of the VO2 films. In addition to the overall optical contrast that is achievable in these films, with increasing film thickness (300 nm) it is observed that the film exhibits zero transmittance in the switched state in the near-IR region. By comparison, in a 100-nm-thick VO2 film the switched state transmittance is around 5% in this region. In Fig. 4.8 are given the typical hysteresis curves of the sputter-deposited VO2 thin films [39]. The hysteresis curves can be obtained either in terms of the optical (transmittance and/or reflectance) parameters or in terms of the resistance change occurring as a function of temperature during the heating and cooling cycles. Shown in Fig. 4.8A is the transmittance hysteresis of a sputter-deposited VO2 film of 300 nm thickness. As seen in this figure, the film exhibits an abrupt decrease in transmittance around 62 C. This optical transition during the cooling cycle is not as abrupt but a more gradual one. The transmittance decrease during the cooling cycle does not commence until 57 C. It is more common to use the resistance hysteresis to characterize the TC behavior of VO2. The resistance hysteresis shown in Fig. 4.8B is more typical of many VO2 films although the onset of the transition during the cooling and heating cycles, the width of the hysteresis, and the actual temperature at which these occur can differ. The order of change in resistance can also differ from one to another depending on the micro- and

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(B) 100 Sheet Resistance Rs(kΩ/sq)

Transmission (%)

80 60 Cooling 40

Heating

20 0 20

30

40 50 60 Temperature °C

70

80

Cooling

10

1

d(log (-Rs))/dT

(A)

20

8 7 6 5 4 3 2 1 0 -1

61.4 °C

Heating

2.61 °C

50 60 70 Temperature °C

30

40 50 60 Temperature °C

70

80

Figure 4.8 (A) Transmittance hysteresis as a function of temperature of a 300-nm VO2 film. (B) Hysteresis of sheet resistance (RS) of the VO2 film with temperature (inset: temperature derivative of RS vs. temperature). (From R. Balu, P.V. Ashrit, Applied Physics Letters 92 (2008) 021904.)

nanostructure imparted to the film by the deposition conditions. The sputter-deposited VO2 film shown in Fig. 4.8B exhibits two orders of change in resistance. The resistance hysteresis is narrower than the optical one (Fig. 4.8A) and the transition slopes during both the heating and the cooling cycles are nearly the same. It is a common practice to work with the derivative of the resistance, R (or sheet resistance RS, d[log (R)]/dT ), to define the exact Tt as the position of the highest value of this derivative [40]. Shown in the inset of Fig. 4.8B is this plot of the derivative of the RS of the sputter-deposited film as a function of the temperature. The peak of the derivative gives the switching temperature and the full width at halfmaximum gives the abruptness or sharpness of the transition. According to this plot, the transition peak for this film is located at 61.4 C and the transition between the cooling and the heating cycles is spread over 2.61 C. As discussed before, in addition to these optical and electrical changes are the structural changes occurring simultaneous with the TC switching of the VO2. Given in Fig. 4.9 are the X-ray diffraction (XRD) patterns of the sputter-deposited VO2 film with 100 and 300 nm thickness in their normal states (room temperature) and in the switched states (high temperature). The XRD patterns of the normal states are indexed with ICDD file 44-0252 and those of the switched states are indexed with ICDD file 044-0253. All the peaks were assigned to VO2 hence indicating the absence of any other VxOy phases. However, quite a bit of difference in the XRD patterns is seen between the VO2 films of 100 and 300 nm thickness. The thicker VO2 film shows multiple peaks corresponding to multiple

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20

30

50 2θ

60

80

(011)

(402)

(113) (022)

(210)

(200)

T=25 °C

(B)

20

(110)

30

40

(121) (220)

50 2θ

60

(002)

T=80 °C (011) (020)

(204)

70

Intensity (arb. units)

(002) (121)

(020)

40

T=80 °C (031)

(402) (113)

(302)

(002)

(102)

(200)

(B) (A) T=25 °C

(011)

(011)

(B) (110)

Intensity (arb. units)

(A) (A)

70

80

Figure 4.9 X-ray diffraction spectra of VO2 films in both normal and switched states for films with thickness of (A) 300 nm and (B) 100 nm. (From R. Balu, P.V. Ashrit, Applied Physics Letters 92 (2008) 021904.)

orientations of the crystals, i.e., a polycrystalline phase, while the thinner film prepared under similar conditions shows a preferential orientation along the (011) plane. Despite this difference in the normal state, both films switch from an initial monoclinic phase at room temperature to a tetragonal phase at high temperature. The crystallite size calculated using the Scherer formula is around 20 nm in both cases. The difference in their structure in the as-deposited state indicates the importance of film thickness and the annealing time. At smaller thicknesses the activation energy provided seems to be more effective in uniformly orienting the crystals, while higher energy may be required to do the same in films with higher thickness. From these initial examples it is clear that the deposition parameters, which strongly influence the micro- and nanostructure of the VO2 films, also lead to a wide range of optical, electrical, crystallinity, and their switching properties, including the hysteresis. Hence, an in-depth study of all these properties needs to be carried out to understand the VO2 thin films on hand.

4.1.2 Thermal Hysteresis in Vanadium Dioxide As mentioned earlier, the thermal hysteresis of VO2 electrical (resistance) and optical (transmittance and reflectance) properties is one of the most important tools of characterization of the TC switching behavior of this material. All the aspects of these hysteresis cycles, such as the width of the

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hysteresis (DT ), the transition slopes during heating and cooling cycles, and more, are found to vary quite a bit depending on the perfection of the VO2 crystals in bulk and even more sensitively on the micro- and nanostructure in VO2 thin films. This micro- or nanostructure or even the atomic arrangement in VO2, in turn, depends on the thin film deposition conditions [41]. A general observation is that a large resistivity change (Dr) is accompanied by a smaller hysteresis width (DT ) [36]. For a quantitative analysis of the VO2 thermal hysteresis loops as shown in Fig. 4.8B, it is common to rely on the mathematical models generally developed for explaining the hysteresis loops of magnetic materials. Ferromagnetic materials such as iron, for example, are formed of domains in which an enormously large number of atoms are in an ordered state with respect to the orientation of their magnetic moments. This parallel orientation of magnetic moments within a domain is due to the strong coupling of moments of neighboring atoms. However, the overall ferromagnetic sample does not exhibit a net magnetic field because of the random orientation of the domains as shown in Fig. 4.10A. When an external magnetic field is imposed on this material, more and more domains orient themselves in the direction of the magnetic field, thus increasing the magnetic field in the material. It happens in such a way that the domains already aligned in the direction of the external field begin to grow by

Figure 4.10 (A) Schematic of the existence of domains with their randomly oriented magnetic moments. (B) Schematic of the domain growth with increasing external magnetic field, B0. (From K. Stowe, Introduction to Statistical Mechanics and Thermodynamics, Johnson Wiley & Sons, 1984.)

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engulfing the non-field-oriented domains as the externally applied magnetic field is increased. Such a growth occurs by the movement of the domain walls or boundaries. This magnetization is as shown in Fig. 4.10B. However, when the external magnetic field is reduced gradually the path along which the internal magnetic field of the material is returned to its original state is not reversible and exhibits a hysteresis loop. This is due to the fact that once the external field is reduced the domains do not become disoriented but retain a certain degree of orientation parallel to the external field. The strong interaction between the atoms in these materials tends to hold the transformed orientation even after the removal of the external magnetic field, thus producing a “magnetic inertia” [42]. Therefore, the return path is not retraceable, leading to a hysteresis. A parallel situation can be envisioned in the thermal hysteresis of VO2. Around the dielectric (semiconductor)-to-metal transition VO2 can be imagined to be made up of two types of domains, the metallic domains and the dielectric (semiconductor) domains, which exhibit their constituent optical and electrical properties as a function of temperature. At temperatures just below the transition (Tt) VO2 stabilizes into a tetragonal rutile structure as shown in Fig. 2.20A. At low temperature ( Tt) as shown earlier in Fig. 4.7. Another interesting aspect of TC switching in contrast with electrochromic switching discussed in Chapter 3 is that the bulk of the change in optical properties upon switching (>Tt) occurs in the IR region and the change in the lower

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(visible) wavelengths is small. Hence, the smart window application through the TC approach is more specifically suitable for the management of the heat part (700e2000 nm) of the solar spectrum while leaving the visible part (400e700 nm) nearly intact. The natural Tt of around 68 C of near-stoichiometric VO2 is quite high for a passive approach to smart window application. However, as discussed earlier, the Tt can be brought down quite effectively by either a suitable doping or through nanostructuring of the VO2 films. The active control through external joule heating, by preparing the VO2 thin films on suitable substrates such as transparent conductive coated glass [66], for example, can also be very effective. Similarly, optical pumping-based TC switching can also be of great interest [67]. However, such active control may not be efficient for an energy management application point of view but can be of great interest in other applications that rely on such active optical or electrical switching of VO2 [68]. As far as the smart window application for energy control is concerned the VO2 thin film-based devices’ switching performance needs to be examined in the light of the important optical parameters, the integrated solar (Tsolar) and photopic (Tphotopic) transmittances discussed in Section 4.2.2 and given by Eqs. (4.3) and (4.4). An important aspect to note in the case of as-deposited VO2 thin films is that they exhibit a lower transmittance in the visible (photopic) region of the spectrum compared to the near-IR region. This visible transmittance can be increased to some extent by working with thinner films. However, as seen earlier in Fig. 4.7, the TC switching efficiency also depends on the film thickness. As seen in that figure, the thicker films exhibit a more efficient switching than the thinner ones. Hence, a compromise needs to be made between maintaining high transmittance in the visible region and a high switching efficiency in the near-IR region for smart window applications. Most of the work related to smart window applications of VO2 thin films focuses on these three aspects: (1) bringing down the Tt of VO2 thin films, (2) increasing the Tphotopic, and (3) increasing the TC switching of Tsolar. VO2-based nanocomposite TC films have been prepared by Manfredi Saeli et al. [69] using a hybrid atmospheric pressure and aerosol-assisted chemical vapor deposition (AP-AACVD) method based on vanadyl acetylacetonate as precursor. The VO2 thin films prepared by this AP-AACVD method were embedded with gold nanoparticles, and tetraacetylammonium bromide (TOAB) was used in some cases as the surfactant. Scanning electron microscopy (SEM) micrographs of the nanocomposite films with

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(A) (A)

(B)

(B) (A)

(B)

(C)

Figure 4.23 (A) Scanning electron microscopy (SEM) pictures of VO2/Au nanocomposite films produced by the atmospheric pressure and aerosol-assisted chemical vapor deposition (AP-AACVD) method with (image A) and without (image B) the tetraacetylammonium bromide (TOAB) surfactant (original image size approx. 3.5  2.5 mm). (B) SEM pictures of VO2 nanostructured films produced by AP-AACVD with (image A) and without (image B) the TOAB surfactant (original image size approx. 1.8  1.3 mm). (C) Photographs of the nanocomposite films in comparison with a regular W-doped VO2 film (image A) and with varying Au/V ratio of 0.09 (image B), 0.15 (image C), 0.3 (image D), and 0.36 (image E) and a gold nanoparticle film (image F). (Adapted from M. Saeli, et al., Solar Energy Materials and Solar Cells 94 (2010) 141e151.)

and without the surfactant and with and without the gold nanoparticles are compared in Fig. 4.23. In Fig. 4.23A are shown the nanocomposite VO2/ Au films with and without the TOAB surfactant. These films show a highly granular structure with larger grains (150 nm) in the absence of the TOAB surfactant and smaller grains (75 nm) in the films with TOAB. In the

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nanostructured VO2 samples (without Au nanoparticles) the grain sizes are 50 and 100 nm for films with and without TOAB, respectively, as shown in Fig. 4.23B. The nanocomposite films grown with TOAB were spherical with a size of 70  10 nm. However, the gold nanoparticles in the nanocomposite films without the TOAB were irregular in shape and smaller in size (40 nm), with a very good interconnectivity. The predominant presence of VO2 and Au was confirmed by XRD and x-ray photoelectron spectroscopy studies, which also indicated the small presence of the V2O5 phase at the surface. The visible appearance of the nanocomposite films was as shown in Fig. 4.23C. The normally brownish VO2 films display a wide range of blue and green colors depending on the Au content. The optical behavior of these films is shown in Fig. 4.24. The transmittance and reflectance spectra for each kind of film are shown here at room temperature (dielectric state) and at high temperature (metallic state). The optical switching behavior of the nanocomposite films with and without the TOAB is shown in Fig. 4.24A and that of VO2 film in Fig. 4.24B for both types. In all cases the transmittance decreases and reflectance increases, in general, with the heating of the films. The films with Au nanoparticles exhibit a lower overall transmittance at room temperature, which leads to their lower switching efficiency. Very minimal changes are seen in the visible region of the spectrum upon TC switching. The most efficient change is seen in the case of VO2 film with TOAB with a drop of transmittance from around 80% to 40% in the high-wavelength (>2000 nm) extreme of the spectrum. This particular film also has a fairly high visible transmittance (60%) with no change after switching. Such films are interesting for the heat management of solar energy. Sihai et al. [70] have studied the TC behavior of nanostructured VO2 thin films prepared by reactive ion beam sputtering and postannealing. Their most optimized film had a well-controlled nanostructure as shown by the micrographs in Fig. 4.25A. The two most interesting traits induced in these optimized films possessing this nanostructure are the low Tt of 35 C and the transmittance switch that occurs in these films between the hightemperature (metallic) and the low-temperature (dielectric) states. The transmittance spectra of the film are shown in Fig. 4.25B in the two states. An excellent degree of transmission switching takes place in this film as a function of temperature. Although not reported in this work, the transmittance in the lower wavelengths (Tt) as required for smart window application, the optical

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modulation in the near-IR region of the solar spectrum is very inefficient, changing by only 3%. Another problem associated with this device is the high Tt at which the modulation occurs. The thermal hysteresis of transmittance of the single VO2 film is compared to that of the sandwich structure in Fig. 4.27C. These hysteresis results recorded at the wavelength of 2000 nm indicate the reduction in transmittance switching efficiency of the sandwich structure compared to that of the single film. Even more important is the fact that although a slight reduction in Tt (58.5 C) is reported, it still remains quite high for practical applications. Despite these shortcomings, the work paves the way to addressing at least one of the important problems related to VO2-based smart windows, i.e., improving the luminous efficiency through multiple AR coatings. Similarly, more advanced multilayer structures based on AR coatings were prepared by Mlyuka et al. [73] to further improve the Tlum of VO2based TC devices. The five-layer configuration, TiO2/VO2/TiO2/VO2/ TiO2, was prepared by reactive sputtering on glass, carbon, and silicon substrates. The experimental and calculated TC performance of a 90-nmthick-base single VO2 layer is shown in Fig. 4.28A and B. As can be seen in this figure, a reasonable visible transmittance of around 40% is exhibited by the VO2 film, which actually improves at high temperature owing to slightly higher values due to the optical effects in the film. An efficient TC modulation in the near-IR region is exhibited. The calculated Tlum and Tsolar of the five-layer structure on glass with a fixed VO2 film thickness of 40 nm are shown here as functions of the angle of incidence and for various thicknesses of TiO2 film. It can be seen from this figure that at all the TiO2 thicknesses a higher degree of Tsolar switch with a lower change in Tlum is seen at temperatures above and below the Tt. At normal incidence and at low temperature ( sc) the Tt. R, reflectance; T, transmittance; Tlum, luminous transmittance; Tsol, solar transmittance. (From N.R. Mlyuka, G.A. Niklasson, C.G. Granqvist, Physica Status Solidi A 206 (9) (2009) 2155e2160.)

and 42% at low temperature (Tt), respectively, i.e., higher than the calculated values from Fig. 4.28B. DTsolar was 8% for normal incidence. In addition to this insignificant improvement even with a five-layer structure, the Tt of the base VO2 layer remains very high

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Figure 4.29 (A) Measured transmittance spectra of the TiO2/VO2/TiO2/VO2/TiO2 structure at a temperature below and above the transition temperature at normal incidence and in two states of polarization (130-nm-thick TiO2 and 40-nm-thick VO2). (B) Hysteresis of sheet resistance of a 50-nm-thick VO2. (From N.R. Mlyuka, G.A. Niklasson, C.G. Granqvist, Physica Status Solidi A 206 (9) (2009) 2155e2160.)

from a solar application point of view. As shown in Fig. 4.29B in the hysteresis cycle of the resistance of the VO2 film the Tt is around 68 C similar, to the bulk value. Another approach to improvement of the visible state transmittance of TC VO2 films is suggested by Khan et al. [74] from their VO2 films prepared by RF sputtering followed by annealing. Their 170-nm-thick VO2

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had the typical high absorption in the visible region of the spectrum, imparting a deep brownish color to the sample, making the transmittance low. It is well known that this optical absorption edge commonly found in vanadium oxides can be moved to lower wavelengths with the insertion of suitable ions (Liþ, Hþ, Naþ, etc.) and electrons as discussed in the case of vanadium pentoxide (V2O5) films’ electrochromic transformation in Chapter 3 (Eq. 3.3). Similar electrochemical reactions ensue in the case of VO2 films [75,76]: VO2 þ xLiþ þ xe 4 Lix VO2 Brown

(4.4)

Transparent

This reaction has been exploited to electrochemically insert various amounts of lithium into the VO2 films. Transmittance spectra of VO2 film with various amounts of lithium are shown in the solar spectral region in Fig. 4.30A. As can be seen here a very significant improvement in the transmittance occurs in the solar spectral region with the increased insertion of lithium. The change is highest in the visible region of the spectrum. After a lithiation of x ¼ 0.43, the VO2 films exhibit a transmittance of around 60% around the midvisible wavelengths, which is quite significant. The TC performance of these films is shown in Fig. 4.30B and compared to that without any lithium. It is seen that while the lithiated films do maintain a reasonable amount of TC switch in the spectral region studied, it is not as efficient as in the case of the unlithiated VO2 film. The same diminishing effect with increased lithium can also be seen in these films through the thermal hysteresis of conductivity. These results are shown in Fig. 4.30C in which the conductivity switch (Ds) after the transition is seen to decrease gradually with increased lithium. Although this approach to increasing the luminous transmittance is very interesting, such films are marred with decreased TC performance and also maintain a high Tt that makes them inefficient for solar management applications. It is perhaps because of this reason that further work related to lithium insertion into VO2 film has not continued. More tangible results in this regard were obtained by Mlyuka et al. [77] by doping the VO2 films with magnesium. As discussed earlier doping with various metal atoms such as W, Mo, and Nb has successfully lowered the Tt. However, these attempts have invariably lowered also the luminous transmittance and the TC efficiency of VO2 films. In this work [77] both these limitations seem to have been overcome by doping VO2 with divalent Mg atoms. TC films with the composition MgxV1xO2 were

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Figure 4.30 (A) Transmittance spectra of a VO2 film in the solar spectral region with the insertion of various amounts of lithium. (B) Transmittance spectra of the lithiated (Li0.43VO2) and as-deposited VO2 film at temperatures above and below the transition temperature. (C) Thermal hysteresis of the VO2 film with different amounts of lithium inserted. (From M.S.R. Khan, K.A. Khan, W. Estrada, C.G. Granqvist, Journal of Applied Physics 69 (5) (1991) 3231e3234.)

reactively sputter deposited on glass substrates varying the x value. In Fig. 4.31A (a) is shown the TC performance of the 7.2% doped and the undoped VO2 sample in the form of transmittance spectra below and above Tt. As can be seen here, in the as-deposited state (22 C) an overall improvement in transmittance is seen throughout the spectral region studied by Mg doping. The doped films exhibit an effective TC switch, albeit with reduced efficiency compared to the undoped films. A significant result of this work is a nearly linear increase in the luminous transmittance with Mg doping as shown in Fig. 4.31A (b). The Tlum value is seen to vary

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Figure 4.31 (A) Transmittance spectra of a Mg-doped and undoped VO2 sample above and below the transition temperature (Tt) (a) and variation in the luminous transmittance as a function of the Mg doping (b). (B) Hysteresis curves of the Mg-doped and undoped VO2 films (a) and Tt as a function of the Mg doping. (From N.R. Mlyuka, G.A. Niklasson, C.G. Granqvist, Mg doping of thermochromic VO2 films enhances the optical transmittance and decreases the metal-insulator transition temperature, Applied Physics Letters 95 (17) (2009), 171909/1e3.)

between 39% and 51% within the doping range studied, while the Tsolar is reported to vary by about 4% in the same range. Another important contribution of this work is the reduction in the Tt. with doping. These results are shown in Fig. 4.31B. The hysteresis loop of the transmittance recorded at the wavelength of 2300 nm seen in Fig. 4.31B (a) shows that the Tt drops to lower values with increased doping of Mg. However, this decrease in Tt is also accompanied by a decrease in the degree of transmittance change. In Fig. 4.31B (b) is shown the decrease in Tt with doping. This linear drop with Mg doping is compared to that of

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W doping of VO2 carried out by the same authors. Though not as effective as W doping (w25 C/W atomic %), the Mg doping corresponds to a mere decrease of around 3 C/Mg atomic %. Despite the moderate success achieved by this approach of combined effect of an increase in luminous transmittance (DTlum w 10%) and a decrease in Tt, the problem of low TC efficiency persists. Although a fairly effective reduction in transmittance as seen in Fig. 4.31A (a) is seen in the near-IR region, the effectiveness of this change as far as the heat part of the solar spectrum (800e1500 nm) is concerned is meager. This is due to the inherent high absorption of reasonably thick and continuous VO2 films and their natural TC modulation being effective only at higher wavelengths beyond the solar spectrum. In these continuous VO2 films in the metallic state rendered at high temperature (>Tt) the plasma frequency (up) is invariably situated at higher wavelengths, making the metallic reflectance of VO2 higher only in this region. Hence, a well-designed VO2-based smart window for energy management application has to address this problem along with a lower Tt and higher Tlum. In some sense, the requirements of higher Tlum and higher IR modulation (Tsolar) within the solar spectrum are contradictory as was seen in Fig. 4.7B. To increase the Tlum one has to work with thinner VO2 films. With decreasing thickness of the film the TC modulation also decreases. Nanothermochromics is a fairly recent idea proposed by Li et al. [78] specifically to overcome these limitations posed by the conventional VO2 thin film-based approach to smart windows for energy efficiency. Hence, the ultimate goal of achieving the highest efficiency in this regard is to maintain a high luminous efficiency in the visible range and to maximize the dynamic modulation in the near-IR region of the solar spectrum. The other stringent requirement along these lines is to induce this modulation at a temperature as close to room temperature as possible. Nanothermochromics refers to working with composites containing VO2 nanoparticles suspended in a dielectric host matrix. Compared to the VO2 single film approach to thermochromics, the VO2 nanoparticle composites have the potential for a high luminous efficiency (Tlum) due to the diminished scattering of these particles, which at high temperature (metallic state) can lead to reflection modulation. The nanoparticle size can be tuned to adjust the plasma resonance frequency (up) to be located in the near-IR range within the solar spectrum. Hence, such composite films are expected to be highly transparent at room temperature (Tt), while in the heat part of the solar region they exhibit

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high transparency at room temperature (Tt). Hence, the nanothermochromics approach to VO2-based smart windows has a high potential for energy efficiency applications. The theoretical feasibility of this approach has been amply demonstrated by Li et al. [78e80]. By considering the optical constants of a solid (continuous) VO2 film in its low-temperature (semiconducting) state and high-temperature (metallic) state, the spectral transmittance as well as the integrated luminous and solar transmittance were calculated. These results are shown in Fig. 4.32A and B. The spectral transmittance for the VO2 film is shown in Fig. 4.32A for the semiconducting and metallic states for different film thicknesses. In accordance with the experimental results seen earlier, these results indicate a decreasing transmittance throughout the spectral region studied. It can also be noted that the TC efficiency at the higher wavelengths increases gradually with increasing VO2 film thickness.

Figure 4.32 (A) Calculated transmittance for a solid (continuous) VO2 film with different thicknesses in the semiconducting (a) and metallic (b) state. (B) Calculated luminous transmittance (a) and solar transmittance (b) of VO2 film as functions of the film thickness in the semiconducting and metallic states. (From S.-Y. Li, G.A. Niklasson, C.G. Granqvist, Journal of Applied Physics 108 (2010) 663525e663528.)

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The luminous and the solar transmittances calculated from these values are given in Fig. 4.32B. As already mentioned it can be seen that the Tlum falls off very rapidly with increasing film thickness. At very low thicknesses (0.05 mm) the Tlum falls to even smaller values in the metallic state. A similar trend is seen in the case of Tsolar, limiting the modulation to a maximum of 14% at higher thicknesses. Even with the best compromise between the two parameters (Tlum, Tsolar) achievable at the thickness of 0.05 mm (Tlum w 40%), only 10% Tsolar modulation can be expected. As per the nanothermochromics approach, the optical constants were then used to calculate the optical behavior of nanocomposites in which VO2 nanoparticles were embedded in a dielectric host material and where these nanoparticles were assumed to be of different shapes (spherical as well as ellipsoid oriented vertically, horizontally, and randomly) as shown in Fig. 4.33. The interaction of these particles with light of pertinent wavelengths much larger than the particle size was considered. This interaction was considered in terms of the aspect ratio (m) expressed as the ratio of the major and minor axes of the particles as shown in the figure. Considering a dilute dispersion of the particles in the dielectric host matrix, the MaxwelleGarnett theory was applied to calculate the dielectric constant of the composite medium, ε: 2 1þ fa 3 ε ¼ εm 1 1 fa 3

(4.5)

with a¼

εp  εm εm þ Lðεp  εm Þ

(4.6)

where εm and f are the dielectric constant of the host medium and the VO2 particle filling factor in the host medium, εp is the dielectric constant of the VO2 particles, and L is the depolarization factor that is dependent on the shape of the particle under consideration. The spectral transmittances were calculated for the different shapes of the VO2 particles considering a filling factor, f, of these particles to be 0.01. The refractive index of 1.5 corresponding to a dielectric constant of the matrix, εm, of 2.25 was considered for these calculations. The substrate was also assumed to be of the same dielectric material. The calculated transmittance

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Figure 4.33 Schematics of light interactions with the assumed nanocomposites formed by the dilute dispersion of VO2 nanoparticles of various shapes embedded in the dielectric matrix. The different shapes were defined by the aspect ratio (m ¼ c/a or a/c) as shown. (A) Spheroids, (B) oriented prolate (major axis oriented perpendicular to the substrate), (C) oblate (major axis parallel to the substrate), (D) prolate randomly oriented. (From S.-Y. Li, G.A. Niklasson, C.G. Granqvist, Journal of Applied Physics 108 (2010) 663525e663528)

values of the composite system with the VO2 particles in their semiconducting (low temperature) and metallic (high temperature) states are shown in Fig. 4.34A (a) and (b), respectively. These transmittance values are calculated for different-sized VO2 particles embedded in the dielectric matrix and for different shapes (spheroid, oblate, and prolate) with different sizes. In the spectral region presented, there is an effective transmittance decrease (increase in reflectance) with semiconducting to metallic state in all the cases. In the semiconducting (low temperature) state all the films exhibit a high transmittance in the spectral region studied, with the highly prolate (m ¼ 10) particles exhibiting the highest transmittance value. The transmittance in this state decreases gradually as the value of m decreases. The decrease in transmittance band with transition to metallic (high temperature) state that is initially centered around 1000 nm broadens to higher wavelength with

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Figure 4.34 (A) Calculated spectral transmittance of nanocomposite films in the semiconducting (a) and metallic state (b) for different values of the aspect ratio, m, for the VO2 particles and a filling factor, f, of 0.01. (B) Luminous (a) and solar (b) transmittance values calculated from the spectra in (A) as functions of the aspect ratio, m. (From S.-Y. Li, G.A. Niklasson, C.G. Granqvist, Journal of Applied Physics 108 (2010) 663525e663528.)

rapidly decreasing transmittance with decreasing value of m. Hence, the most effective change in transmittance with transition seems to be exhibited by the film with oblate particles of m ¼ 0.1. The luminous and solar transmittances calculated from these spectral values are shown in Fig. 4.34B (a) and (b), respectively. The luminous transmittance in both the semiconducting and the metallic state increases gradually with increasing m value [Fig. 4.34B (a)]. A decrease of around 10% between the semiconducting and the metallic state is seen for all the particle shapes and sizes. The maximum values of luminous transmittance are exhibited by films with prolate particles with m ¼ 10 whose values are nearly 70% and 80% in the metallic and semiconducting states, respectively. As far as the solar transmittance for these films is concerned, the nature of increase in transmittance is similar with increasing m. However, the decrease in solar transmittance with semiconductor-to-metal transition is slightly higher (15%e20%). The highest transmittance is once again

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exhibited by films with prolate particles with m ¼ 10 being on the order of 80% and 65% in semiconducting and metallic states, respectively. Hence, the best calculated results correspond to 10% of the change in luminous transmittance with the TC effect in VO2 particles, which is an acceptable change from a smart window application point of view as a transmittance of around 70% is maintained by these films with prolate VO2 particles with m ¼ 10. However, the best modulation achievable by the same films in the solar region is only 15%, which seems to be inadequate. The results for the randomly oriented VO2 particles of different sizes in a dielectric matrix are given in Fig. 4.35A and B. All the films show a fairly high transmittance in the semiconducting state and a substantial decrease in transmittance in the higher wavelength range studied in the metallic state as seen in Fig. 4.35A. In the calculated luminous and solar transmittances shown in Fig. 4.35B it is seen that the tendency in the two films is similar with variation of m value.

Figure 4.35 (A) Calculated spectral transmittance of nanocomposite films in the semiconducting (a) and metallic state (b) for different values of the aspect ratio, m, for the randomly oriented VO2 particles and a filling factor, f, of 0.01. (B) Luminous (a) and solar (b) transmittance values calculated from the spectra in (A) as functions of the aspect ratio, m. (From S.-Y. Li, G.A. Niklasson, C.G. Granqvist, Journal of Applied Physics, 108 (2010) 663525e663528.)

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Interestingly, here, and unlike the specifically oriented particle films, it is the spheroid films (m ¼ 1) that show the highest transmittance values in all states. They exhibit a luminous transmittance value of slightly over 70% and 60% in the semiconducting and metallic states. The solar transmittance values for the same particulate films are 73% and 57% in semiconducting and metallic states. Once again, despite good transmittance values in the visible wavelengths, the modulation of around 16% seems to be insufficient from a smart window application point of view. In view of further refining the nanothermochromics approach, the same authors [78,79] have made similar calculations for nanocomposite TC films in which the embedded particles in the dielectric matrix were made up of VO2 coreeshells, i.e., the TC VO2 spherical nanoparticles were surrounded by another dielectric material as a thin shell and embedded in a host dielectric matrix. The schematic of such a TC nanocomposite interacting with light is shown in Fig. 4.36A. The a value of Eq. (4.5) for ε in this case is given by: a¼3

ðεs  εm Þðεc þ 2εs Þ þ Uð2εs þ εm Þðεc  εs Þ ðεs þ 2εm Þðεc þ 2εs Þ þ 2Uððεs  εm Þðεc  εs Þ

(4.7)

where, along with the dielectric constants of host matrix (εm ), the new dielectric constants εs and εc of the core VO2 particles and the outer shell material are used to characterize the optical behavior. U, the volume ratio of the inner sphere and outer sphere, is given by:  x 3 U¼ (4.8) 2t þ x where x and t are the diameter of the core and the thickness of the shell as shown in Fig. 4.36A. Using this consideration, the authors [79] calculated the spectral transmittance of the composite films in both the low-temperature (semiconducting) and the high-temperature (metallic) state of VO2. The shell of thickness t surrounding the core was considered to be made of VO2 and changing between the semiconducting and the metallic states. Hence, εs in the Eq. (4.7) was applied to VO2 in these two states. The host matrix was considered to be made up of a dielectric constant εm through the refractive index nm of 1.5. The dielectric constant of the core material εc was considered through its refractive index nc, which was varied between 1 and 2.5 to represent different core media. In Fig. 4.36B are given these calculated results for the spectral transmittance with varying core

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diameter-to-shell thickness ratio. In the semiconducting state a fairly high transmittance is seen throughout the spectral region studied with a much higher transmittance at the higher wavelengths. For the nc value of 1 the transmittance decreases with increasing x/t ratio, while at nc ¼ 2, the

Figure 4.36 (A) Schematic of the light interaction with a coreeshell nanocomposite film based on VO2. (B) Calculated spectral transmittance of the coreeshell nanocomposite in the semiconducting (left) and metallic (right) states of the VO2 shell for various values of nc and x/t ratio. (From S.-Y. Li, G.A. Niklasson, C.G. Granqvist, Journal of Applied Physics 109 (2011) 113515e1-5.)

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variation is negligible. In all the cases of the semiconducting state the composite exhibits maximum transmittance when x/t is 0, i.e., when the core diminishes to zero and VO2 particles are embedded in the host matrix. In the high-temperature (metallic) state a very effective reduction in spectral transmittance takes place at the higher wavelengths of the spectral region studied. The transmittance reduction band initially centered around 1200nm wavelength deepens and spreads to higher wavelength regions with increasing x/t as seen in Fig. 4.36B (right). Hence, when the entire spectral region is considered, it is the film with x/t value of 20 that seems to exhibit the maximum modulation with the TC transition of VO2. The luminous and solar transmittance derived from these spectral transmittance results are shown in Fig. 4.37. The dependence of these parameters on the x/t ratio for different values of core refractive index is shown in this figure for both the semiconducting and the metallic states. The figure also shows the actual solar modulation undergone by the composite film of 5 mm thickness through the TC transition of VO2 film and the dependence of this parameter on the x/t ratio. As seen from this figure, in all the cases the maximum solar and luminous transmittance is exhibited by the film made up of solid VO2 particles (x/t ¼ 0). In the semiconducting state such films exhibit a maximum transmittance of over 70% in both the visible and the solar spectral regions. These values decrease by around 8% in the luminous region and by around 15% in the solar spectral region. The solar modulation values calculated from these are also shown in the figure for various values of x/t. As seen in the last (bottom) panel of Fig. 4.37, the maximum modulation is seen in the case of a film of which the core is made up of air (hollow shell) with x/t of 10. Such film made up of a thin VO2 hollow core exhibits a solar modulation of around 21%. Although this is a significant improvement from the earlier seen nanocomposites (Fig. 4.33), the improvement comes with a compromise of the luminous transmittance because the same hollow core VO2 nanocomposite exhibits a diminution of around 15%. Although the nanothermochromics approach proposed by Li et al. [78e80] does provide a tangible theoretical approach to the design of VO2 thermochromics-based smart windows, there seems to be an ultimate limit to the luminous transmittance and the solar modulation achievable through this approach. This is due to the inherent nature of the VO2 thermochromics. However, nanothermochromics and the theoretical approach proposed provide an excellent platform for further improving the performance of such devices. In this approach [78e80] no consideration is given

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to variation in Tt, which can be an important part of smart window applications. Hence, for further work it may be necessary to consider a combination of all the factors involved in the improvement of the TC performance to make nanothermochromics a viable approach. Nanocomposites as proposed earlier with doped species to bring down the Tt need to be considered. In addition, the experimental preparation of these special structured nanocomposites can be quite a challenge considering the

Figure 4.37 Luminous (top) and solar (middle) transmittance of coreeshell nanocomposites based on VO2 in the semiconducting and metallic state as functions of x/t ratio for different values of nc. Solar modulation (bottom) of the nanocomposite film as a function of the x/t ratio is also shown. (From S.-Y. Li, G.A. Niklasson and C.G. Granqvist, Journal of Applied Physics, 109 (2011) 113515e1-5.)

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large-scale application eventually envisioned. Many types of doped and undoped VO2-based nanocomposites have been fabricated using different techniques that may be employed to take the nanothermochromics approach [81e85]. Unlike the WO3 electrochromics-based smart windows discussed in Chapter 3, the performance of the VO2-based smart windows is substantially lower. This is due to the inherent nature of the electrochromic coloration of WO3, especially in the polycrystalline phase, which is centered in the near-IR region and can be tailored to have a good optical modulation in the solar spectral region while maintaining an acceptable level of Tlum. However, the advantage of the VO2 thermochromics-based smart window is that it is, generally, a single-layer approach and is activated passively by ambient heat. The WO3 electrochromics-based smart windows are multilayered structures that have to be actively operated by applying an external electric field. Hence, the route to smart windows as far as the energy management application is concerned has to be taken by weighing the pros and cons of each approach. Table 4.4 lists some of the Table 4.4 Best luminous transmittance and solar modulation optimization results through different approaches Method DTsolar Tlum Tt, 8C References By doping

Experimental Experimental

6.7% 4%

53.8% 51%

47.5 45

[86] [87]

14.7%

50%

e

[88]

48% 59.1% 51.5%

e e e

[78] [89] [90]

70.3% 44.5%

e e

[91] [92]

47.2%

e

[93]

By porosity

Experimental

By nanothermochromism

Simulation Experimental Experimental

20% 12% 11.7%

By special nanostructures

Simulation Experimental

23.1% 7.1%

By antireflection coating

Experimental

15.1%

DTsolar, change in solar transmittance (solar modulation); Tlum, luminous transmittance; Tt, transition temperature. Adapted from Y. Zhou, Y. Cai, X. Hu, Y. Long, Journal of Materials Chemistry A 3 (2015) 1121e1126.

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best results obtained as of this writing by various methods in the endeavor to achieve TC efficiency for smart window applications, i.e., to optimize Tlum and solar modulation (DTsolar). As seen from this table, the best experimental results correspond to only 15% solar modulation and 60% Tlum although much higher theoretical values can be expected by special nanostructures such as moth eyes [91], which predict over 70% and 23% values for Tlum and DTsolar, respectively. However, the experimental reproduction of these values is far from being achieved [92]. More recently, Yang Zhou et al. [94] have fabricated nanothermochromic devices from VO2-based organiceinorganic hybrid films. These authors combined two TC effects, one occurring at high temperature in the IR wavelengths and the other at low temperature in the luminous region, to make an effective modulation throughout the solar spectral region. The hybrid structure was fabricated by dispersing the VO2 nanoparticles with poly(N-isopropylacrylamide) (PNIPAm). This polymer is one of the well-known organic hydrogels that exhibit a temperaturedependent and reversible hydrophilic-to-hydrophobic transition at a low critical temperature of 32 C [95,96]. These hydrogels exhibit a thicknessdependent transmittance in the visible and near-IR spectral region at room temperature. Sandwich structures encompassing the PNIPAm hydrogel between two glass sides were prepared with different thicknesses and studied for their temperature-dependent optical properties. The mechanism of the change is shown in Fig. 4.38A. At temperatures below 32 C the laminated structure is transparent to the solar wavelengths. At higher temperatures (>32 C) the solar wavelengths are partially blocked, rendering the structure less transparent. This effective and sensitive change in optical properties with temperature is gradual. The transmittance spectral change occurring in these devices is shown in Fig. 4.38B for two different PNIPAm thicknesses. Laminate structures with higher thickness (200 mm) show high transmittance in the solar spectral region as well as a higher optical modulation with increasing temperatures up to 60 C as shown in Fig. 4.38B (a). The structure with lower thickness (52 mm) shows [Fig. 4.38B (c)] a high transmittance at low temperature and only a modest change in transmittance with increasing temperature. The optical performance as a function of temperature is also shown in Fig. 4.38B (b) and (d) for the thick and thin structures in the form of spectral transmittance as well as IR, solar, and luminous modulation. The thick film structure exhibits a high degree of modulation for all the optical parameters of interest with increasing temperature above the critical temperature. However, the

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Figure 4.38 (A) Schematic of the mechanism of transition of a laminate structure encompassing poly(N-isopropylacrylamide) (PNIPAm) hydrogel above and below the transition temperature. (B) Spectral transmittance of the laminate structure with thick (200 mm) (a) and thin (52 mm) (c) PNIPAm hydrogel at different temperatures; luminous transmittance (Tlum), change in Tlum (DTlum), change in solar transmittance (DTsolar), and change in infrared transmittance (DTIR) as functions of temperature of the thick (b) and thin (d) laminate structures. (From Y. Zhou, Y. Cai, X. Hu, Y. Long, Journal of Materials Chemistry A 2 (2014) 13550e13555.)

large and simultaneous changes occurring in the visible and solar spectra are an impediment for smart window applications in which high Tlum is expected at high temperatures. The large-thickness structure exhibits a nearzero Tlum, being translucent at the temperature of 60 C. The thin film structure of PNIPAm, however, exhibits a much lesser degree of change as shown in Fig. 4.38B (d). Such a structure exhibits a Tlum of around 54% and a DTsolar of 25.5% at 60 C. These performance parameters are already much superior to all the VO2-based approaches listed in Table 4.4.

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To further enhance the DTsolar of such structures VO2 particles were added to the PNIPAm hydrogels to fabricate laminate structures as shown in Fig. 4.39A [94]. The idea is to overlap the modest IR modulation of the thin structure PNIPAm with the efficient IR modulation of VO2 particles

Figure 4.39 (A) Schematic of solar modulation using a VO2/hydrogel-based system. (B) Transmittance spectra of VO2/hybrid gels of different thicknesses at high (80 C) and low (20 C) temperature (a) and film thickness dependence of integrated solar (DTsol), luminous (DTlum), and infrared (DTIR) transmittance changes between high and low temperatures. (C) Spectral irradiance switch of pure hydrogel, VO2 nanoparticles, and VO2/hydrogel hybrid between high and low temperatures. (From Y. Zhou, Y. Cai, X. Hu, Y. Long, Journal of Materials Chemistry A 3 (2015) 1121e1126.)

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to significantly enhance the solar modulation without impeding the Tlum. This is made possible by choosing the appropriate VO2 particle size for dispersal in the hydrogel. The VO2 particle size was reported to be between 20 and 100 nm. The optical performance of these laminate structures with different VO2/PNIPAm hybrid gels is given for different temperatures in Fig. 4.39B (a). The laminate structures exhibit nearly the same and very high transmittance (80%) through most of the spectral region studied, for all the thicknesses. The transmittance modulation, however, depends strongly on the gel thickness and increases with it as seen in this figure. In Fig. 4.39B (b) are given the other optical parameters (average Tlum,, DTlum, DTsolar, and DTIR) as functions of gel thickness. The corresponding changes are also shown in Table 4.5. The results shown in Fig. 4.39B (b) and Table 4.5 clearly indicate the superior performance of the VO2/PNIPAm hydrogel hybrids, which yield a Tlum of 43.2% and a DTsolar of nearly 35% at a hybrid hydrogel thickness of 52 mm at high temperatures. This performance, especially of DTsolar, is far superior to other experimentally achieved values. Even the performance of the laminate structure with a hydrogel thickness of 26 mm with its Tlum of 52.3% and DTsolar of 26% can be a potential candidate for smart window applications. The TC performance of pure hydrogel, VO2 nanoparticles, and a VO2-based hybrid hydrogel are compared in Fig. 4.39C. Given in this figure is also the solar spectral irradiance as a function of wavelength. A significant improvement in Tlum and DTsolar by dispersing VO2 nanoparticles into the organic PNIPAm hydrogel is quite evident from these curves. Although the authors report the superior performance of the best laminate sample in terms of the average Tlum of 62.6%, for practical purposes the performance at high temperature (>Tt ¼ 68 C) should be considered

Table 4.5 Optical parameters of the hybrid gel with a VO2/Poly(Nisopropylacrylamide) ratio of 1/6 at different thicknesses Tlum (%) Thickness (mm)

208C

808C

DTlum (%)

DTsolar (%)

DTIR (%)

13 26 52

86.8 84.0 62.1

61.8 52.3 43.2

25 31.7 38.9

18.5 26.3 29.9

10.5 29.2 34.7

DTIR, change in infrared transmittance; DTlum, change in luminous transmittance; DTsolar, change in solar transmittance; Tlum, luminous transmittance. From Y. Zhou, Y. Cai, X. Hu, Y. Long, Journal of Materials Chemistry A 3 (2015) 1121e1126.

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for this evaluation. Considering the temperature-dependent performance of the PNIPAm hydrogel alone [95,96], a Tlum of less than 53.6% is plausible. These are by far the best results on TC performance in terms of the smart window application. Even with this relatively superior optical performance of the hybrid devices the limitation of the VO2 TC transition activation at 68 C persists. To make these devices a viable solution for smart window application in all aspects, the Tt (VO2 Tt) of these devices needs to be brought down. Hence, a strategy that combines the hybrid materials approach (PNIPAm þ VO2 nanoparticles) with doping of the VO2 with appropriate dopants to reduce the Tt as well as controlling the size and shape of the doped VO2 nanoparticles may be essential to further improve the application potential of these TC devices. In view of this limitation of high Tt one can also think of other TC applications where the relatively high Tt and its reduction may not be a stringent requirement.

4.2.2 Thermochromic Smart Radiator Devices Another important application of TC materials and devices is for the dynamic control of emissivity of surfaces. Under terrestrial and extraterrestrial conditions where radiation is the main mode of heat transfer between objects, heat loss and heat gain from surfaces subjected to variations in conditions occur. These static surfaces are subjected to a large variations in temperature [97]. If they cannot withstand these variations a dynamic control of their radiative behavior needs to be provided. One such extraterrestrial example is the satellites orbiting around earth geosynchronously. These satellites undergo a regular and severe variation in thermal loads with large temperature variations (150 C to þ150 C) depending on their position with respect to the earth. These extreme temperatures are found between the sun-facing and the sun-obscuring (earth shade) positions. In addition, in the sun-facing position the satellites also receive the sunlight reflected by the earth and are also subjected to the radiation emitted from the earth at all positions. Hence, it becomes essential to maintain the internal temperature of the satellite within a modest range, 20 C to 40 C, to ensure the proper functioning of the electronic components. Static solutions of using a high-emissivity or low-emissivity material lead to problems due to the highly varying conditions of the satellites in orbit. For example, low-emissivity materials such as aluminum, though efficient in keeping the satellite cool below a nominal temperature in the sun-obscuring (earth shade, low thermal load) position, lead to a high

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Figure 4.40 Heat load dependence of radiator temperatures in satellites for low- and high-emissivity materials. (http://esmat.esa.int/materials_news/isme09/pdf/4-New/S4%20-%20Haddad.pdf.)

increase in temperature of the satellite in the sun-facing (high thermal load) position. Similarly, high-emissivity materials are efficient under high thermal load conditions in dissipating the heat to maintain a nominal temperature but cool down fast below the nominal temperature in the earth-shade position. This behavior is shown in Fig. 4.40 for high- and low-emissivity materials along with the expected nominal temperature limits for satellites [98]. A system with dynamic control of satellite temperature is, thus, very essential. The present strategies of heat management (thermal control systems; TCSs) in satellites include mechanical or electrostatic louvers, i.e., plates or windows that open and close in the sunfacing and sun-obscuring (earth shade) positions, respectively, to regulate the emissivity of the satellite surface. In terms of their thermal load management performance these are quite efficient as they provide a 1:7 heat loss ratio between their closed (earth shade) and open (sun facing) positions, respectively. However, they are bulky, expensive, and contain moving parts. Other TCS technologies include specialized paints, micromachined louver systems using microelectromechanical systems, heaters and radiators, heat pipes pumped with cooling fluids, and more. Each of these solutions is marred with some or other limitation concerning their cost, bulkiness, operation, weight, and design complexity. These limitations have given way to thin film-based smart radiator devices (SRDs) with their many advantages. Such thin

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films with dynamic behavior with respect to their optical and, hence, radiative properties have many advantages. Such thin film systems can be applied directly on the satellite surface. Being thin film solutions, they add hardly any weight to the satellites and change the surface properties very effectively, either passively or actively with no moving parts [99]. Two such thin film-based SRD technologies that are being sought after very actively are the electrochromic and the TC devices. Although the electrochromic thin film-based SRD systems [100e103] are very promising, the fact that they are inherently multilayered (five or more) systems and actively (electric field) operated may give them less potential for thermal management applications in satellites than the TC thin film-based SRDs. These systems are, generally, single layered and respond dynamically in a passive manner to the changing thermal environment. Hence, they may be deemed more suitable for the application for thermal management in satellites. As discussed before, the optical properties of vanadium dioxide (VO2) films change very effectively above and below the Tt. These thin films exhibit a semiconductor (low reflectance) to a metallic (high reflection) state transition when this temperature is exceeded. As described later in detail, the emissivity [ε(l)] is intimately related to the optical properties and hence an ambient-temperature-triggered dynamic control of emissivity can be expected. As also seen earlier, the bulk of the optical property change occurring from the TC behavior in VO2 is in the IR region, which is very pertinent to the thermal management application in satellites. Further, the two main factors that blemish or limit the application potential of VO2 thin films, requirement of low Tt and high Tlum, are not relevant to their application in thermal management. The sun-facing and sun-obscuring positions clearly render the temperatures of the satellites above and below the natural Tt (68 C) of VO2. In view of these virtues of TC VO2 thin films a lot of research and development is being conducted to make this SRD approach viable for thermal management application in satellites. 4.2.2.1 Design Considerations Similar to the specific requirements detailed earlier for smart window applications of VO2 thin films, the design considerations for their application in the thermal management of satellites must take into account the related optical and radiative properties of these films in their high-temperature (>Tt, rutile) and low-temperature (Tt) decreases with increasing film thickness, changing from around 75% at a thickness of 30 nm to 25% for film thicknesses above 120 nm. Correspondingly, the switch in emissivity

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Figure 4.44 (A) Optical constants (n, k) spectra of the optimized VO2 film above (380K) and below (300 K) the transition temperature (Tt). (B) VO2 film thickness dependence of the integrated emissivity at temperatures above (380K) and below (300 K) the Tt. (From A. Hendaoui, N. Émonde, M. Chaker, É. Haddad, Highly tunable-emittance radiator based on semiconductor-metal transition of VO2 thin films, Applied Physics Letters 102 (6) (2013), 061107/1e4.)

between the low-temperature semiconductor state and the hightemperature metallic state increases with increasing film thickness as shown in the figure. Hence, the optimized RF-deposited VO2 film of 120 nm thickness prepared at a substrate temperature of 440 C is reported to show the best TC switching behavior and is the candidate with the most potential for IR applications. However, it is to be noted that the emissivity values were obtained by integrating over a limited IR region (8e12 mm),

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while for SRD applications, as discussed earlier, the blackbody radiation extends far beyond this region. However, the efficient optical changes (reflectance, n and k) occurring in these films attest to the rich potential of these films for SRD applications. Ali Hendaoui et al. [109] have reported a highly efficient SRD based on VO2 thin films. This device consists of an 850-nm-thick SiO2 film sandwiched between an outer active coating of VO2 (30 nm) and an inner thick metallic coating of gold (350 nm) as shown in the SEM cross section in Fig. 4.45A (a). The material composition and crystalline structure of these films was confirmed by XRD studies shown in Fig. 4.45A (b). The granular structure of the thin VO2 film was confirmed by SEM and AFM studies shown in Fig. 4.45A (c) and (d), respectively. The average grain size is estimated to be around 23 nm and average height, around 40 nm. The dynamic change in emittance (ε) of such a sandwich structure due to the TC behavior of the active VO2 film with temperature was calculated from the total reflectance measurements on the structure in the wavelength region between 2.5 and 25 mm. The emittance spectrum of the device is shown in Fig. 4.45B at low temperature (Tt). In the low-temperature semiconducting state of the VO2 film, the device exhibits a high reflectance in most of the IR spectral region due to the high transmittance of the VO2 film and high metallic reflectance of the Au film. Hence, the device has low emittance in this region except for the sharp broad peak associated with the SiO2 absorption (emittance) around 10 mm. A very efficient emittance change (Dε) can thus be achieved through most of the spectral region studied. The dynamic change of the near-normal total emittance (integrated over all wavelengths studied) is shown in Fig. 4.45C, which gives a better sense of the application potential of the device as an SRD. The device exhibits an overall emittance of around 22% up to a temperature of 75 C and suddenly rises to around 71% at higher temperatures, as shown in this figure. The hysteresis cycle indicates the full reversibility of this change based in the VO2 film thermochromism. The inset of Fig. 4.45C shows the time derivative of the emittance change. While a sharp increase in emittance is seen in the heating cycle, the return to original state is much broader during the cooling cycle, exhibiting a thermal inertia. From these derivative values fitted to a Gaussian curve, the average Tt is found to be around 66.5 C. An overall emittance switch (Dε) of around 49% is found. The greatest advantage of this device for its application as an SRD is the simplicity of the configuration and its preparation. However, a much improved Dε and a lower Tt

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Figure 4.45 (A) Scanning electron microscopy (SEM) cross section image (a) and X-ray diffraction (b) of the VO2/SiO2/Au sandwich structure. SEM cross section (c) and atomic force microscopy (d) images showing the granular structure of the VO2 film. (B) Nearnormal emittance spectra of the VO2/SiO2/Au structure at temperatures above (100 C) and below (25 C) the transition temperature. (C) Thermal hysteresis of the total emittance of the VO2-based structure along with the temperature derivative of emittance (dε/dt) in the inset. (From A. Hendaoui, Nicole Émonde, M. Chaker, Émile Haddad, Applied Physics Letters, 102 (2013) 061107e1-4.)

may be desirable to make this configuration a viable candidate for this application. Work related to the doping of the VO2 film to decrease the Tt is proposed by the authors [109]. Substituting the dielectric SiO2 layer with a suitable film with no absorption in the mid-IR region around 10 mm may be a great advantage in enhancing the emissivity change (Dε), because this spectral region is of paramount importance for satellite application considerations.

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

Many other doped and undoped VO2-based and multilayered thin film systems have been proposed for the SRD application for satellites by the same group of researchers, Haddad et al. [98,110]. An SEM micrograph of their undoped VO2 film prepared by reactive pulsed laser deposition is shown in Fig. 4.46A. The VO2 films were deposited on dielectric (quartz) as well as aluminum substrates with surface nanostructure. Such systems with VO2 films were reported to show distinctly different TC switching behavior as shown in Fig. 4.46B and C. The best of these undoped VO2based samples show a low-temperature (Tt) emissivity in the range of 55%e75%. The emittance change (Dε) thus provided was between 35% and 45% depending on the sample. These authors carried out the doping of VO2 samples with tungsten (W) and/or titanium (Ti) atoms to decrease the Tt.

Figure 4.46 (A) Scanning electron microscopy image of a reactive pulsed-laserdeposited VO2 film. (B, C) Infrared reflectance spectra of a VO2 film deposited on (B) treated aluminum and (C) quartz substrates. (http://esmat.esa.int/materials_news/ isme09/pdf/4-New/S4%20-%20Haddad.pdf.)

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The Tt was found to decrease up to 20 C with 3% W doping. However, the efficiency of the TC switch decreased with such doping as shown in Fig. 4.47A. A Dε value of around 34% was maintained in 1.5% W-doped VO2 film with a Tt as low as 28 C. Based on this work and the low solar absorptance desired for SRD during sun-facing hours, multilayered stacks were proposed in this work. A schematic of such a multilayered VO2-based SRD is shown in Fig. 4.47B. The configuration is expected to optimize both the IR emittance and the solar absorptance. The thin silver film is added to further increase the solar wavelength reflectance (minimize solar absorptance) of the textured aluminum substrate. The triple interference stacks of W/VO2 and SiO2 of various thicknesses are then added to bring about an efficient emittance tunability at low temperature. The innermost SiO2 interference coating is optimized to provide high reflectance around the peak solar wavelength of 500 nm while the outer thick SiO2 coating provides overall high emissivity. The simulated reflectance behavior of such a multilayer stack is as shown in Fig. 4.47C. An overall high reflectance (low absorptance) is thus expected by this multilayered stack throughout the solar wavelengths. Compared to the single VO2 film on Al substrates, the multilayer stack is reported to exhibit superior performance in both the solar and the IR regions of the spectrum. The Dε value increases from 34% to 36% while solar absorptance decreases from 57% to 32% in the multilayered stack. Both the single VO2-layered and the multilayered structures have been subjected to space and space simulation (thermal cycling, vibrations, atomic oxygen, etc.) tests. The stability of the samples under these conditions and their suitability for the space environment are verified.

4.2.3 Other Applications VO2 thin films and their devices are also studied extensively for other applications because of their efficient and reversible changes in optical and electrical properties as a function of temperature. Following are some such examples. Bonora et al. [68] have studied and proposed the use of the optical switching in VO2 thin films for mid-IR to near-IR conversion and image detection. The experimental setup used in this work is shown in Fig. 4.48A in which mid-IR light (1940-nm Tm fiber laser or CO2 laser at 10.6 mm) was used for the writing by thermochromically induced local changes in the VO2 layers. A near-IR-reading light (850-nm diode laser or 1064 Nd:YAG laser) was sent through a beam splitter to propagate along with the mid-IR beam as shown in the figure. The three-dimensional

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

Figure 4.47 (A) Thermal hysteresis of infrared transmittance of an undoped (a), Wdoped (b), and W/Ti-doped (c) VO2 film. (B) Schematic of a multilayered VO2-based smart radiator device (SRD). (C) Simulated reflectance of a multilayered VO2-based SRD. ESD, electrostatic discharge. (http://esmat.esa.int/materials_news/isme09/pdf/4New/S4%20-%20Haddad.pdf.)

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profile of the writing beam as recorded by the detector is shown in Fig. 4.48B. This clearly indicates the kind of setup that can be used for beam profile characterization by using lasers from visible to mid-IR, where at room temperature a transmittance of around 25% (higher visible) to 50% (mid-IR range) is shown by the VO2 film. In the switched state (>Tt), the VO2 film transmittance decreases in the lower near-IR wavelengths and falls to near zero beyond the wavelength of 1.5 mm (as shown in the inset of Fig. 4.48A). The response sensitivity of the experimental setup was tested for such beam detection experiments by decreasing the temperature gap between the VO2 sample surface and its Tt. A Peltier temperature control, as shown in Fig. 4.48A, was used to achieve this.

Figure 4.48 (A) Schematic of the experimental setup used for the beam profile characterization, mid-IR to near-IR (NIR) conversion, and high-resolution imaging. Inset: The transmittance switch of the VO2 sample between low (25 C) and high (75 C) temperatures. (B) Beam profile of a Tm fiber laser (mid-NIR) detected with an 850-nm diode laser (NIR). (From S. Bonora, et al., Optics Letters 103 (2010) 103e105.)

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

In Fig. 4.49A and B are compared the response times of the VO2-based device held at different temperatures. As can be seen in this figure, at room temperature, under the application of a square Tm laser pulse of 5 ms duration, the change in transmittance is rather slow and occurs only with a threshold power of 54 W/cm2. This response increases in terms of both response speed and change in transmittance with increasing laser power impinging on the VO2 surface. The change was tested up to a 10-fold decrease in transmittance achieved with 109 W/cm2. However, increasing the VO2 sample temperature to near Tt, the same change can be achieved with saturation with a laser power of a mere 12.15 W/cm2. Hence the response speed and the extent of transmittance change can be achieved at much lower laser power. The threshold power required to induce the TC change in VO2 was as low as 144 mW/cm2 in this situation. Work was also carried out to demonstrate the potential of this basic idea of using VO2 optothermochromism at optimized laser power for high-resolution imaging. A target mask was placed in front of the VO2 sample held at 53 C in the path of the illuminating Tm laser beam. The fine details of the mask were detected up to 35 mm as shown in Fig. 4.49C (a). The image resolution was limited by the CCD imaging system. High-resolution image conversion was also demonstrated by combining various pumpeprobe lasers for writing and reading processes. These images are shown in Fig. 4.49C (bed). This work clearly proves the rich potential of such VO2based devices for conversion of mid-IR to near IR and their detection for imagery. This VO2 approach has a great advantage over expensive and elaborate image detecting technologies currently available, especially in the spectral region between 1.7 and 8 mm. Based on the same approach, Bonora et al. [111] have proposed the method of wavelength conversion from the IR to the visible based on the refractive change undergone by the VO2 film during the TC switching. Beam characterization of IR lasers in terms of both beam profile and wavefront using visible detectors is demonstrated. The optical switching of VO2 films between the dielectric and the metallic states was induced not only by thermal effects but also by injection of electric charge or application of voltage. Hence, the VO2 film-based optical valves can be closed or opened to control the transmittance of light beams [98]. They were employed as multiple slits in a Hadamard mask to replace mechanically operative slits in spectrophotometers. Such VO2based Hadamard masks are very advantageous owing to the reduction in signal-to-noise ratio and the absence of moving parts. An example of this is

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(A)

(B)

(C) (a)

(b)

(c)

(d)

Figure 4.49 Comparison of the response of the VO2-based detection device at (A) low (22 C) and (B) high (62 C) temperatures for various writing laser powers (midinfrared). (C) High-resolution images of masks held in front of the impinging beams: (a) Tm laser detection through mask with VO2 held at 53 C; (b) Tm laser detection through another mask with VO2 held at 60 C; and detection of a CO2 laser through a cross (c) and a circular (d) aperture with VO2 held at 40 C. (From S. Bonora, et al., Optics Letters 103 (2010) 103e105.)

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Figure 4.50 (A) A Hadamard shutter based on electrically activated VO2 thin films. (B) VO2 thin films on flexible substrates such as a Ge-coated Kapton sheet. (http://esmat. esa.int/materials_news/isme09/pdf/4-New/S4%20-%20Haddad.pdf.)

shown in Fig. 4.50. Similarly, the TC switching of VO2 films can be used for temperature swing control of astronaut suits if they can be deposited on flexible substrates. The performance of such VO2 films on Ge-coated Kapton sheets has been studied by Haddad et al. [98]. An example of such a flexible film is shown in Fig. 4.50B.

4.3 SUMMARY TC materials showing a reversible phase transition at a critical temperature (Tt) are very interesting from the application point of view as well as from the point view of the study of the fundamental physics and chemistry involved.

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TMOs exhibit a very convenient form of TC change whereby the optical, electrical, and other properties undergo a reversible change. Hence, they are very conducive to building interactive smart systems whose behavior can be controlled via the temperature to which they are exposed. Among the TMOs it is the vanadium oxide (VxOy) series that exhibits a very efficient TC change at various Tt’s. Vanadium dioxide (VO2), for which the Tt is around 68 C in bulk form, exhibits a large change in optical and electrical properties due to the phase change at this temperature from a semiconducting state to a metallic state. At low temperature (Tt) range, it exhibits a high reflectance and a low resistivity. The material exhibits a thermal hysteresis of these properties. It is the proximity of the Tt (68 C) of VO2 to room temperature as well as the possibility of further decreasing this temperature that has made it a very valuable material for scientific research and application. The optical, electrical, and structural changes occurring simultaneous with temperature have been examined for a long time in the light of two principal models, (1) the MotteHubbard insulator (changes occurring are due to the electronic structure) and (2) the Peierl insulator (changes occurring are due to the periodic variation of the crystal structure), and the combination of the two. In view of the large discrepancies experienced in fitting the theory and the experimental results, some authors [27] suggest a judicious choice of theoretical model in the light of the experimental conditions used. More recently, the TC phenomenon has been explained more concretely on the basis of the large phonons created at high temperature (>Tt) that drive VO2 into a stable rutile phase. Despite the phenomenal success of this approach in fitting the theory and the data, it has not received the attention it deserves. A major opportunity exists for researchers in this field to take up this approach based on the electron and phonon contributions to the MIT. Hence, the future work in this area will necessarily have to be examined in the light of this new thermodynamic approach, taking into account the contribution to the total entropy in any system of the electrons and the phonons in VO2 below and above the Tt. The thin film and nanostructured form of VO2 gives the possibility of altering the phase transition behavior. The degree of the optical and electrical change associated with the semiconductor-to-metal transition can vary depending on the film thickness and the nanostructure. The thermal hysteresis exhibited by VO2 also depends very sensitively on the film structure. In addition, the natural Tt of 68 C found in bulk VO2 can be

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altered both by preparing the film in nanostructured form and by doping the films with compatible atoms. The factors affecting TC behavior are doping, nanostructure, nature of the substrate, defects, and impurities. Hence, an enormous opportunity exists for research on VO2 in thin film form by varying these factors. The analysis of thermal hysteresis loops through FORC diagrams itself forms an interesting study on VO2 films. Simultaneous measurement of optical and electrical hysteresis loops can throw a lot of light on the MIT. Multiple steps in the hysteresis loops can reveal multiple phase transitions within the same loop. Nanostructuring and doping of VO2 by Mg atoms is the most effective way of reducing the Tt. Nanothermochromics is a relatively new approach in which VO2 nanoparticles of various shapes and sizes are integrated into other dielectric materials to form nanocomposites to enhance the TC performance. Although theoretically very appealing, the nanothermochromics approach is very challenging from an experimental point of view. From the experimental point of view more success is achieved with hybrid VO2-based structures with polymeric hydrogels to simultaneously enhance the visible transmittance and near-IR switching. An opportunity exists to work with doped VO2 films in these hybrid laminates to further enhance their application potential. Despite the vast amount of research on VO2 thin films, one of the main challenges that persists is the deposition of stoichiometric thin films, as VO2 forms in a very narrow and strict range of deposition parameters. This needs to be overcome if large-scale application of VO2 thin films is envisioned. Several decades of work on VO2 have led to a wide range of methods by which the thin films can be prepared now. The thermally activated reversible changes in optical and electrical properties occurring in these films have become the focus of a wide range of smart devices such as smart windows (solar energy management), SRDs (thermal management in satellites), optical switches, and more.

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[69] M. Saeli, et al., Nano-composite thermochromic thin films and their application in energy-efficient glazing, Solar Energy Materials and Solar Cells 94 (2010) 141e151. [70] S. Chen, H. Ma, J. Dai, X. Yi, Nanostructured vanadium dioxide thin films with low phase transition temperature, Applied Physics Letters 90 (2007), 101117/1e3. [71] M.-H. Lee, J.-S. Cho, Better thermochromic glazing of windows with anti-reflection, Thin Solid Films 365 (2000) 5e6. [72] P. Jin, G. Xu, M. Tazawa, K. Yoshimura, Design, formation and characterization of a novel multifunctional window with VO2 and TiO2 coatings, Applied Physics A 77 (2003) 455e459. [73] N.R. Mlyuka, G.A. Niklasson, C.G. Granqvist, Thermochromic VO2-based multilayer films with enhanced luminous transmittance and solar modulation, Physica Status Solidi A 206 (9) (2009) 2155e2160. [74] M.S.R. Khan, K.A. Khan, W. Estrada, C.G. Granqvist, Electrochromism and thermochromism of LixVO2 thin films, Journal of Applied Physics 69 (5) (1991) 3231e3234. [75] D.W. Murphy, F.J. DiSavo, J.N. Caridas, J.V. Waszezak, Topochemical reactions of rutile related structures with lithium, Materials Research Bulletin 13 (12) (1978) 1395e1402. [76] D.W. Murphy, et al., Lithium incorporation by V6O13 and related vanadium (þ4, þ5) oxide cathode materials, Journal of Electrochemical Society 128 (10) (1981) 2053e2060. [77] N.R. Mlyuka, G.A. Niklasson, C.G. Granqvist, Mg doping of thermochromic VO2 films enhances the optical transmittance and decreases the metal-insulator transition temperature, Applied Physics Letters 95 (17) (2009), 171909/1-3. [78] S. Li, VO2-Based Thermochromic and Nanothermochromic Materials for EnergyEfficient Windows (Ph.D. thesis), Uppsala University, 2013. [79] S.-Y. Li, G.A. Niklasson, C.G. Granqvist, Nanothermochromics: Calculations for VO2 nanoparticles in dielectric hosts show much improved luminous transmittance and solar energy transmittance modulation, Journal of Applied Physics 108 (2010) 663525e663528. [80] S.-Y. Li, G.A. Niklasson, C.G. Granqvist, Nanothermochromics with VO2-based core-shell structures: Calculated luminous and solar optical properties, Journal of Applied Physics 109 (2011), 113515/1-5. [81] J. Zhang, et al., Nanoscience and Nanotechnology 10 (2010) 2092. [82] S.A. Pauli, et al., X-ray diffraction studies of the growth of vanadium dioxide nanoparticles, Journal of Applied Physics 102 (2007) 073527. [83] L. Chen, et al., Synthesis of thermochromic W-doped VO2 (M/R) nanopowders by a simple solution-based process, Journal of Nanomaterials (2012) 8, 491051. [84] H. Miyazaki, et al., Thermochromic tungsten doped VO2-SiO2 nano-particle synthesized by chemical solution deposition technique, Journal of the Ceramic Society of Japan 117 (2009) 970e972. [85] E. Strelcov, Y. Lilach, A. Kolmakov, Gas sensor based on metal-insulator in VO2 nanowire thermistor, Nano Letters 9 (6) (2009) 2322e2326. [86] Y.-Q. Wang, et al., Nanostructured VO2 photocatalysts for hydrogen production, American Chemical Society (ACS) Nano 2 (7) (2008) 1492e1496. [87] X. Cao, et al., Science of Advanced Materials 6 (2014) 558e561. [88] X. Cao, et al., Nanoporous thermochromic VO2 (M) thin films: Controlled porosity, largely enhanced luminous transmittance and solar modulating ability, Langmuir 30 (2014) 1710e1715. [89] C. Liu, et al., VO2/Si-Al gel nanocomposite thermochromic smart foils: Largely enhances luminous transmittance and solar modulation, Journal of Colloid Interface Science 427 (2014) 49e53.

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[90] Y. Gao, et al., VO2-Sb:SnO2 composite thermochromic smart glass foil, Energy & Environmental Science 5 (8) (2012) 8234e8237. [91] A. Taylor, et al., A bioinspired solution for spectrally selective thermochromic VO2 coated intelligent glazing, Optics Express 21 (2013) A750eA764. [92] X. Qian, et al., Bioinspired multifunctional vanadium dioxide: Improved thermochromism and hydrophobicity, Langmuir 30 (2014) 10766e10771. [93] Z. Chen, et al., Solar Energy Materials and Solar Cells 95 (2011) 2677e2684. [94] Y. Zhou, Y. Cai, X. Hu, Y. Long, VO2/hydrogel hybrid nanothermochromic material with ultra-high solar modulation and luminous transmission, Journal of Materials Chemistry A 3 (2015) 1121e1126. [95] Y. Chen, X. Tang, B. Chen, G. Qiu, Low temperature plasma vapor treatment of thermo-sensitive poly(N-isopropylacrylamide) and its application, Applied Surface Science 268 (2013) 332e336. [96] Y. Zhou, Y. Cai, X. Hu, Y. Long, Temperature-responsive hydrogel with ultra-large solar modulation and high luminous transmission for “smart window” applications, Journal of Materials Chemistry A 2 (2014) 13550e13555. [97] D.G. Gillmore, Satellite Thermal Control Handbook, The Aerospace Corporation Press, El Segundo, CA, 1994. [98] E. Hadded et al http://esmat.esa.int/materials_news/isme09/pdf/4-New/S4%20-%20Haddad.pdf. [99] E. Haddad, et al., Passive dynamically dynamically variable Thin-Film Smart Radiator Device, in: 4th Round Table on Micro/Nano Technologies for Space, 2003, p. 13. [100] H. Demiryont, D. Moorehead, Electrochromic emissivity modulator for spacecraft thermal management, Solar Energy Materials & Solar Cells 93 (12) (2009) 2075e2078. [101] H. Demiryont, Electrochromic heat modulator successfully tested in space, SPIE e Newsroom (2008) 3, http://dx.doi.org/10.1117/2.1200805.1147. [102] J.S. Hale, M. De Vries, B. Dworak, J.A. Woollam, Visble and infrared optical constants of electrochromic materials for emissivity modulation applications, Thin Solid Films 313-314 (1998) 205e209. [103] E.B. Franke, et al., Infrared switching electrochromic devices based on tungsten oxide, Journal of Applied Physics 88 (10) (2000) 5777e5784. [104] A. de Rooji, http://esmat.esa.int/publications/published_papers/corrosion_in_space. pdf. [105] http://www.fp7-spacecast.eu/help/bg_sa.pdf. [106] Tsoline Mikaelian, https://arxiv.org/ftp/arxiv/papers/0906/0906.3884.pdf. [107] R.O. Dillon, K. Le, N. Ianno, Thermochromic VO2 sputtered by control of a vanadium-oxygen emission ratio, Thin Solid Films 398-9 (2001) 10e16. [108] F. Guinneton, L. Sauques, J.-C. Valmalette, F. Cros, J.-R. Gavarri, Optimized infrared switching properties in thermochromic vanadium dioxide thin films: role of deposition process and microstructure, Thin Solid Films 446 (2004) 287e295. [109] A. Hendaoui, N. Émonde, M. Chaker, É. Haddad, Highly tunable-emittance radiator based on semiconductor-metal transition of VO2 thin films, Applied Physics Letters 102 (6) (2013), 061107/1-4. [110] https://escies.org/download/webDocumentFile?id¼1110. [111] S. Bonora, G. Beydaghyan, A. Haché, P.V. Ashrit, Mid-IR laser beam quality measurement through vanadium dioxide optical switching, Optics Letters 38 (9) (2013) 1554e1556.

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FURTHER READING [1] M. Soltani, M. Chaker, E. Haddad, R. Kruzelecky, 1  2 optical switch devices based on semiconductor-to-metallic phase transition characteristics of VO2 smart coatings, Measurement Science & Technology 17 (5) (2006) 10.

CHAPTER 5

Photochromic Thin Films and Devices 5.1 BASICS OF PHOTOCHROMIC MECHANISM Photochromic materials are those that show a reversible change in optical properties (color) through the action of light, i.e., electromagnetic radiation. They belong, therefore, to the chromogenic class of materials. Similar to the electrochromic and thermochromic materials seen earlier, photochromic materials have become increasingly popular owing to the interactive nature of the coloration they offer. Historically, a wide range of photochromic effects have been seen in various materials since the 18th century [1e4]. Both organic and inorganic materials exhibiting photochromism exist, although the coloration mechanism differs from one to the other. The most well-known example of the application of photochromic materials are photochromic sunglasses. Inorganic component-based photochromic sunglasses are made with silver halide inclusions in glass lenses. Under the incidence of UV light between the wavelengths of 320 and 400 nm, the electrons in the glass are transferred to the silver halide grains, precipitating elemental silver. It is the elemental silver that blocks the light through absorption, turning the sunglasses from transparent to dark. When the photoactivation is removed (under darkness), the process is reversed and silver returns to its transparent halide state, turning the glass transparent. In lenses using organic materials, however, the coloration process is due to organic dyes such as pyridobenzoxazines or naphthopyrans or indenonaphthopyrans [5]. In these dyes the incidence of UV radiation causes the breaking down of certain bonds that create new organic bonds that absorb light of higher wavelengths such as the visible ones. This being a reversible transformation, the light action can be interactively used for coloring and bleaching the lenses. For efficient optical transformation these organic dyes need to be used with compatible monomers. World-renowned photochromic lens manufacturers such as Transitions Optical have been working constantly to improve the performance of nonglass and dye-based lenses and Transition Metal Oxide Thin FilmeBased Chromogenics and Devices ISBN 978-0-08-101747-0 http://dx.doi.org/10.1016/B978-0-08-101747-0.00005-2

© 2017 Elsevier Ltd. All rights reserved.

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significant improvements have been achieved in terms of response speed of optical transition and their durability. In the context of inorganic material-based photochromism, transition metal oxides (TMOs) are a viable alternative. The multiple oxidation states of some of these oxides offer an interesting and efficient way to achieve photochromism. Especially, in thin film form, and like other chromogenic species, a wide range of photochromism can be envisioned in the TMOs. The photochromic effect in TMOs is similar to the electrochromic effect occurring under the injection of ions and electrons. In the electrochromic coloration of TMOs the ions and electrons needed for the phenomenon come from a storage layer and, under the action of an applied external electric field, are injected into the TMO. In photochromic coloration, electronehole pairs are created under the action of light and it is these electrons and holes that in turn trigger the coloration. Subsequent to the creation of the charge species the insertion and extraction of these species as well as the intervalence charge transfer (IVCT) between the transition metal ions takes place, similar to the electrochromic coloration as described in Chapter 3 [6]. Among the TMOs, molybdenum trioxide (MoO3) exhibits the most efficient coloration under UV light irradiation. The photochromic coloration mechanism in TMOs and the underlying theoretical models can be well discussed in terms of the MoO3 thin films. Under UV light irradiation the initially transparent thin films of MoO3 undergo a deep blue coloration. The degree of coloration increases with intensity and duration of the incident UV light. An example of this is shown in Fig. 5.1A for a 200-nm-thick MoO3 film irradiated with UV light of 121 mW/cm2 over different durations. The transmittance change corresponding to this photochromic coloration is shown in Fig. 5.1B. A deep blue coloration whose peak absorption is centered around 900 nm is seen. The photochromic coloration was also found to be dependent on the film’s nanostructure [7]. Similar to the electrochromic effect seen in tungsten trioxide (WO3) and other TMO thin films, the optical absorption in MoO3 is suggested to emanate from different sources: (1) due to substoichiometric MoO3 (MoO3
x) and/or (2) due to
 the formation of hydrogen molybdenum bronze V Hx MoVI Mo O 1
x x 3 , as well as (3) from small polaron absorption. Accordingly various theoretical models are put forth to explain this coloration. In the case of substoichiometric thin films, depending on the method of preparation, they contain various types of defects, including oxygendeficient sites. This possibility is even more in the case of amorphous molybdenum oxide films. This high density of defects is suggested to create

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(A)

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1300 1800 Wavelength ( nm )

2300

Figure 5.1 (A) Schematic (top) and actual coloration of a 200-nm-thick MoO3 film (bottom) irradiated with UV light for various durations: 1 ¼ 1 min; 2 ¼ 5 min; 3 ¼ 10 min; and 4 ¼ 15 min. (B) Transmittance spectra of MoO3 film corresponding to different UV irradiation times. (From G.-G. Allogho, P.V. Ashrit, Thin Solid Films 520 (2012) 2326e2330. Copyright 2008: Elsevier.)

a defect band within the forbidden band in this semiconductor [8]. Under UV radiation illumination of these films, some of the photogenerated valence band electrons can be excited to these intermediate defect bands and can be easily excited to the conduction band. Hence, F-center formation is thought to be the origin of the mechanism of photochromic coloration. Some of the other aspects that support this theory are the blue coloration and high conductivity of the films that disappear under thermal treatment. The disappearance of the coloration under thermal treatment is, generally, not a reversible phenomenon, as the thermal treatment, especially in the presence of oxygen, restores the stoichiometry of the film in addition to removing the other defect-based color centers. Multiple band coloration can also be seen in some cases, which is expected to arise from different types of defects acting as color centers. The absorption behavior

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arising from these photoelectrons trapped in defects follows Smakula’s equation [9]: n NS fS z 0:87  107 ap d U12 2 ðn þ 2Þ2 where the various parameters are as follows: Ns is the number of color centers per square centimeter, fs is the oscillator strength, n is the refractive index of the uncolored material, ap is the peak absorption coefficient (cm
1), d is the film thickness, and U12 is the full width at half maximum of the generally Gaussian-shaped absorption curve. This explanation of photochromism on the basis of the color center formation, which is purely an electronic behavior approach, is marred by many doubts, similar to its application to explain the electrochromic behavior in TMOs. The theory is applicable in a very narrow range of experimental conditions and has a limited success [11]. A more successful and widely accepted theory is the double insertion/ extraction model, which is associated with IVCT. According to this model, the photochromic phenomenon is similar to the electrochromic coloration occurring in the TMOs. In the electrochromic materials, the double insertion occurs through the electrons and ions from external sources under the application of an electric field. However, in photochromic materials such as MoO3 these are created within the films under study with the UV irradiation. Electrons and holes are created in the film under the impingement of the UV radiation when the light energy is superior to the bandgap (hn > Eg). The holes thus created interact with the water molecules invariably present in the films to form protons (Hþ). The film is now an inherent source of ions and the electrons. The coloration of the film takes place in the presence of these double ion/electron species. The mixed oxidation states of TMOs such as MoO3 and WO3 are conducive to this coloration as the electrons localized on the TMO sites are transferred from one site to another through the absorption of light in the visible and near-infrared region, creating the absorption band and transmittance decrease in this region as shown in Fig. 5.1B. These processes initiated by the UV radiation and leading to the ultimate “photochromic coloration” are as shown below [12]: 1. creation of electronehole pairs under UV irradiation: =

=

MoO3 þ hnðUVÞ / MoO*3 2. formation of ions (protons) and electrons from structural water molecules: H2O þ 2hþ / 2Hþ þ O

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3. double insertion of ions and electrons into MoO3: V MoO3 ðtransparentÞ þ xHþ þ xe
/ Hx MoVI 1
x Mox O3

4. intervalence transfer of electrons localized on one site (MoVI) to another (MoV) through the absorption of light in the visible and near-infrared region: V V VI MoVI A þ MoB þ hnðvis
nirÞ/MoA þ MoB

Some of the major problems associated with photogeneration of electronehole pairs are the following: 1. electronehole pair recombination releasing heat: hþ þ e
/ heat 2. diminishing photochromic coloration efficiency due to formation of oxygen and hydrogen molecules: 1 þ MoO3 þ 2hþ þ 2H2 O / MoO2
4 þ 4H þ O2 2 1 þ
H þ e / H2 2 These side reactions are very unfavorable, especially in the design and proper functioning of the photochromic devices. In addition to the aforementioned water-molecule-dependent photochromism, MoO3 films are also shown to color in vacuum under UV irradiation. This coloration is expected to happen by the creation of a substoichiometric oxides under irradiation as shown below [13]: 1 MoO3 þ hnðUVÞ/MoO3
x þ xO2 2 Another successful model closely related to the IVCT model is the small polaron absorption model. This model is also used in explaining the highenergy tail of the absorption [14] as shown in Fig. 2.17 and discussed in detail in Chapter 2 regarding the electrochromic coloration of amorphous tungsten trioxide films. In photochromically activated MoO3 films, for example, the photogenerated electrons can get trapped on a low-energy Mo6þ site and create a polaron because of the excess charge in the lattice around this region. Under the incidence of light in the visible and nearinfrared region the polaron hopping takes place from site to site as in

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case (4) mentioned earlier, bringing about the blue coloration. The polaron absorption spectrum, which resembles a Gaussian curve, is given by: " # 2
1 ðZu
n
4Up Þ apolaron ðuÞfu 8Up Zu0 where Zu0 is the excess phonon energy given off to the lattice in the photon-assisted hopping, and Up is the energy gained by the lattice polarization. The various parameters are as indicated in Fig. 2.16 for a similar situation for the transition between nonequivalent W sites.

5.2 PHOTOCHROMIC CHARACTERIZATION METHODS Similar to the electrochromic effect, a number of changes occur simultaneously in photochromic thin films and materials upon UV irradiation. Optical (transmittance, absorptance, optical constants) and electrical (electrical conductivity) properties and chemical composition change quite drastically. Hence, it is important to follow all the changes to characterize the effect.

5.2.1 Optical Measurements Spectrophotometry is one of the most common characterization methods used in this work. Measurement of the change in transmittance or absorption of the films as functions of (1) the duration and/or (2) the intensity of UV irradiation is most commonly carried out. This work is mostly carried out on samples prepared under various conditions to correlate the photochromic performance with the preparation conditions and the microstructure of the film. These optical measurements also yield information on various other aspects of the samples, such as the semiconductor bandgap energy, Eg. This intrinsic bandgap of a semiconducting film is related to its absorption coefficient (a ¼ 4pk/l) by: ahn f ðhn
Eg Þh where hn is the photon energy. For amorphous semiconductors, the1 value of h being 2 [9], Eg corresponds to the x axis intercept in a plot of ðahnÞ2 as a function of the photon energy (hnÞ. The value of h is 0.5, 1.5, 2, or 3 according to whether the optical transitions are directeallowed, directeforbidden, indirecteallowed, or indirecteforbidden in crystalline semiconductors [10]. An example of this is shown in Fig. 5.2A and B for 100-nm-thick MoO3 films deposited at different temperatures. The absorption spectra for these films deposited at different substrate temperatures were initially measured. From

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Figure 5.2 (A) Optical absorbance spectra of 100-nm-thick MoO3 film deposited at substrate temperatures of (curve 1) 20 C, (curve 2) 100 C, (curve 3) 230 C, and (curve 4) h i 1 1 275 C. (B) Plot of ðaZuÞ2 orðahnÞ2 versus photon energy, Zu (or hn) for these samples deposited at different substrate temperatures. (From C.G. Granqvist, Handbook of Inorganic Electrochromic Materials, Elsevier, Amsterdam, 1995.)

these data the plot as shown in Fig. 5.2B can be obtained and the bandgap 1 determined from the intercept on the photon energy axis of the ðahnÞ2 plot. Such plots are generally used to discern the dependence of the bandgap (Eg) on the film preparation conditions.

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Ellipsometry is also a powerful tool for characterizing the fundamental optical parameters of the films, such as the refractive index (n) and the absorption coefficient (k). This technique is also used quite often in measuring the film thickness. All the measurements are carried out on the films before and after the exposure to UV radiation. The optical constants obtained from this technique can also be used in estimating the porosity of the film. The porosity p can be estimated by comparing the refractive index of the film (nf ) at a given wavelength with that of the bulk value (n) through the relation [14]: 
2 nf
1 ¼1
p ðn2
1Þ As far as the optical measurements are concerned, the photochromic response can also be measured in some cases by the relative changes in transmittance (T) and/or reflectance (R) as [15]: DT Ta
Tb DR Ra
Rb and ¼ ¼ T R Tb Rb where the subscripts a and b represent, respectively, the optical parameters after and before the photochromic coloration.

5.2.2 Chemical Composition Measurements The determination of the chemical composition of the photochromic film before and after the UV irradiation is also very important as the photochromic efficiency depends on many factors in TMO films. In TMOs such as MoO3, in addition to multiple oxidation states such as Mo5þ and Mo6þ, the structural and adsorbed water molecules, various defects such as oxygen vacancies also play important roles in inducing the photochromic efficiency. As discussed before it is important to discern the intrinsic coloration of the “as-deposited” films arising from the oxygen vacancies and the coloration arising from the UV irradiation. Under the action of UV light and with the insertion of photogenerated electrons the concentration of Mo5þ increases. The ensuing increase in blue coloration is due to the increased formation of Mo5þ. Hence, the measurement of the relative concentration of Mo5þ and Mo6þ before and after irradiation is an important aspect in the photochromic characterization. X-ray photoelectron spectroscopy (XPS) and Raman spectroscopy are powerful techniques in this regard. While the XPS measurements give the chemical composition of the film only at the surface, Raman spectroscopy is preferable for more precise evaluation of the chemical nature of the entire film. This factor is even more important in

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Figure 5.3 X-ray photoelectron spectroscopy spectra of Mo 3d structure of MoO3
x films with (A) x ¼ 0.08, (B) x ¼ 0.18, (C) x ¼ 0.3, and (D) x ¼ 0.4, before and after UV irradiation. In each case, solid line, raw data; dashedot line, Gaussian peaks associated with Mo6þ oxidation state; dashed line, Gaussian peaks associated with Mo5þ oxidation state; dotted line, fitted envelope of the Mo6þ and Mo5þ spectra. (Adapted from M. Rouhani, et al., Applied Surface Science 273 (2013) 150e158.)

films such as MoO3 showing surface-dependent photochromism and other sensing activities related to atmospheric humidity. Examples of such XPS measurements on MoO3 samples prepared under different conditions are shown before and after UV irradiation in Fig. 5.3 [15]. From the deconvoluted Gaussian peaks in the XPS spectra of the 3d core relating to the   different oxidation states Mo6þ ; Mo5þ , the variation of peaks relating to Mo5þ with UV irradiation is quite evident. The single spin-orbit doublet peak in binding energy corresponding to the presence of Mo6þ is located around 232.3 and 235.4 eV. The doublet peak corresponding to the

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presence of Mo5þ is found with binding energies of around 231.2 and 234.3 eV. Hence, by fitting the raw data to these binding energy spectra as shown in Fig. 5.3 one can obtain valuable information. The analysis of the data presented here, as an example, shows not only the variation occurring in the film with the photochromic change but also the variation within the as-deposited film as a function of the preparation conditions. Similarly, a more in-depth analysis of the entire film can be carried out through Raman spectroscopy of the films before and after the irradiation. This is shown, as an example, in Fig. 5.4, for the same MoO3 films deposited under different conditions as in Fig. 5.3. The Gaussian peaks corresponding to the three vibration modes of Mo6þ and Mo5þ are shown to be located at 735, 847, and 956 cm
1 and at 670, 905, and 990 cm
1, respectively. These peaks have been fitted with the raw data to analyze the chemical changes occurring in the bulk of the film with the photochromic effect. Based on these measurements and in-depth analyses, the authors [15] concluded the origin of photochromic coloration and tied it to a predominant role in the formation of different types of defects and oxygen vacancies depending on the deposition conditions.

5.2.3 Electrical Measurements As seen earlier, photochromic coloration ensues from the creation of electronehole pairs under UV irradiation. Hence, along with the change in optical properties, electrical properties also change drastically as the film changes from an insulating state to a conducting state. Hence, measurement of electrical resistivity or IeV characteristics of the films before, during, and after UV irradiation can also be an important tool in analyzing the films. Measurements such as the one shown in Fig. 5.5 for photochromic WO3 films can be carried out [16]. The as-deposited WO3 films show no conductivity, while the UV-irradiated films show a nonlinear and drastic increase in conductivity with increasing electric field. Even more interesting is the exponential increase in conductivity with electric field if measured during the irradiation. The authors [16] base their conclusion on the double insertion model. The increase in conductivity in the UV-irradiated state is attributed to the contribution of photogenerated holes while the electrons are trapped on the W sites. The even higher degree of conductivity observed during the process of irradiation is attributed to the contributions of both the electrons and the holes to the conductivity. The measurement of conductivity as a function of time of irradiation can also be very valuable in discerning the photochromic mechanism.

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Figure 5.4 Gaussian peak fitting of the Raman spectra of MoO3
x films with (A) x ¼ 0.08, (B) x ¼ 0.18, (C) x ¼ 0.3, and (D) x ¼ 0.4, before and after UV irradiation. In each case, dotted line, experimental data after subtracting the baseline; solid line, Gaussian peaks associated with Mo6þ vibration bands; dashed line, Gaussian peaks associated with Mo5þ vibration bands. The solid line fitted to the experimental data corresponds to the fitted envelope of the Gaussian peaks. (Adapted from T. He, Fabrication and Properties of Novel Materials with High Photo-electro-chromic Performance (Ph.D. thesis), Institute of Chemistry, Chinese Academy of Sciences, 2002.)

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Figure 5.5 Currentevolt (IeV) curves measured for a thermally evaporated photochromic WO3 film, before, after, and during UV irradiation. (From K. Kim, C. Seo, H. Cheong, Journal of the Korean Physical Society 48 (6) (2006) 1657e1660.)

5.3 PHOTOCHROMIC EFFECTS IN MOLYBDENUM TRIOXIDE- AND TUNGSTEN TRIOXIDE-BASED FILMS Among the TMOs, both tungsten trioxide (WO3) and molybdenum trioxide (MoO3) exhibit very effective electrochromic and photochromic colorations. WO3, with its most efficient and versatile electrochromism, has attracted much more attention than MoO3, although the latter exhibits a nearly equivalent electrochromic efficiency.

5.3.1 Molybdenum Trioxide As far as the photochromic effect is concerned, MoO3 exhibits by far the most efficient coloration under UV irradiation, making it the most attractive material. The semiconductor bandgap of MoO3 films is much lower, being around 2.6 eV [17], although values between 2.8 and 3.2 eV are also reported [18]. The photochromic absorption band is located in the visible to near-infrared region of the spectrum. Hence, the material is of great interest for various applications. The photochromic coloration efficiency in these films depends very strongly on the thin film preparation conditions as well as the film nanostructure. This is due to the strong dependence of the coloration on the availability of the water molecules in these films. These water molecules, either adsorbed at the surface or in the

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bulk of the film, act as the source of ions needed for the coloration occurring via double insertion as discussed earlier. The chemisorbed water molecules are an integral part of the chemical bonding within MoO3 and do not contribute to the photochromic process. Physisorbed water, on the other hand, is present between the grains and the crevices of MoO3 and, hence, easily available for the light-activated breakdown [19,20]. The presence of such water molecules, in addition to providing the ions needed for the coloration, also promotes the rapid mobility of the created protons within such films [21]. The photochromic behavior of nanostructured MoO3 films deposited by the glancing angle deposition (GLAD) method using e-beam evaporation has been studied by Beydaghyan et al. [22]. The nanostructure was varied in these films by deposition at various tilt angles. Atomic force microscopy (AFM) images of the MoO3 films are shown in Fig. 5.6A. An increase in film porosity and nanostructure with increasing tilt angle of deposition is seen in these films. Ellipsometric studies were carried out on the MoO3 films before and after the UV irradiation to determine the optical indices (n and k). These results are shown in Fig. 5.6B. The subscripts xy and z in this figure refer to refractive indices of light polarized parallel and perpendicular to the substrate, respectively. An increase in absorption is seen in all the MoO3 films studied with irradiation. The coloration is concentrated in the lower visible to near-infrared region. The average change in extinction coefficient (Dkx) occurring at different wavelengths and with different tilt angles is shown in Fig. 5.7A. The dependence of the photochromic response on the substrate tilt and the ensuing nanostructure are clearly evident from these figures. The MoO3 film deposited with a 60-degree tilt is the most photochromically efficient one. This photochromic behavior in GLAD-deposited film is attributed to optimization of the surface area available for water molecule adsorption, which contributes to the increased photochromic activity. The intrinsic 1

MoO3 bandgap (Eg) is determined through the plot of ðahnÞ2 as a function of the photon energy ðhnÞ as shown in Fig. 5.7B (a) for a MoO3 film deposited at a tilt angle of 60 degrees. The a values were determined from the corresponding kx values for a film deposited at this tilt. The Eg value in this case was 3 eV. These plots for MoO3 films deposited at different tilt angles yielded the dependence of Eg on the substrate tilt angle as shown in the inset of Fig. 5.7B (a). With increasing tilt angle between 0 and 80 degrees, the bandgap (Eg) is found to vary between 2.82 and 3.05 eV. The UV irradiation in this case corresponds to 3.26 eV (380 nm).

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

Figure 5.6 (A) Scanning electron microscopy cross-section (left) and atomic force microscopy (right) images of MoO3 films deposited by the glancing angle deposition method at various tilt angles. (B) Optical indices of a MoO3 film deposited at 60-degree substrate tilt before (a) and after (b) UV irradiation. (From G. Beydaghyan, et al., Applied Physics Letters 95 (2009) 051917e1-3.)

The absorption coefficient (a) at this wavelength is shown as a function of various tilt angles of MoO3 deposition in Fig. 5.3B (b). From these results it is concluded that more than the bandgap (Eg), it is the surface area (nanostructure) available for water molecule adsorption that leads to enhanced photochromic effect. This work also shows the ease with which the nanostructure control can be achieved using the GLAD method.

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Figure 5.7 (A) Average change in extinction coefficient (Dkx) at different tilt angles as a function of wavelength (a) and as a function of tilt angle at the wavelength of 800 nm (b). (B) (a) Semiconductor bandgap (Eg) of a MoO3 film deposited at the tilt angle of 60 degrees (inset: variation of Eg with tilt angle), and (b) variation of absorption coefficient (a) with substrate tilt. (Adapted from G. Beydaghyan, et al., Applied Physics Letters 95 (2009) 051917e1-3.)

The enhancement of photochromic effect can also be brought about through various other means. One such method is to deposit a thin metal film on top of the MoO3 or WO3 film or to dope it with metal particles. One of the main problems associated with the photochromic effect that originates from the electronehole pair creation under UV irradiation is their recombination. It has been found that the addition of metal film or particles has multiple effects in enhancing photochromic efficiency [23e26]: 1. The work function of metals such as gold and platinum being 5.1 and 5.64 eV, respectively, and the Fermi level (EF) of TMOs such as MoO3 and WO3 being far below that (4.3e4.9 eV), a Schottky barrier is formed at the interface of the metal and the oxide. Hence,

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the electronehole pairs formed through UV irradiation drift in different directions under the influence of this electric field. 2. The Schottky barrier thus formed also reduces the photocorrosion reactions. 3. The metal particles also enhance the photochromism by adsorbing water molecules at the nanoparticle surface. Yao and his group carried out an in-depth study of the photochromic effects of MoO3 and WO3 thin films prepared under various conditions [5,23e26]. The results of their study on the effect of depositing a 20-nm thin film of gold on a 1-mm-thick MoO3 film on an indium tin oxide (ITO)-coated glass are shown in Fig. 5.8. A significant enhancement of absorption is seen in the MoO3/Au film compared to the pure MoO3 film.

Figure 5.8 (A) Absorption spectra of a MoO3 sample on indium tin oxide (ITO) in the as-deposited state (curve a) and after UV irradiation for 10 min (curve b). (B) Absorption spectra of a MoO3 sample on ITO overlaid with a thin film of Au, in the as-deposited state (curve a) and after UV irradiation for 10 min (curve b). ABS, absorption. (From J.N. Yao, Y.A. Yang, B.H. Loo, Journal of Physical Chemistry B 102 (1998) 1856e1860.)

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A change in absorption (DABS) of 56% at the wavelength of 900 nm is seen in the film coated with Au, while it is on the order of 21% at the same wavelength in the pure film. An extremely important aspect of their work is the demonstration of enhancement of photochromism by combining the effect with electrochromism in these TMOs. The bandgap (Eg) of the wideband semiconductor MoO3 films being on the order of 3 eV, the photochromic effect cannot be induced by lower energy photons. Hence, visible light photochromism cannot be induced in these films. However, it has been shown that it is indeed possible to color the MoO3 films by visible light if the films are pretreated electrolytically. This breakthrough work of Yao et al. [27] of inducing visible light photochromism is a major progress in this field. This is because of the large interest in visible-light-related applications such as smart windows for solar control, display devices, and more. Further, commercial lasers that are economically available are those that operate in the visible and near-infrared region. Hence, from a commercial application point of view it is highly desirable to have photochromic materials and devices that operate in a higher wavelength region than the UV. MoO3 films deposited on ITO-coated substrates were inserted in an electrochemical cell containing 0.2 M LiClO4 in a polypropylene carbonate solution [23]. By cathodically polarizing the MoO3/ITO electrode, 10 mC of lithium was inserted into the film to lightly color the samples. The films were then exposed only to visible light for 10 min. The results are shown in Fig. 5.9A. A significant coloration ensues with the visible light irradiation as seen in this figure. The MoO3 samples coated with 20 nm of Au on ITO samples were also subjected to the same steps of initial light coloration through electrochromic insertion and then exposed to visible light irradiation for 10 min. These results are shown in Fig. 5.9B. It can be seen that a much higher degree of photochromic coloration through absorption is seen in MoO3/Au composite films. The DABS values corresponding to the initial coloration by polarization of the MoO3 and MoO3/Au films on ITO were, respectively, 13% and 12% at the wavelength of 900 nm. DABS values for these films after an exposure of 10 min to visible light were, respectively, 13% and 30% at the same wavelength. These results clearly show not only the effectiveness of initial electrochromic coloration in bringing about the visible photochromism, but also its enhancement by the inclusion of a thin metallic layer as in the case of UV light photochromism. Similar work has been carried out on the photochromic enhancement of MoO3 films overcoated with platinum film. MoO3/ITO films overlaid

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Figure 5.9 (A) Absorption spectra of a MoO3 sample on ITO in the as-deposited state (curve a), after the electrochromic insertion of lithium (curve b), and after UV irradiation for 10 min (curve c). (B) Absorption spectra of a MoO3sample on ITO overlaid with a thin film of Au, in the as-deposited state (curve a), after the electrochromic insertion of lithium (curve b), and after UV irradiation for 10 min (curve c). ABS, absorption. (Adapted from J.N. Yao, Y.A. Yang, B.H. Loo, Journal of Physical Chemistry B 102 (1998) 1856e1860.)

with a thin film of Pt were initially polarized cathodically to insert lithium ions. The work function of Pt being higher (5.64 eV) than that of Au (5.1 eV), an even more significant enhancement in visible light photochromism is observed for Pt/MoO3/ITO film. The DABS value for films colored lightly by lithium insertion, i.e., electrochromically, and then irradiated with visible light for 10 min was around 49% at 900-nm wavelength. This is a significant jump from even the Au-overlaid MoO3 films,

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Figure 5.10 (A) Comparison of the change in absorption (DABS) of the visible light photochromism of a MoO3 sample on indium tin oxide overlaid with a thin film of Au and Pt, after the initial electrochromic coloration (DABS1) and after the visible light irradiation (DABS2). (B) Schematic of the formation of the Schottky barrier and the charge-transfer process at the Au/MoO3 (or WO3) layer during the photochromic process. TMO, transition metal oxide. (Adapted from T. He, J. Yao, Journal of Photochemistry and Photobiology C 4 (2003) 125e143; J.N. Yao, Y.A. Yang, B.H. Loo, Journal of Physical Chemistry B 102 (1998) 1856e1860.)

which showed a DABS of 30%. The degree of visible photochromic enhancement occurring with the inclusion of gold and platinum on MoO3 films and by cathodic polarization is shown in Fig. 5.10A. The enhancement phenomenon is attributed to the formation of a Schottky barrier at the metal (Au, Pt)/TMO (MoO3, WO3) interface [10]. The UVphotogenerated electronehole pairs are subjected to this built-in electric field in the film. This electric field drives the electrons into the bulk of the TMO along the conduction band while the holes drift to the interface via the valence band, thus minimizing the electronehole recombination. This is the reason for the enhancement of the coloration in MoO3 films with metal inclusions. The schematics of these processes are shown in Fig. 5.10B.

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As far as the visible light photochromism is concerned, the initial electrochromic coloration by lithium insertion and formation of bronze (e.g., LixMoO3) facilitates the process. Normally, in the broadband semiconductors such as MoO3 with a bandgap around 3 eV, the photons associated with UV light will not be sufficient to excite the electrons over this gap. However, the electrochromic coloration and formation of LixMoO3 leads to stable trap states that are located within the bandgap, thus reducing the energy required to excite the photogenerated electrons. Hence, the energy associated with the visible light photons is sufficient to excite the electrons into the conduction band. The presence of these trap states is verified by surface photovoltage spectroscopy. The enhancement of visible light photochromism is explained again by the formation of a Schottky barrier and the diminution of electronehole recombination. The work function of Pt being higher than that of Au, this charge separation is even more efficient, leading to much higher photochromic coloration.

5.3.2 Tungsten Trioxide WO3 thin films also exhibit quite efficient photochromic coloration under UV light irradiation. Yao et al. [28] made an in-depth study of the photochromic behavior of WO3 thin films. The photochromic mechanism and the factors influencing it are similar to those discussed earlier for molybdenum oxide (MoO3) thin films. The different theoretical models used to explain the phenomenon [color center formation, IVCT, defects (oxygen vacancies)] are also the same. An example of their work on nanostructured photochromic WO3 films is given in Fig. 5.11. These films were prepared by colloidal chemistry using oxalic acid to control the nanostructure induced in the film as well as to prevent the WO3 particle aggregation. The nanostructure control with various molar concentrations of oxalic acid is as shown in Fig. 5.11A. These authors reported several interesting observations that occurred as a function of the oxalic acid concentration: (1) as seen in this figure, with increasing oxalic acid concentration the particle size decreased rapidly, leading to higher amorphousness from an initial polycrystalline state along with the increase in defect states; (2) the decreasing particle size also led to the increase in the bandgap (Eg) as the electron energy state spacing increased with decreasing volume or diameter of the particles; and (3) the presence of oxalic acid itself played an important role in the photochromic process by acting as a scavenger of the photogenerated holes and electrons.

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Figure 5.11 (A) Transmission electron microscopy micrographs and electron diffraction patterns of nanostructured WO3 films prepared by the colloid chemistry method using various concentrations of oxalic acid: (a) 0.1, (b) 0.3, (c) 0.4, and (d) 0.8 M. (B) Absorption spectra of nanostructured WO3 films prepared with an oxalic acid concentration of 0.3 M. (Curve a) As prepared, (Curve b) UV-irradiated. (Adapted from T. He, J. Yao, Journal of Materials Chemistry 17 (2007) 4547e4557.)

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The photochromic comportment of the nanostructured WO3 films is explained in the light of these factors. Shown in Fig. 5.11B are the absorption spectra of nanostructured WO3 films prepared with an oxalic acid concentration of 0.3 M. Compared to the transparent film of WO3 with no absorption structure in the visible and near-infrared region of the spectrum in the as-prepared state, a strong absorption peak centered around 900 nm develops in the UV-irradiated film. The broadband absorption spans the entire visible and near-infrared region, rendering the film blue in appearance owing to the absorption tail in the visible region. This photochromic coloration depends very strongly on the oxalic acid concentration. As seen in Fig. 5.12, the absorption increases throughout the spectral region studied rapidly and more effectively with increasing molar concentration of oxalic acid. It was also observed that there is a slight blue shift in the absorption spectra with increasing oxalic acid concentration. The enhancement of the photochromic coloration and the blue shift are attributed to the three factors discussed earlier occurring with increasing oxalic acid concentration. The decreasing particle size leads to an increase in Eg and thus more energetic photons are required for the photochromic excitation of electrons resulting in the blue shift. The increasing defect states with increasing amorphousness facilitate the phototransition and, thus, the photochromism. The presence of increasing oxalic acid inhibits the free movement of the photogenerated holes and electrons also hinder their recombination, thus enhancing the photochromic effect. The maximum DABS of these films around the 900-nm wavelength corresponds to 0.2.

Figure 5.12 Absorption spectra of WO3 thin films prepared by colloidal chemistry with various amounts of oxalic acid as shown. (Adapted from T. He, J. Yao, Journal of Materials Chemistry 17 (2007) 4547e4557.)

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This value is much higher than the DABS of the pure MoO3 film around the same wavelength, which is reported to be 0.12 (Fig. 5.8). Another important difference between the nanostructured WO3 films with oxalic acid [28] and the thermally evaporated MoO3 films [23] is their optical nature in the as-prepared state. The WO3 films are highly transparent in the visible and near-infrared region, whereas the thermally evaporated MoO3 films are bluish with an average absorptance of nearly 0.2 at the wavelength of 900 nm. This initial absorptance in pure molybdenum oxide films might be a result of nonstoichiometry of the films in the as-prepared state. It can also be seen that the molybdenum oxide films in the as-prepared state exhibit a much broader band absorptance than as-prepared films of WO3 with oxalic acid. From these initial states, the two types of films show different DABS values. With the reported photochromic effect in the two cases, the value of DABS is, respectively, 0.12 and 0.2 for thermally evaporated MoO3 films and WO3 films with oxalic acid. Hence, the photochromic performance of WO3 films seems to be superior to that of MoO3 films, although a direct comparison of the two types of films reported here is not justified, as the natures of the two films are entirely different. The methodology of preparation and the nanostructure induced play extremely important roles in the photochromic process. For a tangible comparison the two films have to be prepared under exactly the same conditions. This is even more important as several factors come into play in bringing about the photochromic effect, especially the nanostructure. Similar to MoO3, WO3 is also an n-type wideband semiconductor and a TMO with a bandgap around 3.2 eV. Hence, high-energy photons in the UV region are required to generate electronehole pairs that induce the photochromic effect in these films. Following the breakthrough work of inducing visible light photochromism, i.e., with lower energy photons in MoO3 films, the same was shown to take place in WO3 films by Yang et al. [29]. The method of electrolytically polarizing the WO3 films in a nonaqueous electrolyte prior to irradiating them with visible light was applied in this work. One-micrometer-thick WO3 films were prepared by thermal evaporation on SnO2-coated glass substrates and then immersed in an electrochemical cell containing potassium nitrate (KNO3) in polycarbonate solution. The electrodes were polarized cathodically to insert potassium ions and electrons to initially color the films electrochromically. Such films were then exposed to visible radiation for 40 min. The photochromic behavior of such cathodically polarized WO3 films is shown in Fig. 5.13A. The optical density spectra of the films are given in the

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Figure 5.13 Optical density spectra of (A) a thermally evaporated WO3 film and (B) a MoO3 film in the as-deposited (curve a), cathodically polarized (curve b), and subsequently visible-light-irradiated (curve c) states. (From Y.A. Yang, et al., Journal of Physical Chemistry of Solids 59 (9) (1998) 1667e1670.)

as-deposited, cathodically polarized, and visible-light-irradiated states. It is seen that the as-deposited films exhibit a fairly high coloration corresponding to an optical density of 0.2 at lower wavelengths studied and gradually increasing to more than 0.35 in the near-infrared region. With polarization the optical density increases quite effectively with increasing change toward the higher wavelengths. With subsequent exposure to visible light it is seen that the optical density increases further, thus clearly exhibiting visible light photochromism in WO3 films. By comparison to similarly treated thermally evaporated MoO3 films shown in Fig. 5.13B, it is seen that visible light photochromism in WO3 films is much more effective

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than in MoO3 films of the same type. The ultimate change in optical density (DOD) from the polarized state to the visible-light-irradiated state is 0.15 and 0.3 for MoO3 and WO3 films, respectively. However, it is to be noted that the visible light exposure times in the two films compared are quite different, being 10 min for MoO3 films vs. 40 min in the case of WO3 films. However, both studies are important steps in how to induce visible light photochromism. The observation of visible light photochromism only after electrolytically polarizing samples in both the WO3 and the MoO3 films is attributed to the formation of defect states inside the bandgap region upon the electrochromic double insertion of ions and electrons and the formation of the respective bronzes (KxWO3 and LixMoO3). These defect states formed seem to be more stable in the case of MoO3 films [23] than in the WO3 films. The photochromic coloration is reported to be stable over a long period and the coloration/ bleaching more repetitive in the case of the MoO3 films. In the case of the WO3 films the samples themselves and the coloration/bleaching cycles seem to be very short-lived. Hence, most aspects related to the photochromic effect seem to point to MoO3 as a better candidate for this work, both for UV light photochromism and for visible light photochromism. Under similar conditions of preparation and treatment the MoO3 films seem to exhibit superior optical change, stability of operation, and cycling.

5.3.3 Electrochromic and Photochromic Colorations In addition to the capacity to induce visible light photochromism, the process of slightly coloring the photochromic films by cathodic polarization (electrochromic coloration) is also advantageous from the application point of view. Generally, the purely photochromic process in MoO3 and WO3 films is not interactively reversible. Once such films are photochromically colored, the coloration is metastable and in some cases can last for several months to several years. Given in Fig. 5.14A is an example of a Bragg mirror (multiple layers of high- and low-refractive index) made out of radio-frequency-sputtered MoO3 and SiO2 films [30]. The film shown in Fig. 5.14A is a 20-pair alternating stack of these films with quarter-wave stacks with a photonic stop band at 650 nm. Such a multilayer device was exposed to UV light of 380 nm with a 50-nm bandwidth for 1 h at an irradiance of 20 mW/cm2. An aluminum mask with an Université de Moncton logo was held in front of the sample during irradiation. The UV photochromic coloration has been stable for over 5 years, demonstrating the application potential of such devices for memory storage. The Bragg mirror has been constantly exposed to air with no degradation of the

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Figure 5.14 (A) Photo of a MoO3/SiO2-based Bragg mirror with UV-irradiated darker coloration in the middle showing the logo of the Université de Moncton. (B) Scheme of the proposed display device combining electrochromic and photochromic effects. (Adapted from T. He, J. Yao, Journal of Photochemistry and Photobiology C 4 (2003) 125e143.)

photochromic coloration. One of the reasons for this high degree of optical memory seems to be the outermost layer of SiO2, which acts as a protective layer, too, against any interaction of the inner colored MoO3 films. Despite this interesting and desirable optical memory, the quasi-irreversible coloration is a shortcoming for many applications in which an interactive reversibility of coloration is necessary. From an application point of view the devices in which a combination of electrochromic and photochromic coloration can be induced are more interesting. The electrochemical cell used to induce the initial coloration through cathodic polarization (electrochromic coloration) for bringing about the visible light photochromism can also be used to “erase” (reverse) the photochromic coloration. An anodic polarization needs

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to be applied to the TMO (MoO3, WO3) film in the electrochromic cell to discolor the film. Such a device would provide interactive control of the optical change occurring through electrophotochromism. An example of this arrangement is shown in Fig. 5.14B for use as an electronic display device [10]. The initially transparent sample as shown in stage A of this figure is initially colored electrochromically in an electrochemical cell through cathodic polarization to bring it to stage B. After this pretreatment, the sample is exposed to visible light irradiation with a desired pattern mask in front of the sample to further photochromically color the film, shown as stage C in the figure. From this stage C, the sample can be brought back to the initial stage A (complete discoloration) or to stage B (partial discoloration to the initial electrochromic level) by controlling either the duration or the level of anodic voltage applied to the sample for discoloration. Such a coloration and bleaching can be achieved over several hundred cycles.

5.3.4 Photochromic Effects in Molybdenum Trioxide and Tungsten Trioxide Films Modified by Metal Nanoclusters An enhancement of the UV-related photochromism as seen in MoO3 films can also be brought about in WO3 thin films by the inclusion of metallic components. The creation of the Schottky barrier at the WO3 metal interface separates the photoexcited electrons and holes, which contribute more effectively to the coloration, thus leading to an enhancement of the photochromic effect. If these gold nanoparticles are at the surface of the WO3 film they will also be effective in attracting water molecules to the surface. These water molecules further contribute to the photochromic effect by consuming the photogenerated holes and thus diminishing the electronehole recombination [31]. In Fig. 5.15A is shown the effect of such photochromic enhancement in WO3 thin films overlaid with gold nanoparticles [32]. The WO3 thin films were deposited by thermal evaporation, whereas the gold nanoparticles were deposited by spin coating. The coloration shown in Fig. 5.15A was effected by UV irradiation of 3 min using a 500-W mercury lamp. Absorbance spectra of the WO3 films with (WO3/Au) and without (WO3) gold are compared in this figure before and after the UV irradiation. As can be seen in this figure, both films are fairly transparent in the as-prepared state with a slight absorption probably coming from the nonstoichiometry and light coloration due to the presence of water. Upon UV irradiation a fairly high absorbance is

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Figure 5.15 Absorption spectra of (A) WO3 and (B) WO3/Au thin films before and after UV light irradiation for 3 min in air. (From T. He, et al., Physical Chemistry and Chemical Physics 4 (2002) 1637e1639.)

induced in the film throughout the spectral region studied. The coloration is significantly higher in the case of the WO3 film overlaid with gold nanoparticles. This enhancement in coloration is attributed to the various factors discussed earlier, similar to the case of metal-particle-overlaid MoO3 films. The DABS due to the photochromic effect at the wavelength of 900 nm is around 21% for the pure WO3 film, while it is around 37.5% for the film overlaid with Au nanoparticles. Hence, an enhancement of around 1.8 is induced by the presence of the metal nanoparticles.

5.3.5 Photochromic Effects in Heterostructure-Based Films Other types of photochromic enhancements have also been studied in both WO3 and MoO3 films by putting them in contact with other oxide semiconductors. Amorphous MoO3 films were deposited on glass substrates

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Figure 5.16 Optical density spectra of MoO3 films on glass (A) and of CdS/MoO3 on glass (B) as functions of irradiation time. (From M.A. Quevedo-Lopez, et al., Journal of Materials Science 11 (2000) 151e155.)

precoated with polycrystalline cadmium telluride (CdS) films. The irradiation of such double layers enhances the formation of color centers, thus contributing to enhanced photochromic coloration. Shown in Fig. 5.16 is an example of such enhancement in the MoO3/CdS double layer. A comparison of the optical density change of pure MoO3 and the double layer irradiated with a 100-W tungsten lamp is shown here [33]. In the pure MoO3 films (Fig. 5.16A), there is an optimum energy and/or duration during which color center formation is effective. At the constant power level used in this study, the maximum photochromic coloration seems to take place at 45 min of irradiation. The peak of this coloration is around 850-nm wavelength with an optical density of 1. However, in the films with the MoO3/CdS double layer there is a significant improvement in the

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photochromic coloration. In Fig. 5.16B are shown the optical density spectra of MoO3/CdS films before and after UV irradiation of varying duration. The photochromic coloration in this case increases nearly systematically with increasing duration of irradiation for wavelengths above 500 nm. At wavelengths below 500 nm, there is a static absorption band emanating from the CdS film. There is hardly any change in optical density with UV exposure of 45 min. However, for other prolonged UV exposures, the optical density increases, giving a maximum optical density of around 2.75 at the wavelength of 850 nm for 180 min of exposure. Hence, such double-layered films with other semiconducting oxides induce a significant enhancement of photochromic performance. The photochromic coloration and its enhancement in the double layer are attributed to the formation of color centers in the film. The color center concentration per unit volume, N, in the two types of films, MoO3 on glass and MoO3 on glass coated with CdS, has been calculated using the Smakula equation [34]:  2   Z Eeff mc N fij ¼ n aðuÞdu 2p2 e2 h E  2 E where Eeff is the local effect of polarization, fij is the oscillator strength, and n is the refractive index along with the fundamental constants, m, c, e, and h. It was found that the color center concentration varied with the duration of UV exposure and the type of MoO3 film dealt with. In MoO3 film deposited on glass, a maximum concentration (10  1017/cm3) was found for a UV exposure of 45 min in accordance with the result seen in Fig. 5.16A. Similarly the aforementioned model indicated the increasing formation of color centers in the case of MoO3 deposited on glass coated with CdS with increasing duration of UV exposure. The maximum number of color centers in this case corresponded to the CdS/MoO3 double layer exposed to UV irradiation for 180 min, giving a value of 25  1017/cm3 of color centers. This enhancement in the color center formation and the eventual enhancement of photochromic activity are attributed to the presence of the semiconductor CdS film. Under the action of light, the electrons from the semiconductor’s valence band are excited to the conduction band and eventually into the MoO3 film. The color centers are formed through the trapping of these electrons in the oxygen vacancies in the film. The photochromic effect of a nanostructured MoO3 double layer with CdS was studied by Srinivas Rao et al. [35]. The MoO3 films were

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Figure 5.17 (A) Atomic force micrographs of MoO3 films deposited on glass (left) and on glass coated with CdS film (right). (B) Plot of ln a versus hn (b) and of (ahn)2 (a) as functions of photon energy. (From K. Srinivas Rao, et al., Materials Science and Engineering B 100 (2003) 79e86.)

deposited by thermal evaporation under high-pressure oxygen plasma facilitating the nanostructuring of the films. Films were deposited on glass substrates as well as glass substrates coated with polycrystalline CdS films. The films deposited on bare glass substrates were amorphous and had a uniform nanostructure with an average grain size of around 50 nm. The films deposited on CdS-coated glass, on the other hand, had elongated grains with an average grain size of around 100 nm. In Fig. 5.17A are shown AFM images of the two films for comparison. The absorption

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coefficient, a; was obtained from the measured transmittance (T), reflectance (R), and thickness (t) of the film using the following relation [36]: ! T ln ð1
RÞ2 a¼ t Applying Urbach’s rule about the dependence of the absorption coefficient, a, a; on photon energy through the relation,   hn a ¼ a0 exp Ee where Ee is the energy of the localized states, the value of Ee was obtained from a plot of ln a as a function of the photon energy. As shown in Fig. 5.17B (b), the energy of bound states, Ee, was 0.22 eV. The bandgap (Eg) for the MoO3 film was obtained from plot of ahn versus ðhn
Eg Þh as discussed before. This plot is shown in Fig. 5.17B (a). The intercept on the energy axis obtained from the best fit to the linear part of this dependence for a h value of 0.5 (allowed direct transitions) was 3.18 eV. The photochromic effect on the two types of MoO3 films, i.e., on glass and on glass coated with CdS, was induced by exposing the films to 120 mW/cm2 radiation using a tungsten lamp over varying duration from 15 to 200 min. The optical density spectra are as shown in Fig. 5.18. In the nanostructured MoO3 film on bare glass substrate shown in Fig. 5.18A, there exists an optimum irradiation duration for which the film exhibits the maximum change in optical density. This film upon an exposure of 40 min shows the maximum photochromic change, which peaks around the wavelength of 850 nm. However, as also seen in the earlier work, with prolonged irradiation the photochromic coloration diminishes. The photochromic coloration is attributed to the formation of color centers through the trapping of the photogenerated electrons by the oxygen vacancies. The peak optical density value for these nanostructured films is around 1.25 at the wavelength of 850 nm. The MoO3 films deposited on the CdS-coated glass, on the other hand, exhibit a systematic increase in optical density with irradiation duration, as shown in Fig. 5.18B. The peak value of optical density in all cases is around 820 nm, although a slight shift toward higher wavelength with irradiation duration is also evident. The peak optical density in this case is for an exposure of 150 min, which corresponds to a value of 2.5. Hence, the optical density value is doubled in

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Figure 5.18 Optical density spectra with varying irradiation duration of (A) a MoO3 film on glass and (B) a MoO3 film on CdS-covered glass. (From K. Srinivas Rao, et al., Materials Science and Engineering B 100 (2003) 79e86.)

the case of MoO3 film in contact with the semiconducting CdS layer compared to the film on bare glass substrate. This enhancement in photochromic effect in the MoO3/CdS layer is attributed to the excitation of electrons to the conduction band and their transfer into the MoO3 film creating color centers through their capture by oxygen ion vacancies. A schematic representation of this transfer is shown in Fig. 5.19A. As shown in this schematic the bandgaps of CdS and MoO3 are, respectively, 2.42 and 3.18 eV. The conduction band of MoO3 is lower than that of CdS, making the transfer of photogenerated electrons to MoO3

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Figure 5.19 (A) Schematic representation of the electron transfer between CdS and MoO3 semiconducting layers. (B) Optical density spectra of MoO3film on In:CdS/glass substrate with varying duration of irradiation. (C) Color center concentration (N) of MoO3 film on different substrates as a function of irradiation duration. (From K. Srinivas Rao, et al., Materials Science and Engineering B 100 (2003) 79e86.)

possible. Indium (In)-doped CdS films on glass were also studied in this work to examine the photochromic behavior. The optical density spectra of such films with various durations of irradiation are shown in Fig. 5.19B. Similar to the MoO3 layers, the MoO3/In:CdS layers exhibit an increasing photochromic coloration with increasing irradiation duration. The photochromic effect in these films is even more efficient, though slightly, than in the case of MoO3 on CdS film. The maximum optical density corresponding to 150 min of irradiation in this case is 2.7 at the wavelength of around 850 nm. Compared to the optical density of 2.5 exhibited by the MoO3/CdS double layers, there is a slight enhancement in the MoO3/ In:CdS layers. This enhancement in MoO3 photochromism in In-doped

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CdS film is attributed to the increasing bandgap of the CdS through In doping. It is also assumed that In provides more photogenerated electrons to the MoO3 conduction band. This effect further facilitates the transfer of photoexcited electrons, enhancing the photochromic effect. The color center concentration (N) has been calculated for the three types of MoO3 films, i.e., bare glass, glass coated with CdS, and glass coated with In-doped CdS, using Smakula’s relation discussed earlier. In Fig. 5.19C is given the color center concentration with irradiation duration. This dependence is in accordance with the photochromic effect observed in the three cases. The MoO3 film on bare glass substrates exhibits a maximum concentration of 16  1017/cm3 for an optimum light exposure of 40 min. The color center concentration increases nearly linearly in MoO3 film deposited on CdS and In-doped CdS, though much more efficiently in the latter case. At the longest irradiation time studied, 200 min, an estimated color center concentration of 33  1017 and 47  1017/cm3 is found for MoO3 film on CdS and on In-doped CdS, respectively. In the scheme of photochromic enhancement by combining with other oxides, one of the popular combinations is the mixed MoO3/WO3 films, which yield a better photochromic effect than either of the individual layers. For example, shown in Fig. 5.20 is a comparison of the photochromic coloration in the form of DOD at 800 nm of a MoO3 and a WO3 film as well as their composite film [28]. The individual films were prepared by thermal evaporation and the composite films were prepared by coevaporation. The figure shows the DOD of these WO3 (92%)/MoO3 (8%) films under UV irradiation and as tested under various reducing agents. As can be seen in this figure, the performance of the composite film in terms of DOD and under similar conditions of testing is 2.5 and 1.5 times that of the individual layers of MoO3 and WO3, respectively. The photochromic performance of the composite films also depends on the type of reducing agent under which the tests are conducted, as shown in the inset of Fig. 5.20. In addition, as seen in the previous discussion, the peak photochromic absorption in the case of individual layers of MoO3 is situated in the near-infrared region, while that of the composite film is located at a lower wavelength bringing it into the middle of the visible region. Owing to the highest sensitivity of human eyes around these wavelengths (550 nm), the composite films provide a distinct advantage from the application point of view in display devices and more. Ahmed et al. [36] reported on the visible-light-induced photochromic coloration of MoO3/WO3 nanocomposite films. In Fig. 5.21A are shown

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Figure 5.20 Comparative study of the photochromic response (DOD) at 800 nm of a WO3 (92%)/MoO3 composite film with that of individual WO3 and MoO3 films after 5 min of UV irradiation in N2 in the presence of different vol% of C2H5OH vapor: (column 1) 0.25%, (column 2) 0.5%, and (column 3) 1%. Inset: Photochromic response of the composite films in 0.5% vol% of (column a) ethanol, (column b) methanol, (column c) n-propanol, and (column d) isopropanol. (Adapted from T. He, J. Yao, Journal of Materials Chemistry 17 (2007) 4547e4557.)

scanning electron microscopy images of liquid-phase-deposited and hightemperature-annealed MoO3/WO3 composite films with 10% and 20% WO3. The nanostructure is made up of needle-like structures in the 100e500 nm range. An interesting feature reported is the development of a crystalline nature of the MoO3, while the WO3 component is amorphous. In Fig. 5.21B are shown the optical spectra of the two films as well as the photochromic absorption spectra with increasing visible light irradiation. In the as-deposited film the overall transmittance of the composite films decreases with increasing WO3 composition. It is seen that the absorption in the mixed oxide film increases with increasing visible light irradiation.

5.3.6 Photochromic Effects in Doped Molybdenum Trioxide and Tungsten Trioxide Films The photochromic performance of the MoO3 and WO3 thin films and enhancement effects have also been studied by doping these materials with other dopants such as titanium (Ti), niobium (Nb), tantalum (Ta), and zirconium [38]. A solegel method of preparation of WO3 thin films and their doping has been carried out and used to study the effects of these materials on the photochromic performance of the films. The photochromic effect was also found to depend strongly on the thermal treatment.

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Figure 5.21 (A) Scanning electron microscopy images of liquid-phase-deposited and annealed MoO3/WO3 nanocomposite films with 10% (left) and 20% (right) WO3. (B) Transmittance spectra of MoO3/WO3 nanocomposite films with 10% (curve a) and 20% (curve b) WO3. (C) Absorption (Abs) spectra of MoO3/WO3 nanocomposite films with different irradiation times: (curve a) before irradiation and after (curve b) 10, (curve c) 20, (curve d) 30, (curve e) 40, (curve f) 50, and (curve g) 60 min of irradiation. (Adapted from H.M. Farveez Ahmed, N.S. Begum, Bulletin of Materials Science 36 (1) (2013) 45e49.)

An example of the photochromic change undergone by a Zr-doped WO3 film is shown in Fig. 5.22A. A highly transparent as-deposited film thermally treated at 100 C shows a very high transmittance. The film transmittance drops quite drastically after visible light irradiation, with peak absorption around 900 nm. Compared to all other dopants, the films with Zr seem to show the maximum coloration. The coloration responses of the different dopants are shown in Fig. 5.22B in the form of DOD. As seen in this figure, quite a drastic enhancement of the photochromic performance can be expected from these dopants. Molybdenum (Mo)-doped tungsten oxide (WO3) composites with urethane resin have been studied for their photochromic properties [39]. The composites were formed using peroxoisopolytungstic acid (W-IPA), peroxoisopolymolybdenum acid (Mo-IPA), and liquid urethane resin (M-40) as starting materials. Undoped samples were prepared by mixing

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Figure 5.22 (A) Transmittance spectra of solegel-prepared WO3:Zr thin film with two layers and heat treated at 100 C for 1 h in the bleached (solid line) and colored (dotted line) state. (B) Comparison of the optical density change (DOD) of WO3 films with different dopants. (From C.O. Avellaneda, L.O.S. Bulhoes, Solid state Ionics 165 (2003) 117e121.)

W-IPA in liquid resin. Doped samples were obtained by initially mixing W-IPA and Mo-IPA; eventually this solution was mixed with the urethane resin. The resin was cured in each case through UVevisible light radiation. The initial coloration induced in the composite due to the exposure to UVevisible light was removed by keeping the samples in the dark for 7 days. Shown in Fig. 5.23 is the transmittance of an undoped WO3-based composite as a function of the duration of UVevisible light irradiation. With increasing duration of irradiation, it is seen that the transmittance in the midregion of the spectrum diminishes with peak coloration around 600 nm. A much weaker peak is also reported around 900 nm, as shown in the inset of Fig. 5.23A. As also can be seen in the inset of this figure the initially transparent film turned increasingly blue with photochromic coloration. The coloration rate constant (k) was calculated from the absorption values using the relation:   A
ln ¼ kt A0 where A and A0 are the absorption values after and before photochromic coloration and t is the time. The reaction rate constant was

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Figure 5.23 (A) Transmittance spectra of WO3/resin composites with increasing duration of UVevisible light irradiation (inset shows the samples before and after 120 min of coloration). (B) Transmittance spectra of Mo-doped WO3/resin composites with different Mo/W ratios (inset shows the samples after 120 min of coloration). (Adapted from H. Miyazaki, M. Inada, H. Suzuki, T. Ota, Journal of the Ceramic Society of Japan 121 (1) (2013) 106e108.)

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Figure 5.24 Absorbance spectra of (A) WO3 and (B) ZnO/WO3 colloids before (curve a) and after (curve b) exposure to UV irradiation for 1 min (inset in B shows the optical density change (DOD) as a function of the ZnO concentration. (Adapted from T. He, J. Yao, Journal of Materials Chemistry 17 (2007) 4547e4557.)

calculated to be 2.68  10
2/min for the undoped WO3 composite. In Fig. 5.23B are shown the photochromic coloration spectra of the Mo-doped WO3/resin composites as a function of Mo-to-W ratio. The film with the smallest Mo/W ratio of 0.1 shows the maximum photochromic change. The blue coloration of this film with an irradiation of 120 min is similar to that of the undoped composite. With increasing Mo/W ratio in the composite the overall coloration decreases and exhibits a blue shift as seen in Fig. 5.23B. The doped composite samples appear light brown after the photochromic coloration. Using the peak position in each case, the coloration rate constant was calculated for the Mo-doped

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composites. With values of k being 3  10
2, 4.23  10
2, and 3.53  10
2/min for composites with Mo/W ratios of 0.1, 0.5, and 1, respectively, the presence of Mo in the composite seems to increase the rate of coloration. Similarly, the photochromic behavior of WO3/ZnO-based colloids has also been studied, which showed an extremely enhanced photochromic effect compared to the thin films [40]. Depending on the concentration of ZnO inserted into these colloids, a 200-fold improvement in DOD was reported in these colloids compared to the WO3 colloids. The absorbance spectra of the two types of colloids are shown in Fig. 5.24. In addition to the severe photochromic enhancement of these ZnO-inserted colloids, the UV coloration can be brought about within a minute, coloring the colloids deep blue. The photochromism in such colloids is attributed to the electrons trapped in the energy levels within the forbidden gap of WO3. No formation of hydrogen bronze is associated with this coloration.

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[20] Y.A. Yang, Studies of Photo-electronic Functional Properties of Transition e MetalOxide Semiconductor Thin Films (Ph.D. thesis), Institute of Photographic Chemistry, Chinese Academy of Sciences, 1999. [21] J. Vondrak, J. Baluska, Solid State Ionics 68 (1994) 317e323. [22] G. Beydaghyan, et al., Applied Physics Letters 95 (2009), 051917e1-3. [23] J.N. Yao, Y.A. Yang, B.H. Loo, Journal of Physical Chemistry B 102 (1998) 1856e1860. [24] J.N. Yao, B.H. Loo, Solid State Communications 105 (1998) 479. [25] J.N. Yao, Y.A. Yang, Progress in Natural Science 9 (1999) 153. [26] T. He, et al., Langmuir 17 (2001) 8024. [27] J.N. Yao, K. Hashimoto, A. Fujishima, Nature 355 (1992) 624. [28] T. He, J. Yao, Journal of Materials Chemistry 17 (2007) 4547e4557. [29] Y.A. Yang, et al., Journal of Physical Chemistry of Solids 59 (9) (1998) 1667e1670. [30] T. Ben-Messaoud, et al., Applied Physics Letters. 94 (2009) 111904. [31] C. Bechinger, G. Oefinger, S. Herminghaus, Leiderer, Journal of Applyied Physics 74 (1993) 4527e4533. [32] T. He, et al., Physical Chemistry & Chemical Physics 4 (2002) 1637e1639. [33] M.A. Quevedo-Lopez, et al., Journal of Materials Science 11 (2000) 151e155. [34] D.C. Dexter, in “Solid State Physics: Advances in Research and Applications”, vol. 6, Academic Press, p-3531. [35] K. Srinivas Rao, et al., Materials Science and Engineering B 100 (2003) 79e86. [36] F. Urabch, Physical Review 92 (1966) 627. [37] H.M. Farveez Ahmed, N.S. Begum, Bulletin of Materials Science 36 (1) (2013) 45e49. [38] C.O. Avellaneda, L.O.S. Bulhoes, Solid State Ionics 165 (2003) 117e121. [39] H. Miyazaki, M. Inada, H. Suzuki, T. Ota, Journal of the Ceramic Society of Japan 121 (1) (2013) 106e108. [40] R.G. Xie, et al., Chemical Journal of Chinese Universities 24 (2003) 2086.

CHAPTER 6

Chromogenic Thin Film Photonic Crystals 6.1 BASICS OF PHOTONIC CRYSTALS Photonic crystals are one-, two-, or three-dimensional, periodically ordered structures on a light wavelength scale, i.e., on a nanometric scale [1,2]. These materials are made of regions of alternately high- and low-refractive-index units. An incoming electromagnetic radiation impinging on such materials would see this periodic variation in the dielectric constant of the medium during its propagation in different directions. This results in their constructive or destructive interference depending on the dielectric variation encountered in a certain direction and the wavelength of the incident light. The simplest example of this is the thin film Bragg mirror, which can be considered as a photonic crystal of one dimension as shown in Fig. 6.1A. As the incident light goes through this medium of an alternating high (nH)- and low (nL)-refractive-index thin film stack, the wavelengths correspond to the periodicity and the thickness of the high- and low-refractive-index media via the relation [3]: ðdH nH cos qH þ dL nL cos qL Þ where dH and dL are the thicknesses of the high- and low-index media and 12 12   1n2a sin2 qa 1n2a sin2 qa cos qH ¼ and cos qL ¼ , with the parameters n2 n2 H

L

with the subscript a indicating the refractive index of the incident medium and the angle of incidence. The wavelengths corresponding to this relation suffer destructive interference in the forward direction (transmittance) and are completely or partially forbidden from propagating through the medium. Thus a photonic bandgap (PBG) is created for these particular wavelengths. Shown in Fig. 6.1B are the transmittance spectra of a Bragg mirror fabricated in our laboratory with 42 layers of high-and low-refractive-index materials. This Bragg mirror was designed to have a PBG centered around 530 nm. Transition Metal Oxide Thin FilmeBased Chromogenics and Devices ISBN 978-0-08-101747-0 http://dx.doi.org/10.1016/B978-0-08-101747-0.00006-4

© 2017 Elsevier Ltd. All rights reserved.

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Figure 6.1 (A) Schematic diagram of a one-dimensional photonic crystal (Bragg mirror). (B) Transmittance spectrum of a 42-layer high- and low-refractive-index material showing the photonic bandgap at 630 nm. (A) http://www.batop.de/information/r_ Bragg.html.)

If the dielectric constant variation and the wavelength match perfectly for destructive interference, these wavelengths cannot propagate through the material, thus creating a complete PBG. The most well-known examples of naturally occurring photonic crystals are the opals, which exhibit an array of beautiful colors. The colors arise from the interference and diffraction of light passing though these periodically ordered threedimensional (3-D) structures made up of SiO2 beads 150e300 nm in size. A schematic of such an opal is shown in Fig. 6.2A. The incident light goes through a medium of high refractive index (SiO2) interspersed with a low-index (air) medium. Only a small bandgap is created because of the low-refractive-index contrast. The degree of effectiveness of the PBG created depends on the match between the incident wavelength and the scale of the periodicity as well as the refractive index contrast. Hence,

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Figure 6.2 (A) Idealized diagram of the periodic nanostructure of an opal. (B) Schematic of one-, two- and three-dimensional photonic crystals. ((A) From https://en.wikipedia. org/wiki/Opal#/media/File:Opal_molecular_structure2.jpg; (B) From J.D. Joannopoulos, S.G. Johnson, J.N. Winn, R.D. Meade, Photonic Crystals: Molding the Flow of Light, second ed., Princeton University Press, 2008.)

depending on the wavelength region of interest, 1-D, 2-D, and 3-D photonic crystals can be artificially created by adjusting these parameters and tailoring the PBG. Shown in Fig. 6.2B are schematic representations of these artificial photonic crystals. Photonic crystals offer a unique way to control or manipulate light propagation. Photonic crystals have thus become increasingly important since the late 1990s both from a fundamental research point of view and from an application point of view. This is because innumerable industrial applications in areas such as communications, signal processing, and more are tending more and more toward the photonic mode of operation, which offers distinct advantages compared to the hitherto electronic or optoelectronic modes. These advantages stem from the fact that photonic processing and devices based on this mode are expected to be faster and

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more efficient in terms of signal-to-noise ratio, low loss, and the spectral bandwidth that can be handled. However, photons, unlike particles such as electrons, do not carry any charge. Thus the manipulation and control of their flow in a medium via external forces becomes challenging. The flow of electrons in a material, for example, can be manipulated and/or controlled by the application of an external electric field because of their inherent charge. Thus the creation of a PBG in photonic crystals offers this unique possibility of manipulating light flow similar to controlling the flow of electrons in atomic crystals. The familiar word crystal is applied to the periodic arrangement of atoms and molecules in solid materials. The electrons in these solids have to navigate around the periodic potential during their transport because of the coulomb force existing at the nuclear sites. From quantum mechanical considerations it is established that electron transport in the crystal takes place owing to their wave nature. This interaction of the electrons with the coulomb force creates regions of allowed and forbidden energy states. In metals, for example, owing to the overlapping of the wave functions associated with the nuclear sites, the free electrons behave as a gas, leading to large conductivity as there are no forbidden energy states. However, when the crystal is not perfectly periodic, with either missing atoms or the presence of atoms of a kind different from the native one of the crystal or owing to the presence of imperfections in the crystal (defects, impurities, etc.), an electronic bandgap arises in the crystal. Electrons with energies corresponding to the bandgap cannot propagate through the crystal. Thus in semiconductors, forbidden regions of energy are created. Electrons with energies corresponding to the energies in the forbidden region cannot propagate through the crystals, thus creating an electronic bandgap. Analogously, for photons traveling through photonic crystals, the periodic variation (contrast) of dielectric constants (refractive index) is similar to the electric potential experienced by the electrons in atomic crystals. If this contrast in refractive index is too large, then the incoming light gets trapped within that region and cannot propagate through the crystal. Hence, a PBG is created for the light wavelength corresponding to that bandgap similar to the electronic bandgap in atomic crystals. Only the light wavelengths corresponding to the PBG are blocked from propagation through the photonic crystal, while other wavelengths are not affected. Although 1-D photonic crystals such as Bragg mirrors made up of multilayered thin or thick film structures as shown in Fig. 6.1 have been known for a long time, the actual term “photonic crystals” and their potential for

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light manipulation were first introduced in 1987 by Eli Yablonovitch and Sajeev John [6e8]. A 3-D photonic crystal was built and tested in 1991, in the microwave range, by drilling millimeter-sized holes in a material [9]. Being made of material with refractive index of 3.6 and air-filled holes, the photonic crystal provided a very high contrast in refractive index and blocked microwaves from all directions, providing a complete PBG. Since this pioneering work, efforts have been under way all over the world to reproduce the same at near-infrared and visible wavelengths, at which an enormous number of applications for photonic crystals can be envisioned, from optical communications to sensors to the energy sector. The utilization of photonic crystals for various applications has opened up a new and important domain in the optics and photonics industry. The various possibilities opened up by the domain of photonic crystals have also prompted a vast amount of theoretical work in this field. Analogous to the use of Schrödinger’s equations and Bloch functions to account for the electron behavior in atomic crystals, Maxwell’s equations are used to deal with the propagation of light in photonic crystals. The band structure calculations relating to photonic crystals are usually made through numerical methods. These calculations are quite tedious and challenging as the representation of the entire structure with all the unit cells and their shapes and positions needs to be taken into account. The three main approaches used for these calculations are [2]: 1. Band structure calculation assuming an infinite periodic crystal Here the entre crystal is represented through its infinitely repetitive single unit cell. The interaction of light with this single unit cell is calculated to arrive at the frequencies or wavelengths that can pass through or be stopped by the crystal. Although a convenient approach, it cannot be applied to practical photonic crystals because of their invariable imperfectness of periodicity and finite nature. The various algorithms and methods used for the photonic band structure calculations are the plane-wave expansion method [10], finite differences in time domain (FTDT) [11], transfer matrix method [12], KorringaeKohneRostoker method for photonic crystals [13,14], and more. 2. General electromagnetic approach In this approach, transmittance and reflectance are calculated for any arbitrary structure, not essentially taking into consideration the periodic nature of the structure. Although this method gives a lot of flexibility, the fact that the photonic crystal periodicity is ignored makes the calculations more tedious. The method is very appropriate for certain

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photonic crystal devices such as waveguides and resonators. Here too, the FTDT method is applied for the treatment of Maxwell’s equation on a finite grid [15]. 3. Transfer matrix methods In these methods, the band structure of one- and higher-dimensional photonic crystals is calculated by using the transfer matrix equations with the Bloch theorem to solve the crystal structure-associated eigenvalue problem [16e19]. Realizing the importance of photonic crystals for the control and manipulation of light propagation and its consequence for photonic applications in many areas, an enormous amount of theoretical and experimental work has been taking place since the late 1900s. On the experimental side, the biggest challenge in this domain is associated with fabrication of these devices on a large scale. A range of fabrication methods are proposed based on the wavelength region of interest and application. One can divide the photonic crystals very broadly as photonic crystal heterostructures and their devices. The photonic crystal heterostructures are individual components in one, two, and three dimensions, while the devices are multicomponent devices incorporating the photonic crystals. Hence, these devices function as integrated photonic circuits. The photonic heterostructures are fabricated, generally, using one of the three main methods, although other variations of these methods are also reported. These methods are: (1) self-assembly, (2) micromachining, and (3) holographic exposure. The self-assembly method is simple and easily amenable to the large volume growth of 2- and 3-D photonic crystals. A wide range of colloidal crystals are grown by the self-assembly method. This method is most commonly applied in the case of silica or polystyrene sphere-based photonic crystal growth [20,21]. This is achieved by introducing these nanometric spheres into a suitable solvent and letting them organize in a periodic manner under gravitational sedimentation. These so-called colloidal crystals are formed by the 3-D assembly of monodispersed spheres of nanometric size. The crystalline assembly is also known to take place through attractive capillary forces as the solvent of the suspension evaporates slowly. A third method by which the crystalline assembly can take place is repulsive electrostatic interactions. Shown in Fig. 6.3A and B are the schematics of the two forms of the gravitational sedimentation method. In one of these (Fig. 6.3A) the assembly of light and small spheres such as those of polystyrene takes place through the continuous evaporation

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Figure 6.3 (A) Schematic of self-assembly process for the growth of colloidal crystals on a horizontal substrate driven by surface tension and capillary force. (B) Selfassembly driven by interactions between particles and substrate in a vertically held substrate. (C, D) Scanning electron microscopy images of a polystyrene photonic crystal grown by self-assembly on a vertically held glass substrate: (C) 300-nm spheres and (D) 600-nm spheres. (A) Adapted from http://www.tytlabs.com/english/ review/rev394epdf/e394_033nakamura.pdf; (B) S.-L. Kuai, G. Bader, P.V. Ashrit, Tunable electrochromic photonic crystals, Applied Physics Letters 86 (22) (2005) 221110/1e3; (D) S.-L. Kuai, G. Bader, P.V. Ashrit, Tunable electrochromic photonic crystals, Applied Physics Letters 86 (22) (2005) 221110/1e3.)

of the solvent from the horizontally placed substrate surface. The organization of the spheres takes place under surface tension and capillary force. If the substrate is placed in a vertical position, the crystal growth takes place through the interparticle interactions with the evaporation of the solvent, as shown in Fig. 6.3B. Each of these methods has its advantages and disadvantages. In Fig. 6.3C is shown an example of crystals of 300 and 600 nm grown on a glass substrate [22]. Two- and three-dimensional photonic crystals are also grown by depositing different materials with varying refractive index on an already patterned substrate [23]. Shown in Fig. 6.4A is this technique, called autocloning. A suitable substrate is first etched by electron beam lithography and dry etching. Subsequently the two different materials are sputter deposited and sputter etched to define a well-corrugated and repeating

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Figure 6.4 (A) Schematic of the growth of a two-dimensional photonic crystal by autocloning. (B) Fabrication apparatus for the autocloning. (C) Sputter deposition of two materials and sputter etching are carried out when radio-frequency (RF) power is applied to the targets and the substrate electrode, respectively. (From T. Sato, et al., Photonic crystals for the visible range fabricated by autocloning technique and their application, Optical and Quantum Electronics 34 (1) (2002) 63e70.)

pattern by using an apparatus as shown in Fig. 6.4B. In Fig. 6.4C are shown two such patterned 2-D photonic crystals based on the SiO2/TiO2 and SiO2/Ta2O5 film combinations. Such photonic crystals display a range of interesting properties important for a wide range of applications. In addition to the PBG, these films have birefringence, which is useful in controlling the state of polarization of light going through such crystals. A combination of a regular alternating flat multilayer and a corrugated multilayer as shown in Fig. 6.4A in the same photonic crystal device can be useful in creating polarization-selective grating. Shown in Fig. 6.5B is the actual multilayer device combining these two effects and the result of the optical power transmitted through such a grating. Similarly, actual devices embedded with the photonic crystals within the material are also fabricated. These so-called photonic crystal heterostructure

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Figure 6.5 (A) Schematic of a polarization-selective grating made from the combination of a flat multilayer and a corrugated multilayer [two-dimensional (2D) photonic crystal]. (B) Cross-sectional scanning electron microscopy photograph of the polarization-selective grating composed of TiO2/SiO2. (From T. Sato, et al., Photonic crystals for the visible range fabricated by autocloning technique and their application, Optical and Quantum Electronics 34 (1) (2002) 63e70.)

devices are useful for various applications such as resonant cavities, waveguides, and arrays of photonic crystals [2]. The easiest way to make the resonant cavities is to drill holes periodically in a dielectric slab. The refractive index contrast is provided by the material dielectric constant and that of the air filling the holes. Such a structure can block the propagation of certain wavelengths if the lattice constant (periodicity) of the structure is equal to half the wavelength of the guided light. Also, the introduction of various types of defects in the periodic structure can create states with frequencies within the bandgap. Modes corresponding to these frequencies can be confined within the defective region surrounded by the photonic crystal. Such slabs are used to create localized modes that can have frequencies inside the PBG. Such photonic crystals are

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Figure 6.6 (A) Schematic of a resonant cavity realized using photonic crystal heterostructure showing the refractive index profile. (B) A graded photonic crystal resonator. (Adapted from E. Istrate, E.H. Sargent, Photonic crystal heterostructures and interfaces, Reviews of Modern Physics 78 (2) (2006) 455e481.)

used in integrated laser and narrowband filters for wavelength division multiplexing applications [24]. It is also possible to construct various other types of photonic crystal heterostructure devices. Shown in Fig. 6.6A is a resonant cavity device that consists of two photonic crystals presenting a PBG at a particular wavelength of interest separated by a region that has a passband for the wavelength. Hence, light will be confined to the middle region creating resonant states. In addition to this type of abrupt structure, graded heterostructured photonic crystals such as the ones shown in Fig. 6.6B have been fabricated, which are shown to have less radiation loss in the resonant cavities [25].

6.2 TUNABLE PHOTONIC CRYSTALS Hence, it is clear that a wide range of photonic crystals and their integrated devices can be created to control and manipulate the propagation of light in materials. The versatility of such a control stems from the fact that the important parameters of interest relating to photonic crystals, such as the PBG, the quality factor, the mode volume, and more, can be controlled by a judicious choice of the materials of which the crystals are made (dielectric

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constant contrast) and their lattice constants (periodicity). However, the discussion hitherto was limited to passive modes of photonic crystals, i.e., crystals in which the PBG, for example, cannot be dynamically altered. If such a possibility of dynamically altering or “tuning” the behavior of a single photonic crystal or an integrated device can be realized, the photonic crystal application would become even more versatile. Such tunable photonic crystals can be controlled via external forces to manipulate the light flow, and a wavelength or frequency selectivity can be introduced into such a control in a single photonic crystal or its device. This possibility would open up numerous dynamically controllable photonic components such as tunable filters, optical switches, dynamic waveguides and cavities, reconfigurable optical networks, and more. To this end, a lot of research and development effort is concentrated to realize such tunable photonic crystals by varying the two important inherent parameters of the photonic crystals: (1) the lattice constant (periodicity) and (2) the refractive index (dielectric constant). The first category, of bringing about a reversible change in the lattice constant (periodicity) of the photonic crystals, has been carried out by using various external parameters such as mechanical force, electric or magnetic field, or light. Sumioka et al. [26] have fabricated polymethylmethacrylate (PMMA) inverse opals using silica opals with silica microspheres with a diameter of 285 nm. Such a PMMA-based inverse opal shown in Fig. 6.7A with air voids presents a PBG around 540 nm. It has been demonstrated that such a structure can be deformed reversibly, providing a tunability of the PBG, as shown in Fig. 6.7B. As can be seen in the transmittance spectra, for a stretch ratio of 1.5, the peak of the bandgap shifts from around 540 nm to around 480 nm. A distinct color change from green to blue associated with this stretch takes place as shown in Fig. 6.7C. The PBG tunability of such PMMA inverse opals is shown to be quasilinear with the stretch ratio of up to 1.6. Hence, mechanical force is the external control parameter used for the reversible tunability of the bandgap properties. Sheng et al. [27] have used electric and magnetic fields to transform the internal structure of electromagnetorheological (EMR) microspheres from body-centered tetragonal to face-centered cubic. The EMR spheres were made up of core microspheres of glass multiply coated with layers of Ni, Lead Zirconate Titanate (PZT), and TiO2. The variation in the dielectric constants as a function of the applied electric and magnetic fields is shown in Fig. 6.8. As can be seen in this figure a large variation in the dielectric constants can be induced by a suitable combination of electric and magnetic fields. Similarly,

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Figure 6.7 (A) Cross-sectional scanning electron microscopy image of polymethylmethacrylate (PMMA) inverse opal film fabricated from silica opal. (B) Transmission spectra for PMMA inverse opal film (dotted line) and stretched inverse opal film (stretch ratio ¼ 1.5). (C) Optical photograph of PMMA inverse opal film (left) and stretched PMMA opal film (right). (From K. Sumioka, H. Kayashima, T. Tsutsui, Tuning the optical properties of inverse opal photonic crystals by deformation, Advanced Materials 14 (18) (2002) 1284e1286.)

photoinduced phase transition has been achieved in photoresponsive photonic crystals to reversibly tune the optical properties by changing the lattice parameters [27]. Although a good variation in the PBG can be achieved with these methods, the structural change on the order of micrometers that is necessary may put a stringent limitation on their use in device form. The variation in refractive index (or dielectric constant) to reversibly tune the photonic properties is largely applied in the case of liquid crystals as it is very well known that these crystals show a wide range of optical property change with orientation, phase, and temperature. PBG tunability has been achieved by infiltrating birefringent liquid crystal into the voids of an inverse opal [28]. The optical property change, i.e., change in refractive index, in such a photonic crystal is achieved by applying a small

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Figure 6.8 Dielectric constant of the multicoated electromagnetorheological microspheres measured along the y direction (perpendicular to both the electric and the magnetic field) as a function of applied magnetic field under four fixed electric fields along the z direction. The dielectric constant is seen to exhibit a 3% dip. The position of the minimum is a linear function of the applied electric field (upper right inset). The volume fraction of the sample is 20%. The arrow indicates the magnetic field value where the face centered cubic structure was observed. (Adapted from P. Sheng, et al., Multiply coated microspheres. A platform for realizing fields-induced structural transition and photonic bandgap, Pure and Applied Chemistry 72 (2000) 309e315.)

electric field. Similarly, Martens et al. [28a] have achieved small PBG tunability in their silicon-based photonic crystals filled with a nematic liquid crystal (4-cyano-40 -pentylbiphenyl), which shows a clearing temperature of 34 C. In Fig. 6.9 are shown the results of such photonic crystals. As can be seen in this figure, a shift in the band edge can be brought about by changing the temperature as well as the plane of polarization of the light. Ferroelectric materials have also been proposed for such PBG tuning because of the phase transition they undergo with an electrooptical effect [29]. In both these categories of tunable photonic crystals either the PBG shift is not large enough or the control is not fine and versatile enough for practical applications. Transition metal oxides (TMOs) with their highly efficient, rapid, and versatile chromogenic coloration are an excellent choice for the fabrication of tunable photonic crystals. Various types of such photonic crystals based on TMO thin films have been realized, opening up

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Figure 6.9 (A) Scanning electron microscopy images (top and side view) of a threedimensional photonic crystal made of macroporous silicon with modulated pore radius. (B) Transmission spectra of a liquid crystal (LC)-filled photonic crystal in the G-A direction (solid line, pores filled with air; dotted lines, pores filled with 4-cyano-40 pentylbiphenyl). The shift in the band edge by 144 nm toward larger wavelengths is induced by the increase in temperature from 24 Ce to 40 C. (C) Polarization dependence. (From K. Busch, S. John, Liquid-crystal photonic-band-gap materials: the tunable electromagnetic vacuum, Physical Review Letters 83 (5) (1999) 967e970.)

a new vista of photonic application possibilities from waveguides to chemical sensors to high-efficiency lasers to solar cells and more. A wide variety of TMOs based 1-D, 2-D, and 3-D chromogenically tunable photonic crystals can be envisioned, which can be interactively operated by various external forces such as light (photochromic), electric field (electrochromics), heat (thermochromics), and more. Their optical switching properties can be tailored to different wavelengths by controlling the periodicity scale. One can envision either direct photonic crystals or inverse opals based on the TMOs. Given this wide range of possibilities, the field of tunable photonic crystals, especially the TMO-based ones, is just beginning to emerge and can afford a vast amount of scientific research and development work, which will open up the potential for new applications in photonics. In the following sections a discussion of the various types of TMO chromogenic

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properties based on photonic crystals is undertaken. These sections are divided according to the type of chromogenic property they are based on, i.e., electrochromic, thermochromic, and photochromic devices.

6.2.1 Electrochromically Tunable Photonic Crystals As discussed in Chapter 3, the phenomenon of reversible electrochromic coloration occurs via the double injection of electrons and ions bringing about a high-efficiency coloration in some of the TMOs such as WO3, MoO3, V2O5, TiO2, Cr2O3, and others. With this double insertion under the influence of a small electric field a strong coloration in the visible and near-infrared region occurs in these materials due to either an enhanced reflectance (free electrons) or an increased absorption (localized electrons) depending on the phase of the TMO thin film (amorphous or crystalline). Hence, a strong change in the optical (dielectric) constant ensues with the electrochromic coloration [30]. These films, therefore, have a great application potential in various devices such as energy-efficient smart windows, display devices, switchable mirrors, and more. Owing to these reasons the fabrication of photonic crystals based on these electrochromic materials is a highly desirable proposition. In the field of electrochromics, tungsten trioxide (WO3) is known to be the best and most versatile material, showing a highly efficient coloration in both polycrystalline (reflection modulation) and amorphous (absorption modulation) states. Further, WO3 thin films, as discussed in Chapter 3, can be prepared by a wide range of methods. The electrochromic modulation in these films occurs conveniently in the solar energy range, making them extremely attractive candidates for various applications. Hence, fabrication of tunable photonic crystals based on WO3 would be highly desirable. Sulan Kuai et al. [22,31] have fabricated inverse opals based on polycrystalline colloidal templates. The templates were grown on glass microslides through a convective assembly process. Photonic crystal templates were fabricated using polystyrene spheres of 300 and 760 nm. These templates were then dipped into a WO3 sol [32] to introduce WO3 by a dip-infiltrating solegel technique [33]. The resulting polystyreneegel composites were initially dried at 80 C and heated eventually at 460 C to sinter the polystyrene template, leaving behind a periodic framework of WO3-filled voids and airfilled spheres. Scanning electron microscopy images of these structures leading to the WO3 opals prepared with 300- and 760-nm polystyrene spheres are shown in Fig. 6.10. As can be seen in Fig. 6.10C and D the opals are made up of periodic macroporous hexagonal structures with three

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Figure 6.10 (A and B) Scanning electron microscopy (SEM) surface images of polystyrene (PS) photonic crystal template of (A) 300 nm and (B) 760 nm. (C and D) SEM surface images of WO3 inverse opals prepared with templates of PS spheres of (C) 300 nm and (D) 760 nm. (E and F) SEM surface images of WO3 inverse opals prepared with templates of PS spheres of (E) 300 nm and (F) 760 nm on a larger scale. (From S.-L. Kuai, G. Bader, P.V. Ashrit, Tunable electrochromic photonic crystals, Applied Physics Letters 86 (22) (2005) 221110/1e3.)

“dark” windows in each pore, corresponding to air spheres of underlying layers. The opals corresponding to the polystyrene spheres of 300 and 760 nm showed a center-to-center distance between air spheres of 283 and 702 nm, respectively, i.e., corresponding to a shrinkage of 6% and 8% with respect to the initial polystyrene spheres. According to X-ray diffraction analysis, the WO3 porous framework is well crystallized and mainly consists of hexagonal phase with an average crystal size of 26 nm. The presence of a PBG is clearly evident, as shown in Fig. 6.11, where the strong bands are present in the reflection spectra.

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Figure 6.11 Reflection spectra at near-normal incidence of the WO3 inverse opals prepared with templates using polystyrene spheres of 300 nm (curve a) and 760 nm (curve b). (From S.-L. Kuai, G. Bader, P.V. Ashrit, Tunable electrochromic photonic crystals, Applied Physics Letters 86 (22) (2005) 221110/1e3.)

These peaks in the spectra correspond to the pronounced Bragg diffraction peaks indicating the first and second stop bands in each of the samples. The band position depends (peak reflectance) strongly on the lattice constant (periodicity) with the first and the second PBGs at 532 nm for the WO3 inverse opal produced with 300-nm polystyrene spheres. In the case of the inverse opal fabricated with 760-nm spheres, these peaks are located at 1302 and 700 nm. Thus, by working with various-sized polystyrene or other nanospheres one can tailor the position of the bandgap. A significant result of this work is the demonstration of the continuous tunability of this bandgap by the electrochromic phenomenon. Lithium was inserted into these WO3 inverse opals by a dry lithiation process [34] to bring about electrochromic coloration according to the following reaction: WO3 ðtransparentÞ þ Li0x / Lix WO3 ðblueÞ In Fig. 6.12 are shown the optical results of this insertion into the WO3 inverse opal prepared with 300-nm spheres. As can be seen from the results in this figure, a systematic shift of the first bandgap and the second bandgap of this sample in reflectance toward shorter and longer wavelengths, respectively, occurs as a function of the inserted quantity of lithium. These changes stem from the complex changes occurring in both the real (N) and the imaginary (K) parts of the refractive index of WO3 with the insertion of lithium [35]. Although not reversible by design, this work demonstrates the

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Figure 6.12 (A) Reflection spectra of a WO3 inverse opal prepared with a template using 300-nm polystyrene spheres as a function of the inserted quantity of lithium (shown in the key). (B) Shift of first reflection peak located at 532 nm. (C) Shift of second reflection peak located at 344 nm. (D) Shift of the peak position of the first reflection. (E) Shift of the peak position of the second reflection. (From S.-L. Kuai, G. Bader, P.V. Ashrit, Tunable electrochromic photonic crystals, Applied Physics Letters 86 (22) (2005) 221110/1e3.)

potential for this approach and paves the way for building reversibly tunable photonic crystal devices. The experimental results shown in Fig. 6.12 were examined in the light of the theoretical models. The FTDT [11] method was applied to calculate the band structure of the initial polystyrene (PS) photonic crystal template and later the tunable PBG. The Brillouin zones of FCC structure in

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(A) K

L

X

W

L

Γ

X

(B) 1.2 1.0

ω a/2 π c

0.8 0.6 0.4 0.2 0.0

X

U

X

W

K

Figure 6.13 (A) The Brillouin zones of FCC structure in reciprocal space: X (1,0,0), U (1,1/4,1/4), L (1/2 ,1/2 ,1/2 ), G (0,0,0), X (1,0,0), W (1,1/2 ,0), K (3/4 ,3/4 ,0). (B) Band structure of a polystyrene crystal template, N ¼ 1.59 and K ¼ 0 (by plane-wave expansion method).

reciprocal space are shown in Fig. 6.13A indicating the different directions. The band structure of the PS crystal template is as shown in Fig. 6.13B. These calculations are done using by the plane-wave expansion method as far as the PS crystal template is concerned, assuming a refractive index (N) value of 1.59 and an extinction coefficient (K) of zero. The (111) planes of the WO3 inverse opal structure being parallel to the surface of the sample, it can be assumed that the first reflectance peak (Fig. 6.13B) corresponds to the stop band at the L point in the FCC reciprocal lattice. For the band structure calculations of the WO3 inverse opal in the as-prepared state and lithium-inserted state it is important to know the optical constants at each level of lithium insertion. Given in Fig. 6.14A and B are the real (N) and imaginary (K) parts of the refractive index as a function of the inserted lithium. The band structure calculated for the WO3 inverse

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(A) 2.4

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Figure 6.14 (A) Real, N, and (B) imaginary, K, parts of the refractive index of a sole gel-derived WO3 film of 190 nm as functions of the inserted lithium. (C) Band structure of an inverse opal structure assuming N ¼ 2.0307 and K ¼ 0 (by plane-wave expansion method).

opal with N and K values of 2.0207 and 0 at the wavelength of 515 nm is shown in Fig. 6.14C. As can be seen in this figure, the stop band widens at the L point. For calculating the band structure for the WO3 inverse opal with the insertion of lithium (colored state), i.e., K s 0, the FTDT simulation method was used. A typical band structure of a WO3 inverse opal structure is shown in Fig. 6.15A calculated with N and K values of 1.993 and 0.0239 obtained with the insertion of 5 nm equivalent thickness of lithium. Compared with WO3 inverse opal without lithium shown in Fig. 6.14C, with K ¼ 0, it can be seen that the introduction of K seems to cause a shift in the band structure and thus the position of the bandgap. In Fig. 6.15B is shown this shift of calculated stop band midposition as a function of the inserted lithium. It can be seen that with lithium insertion, the gap position shifts

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(A) 1.2 1.0

ωa/2πc

0.8 0.6 0.4 0.2 0.0 X

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Γ

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(B) 1.50

1.49

λc/a

1.48

1.47

1.46

1.45

0

2

4

6 8 10 Lithium (a.u.)

12

14

Figure 6.15 (A) Band structure of a WO3 inverse opal with N ¼ 1.993 and K ¼ 0.0239 at the wavelength of 515 nm (by finite differences in time domain method).

toward shorter wavelength, accordance with the experimental results for the first reflectance peak (Fig. 6.11B). However, to get the exact tendency of the entire spectrum, one has to consider the dispersion of both N and K as a function of the wavelength instead of choosing these values at a specific wavelength. The interpretation of the second reflection peak (Fig. 6.12C), which shifts to higher wavelengths with the insertion of lithium, is quite complicated owing to the absorption band edge in this wavelength region with abrupt changes in optical constants. The reflection spectra of such

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samples are influenced by many factors other than the photonic band structure, such as N and K values, surface and interface status, defects, etc. Thus, using photonic band structure to explain the reflection spectra and the changes occurring with the insertion of ions and electrons (electrochromic effect) is limited in some way. However, the calculations presented here give the overall tendency. In a further patented work, an electrochromically tunable photonic crystal device was fabricated by the same group [31,36]. A glass slide coated with an indium tin oxide (ITO) transparent conductor was used as the substrate to grow an initial PS-based template and then a WO3 inverse opal as described earlier. Two types of measurements can be carried out in this structure as shown in the schematic of Fig. 6.16A: (1) light incident from the inverse opal side and (2) light incident from the glass side. The reflectance spectra taken following these modes are shown in Fig. 6.16B. Despite some difference in the natures of the spectra, the two reflection peaks fall around the 500-nm wavelength. The device, as shown in Fig. 6.17A, was then constructed by sandwiching a plastic gasket (made of special plastic paper) between the ITO/WO3 inverse opal and a plain ITO glass (two ITO-film sides facing each other with the WO3 inverse opal in between). This assembly containing a small gap in the plastic gasket was completely immersed in a container filled with 1 M LiClO4 in propylene carbonate. The LiClO4 was backfilled into the sandwich structure by initially pumping a vacuum to take out the air in the structure and then by slowly letting the air into the chamber. The small hole in the plastic gasket was then sealed. Shown in Fig. 6.17B are the transmittance spectra of the electrochromic WO3 inverse opal device before and after the insertion of the LiClO4 solution. The light passing through the various layers in the device presents a very broad band and obscure dip in transmittance around 500 nm. The backfilling of the inverse opal with the LiClO4 solution provides a better defined stop band located at 640 nm. The factors influencing these changes are the refractive index contrast provided and the light diffusion provided by the opal structure. The electrical tunability of the device by applying positive and negative potentials to the ITO electrodes has been carried out. In Fig. 6.18A and B are shown transmittance spectra with the insertion and extraction of the lithium, respectively. With the application of increasing negative potential from 0 to 3.25 V a visible blue coloration takes place with the increasing absorption. These optical changes (refractive index) lead to a shift of the stop band toward the lower wavelengths. The discoloration of the device

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311

(A) 1

ITO Glass 2

(B) 250

Reflectance (a.u.)

200 2

150 100 50 0 300

1 400

500 600 700 Wavelength (nm)

800

Figure 6.16 (A) Two modes of reflection measurements possible with a WO3 inverse opal on indium tin oxide (ITO)-coated glass. (B) Reflectance spectra of the WO3 inverse opals on ITO carried out in the two modes of reflection. (From P. Ashrit, S. Kuai, Chromogenically Tunable Photonic Crystals, (US Patent 7660029 B2), February 9, 2010.)

can be induced by changing the polarity of the applied potential as shown in Fig. 6.18B. The peak shifts toward higher wavelength with the application of positive potential (extraction). In the as-prepared state of the device a mild electrochromic coloration exists owing to the residual insertion of the ions and electrons from the LiClO4 solution, which is removed upon the application of a high positive potential to the device. In Fig. 6.18C are shown the shifts of the peak wavelength (lc) with the double (electrons and ions) insertion and extraction of the charges with the application of a negative and positive potential, respectively, to the device. A total shift of around 15 nm is seen to occur in these devices. The current

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(A)

Glass Slide ITO

LiCIO4/PC ITO Glass Slide

Transmittance (%)

(B) 30 20 b

10

0 300

a

400 500 600 Wavelength (nm)

700

Figure 6.17 (A) Schematic of a WO3 inverse opal-based electrochromically tunable photonic crystal (EC-PC) device. (B) Transmittance spectra of the EC-PC device before (curve a) and after (curve b) infiltrating with the LiClO4 solution. ITO, indium tin oxide; PC, propylene carbonate. (From P. Ashrit, S. Kuai, Chromogenically Tunable Photonic Crystals, (US Patent 7660029 B2), February 9, 2010.)

(I) vs. voltage (V) curves associated with the insertion and extraction of charges are shown in Fig. 6.18D. Although the tuning of the stop band in these devices is not as well defined and effective as the WO3 inverse opals discussed earlier (Fig. 6.12), nevertheless these devices indicate the potential for a better stop band tunability. The main difference between the two approaches is that the stand-alone inverse opals were deposited on glass and the lithium insertion was carried out by a dry method. In the case of the electrochromic (EC) devices the lithium insertion and extraction were carried out electrochemically. This original work and first demonstration of the ability to tune the PBG by intercalating lithium into amorphous [36] and crystalline [37] WO3 inverse opals is expected to open up new and exciting developments for the photonic industry [38].

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-2

-1

0 -1

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1

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Voltage (V)

Lithium Insertion -2

Figure 6.18 Transmittance spectra of an electrochromically tunable photonic crystal device as a function of the applied (A) negative and (B) positive potential. (C) Peak wavelength (lc) of transmittance as a function of the applied voltage. (D) Current vs. voltage with the charge insertion and extraction. (From P. Ashrit, S. Kuai, Chromogenically Tunable Photonic Crystals, (US Patent 7660029 B2), February 9, 2010.)

Lili Yang et al. [39] have studied similar ordered macroporous WO3 films for infrared applications. WO3 inverse opals were prepared using 325-, 410-, and 620-nm PS latex spheres on fluorine-doped tin oxidecoated glass substrates. The scanning electron microscopy (SEM) images of these photonic crystals and the inverse opals are shown in Fig. 6.19A. An excellent closely packed ordering of the PS sphere photonic crystal is seen. Subsequent to the removal of the PS spheres after the infiltration of the WO3 precursor sol, WO3 inverse opals with hexagonal ordering is seen. Similar to the earlier work, three dark regions can be seen in each hole in the opal structure. A better ordering of the structure is reported with increasing pore size. The center-to-center distances of the macroporous holes of the inverse opal and those of the PS sphere crystals are reported to be, respectively, 330, 420, and 615 nm and 325, 415, and 635 nm. For a comparative study of the EC properties, a 520-nm-thick WO3 film and equivalent 2-mm-thick macroporous WO3 inverse opals

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Figure 6.19 (A) Scanning electron microscopy (SEM) images of polystyrene (PS) templates with different-sized spheres (a) and the corresponding macroporous inverse opals (b). (B) SEM images of a dense WO3 film of 520 nm (a) and macroporous WO3 inverse opals prepared using templates of PS sphere sizes of 325 nm (b), 410 nm (c), and 620 nm (d). (From L. Yang, et al., Improved Electrochromic performance of ordered macroporous tungsten oxide films for IR Electrochromic device, Solar Energy Materials and Solar Cells 100 (2012) 251e257.)

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Figure 6.20 (A) Cyclic voltammograms of the dense WO3 film (curve a) and the macroporous WO3 inverse opal (curve b). (B) Coloration efficiency (CE) spectra of a dense WO3 film and a macroporous WO3 inverse opal with a pore size of 325 nm. (C) CE spectra of WO3 inverse opal with different pore sizes, 325 nm (curve a), 410 nm (curve b), and 620 nm (curve c). (From L. Yang, et al., Improved Electrochromic performance of ordered macroporous tungsten oxide films for IR Electrochromic device, Solar Energy Materials and Solar Cells 100 (2012) 251e257.)

were deposited using the three sizes of PS sphere-based templates. Hence, the dense film and the macroporous inverse opals had the same amount of WO3. In Fig. 6.19B are shown the cross-section SEM images of all the samples. The EC insertion of lithium was carried out in these films, giving good coloration efficiency and reversible cycles. In Fig. 6.20A are compared the cyclic voltammetry curves associated with the double insertion and extraction of charges indicating a more efficient electrochemical performance in the inverse opal. The coloration efficiency (CE) spectra of the various types of structures are shown in Fig. 6.20B and C. It can be observed that (1) the macroporous WO3 inverse opals show a better CE compared to the dense WO3 film of equal thickness and (2) the wavelength at which occurs the peak efficiency in the broad CE spectra shifts to a higher wavelength in the inverse opal. Although the tunability of optical properties as a function of inserted

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

lithium is not explicitly verified in this work, it is indeed clear that one can bring about the wavelength selectivity through the quantity of charge inserted. A schematic of the inverse opal is shown in Fig. 6.21A to explain the improvement in EC CE in the WO3 inverse opal. The phenomenon is explained on the basis of light wavelength-to-pore size ratio. Light of lower wavelength radiation undergoes Mie scattering owing to the close correspondence between the pore size and the wavelength and also aided by the highly rigorous surface. Less light thus enters the porous and ordered structure and undergoes Bragg scattering. Thus absorption and hence the CE are low at lower wavelengths. Higher wavelengths do not undergo Mie scattering, as much more light enters, gets multiply refracted, and gets absorbed in the bulk of the structure, thus contributing to high CE at higher wavelengths. The overall EC efficiency also increases in the inverse opal because of the electronic conductivity established along the continuous thin wall structure of the inverse opal. Hence, such ordered structures are conducive to better EC performance because of the enhanced conductivity as well as increased absorption. The shift of peak CE of WO3 inverse opals that can be controlled electrochromically is very advantageous for infrared applications such as aerospace thermal control systems. Such systems can also be switched between colored and bleached states over many reversible cycles. Shown in Fig. 6.21B are transmittance switching cycles recorded at 800 nm of the three WO3 inverse opals studied. Quite efficient and reversible optical changes can be induced in these films within a few seconds. L. Kavan et al. [40] have examined the optical behavior of TiO2 anatase inverse opals under lithium insertion for applications such as dye-sensitized solar cells [41]. However, they found that the electrochemical performance of the inverse opals was poorer than that of the regular dense films. This was attributed to the lack of good connectivity between the particles forming the inverse opals, making the charge transport weak. From these few examples it is quite evident that the electrochromically tunable photonic crystals (EC-PCs), especially based on TMOs, offer an excellent possibility of controlling light propagation. The field of TMObased EC-PCs is in its infancy and offers a wide range of exciting possibilities for future research and applications in photonics. Although there has been tremendous research in the area of photonic crystals and tunable photonic crystals since the late 1990s, most of this work is dedicated to other forms of tunability. The TMOs such as WO3 offer a distinct and versatile form of tunability of optical and electrical properties. Even the few works reported to date have concentrated mostly on WO3 and TiO2 films,

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Figure 6.21 (A) Schematic of the bicontinuous electron/ion-conducting and light transmittance structure of WO3 inverse opal. (B) Coloration and bleaching cycles of the WO3 inverse opal at 800 nm with different pore sizes of 325 nm (a), 410 nm (b), and 620 nm (c). (From L. Yang, et al., Improved Electrochromic performance of ordered macroporous tungsten oxide films for IR Electrochromic device, Solar Energy Materials and Solar Cells 100 (2012) 251e257.)

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

while a large number of known and efficient EC TMOs such as molybdenum oxide (MoO3), chromium oxide (Cr2O3), and vanadium pentoxide (V2O5) need to be explored with a view to fabricating photonic crystals or inverse opals. This might lead to electrical and optical property tunability over a wide wavelength region, especially in the visible and near-infrared regions where rich photonic application possibilities exist.

6.2.2 Thermochromically Tunable Photonic Crystals As discussed in Chapter 4, thermochromism offers the possibility of tuning reversibly the optical and electrical properties in certain materials via the influence of heat. Among the TMOs, vanadium dioxide (VO2) is the most researched material because of its thermochromic effect, which occurs at a temperature nearest to room temperature. Hence, the thermochromic phenomenon in this material in which a semiconductor-to-metal transition occurs at the transition temperature (Tc) of 68 C is of immense interest from the scientific and application points of view. Around the Tc, VO2 is known to go through a conductivity increase of several orders and in the near-infrared region the material changes from a highly transparent state (semiconductor) to a highly reflective state. These changes are accompanied by a continuous or drastic change in their optical constants [41]. The nature of the semiconductor-to-metal transition (transition temperature, abruptness, recovery, and hysteresis) depends on a wide variety of parameters such as film thickness [42], nature of substrate [43], doping [44], and structure [45]. It is possible to prepare the doped and undoped VO2 films by various well-known solegel methods [46], which are conducive to fabrication of photonic crystals based on this material. Thus, this efficient and tailorable phase transition in VO2 opens the doors to the fabrication of highly efficient thermochromic photonic crystals (TC-PCs) in which the special optical and electrical properties of the photonic crystal can be controlled via the influence of heat. Further, the necessary conditions for the thermochromic switching can be provided, for example, by an electric heater (Joule effect) and/or by photothermal effect using an external and pertinent light source. PBG tuning using TC-PCs has become the object of an immense research effort owing to the application possibilities that exist in the area of optical communication. Further, it is possible to bring down the Tc via the doping of other transition metal atoms into VO2. All these possibilities make VO2 thin films and their photonic crystals very attractive for scientific and applied research.

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Figure 6.22 (A) Scanning electron microscopy (SEM) image of cleaved edges of a V2O5/SiO2 composite opal obtained after 6 infiltration cycles. The inset shows the same composite opal after 20 cycles. (B) SEM image of a VO2/SiO2composite opal obtained after 6 infiltration cycles. The inset shows the VO2 inverse opal after 20 cycles. The opals were made with 610-nm silica particles. Scale bars represent 2 mm. (C) Reflectance spectra of the base silica opal (a), V2O5/SiO2 composite (b), VO2/SiO2 composite (c), and VO2 inverse opal. (From M. Ibisate, D. Golmayo, C. Lopez, Vanadium dioxide thermochromic opals grown by chemical vapour deposition, Journal of Optics: Pure and Applied Optics 10 (12) (2008) 125202/1e6.)

M. Ibisate et al. [47] have fabricated large-area high-quality VO2/ SiO2opals. Silica sphere-based opals were first grown on silicon substrates by the vertical deposition method [47]. Vanadium pentoxide (V2O5) was first synthesized on the opals by chemical vapor deposition and then converted to VO2 by reduction. The VO2 inverse opals were then obtained by etching away the silica in HF. Shown in Fig. 6.22A and B are the SEM

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

images of the samples at various stages of preparation. Fig. 6.22C shows the reflectance spectra corresponding to the various stages of the VO2 inverse opal preparation. The Bragg peak arising from the silica photonic crystal’s periodicity is located at 1237 nm. With the infiltration of V2O5 into the voids of the photonic crystal and with the higher refractive index, the peak shifts to higher wavelengths situated around 1750 nm. The reduction process leading to the formation of VO2/SiO2 composite opal broadens and shifts the peak to slightly lower wavelengths. With the dissolution of the silica spheres forming the VO2 inverse opal the peak shifts to lower wavelengths again, reflecting the average lower refractive index of VO2 at these wavelengths. The thermochromic switching characteristics and the tunability of the optical properties of the VO2/SiO2 composite opal and the VO2 inverse opal have been studied in reflectance mode. In Fig. 6.23 are given the

Figure 6.23 (A) Reflectance spectra from the (111) surface of the VO2/SiO2 composite opal at different temperatures, including a spectrum taken after cooling back down to room temperature (RT). (B) Reflectance spectra from the (111) surface of the VO2 inverse opal at different temperatures. (C) Thermal differential reflectance spectra obtained from the spectra shown in (B). (From M. Ibisate, D. Golmayom, C. Lopez, Vanadium dioxide thermochromic opals grown by chemical vapour deposition, Journal of Optics: Pure and Applied Optics 10 (12) (2008) 125202/1e6.)

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reflectance results as a function of temperature. In Fig. 6.23A are given the reflectance spectra of the VO2/SiO2 composite opal at different temperatures from room temperature to above the transition temperature. A distinct difference in the spectral nature of the opals occurs below and above the transition temperature. With increasing temperature there is a blue shift in the peak around 1662 nm up to a temperature (65 C) just below the Tc. At the Tc (68 C), drastic changes occur and a broad peak centered around 1618 nm develops in the spectra. For temperatures above the Tc the phase transition of the VO2 to metallic phase leads to a high reflectance of the opal for wavelengths above 1.5 mm. Similar results are shown in Fig. 6.23B for the VO2 inverse opal. Here, with increasing temperature, the Bragg peak situated at 1294 nm decreases with nearly no shift up to a temperature of 65 C. For temperatures above Tc, with the onset of the metallic state, drastic changes take place, which are very different from the composite opal seen in Fig. 6.23B. The complex changes occurring in the refractive index give rise to a new peak around 105 nm and a high reflectance in the infrared region. Differential reflectance data arising from these thermal variations are presented in Fig. 6.23C to better present the tunability of the optical properties in the VO2 inverse opals. Quite significant oscillations in the differential reflectance can be observed near the transition temperatures. These changes can have important implications in certain devices. Thus a fine control of optical properties using such photonic devices is a distinct possibility. Golubev et al. [48] have studied PBG control via thermochromic phase transition of opal/VO2 photonic crystals. These 3-D crystals were prepared using SiO2-based synthetic opals wherein the diameter of the SiO2 spheres was around 250 nm. These photonic crystals were infiltrated initially with V2O5 by chemical bath deposition and later annealed in an atmosphere at 800 C to obtain opal/VO2 composites. The Bragg reflection peak arising from the periodic planes in the opal/VO2 photonic crystal is situated around 625 nm. Shown in Fig. 6.24A are the reflection spectra of the opal/VO2 composite at temperatures below (semiconductor state) and above (metallic state) the transition temperature (Tc) of VO2. As can be seen in this figure the PBG clearly shifts by 25 nm to lower wavelengths with the thermochromic phase transition. The reflection band becomes narrower and lower in intensity at temperatures above the transition owing to the overall decrease in the dielectric constant of the composite structure and the optical contrast between the two components of the photonic crystal. Fig. 6.24B shows the thermal hysteresis of the peak position of reflectance (PBG) of the composite

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Transition Metal Oxide Thin FilmeBased Chromogenics and Devices

Figure 6.24 (A) Reflection spectra from the (111) surface of the opal/VO2 composite. (Curve 1) VO2 in the semiconductor phase, T ¼ 15 C; (curve 2) VO2 in the metallic phase, T ¼ 75 C. (B) Thermal hysteresis loop of the reflection peak position of the opal/ VO2 composite during the sample heating and cooling. (From V.G. Golubev, et al., Phase transition-governed opal VO2 photonic crystal, Applied Physics Letters 79 (14) (2001) 2127e2129.)

structure. Hence, a good control over the PBG as a function of temperature within this wavelength is clearly achievable in these structures. Among the inorganic materials, VO2 is by far the most studied for thermochromic properties. Hence, even the few works related to photonic crystals focus on VO2-based TC-PCs. These few examples of the VO2-based TC-PCs are cited to illustrate the immense possibilities that exist for controlling the optical and electrical properties by building other TMO-based photonic crystals. Such a temperature-based tunability can also be extended to other TMOs such as oxides of vanadium (VxOy), WO3, MoO3, and more, which show phase transition at other temperatures. As mentioned earlier, the possibility also exists of fabricating other transition metal atom-doped VO2-based photonic crystals or their inverse opals, which would have their Tc even closer to room temperature. The thermochromic switching such as the one in VO2 is, in general, an abrupt phenomenon. The switch occurs precisely around the Tc at which the semiconductor-to-metal transition occurs and hence such TC-PCs are invariably proposed for switching applications such as in optical switches, etc. This is unlike the gradual tuning of properties that occurs in EC-PCs. However, in many VO2 thin films a gradual or multistep transition has been observed depending on the phase, structure, and particle size

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distribution. The fabrication of such TC-PCs with VO2 with controlled phase transition can be envisioned to have a better tuning of the optical properties such as PBG.

6.2.3 Photochromically Tunable Photonic Crystals Similar to photonic crystal-based EC materials such as WO3 and photochromic materials such as VO2, one can envision the fabrication of photonic crystals based on photochromic materials. The photochromic phenomenon relating to TMO thin films is discussed in Chapter 5 in detail. Among the TMOs, MoO3 and WO3 exhibit a very efficient photochromism under the action of UV light and, in some cases, visible light. Hence, if photonic crystals can be fabricated using such photochromic materials, the tuning of optical properties via light irradiation can be brought about. These changes originate from the ensuing chemical reaction under light irradiation [49]. The photochromic mechanism in MoO3, for example, can be explained in a way similar to that of electrochromism with the difference that in the former the inserted positive ions come from within the structure of MoO3 films. Electrons and holes are generated in these under appropriate conditions. The electrons thus generated then react with the adsorbed H2O molecules, forming the protons necessary for the coloration owing to the formation of molybdenum bronze (HxMoO3). This reversible reaction can lead to large changes in the refractive indices under the action of light. Hence, the photochromic mechanism in TMO-based photonic crystals can be of great interest in tuning the optical properties of these structures via an external light source. One-dimensional photonic crystals (Bragg mirrors) based on the photochromic effect in MoO3 are reported by Ben-Messaoud et al. [50]. Alternating layers of high- and low-refractive-index MoO3 and SiO2 by radio-frequency sputtering on glass substrate provide a refractive index contrast of 2.1 to 1.5. The photochromic response of the multilayer stack was measured using a standard pumpeprobe setup. The pump beam was a UV radiation centered at 380 nm with a spectral band width of 50 nm. The optical changes were probed using halogenedeuterium white light. The photonic band structure of MoO3/SiO2 multilayers before and after 1 h of UV irradiation is shown in Fig. 6.25A for 5, 10, and 20 pairs (multilayers). A photonic stop band appears between 625 and 700 nm whose intensity depends on the number of pairs of high/low index in the multilayer system. The 20-multilayer system has a much less pronounced transmittance and stop band as it is marred by high diffusion and nonuniformity. The optical changes occurring with UV irradiation are better represented in the system

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Figure 6.25 (A) Transmittance spectra of MoO3/SiO2 multilayers before and after the films were irradiated with UV light for 1 h. 5, 10, and 20 m indicate 5, 10, and 20 pairs of layers in the system. (B) Relative changes in reflectance (above) and transmittance (below) in MoO3/SiO2 multilayers exposed to UV light for 1 hour. (From T. Ben-Messaoud, et al., Photoactive periodic media, Applied Physics Letters 94 (11) (2009) 111904/1e3.)

with 10 pairs of high/low-index films. To see the changes occurring around the band edges, relative changes occurring in transmittance and reflectance were calculated. These are shown in Fig. 6.25B. These relative optical changes increase with increasing number of pairs in the structure. A sharp peak arising around the lower edge of the stop band becomes very well pronounced in the photonic crystal with 20 pairs of high/low-index films. This change is due to the variation in refractive index of MoO3. Hence, it may be possible to tune this DR/R, for example, by varying the duration and/or intensity of UV irradiation. Although such a tunability is not explicitly studied here, the work indeed confirms this possibility. As discussed in Chapter 5, the photochromic coloration in MoO3 and other TMO thin films is quite stable and not easily reversible. To induce a reversible tuning of optical properties in TMOs and TMO-based photonic crystals, a special device configuration that integrates EC discoloration, for example, as seen earlier in Fig. 5.14B, may be envisioned. An integrated photoelectrochromic photonic crystal (PEC-PC) device based on TMOs such as MoO3 and WO3 is plausible. A schematic of such a device is shown in Fig. 6.26. Such a device can be initially irradiated to induce photochromic

Chromogenic Thin Film Photonic Crystals

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Glass ITO Electrolyte

ITO

TMO based photochromic photonic crystal

Glass

Figure 6.26 Schematic of a proposed photoelectrochromic photonic crystal device. ITO, indium tin oxide; TMO, transition metal oxide.

coloration in the TMO-based photonic crystal due to the photogenerated electron and proton injection. The photonic stop band arising from the periodic structure can be displaced progressively depending on the duration and intensity of the UV or visible light source. These changes occur because of the change in refractive index of the TMO used. The process of photochromic coloration can be reversed progressively electrochromically by applying a small voltage with polarity as shown in Fig. 6.26. Thus, this scheme of photochromic coloration in conjunction with EC discoloration will render possible the reversible tuning of the optical properties in photonic crystals and will increase the technological scope of such devices in photonics. Further, as discussed in Chapter 5, it is reported that a slight EC coloration by cathodic polarization prior to photoirradiation of the TMO samples is expected to render the photochromic phenomenon even more efficient [51]. Hence, the configuration shown in Fig. 6.26 can be advantageous not only in inducing a reversible coloration but also in enhancing it. To the best of author’s knowledge such PEC-PCs have not been realized as of this writing.

6.3 SUMMARY Photonic crystals are periodically structured electromagnetic media possessing a PBG for a certain range of frequencies. In recent years, photonic crystals have received a lot of research attention because of the unique possibility they offer of controlling the propagation of light. Among many possible applications, a tunable PBG is highly desirable as it allows the control of the photonic properties externally. This can lead to various

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application possibilities such as optical switches, tunable filters, reconfigurable optical networks, etc. The PBG depends mainly on two material aspects of the photonic crystals: (1) periodicity (lattice constants) and (2) refractive index. One can apply external mechanical, electrical, and magnetic forces to change the structure of the crystal. However, this approach induces a lot of strain on the crystal structure and is hard to adopt in device form. The reversible and facile coloration of TMOs is very conducive for the fabrication of photonic crystals as they undergo large variations in refractive index. By fabricating photonic crystals of various TMOs, the optical properties such as PBG can be reversibly tuned electrochromically, thermochromically, or photochromically by using an electric field or heat or light, respectively. EC-PCs seem to be the most versatile as they offer a reversible and continuously tunable control of optical properties. Thermochromically tunable photonic crystals are very efficient from the point of view of sudden switching of properties, which happens around the transition temperature (Tc). Photochromically tunable photonic crystals are also very efficient from the point of view of gradual tuning of optical properties. However, at present their scope in stand-alone form is limited by the controlled irreversibility of their operation. However, if integrated with EC discoloration such PECECs can be extremely interesting. The field of chromogenically tunable photonic crystals is in its infancy. Although a lot of theoretical and a substantial amount of experimental work has been carried out on passive photonic crystals and some of their applications have been realized, an enormous scope exists for basic and applied research in the area of TMObased tunable photonic crystals. On the theoretical side the challenge is to model the complicated and dynamic changes occurring in the crystals by appropriate methods. The success of a good modeling hinges on establishing a good correlation between the changes occurring in the dielectric (optical) constants in the crystals with the external force applied, such as heat, light, and electric field. On the experimental side the challenge is to fabricate in a reproducible way the photonic crystals or their inverse opals based on the TMOs. If this challenge is well addressed, with the wide range of TMOs that are available that exhibit chromogenic properties, an enormous scope exists to undertake research in this domain. Further, by varying the size of the particles of which the photonic crystals are made, the PBG and its tunability via different external forces can be tailored to the wavelength region of interest. All these possibilities for research and industrial application in the field of photonics make the field of chromogenically tunable photonic crystals an exciting one for the near future.

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REFERENCES [1] J.D. Joannopoulos, S.G. Johnson, J.N. Winn, R.D. Meade, Photonic Crystals: Molding the Flow of Light, second ed., Princeton University Press, 2008. [2] E. Istrate, E.H. Sargent, Photonic crystal heterostructures and interfaces, Reviews of Modern Physics 78 (2) (2006) 455e481. [3] J. Lekner, Omnidirectional reflection by multilayer dielectric mirrors, Journal of Optics A: Pure and Applied Optics 2 (5) (2000) 349e352. [4] http://www.batop.de/information/r_Bragg.html. [5] https://en.wikipedia.org/wiki/Opal#/media/File:Opal_molecular_structure2.jpg. [6] E. Yablonovitch, Inhibited spontaneous emission in solid-state physics and electronics, Physical Review Letters 58 (20) (1987) 2059e2062. [7] S. John, Strong localization of photons in certain disordered dielectric superlattices, Physical Review Letters 58 (23) (1987) 2486e2489. [8] E. Yablonovitch, Photonic crystals: semiconductors of light, Scientific American 285 (6) (2001) 47e55. [9] E. Yablonovitch, T.J. Gmitter, K.M. Leung, Photonic band structure: the facecentered- cubic case employing nonspherical atoms, Physical Review Letters 67 (17) (1991) 2295e2298. [10] K.M. Ho, C.T. Chan, C.M. Soukoulis, Existence of a photonic gap in periodic dielectric structures, Physical Review Letters 65 (25) (1990) 3152e3155. [11] C.T. Chan, Q.L. Yu, K.M. Ho, Order-N spectral method for electromagnetic waves, Physical Review B 51 (23) (1995) 16635e16642. [12] J.B. Pendry, Calculating photonic band structure, Journal of Physics: Condensed Matter 8 (9) (1996) 1085e1108. [13] N. Stefanou, V. Karathanos, A. Modinos, Scattering of electromagnetic waves by periodic structures, Journal of Physics: Condensed Matter 4 (36) (1992) 7389e7400. [14] N. Stefanou, V. Yannopapas, A. Modinos, MULTEM 2: a new version of the program for transmission and band-structure calculations of photonic crystals, Computer Physics Communications 132 (1e2) (2000) 189e196. [15] K.S. Yee, Numerical solution of initial boundary value problems involving maxwell’s equations in isotropic media, IEEE Transactions on Antennas and Propagation 14 (3) (1966) 302e307. [16] A. Taflove, S.C. Hagness, Computational Electrodynamics: The Finite-Difference Time-Domain Method, Artech House, Norwood, MA, 2000. [17] P. Yeh, Optical Waves in Periodic Media, Wiley, New York, 1998. [18] J.B. Pendry, A. MacKinnon, Calculation of photon dispersion relations, Physical Review Letters 69 (19) (1992) 2772e2775. [19] P.M. Bell, J.B. Pendry, L.M. Moreno, A.J. Ward, A program for calculating photonic band structures and transmission coefficients of complex structures, Computational Physics Communications 85 (2) (1995) 306e322. [20] Y. Xia, B. Gates, Y. Yin, Y. Lu, Monodispersed colloidal spheres: old materials with new applications, Advanced Materials 12 (10) (2000) 693e713. [21] http://www.tytlabs.com/english/review/rev394epdf/e394_033nakamura.pdf. [22] S.-L. Kuai, G. Bader, P.V. Ashrit, Tunable electrochromic photonic crystals, Applied Physics Letters 86 (22) (2005), 221110/1e3. [23] T. Sato, et al., Photonic crystals for the visible range fabricated by autocloning technique and their application, Optical and Quantum Electronics 34 (1) (2002) 63e70. [24] https://www.cst.com/Content/Articles/article687/CST_Whitepaper_Photonic_ CST.pdf. [25] K. Srinivasan, P.E. Barclay, O. Painter, Fabrication-tolerant high quality factor photonic crystal microcavities, Optics Express 12 (7) (2004) 1458e1463.

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[26] K. Sumioka, H. Kayashima, T. Tsutsui, Tuning the optical properties of inverse opal photonic crystals by deformation, Advanced Materials 14 (18) (2002) 1284e1286. [27] P. Sheng, et al., Multiply coated microspheres. A platform for realizing fields-induced structural transition and photonic bandgap, Pure and Applied Chemistry 72 (2000) 309e315. [28] K. Busch, S. John, Liquid-crystal photonic-band-gap materials: the tunable electromagnetic vacuum, Physical Review Letters 83 (5) (1999) 967e970. [28a] G. Martens, et al., Shift of the photonic band gap in two photonic Crystal/liquid Crystal composites, Applied Physics Letters 80 (11) (2002) 1885e1887. [29] B. Li, et al., Ferroelectric inverse opals with electrically tunable photonic band gap, Applied Physics Letters 83 (23) (2003) 4704e4706. [30] C.G. Granqvist, Handbook of Inorganic Electrochromic Materials, Elsevier, Amsterdam, 1995. [31] P. Ashrit, A new approach to manipulating light propagation, SPIE Newsroom (2010) 2, http://dx.doi.org/10.1117/2.1201005.002937. [32] J.P. Cronin, D.J. Tarico, A. Agarwal, L. Zhang, (US Patent No. 5277986), January 11, 1994. [33] S. Kuai, et al., Preparation of large-area 3D ordered macroporous titania films by silica colloidal crystal templating, Advanced Materials 15 (1) (2003) 73e75. [34] P.V. Ashrit, G. Bader, F.E. Girouard, V.-V. Truong, Electrochromic materials for smart window application, Proceedings of SPIE 1401 (1990) 119e129. [35] G. Bader, P.V. Ashrit, V.-V. Truong, Transmission and reflection ellipsometry studies of Electrochromic materials and devices, Proceedings of SPIE 2531 (1995) 70. [36] P. Ashrit, S. Kuai, Chromogenically Tunable Photonic Crystals (US Patent 7660029 B2), February 9, 2010. [37] T. Sumida, Y. Wada, T. Kitamura, S. Yanagida, Electrochemical change of the photonic stop band of the ordered macroporous WO3, films, Chemical Letters 31 (2) (2002) 180e181. [38] S.K. Deb, Opportunities and challenges in science and technology of WO3 for electrochemic and related application, Solar Energy Materials and Solar Cells 92 (2) (2008) 245e258. [39] L. Yang, et al., Improved Electrochromic performance of ordered macroporous tungsten oxide films for IR Electrochromic device, Solar Energy Materials and Solar Cells 100 (2012) 251e257. [40] L. Kavan, M. Zukalova, M. Kalbac, M. Graetzel, Lithium insertion into anatase inverse opal, Journal of the Electrochemical Society 151 (8) (2004) A1301eA1307. [41] S. Nishimura, et al., Standing wave enhancement of red absorbance and photocurrent in dye-sensitized titanium dioxide photoelectrodes coupled to photonic crystals, Journal of American Chemical Society 125 (20) (2003) 6306e6310. [42] G. Xu, P. Jin, M. Tazawa, K. Yoshimura, Thickness dependence of optical properties of VO2 thin films epitaxially grown on sapphire, Applied Surface Science 244 (2005) 449e452. [43] Y. Muraoka, Z. Hiroi, Metal-insulator transition of VO2 thin films grown on TiO2 (001) and (110) substrates, Applied Physics Letters 80 (4) (2002) 583e585. [44] M. Soltani, et al., Optical switching of vanadium dioxide thin films deposited by reactive pulsed laser deposition, Journal of Vacuum Science & Technology A: Vacuum Surfaces and Films 22 (3) (2004) 859. [45] C. Petit, J.-M. Frigerio, M. Goldmann, Hysteresis of the metal- insulator transition of VO2; evidence of the influence of microscopic texturation, Journal of Physics; Condensed Matter 11 (16) (1999) 3259e3264. [46] Y. Ningyi, L. Jinhua, L. Chenglu, Valance reduction process of sold-gel V2O5 to VO2 thin films, Applied Surface Science 191 (1e4) (2002) 176e180.

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[47] M. Ibisate, D. Golmayo, C. Lopez, Vanadium dioxide thermochromic opals grown by chemical vapour deposition, Journal of Optics: Pure and Applied Optics 10 (12) (2008), 125202/1e6. [48] V.G. Golubev, et al., Phase transition-governed opal VO2 photonic crystal, Applied Physics Letters 79 (14) (2001) 2127e2129. [49] T. He, et al., Enhancement effect of gold nanoparticles on the UV-light photochromism of molybdenum trioxide thin films, Langmuir 17 (26) (2001) 8024e8027. [50] T. Ben-Messaoud, et al., Photoactive periodic media, Applied Physics Letters 94 (11) (2009), 111904/1e3. [51] Y.A. Yang, et al., Visible light photochromism in electrolytically pretreated WO3, Journal of Physics and Chemistry of Solids 59 (9) (1998) 1667e1670.

FURTHER READING [1] Z.Z. Gu, A. Fujishima, O. Sato, Photochemically tunable colloidal crystals, Journal of American Chemical Society 122 (49) (2000) 12387e12388.

CHAPTER 7

Emerging Technologies* The marriage between chromogenics and transition metal oxides (TMOs) is becoming increasingly important in view of the various application possibilities. In the previous chapters, we discussed in detail electrochromic, photochromic, and thermochromic effects. We also observed that TMOs lend themselves quite well to fabrication as thin films with established techniques. As thin films they are susceptible to nanostructuring and nanosculpting, which add to their potential morphologies and hence novel or enhanced properties [1e4]. In this chapter we try to summarize and highlight the most interesting, innovative, and important applications of these effects. The choice of emerging technologies in this chapter was based on their potential, as well as their ability to highlight the use of various chromogenic effects. That means that some interesting applications, such as the use of MoO3 as a photoinduced catalyst, its possible use in battery technology, or its potential as a hole-injection layer in organic lightemitting diodes, are omitted, even though the mechanisms for such applications are closely related to those for chromogenic properties [5e8]. The following chosen examples start with the use of chromogenic effects in smart windows, a well-researched field that is already commercialized, albeit there are still areas of improvements to be addressed. On the other side of the spectrum, we have chosen applications such as infrared wave-front mapping, memristive devices, and security features. These effects have so far been demonstrated in research only; however, they represent new venues and a vast potential. Whenever possible, we have given examples of companies working on chromogenic devices.

7.1 SMART WINDOWS With the dire and imminent consequences of global climate change, it is ever more important to reduce the use of fossil fuels and employ energyefficient techniques. So far, perhaps the most important use of chromogenic materials has been in the design and fabrication of energy-efficient smart windows with the ability to automatically or manually control the * The major contribution of Dr. Gisia Beydaghyan in writing this chapter is greatly appreciated. Transition Metal Oxide Thin FilmeBased Chromogenics and Devices ISBN 978-0-08-101747-0 http://dx.doi.org/10.1016/B978-0-08-101747-0.00007-6

© 2017 Elsevier Ltd. All rights reserved.

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amount of transmitted radiation into interior spaces. Such designs could be based on electrochromism, photochromism, or thermochromism, of which the first is the most investigated and furthest developed [9e14]. The electrochromic design is basically an electrochromic cell in which coloration and bleaching are the means by which the amount of transmitted radiation is controlled (see Fig. 3.1). Further, it is possible to employ nanostructured, amorphous, or crystalline TMOs (e.g., WO3) to choose the peak of the energy rejection/transmission band, as it would be desirable to reject heat-generating near-IR and IR radiation while passing the visible portion of the spectrum [3]. As well, embedded nanocrystals such as indium tin oxide (ITO) can increase the electrochromic effect in such designs [15]. According to a report prepared by the Freedonia Group, the size of the worldwide market for flat glass was worth US$72 billion in 2013 and is expected to rise to US$139.9 billion by 2023 [16]. A major segment of the flat glass market is architectural glass and about 40% of global glass utilization is in building refurbishment, which is an ideal area for the use of smart windows. Other areas for use include aircraft windows, automobile sunroofs, displays, and visors. Smart windows based on chromogenic TMOs have already been commercially produced and utilized for energy efficiency in large buildings by Asahi and Saint-Gobain. Research Frontiers has developed and licensed SPD Smart Glass for use in the automotive and aircraft industries. Boeing and Airbus have already installed chromogenic aircraft windows in their fleet. An excellent dossier of companies working on chromogenic smart materials, especially smart glass, has been gathered by C.M. Lampert and in one of the older references, he provides a very useful table of electrochromic device structures along with the device maker and application areas [17e19]. It is interesting to note that the design of an electrochromic (EC) cell bears a great similarity to that of a photovoltaic (PV) cell, and in fact, designs have been proposed to fabricate EC-PV cells either side by side or as part of an integrated monolithic device. As such, the PV cell provides the power required for the operation of the EC device. One such design is proposed by Deb [20] in which a wide-bandgap amorphous SiC:H n-i-p PV cell provides the power for a LixWO3/LiAlF4/V2O5 EC device (Fig. 7.1). The coloration provided by this stand-alone device is comparable to that obtained by applying 0.6 V of external voltage under 1-sun illumination (60% for PV powered vs. 80% with applied external voltage). A middle transparent conductor may be further added to this design to provide better control and possible options for charging [20].

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Figure 7.1 Monolithic photovoltaic (PV) (n-type/intrinsic/p-type (nip) SiC:H)e electrochromic (EC) smart window. (A) Schematic structure and (B) coloration (under short circuit (S.C) and under 0.6V application) and bleaching spectra [20]. a-SiC:H, amorphous SiC:H; ITO, indium tin oxide.

In the same reference, Deb discusses an alternative design for smart windows with a simpler approach than a combined EC-PV cell described earlier. In this approach one-half of a dye-sensitized solar cell is combined with one-half of an EC cell with lithium ions added to the electrolyte [21]. This design is capable of producing large changes between bleached and colored states (Fig. 7.2). However, because there are only two electrodes, there is less user control in its operation. This device could be designed based on purely EC or photochromic effect or a combination of two.

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Figure 7.2 Photochromic device using dye-sensitized TiO2 and electrochromic WO3. (A) Schematic diagram and (B) spectra of colored and bleached states [20].

Smart windows may also be designed based on the gasochromic effect, which is coloration due to the presence of a reducing gas such as hydrogen. The structure consists of a thin layer of WO3 covered by a thin layer of platinum in an isolated, closed cavity [22]. When the hydrogen gas is introduced, the WO3 changes to the colored state. The introduction of oxygen brings the window back to the bleached state. The use of gases in these designs makes them somewhat complicated to install; however, they have the advantage of not requiring electrodes and being fairly simple to construct.

7.2 CHROMOGENIC MIRRORS A concept naturally evolving from chromogenic devices as smart windows is chromogenic mirrors. By replacing the substrate (or end surface) of an EC device with a reflecting surface, the device is transformed from a window to

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Figure 7.3 An electrochromic device used in transmittance (left) and reflection mode (right) [24]. EC, electrochromic; TCE, transparent conducting electrode.

a mirror, as depicted in Fig. 7.3. Hence devices similar to those discussed in the previous section can be envisioned as mirrors with electrochromism, photochromism, or thermochromism as the underlying mechanism. EC rearview mirrors are already commercialized and have been used in the automotive industry since 1982. In such designs, WO3 is universally the EC layer. Chromogenic mirrors are now an option in most major brands of car. Gentex has been a pioneer in this area and still has the dominant market share. It estimated that by 2004, 18% of vehicles had an interior autodimming mirror [23]. Other companies producing chromogenic mirrors for the automotive industry are Magna Donnelly, Schott, Nikon, and Toyota, among others [17,19].

7.3 SATELLITE WINDOWS Among the challenges for successful operation of satellites and space missions is the issue of managing extreme temperatures [25]. It is desirable to protect the astronauts and also the instruments in the interior of a spacecraft or on a satellite from excessive heat and radiation. Any device for temperature management should also have as small a mass as possible. Satellite windows with suitable coatings can provide protection and control. However, a coating with variable emissivity is the most desired. That is because under the condition of high radiation, when the satellite is facing the sun, the window is required to have low absorption and high emissivity, whereas under satellite eclipse (earth shade) conditions, it is required to have

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Figure 7.4 Schematic of a variable emissivity electrochromic device [26].

high absorption and low emissivity. Therefore designs based on electrochromism or thermochromism are very attractive for such coatings. A heat modulator with an EC layer and an IR-reflective electrode is shown in Fig. 7.4. This device is controlled by intercalation or extraction of electroneion pairs. When the ions are in the ion storage layer, the EC layer is transparent and therefore the device has low absorption and high emissivity. When the electroneion pair is intercalated in the EC layer, it becomes absorptive and nonreflective, i.e., a low-emissivity device. This device was tested in space with very favorable results in the 7e12 mm farIR region [26]. Similar IR switching EC devices for thermal emittance modulation were fabricated by Franke et al. [27]. This device is based on tungsten oxide (WO3) with an Al layer or grid as the reflecting surface and operates in the 2e35 mm region. Similar work in this area was reported by Tachikawa et al. [28]. In this work a “smart radiation device” (SRD) was made of thermochromic oxides such as La1 xSrxMnO3 or La1 yCayMnO3. These oxides, which are based on MnO3, have a transition temperature near room temperature, and this transition is exploited to provide the variation in emissivity. The device has low and high emissivity below and above the transition temperature, respectively. A maximum emissivity of 0.60e0.65 was measured for this device; however, the absorption of this device was too large and unacceptable when exposed to solar radiation. To remedy this situation, the

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authors further fabricated SRDs with multilayer coatings to reflect solar radiation. The integrated absorptance decreased from 0.81 to 0.28 upon the addition of multilayer coatings.

7.4 INFRARED LASER WAVE-FRONT MAPPING In almost any industrial or scientific application of a laser beam, characterization of laser wave front and its profile are of fundamental importance. Although simple, compact, and affordable sensors exist in the visible range, producing comparable products in the IR region is much more challenging. In particular, owing to the thermal noise being comparable to the photon energy in the IR region, cooling of detectors to liquid N2 temperatures (77 K) is often required, especially in the 1.7e8 mm range. In the work of Bonora et al. [29,30], this challenge is met by exploiting the change in the refractive index of a VO2 thin film due to thermochromic phase transition. The resulting change in the reflectivity of the VO2 film under the irradiation of the pumping beam can be read at the near-IR or visible range, and hence one is able to convert the wavelength of detection to a region in which compact sensors are available. Fig. 7.5 shows the results of this work, which compares very well to a commercially available sensor. In accordance with the VO2 transition, this technique covers a broad spectral range and is shown to be capable of high resolution and low threshold (w100 mW/cm2).

7.5 MEMRISTIVE DEVICES Memristive devices are systems whose resistance depends on the entire dynamical history of applied voltage and current. Further, they can keep the value of the resistance even when the applied voltage is changed and thus have nonvolatile memory. This differs from ohmic resistors, in which the value of the current depends on the instantaneous value of the applied voltage. Further, they are new circuit elements that cannot be replaced by a combination of capacitors, inductors, and ohmic resistors. Their main potential use is in resistive random access memory. An excellent review of memristive devices and their potential in bringing about new computing logic architecture is given in [31]. TMOs (TiOx, MoOx, WOx, VOx, etc.) exhibiting metal-to-insulator transition (MIT) have been extensively studied for potential use in memristive devices [31,32]. VO2 is a very promising material for use in memristive devices owing to fast switching, having an MIT temperature near ambient and the possibility of inducing said transition by diverse means such

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Figure 7.5 Interferograms measured from the generation of cylindrical wave fronts. The first two columns show (A, D) measurements with VO2 and (B, E) comparison with the results obtained with a commercial device in the visible range. The third column (C, F) shows detection of the far-field spot. (G) Simulation and (H) measurement for the far-field distribution obtained from the application of the spherical aberration illustrated in (I) [29].

as electric field, applied current, and photoexcitation. It has already been demonstrated in such systems [33,34]. An example of a VO2 insulator-tometal transition-based memristive device is shown in Fig. 7.6.

7.6 SENSING APPLICATIONS: GAS SENSORS For many environmental, ecological, industrial, and safety reasons, gas monitoring is of utmost importance. Gas sensing is yet another application in which chromogenic oxides such as WO3 that possess gasochromic properties can be useful.

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Figure 7.6 (A) Resistivityetemperature curves of a memristive device illustrating the hysteretic nature of the insulator-to-metal phase transition. The vertical dotted line shows the (applied) bias temperature. (B) Demonstration of information storage in a memristive vanadium dioxide film. Each 50-V pulse triggers the transition to a new resistivity level [33].

Tungsten trioxide has been extensively investigated for its gasochromic properties because of its high sensitivity and good stability in thin or thick film form, and therefore can be part of a solid-state sensor device. Incorporation of WO3 as a gas sensor can in fact be based on several effects. For example, it could be based on its change in conductivity, similar to the well-known effect in semiconductors, owing to the presence of oxidizing or reducing gases in the surrounding environment. Reducing gases (such as H2, H2S, NH3) increase surface conductivity by creating oxygen vacancies, which also result in the coloration of WO3. Oxidizing gases (such as O3, NO2, or CO2) reduce the number of such vacancies, resulting in more transparency and reduced conductivity [2,35,36]. A plot of impedance variation of a WO3 sensor with ozone concentration is given in Fig. 7.7. A gas sensor could utilize the change in resistivity or the change in the dielectric constants as the sensing mechanism, and both mechanisms are reversible. Keeping WO3 at high temperature improves sensitivity, response, and recovery time. It has been reported that a sensor based solely on absorption changes will not be sensitive enough [20]. However, it is possible to build a device combining WO3 and a metal such as Au or Pt to

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Figure 7.7 Impedance variation of a WO3 gasochromic sensor with O3 concentration [35].

create surface plasmon resonance (SPR) at a certain wavelength and monitor the shift in the SPR wavelength due to change in dielectric constants of WO2 film. Such sensors based on 40 nm Au or Pt/600 nm WO3/3 nm Pd have been reported. The Pd enhances the reaction rate through its catalytic action [20]. Gas sensors are especially vital for H2 sensing. In the reaction of WO3 with hydrogen gas, its spectrum develops an absorption band in the IR but not in the short wavelength range. Therefore one can use the IR portion as the signal and a short wavelength as reference [20]. Molybdenum oxide (MoO3) is another chromogenic oxide investigated for its gasochromic properties. It was found by Yao et al. that a-MoO3 films provide better response to H2 gas than b-MoO3 and that the gasochromic response significantly increased with porous films [37].

7.7 SENSING APPLICATIONS: BIOSENSORS A natural consequence of gasochromic sensors is their use in biological applications. A good example is given in [20] in which a H2 sensor is used for selecting the specific mutant green algae that produces H2 gas while tolerating high levels of O2 presence. This is an important step toward developing biohydrogen production under atmospheric levels of O2. In another study, Santos et al. [38] used WO3 nanoparticles to improve the electron transfer of cytochrome c nitrite reductase with nitrite biosensor electrodes. Compared to the bare ITO electrodes, the exchange rate of WO3/ITO electrodes increased by an order of magnitude.

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7.8 SECURITY The global cost of counterfeiting, including money, credit card fraud, and goods, was estimated as US$1.8 trillion in 2014 [39]. Of this, global credit card fraud was estimated at US$16.31 billion and estimated to reach US$35 billion by 2020 [40]. The US market has the largest loss due to credit card fraud, and a major portion of the fraud (37% in 2014) is due to counterfeiting [41]. Bank notes are, of course, always a major target for fraud. Even though the amount of counterfeit notes in circulation may seem small, e.g., US$220 million in the United Sates in 2014, they cost individuals and businesses money and erode confidence in the currency itself [42]. As a result, extensive measures to prevent fraud are being taken. Smithers Pira predicts that the value of the security printing market will reach US$36.6 billion in 2020 [43]. Interference security image structures (ISISs) are one of various security features in bank notes or credit cards to impede and discourage counterfeiting. They have iridescent colors resulting from interference filters, with the apparent color changing with the viewing angle. Because of this color change, they are harder to reproduce by photographic reproduction methods such as copying or scanning or printing. One way of adding another level of authentication to ISIS security features is to have a noniridescent color (or another filter) that matches only at a certain angle, e.g., 0 degree [44]. In the work of Baloukas et al. EC WO3 was used for the purpose of creating multilevel authentication [45]. In their work, WO3 and SiO2 films are sandwiched between an ITO electrode at the substrate interface and an Au electrode at the top surface. SiO2 acts as an electron barrier and WO3 is the coloration material (Fig. 7.8). As depicted in the figure, this is an active device in which WO3 can change color from transparent to blue upon the application of a small voltage. When WO3 is partially colored, there are two images visible, one of which disappears upon further coloration. This is due to metameric color matching with the background. Similarly, the other image disappears when the device is transparent. Further, the images are iridescent, meaning they change color with the viewing angle [45].

7.9 TUNABLE PHOTONIC CRYSTALS Chromogenic thin film photonic crystals were extensively discussed in Chapter 6. Here we highlight the emerging and promising applications of this technology.

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Figure 7.8 Conceptual example combining interference security image structures (ISISs) with an electrochromic device (ECD). ISIS A is metameric in the bleached state of the active device and is therefore invisible at normal incidence (dashed star). ISIS B is a metameric in the colored state of the ECD and consequently becomes invisible during coloration. When the whole device is tilted, both ISISs change color while the rest of the device essentially remains the same. The layered structure of our ECD is also shown [45]. ITO, indium tin oxide.

Photonic bandgap (PBG) materials are defined by the presence of a gap in their photonic band structure, similar to the presence of an electronic bandgap in semiconductor materials. Photonic crystals are periodic structures in which the lattice periodicity is comparable to a wavelength of light. Interestingly, the periodicity may not be strictly necessary and such gaps may exist in amorphous structures as well. A PBG could be one, two, or three dimensional, with the Bragg mirror being a good example of a 1-D photonic crystal. A 3-D PBG material has a complete photonic gap for a particular wavelength of light and hence inhibits propagation of that wavelength in its interior. Such materials with a complete gap do not exist in nature; however, incomplete photonic gap materials do exist. They are distinguished by the iridescent, structural light arising from interference. Examples of such naturally occurring structures in beetles, the feathers of pigeons, the eyes of the male peacock, and the wings of morpho butterflies are given by Wang and Zhang in [46].

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There has been intense interest in PBG structures since they were proposed by S. John [47] and E. Yablonovitch [48] in 1987. They bring about exciting new possibilities such as optical transistors and amplifiers, waveguides with the capability of diverting light at sharp corners, localized light due to defects or cavities, etc. The first 3-D PBG fabricated in a laboratory was based on a woodpile structure with a complete bandgap at microwave frequencies [49]. Theoretical work on periodic structures generally shows that the face-centered cubic and diamond structures are more amenable to the presence of a complete and relatively broad PBG [49]. It has also been shown that spiral structures such as square spirals can produce robust and broad bandgaps [50e52]. As mentioned, inverse opal structures based on silicon with the bandgap in the IR range (at communication wavelength) have been produced, and further shown to be capable of being tuned by being impregnated with liquid crystals [53,54]. Incorporating chromogenic oxides in the photonic crystals opens an exciting possibility of tuning the position and width of the bandgap, or switching it on and off by intercalation, photoexcitation, heating, or electric field. In the work of Kuai et al. [55] inverse opal WO3 films were fabricated on polystyrene (PS) colloidal crystal templates. WO3 was introduced into this template through a dip-infiltrating solegel technique, and subsequently dried at 80 C for half an hour. The sample was then sintered at 460 C to remove the PS template. The resulting macroporous WO3 periodic structures were measured and were determined to have a PBG at 532 nm (for starting PS spheres of 300 nm diameter) and 1302 nm (PS sphere of 760 nm diameter) with a bandwidth of 13%. More importantly, it was shown that the center of the gap shifts with lithium intercalation. Lithium in the amount of 0e105 nm was introduced into the sample by a dry lithiation method and caused the center wavelength of the gap to shift from 532 to 496 nm. This work clearly showed that reversible tuning of gap position is possible. Further work showed that the ordered structure results in improved EC performance of the macroporous tungsten oxide film in the IR region [56]. Further work in this area includes the fabrication of large-scale thermochromic vanadium oxide in a VO2eSiO2 opal structure with a peak near 1.55 mm (communication wavelength) [57]. VO2 inverse opal structures were also grown in the same work. It was shown that at transition temperature (68 C) the intensity of the reflection peak reduces drastically. Interference filters of WO3 were also fabricated by Baloukas et al. [58], and Kavan et al. reported on the EC properties of TiO2 anatase opals [59].

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These examples illustrate the vast potential of chromogenic oxides as building blocks or as incorporated in the structure of photonic crystals. The possibility of actively tuning the band position and the presence of the gap itself via chromogenic mechanisms opens a vista of possibilities.

7.10 SUMMARY In this chapter we have given examples of chromogenic devices. Some of these devices have already made substantial inroads into their respective markets (switchable mirrors, smart windows, smart glass), but they still have tremendous potential for growth. Other examples are exciting concepts of which prototypes have been demonstrated, but they are still many steps away from commercialization. Even for more mature technologies, cost is one of the major issues to be addressed and lowered. Chromogenic mirrors, glass, and displays are more complex and more costly to manufacture than their traditional counterparts and the cost per unit area reflects that. There are also issues regarding response time and durability, especially for smart windows. Perhaps that is why the smaller devices such as automobile windows, visors, and sunroofs are more attractive from a marketing point of view as they are generally expected to be replaced in less than 10e15 years. Yet another important task would be developing a good solid ion conductor, comparable to liquid electrolytes, for improvement of EC devices. Another area of improvement is developing the capability of making chromogenic devices on flexible substrates, which may lead to fabricating or even printing a complete system (including batteries, device, and sensors) on the same substrate. This ability also brings the possibility of making chromogenic devices on thin, flexible, low-cost substrates that can then be easily put on existing windows thus eliminating the need for costlier replacement. Another idea worth exploring is integrating more functions into the existing devices. An example would be using a smart window as a projection screen, or adding an information display to switchable mirrors. There has even been work in integrating active surfaces to artistic work as an extra dimension of exploration and expression [60]. The field of TMO chromogenics-based photonic crystals is in its infancy and provides a wide scope for future research and development. The wide range of TMOs with varying refractive index available to build these devices and the variability of crystal lattice parameters offer a wide range of wavelengths and tunable PBGs. However, the main challenge associated

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with this field is the limitation of techniques available to fabricate large-area 3-D periodic structures. Hence, the near-future application of chromogenic photonic crystal-based devices can be envisioned only in small-area applications. In conclusion, there are tremendous opportunities in the field of TMObased chromogenics for further research and development, exploring new venues, development of devices, and bringing them to fruition. The possibilities are many and exciting.

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[44] B. Baloukas, L. Martinu, Metameric interference security image structures, Applied Optics 47 (2008) 1585e1593. [45] B. Baloukas, J.-M. Lamarre, L. Martinu, Active metameric security devices using an electrochromic material, Applied Optics 50 (2011) C41eC49. [46] H. Wang, K.-Q. Zhang, Photonic crystal structures with tunable structure color as colorimetric sensors, Sensors 13 (2013) 4192e4213. [47] S. John, Strong localization of photons in certain disordered dielectric superlattices, Physical Review Letters 58 (23) (1987) 2486e2489. [48] E. Yablonovitch, Inhibited spontaneous emission in solid-state physics and electronics, Physical Review Letters 58 (20) (1987) 2059e2062. [49] J.D. Joannopoulos, S.G. Johnson, J.N. Winn, R.D. Meade, Photonic Crystals: Molding the Flow of Light, second ed., Princeton University Press, 2008. [50] O. Toader, S. John, Square spiral photonic crystals: robust architecture for microfabrication of materials with large three-dimensional photonic band gaps, Physical Review E 66 (2002), 016610. [51] S.R. Kennedy, et al., Optical properties of a three-dimensional silicon square spiral photonic crystal, Photonics and Nanostructures e Fundamentals and Applications 1 (2003) 37e42. [52] G. Beydaghyan, Photonic Crystals with Glancing Angle Deposition (thesis), Queen’s University, 2003. [53] A. Blanco, et al., Large-scale synthesis of a silicon photonic crystal with a complete three-dimensional bandgap near 1.5 micrometres, Nature 405 (2000) 437e440. [54] K. Busch, S. John, Liquid-crystal photonic-band-gap materials: the tunable electromagnetic vacuum, Physical Review Letters 83 (5) (1999) 967e970. [55] S.-L. Kuai, G. Bader, P.V. Ashrit, Tunable electrochromic photonic crystals, Applied Physics Letters 86 (22) (2005), 221110/1e3. [56] L. Yang, et al., Improved Electrochromic performance of ordered macroporous tungsten oxide films for IR Electrochromic device, Solar Energy Materials and Solar Cells 100 (2012) 251e257. [57] M. Ibisate, D. Golmayo, C. Lopez, Vanadium dioxide thermochromic opals grown by chemical vapour deposition, Journal of Optics: Pure and Applied Optics 10 (12) (2008), 125202/1e6. [58] B. Baloukas, J.-M. Lamarre, L. Martinu, Electrochromic interference filters fabricated from dense and porous tungsten oxide films, Solar Energy Materials and Solar Cells 95 (2011) 807e815. [59] L. Kavan, M. Zukalova, M. Kalbac, M. Graetzel, Lithium insertion into anatase inverse opal, Journal of the Electrochemical Society 151 (8) (2004) A1301eA1307. [60] M. Ferrara, M. Bengisu, Intelligent design with chromogenic materials, Journal of the International Color Association 13 (2014) 54e66.

INDEX ‘Note: Page numbers followed by “f” indicate figures, “t” indicate tables.’

A Amorphous tungsten trioxide film, 27e29 IVCT model, 33e34 optical density spectra, 34, 35f polaron model, 33e34, 33f Antireflection (AR) coatings, 195e196 Atmospheric pressure and aerosol-assisted chemical vapor deposition (AP-AACVD) method, 191e193 Atomic force microscopy (AFM) images, 118, 119f, 175e176 Autocloning, 295e296

B Biosensors, 340 Bloch functions, 293e294 Bragg reflection peak, 321e322 Brillouin zones, 306e307, 307f

C

Calcium fluoride (CaF), 58e60 Chemical vapor deposition (CVD), 64e65 Chromogenic effects chromogenic mirrors, 334e335 infrared laser wave-front mapping, 337 memristive devices, 337e338 satellite windows, 335e337 security, 341 sensing applications biosensors, 340 gas sensors, 338e340 smart windows, 331e334 tunable photonic crystals, 341e344 Chromogenic mirrors, 334e335 Chromogenics effects chromogenic mirrors, 334e335

infrared laser wave-front mapping, 337 memristive devices, 337e338 satellite windows, 335e337 security, 341 sensing applications, 338e340 smart windows, 331e334 tunable photonic crystals, 341e344 materials chemichromic, 7e8 electrochromic, 1e2 magnetochromic, 6e7 photochromic, 2e4 piezochromic, 5e6, 6f thermochromic, 4e5 TMO, 8e9 origin, 1 physical properties, 1 Chronoamperometry (CA), 78e80 Chronopotentiometry (CP), 78e80 Coarse grain thin films, 8e9, 47e49 Colloidal crystals, 294e295 Coloration efficiency (CE), 315e316 Colorations, 248, 250e251, 271e273 Coulomb repulsion, 16 Counterelectrode layer, 77 Cyclic voltammetry (CV), 78e80

D 1,2-Dimethyl-3-propylimidazolium iodide (DMPII), 133e134 Dip-coating method, 63 Double-insertion methods characterization techniques, 87e88 dry lithiation method advantages of, 82e84 LiNbO3, 82 lithium atoms, source of, 81e82 electrochemical method CA, 78e80

349

350

Index

Double-insertion methods (Continued) ClO4 /AsF6 ions, 78e80 CV technique, 80e81 galvanostatic measurements, 80e81 I V cycles, 80e81 optical characterization atmospheric absorption bands, 84e86 optical ellipsometry, 87 photopic transmittance, 86e87 Planck’s radiation law, 84 solar radiation, 84e86 spectral transmittance, 86e87 transparent windows, 84 Drude theory, 30e31 Dry lithiation method advantages of, 82e84 LiNbO3, 82 lithium atoms, source of, 81e82

E Effective medium approximation (EMA), 158e159 Electrochemical method CA, 78e80 ClO4 /AsF6 ions, 78e80 CV technique, 80e81 galvanostatic measurements, 80e81 I V cycles, 80e81 Electrochromic (EC) anodic TMOs, 73e74 aspects, 74 cathodic TMOs, 73 coloration, 1e2 device configuration counterelectrode layer, 77 EC layer, 75e78 IC layer, 77 TC layers, 78 ECD, nanostructured thin films EC layer, 127 gel electrolytes, 142e146 half cells, 128e130 IC layer, 128e130 IS layer, 127e128 liquid electrolytes, 130e133 solid-state, 133e142 TC layers, 128e130

electric field, 73 electrochromically tunable photonic crystals CE spectra, 315e316, 315f dry lithiation process, 305 electrical and optical property, 316e318 high-efficiency coloration, 303 LiClO4 solution, 310 measurements, types of, 310 PBG, 304e305 polystyrene gel composites, 303e304 PS crystal template, 306e307, 313e315 scanning electron microscopy images, 303e304 WO3, band structure calculations, 307e310, 309f X-ray diffraction analysis, 304e305 ITO, 74e75 materials and devices characterization techniques, 87e88 double-insertion methods, 78e84 optical characterization, 84e87 nanostructured TMO thin films, properties of MoO3, 117e127 V2O5, 108e117 WO3, 88e108 Electrodeposition/electrophoretic deposition systems, 1e2 Electromagnetorheological (EMR), 299e300 Electron orbital, 16 Electrostatic discharge (ESD), 224e225 Ellipsometry, 254

F Fast atom beam (FAB), 136e138 Fermi level, 27e29 Ferroelasticity, 6e7 Ferroelectricity, 6e7 Ferromagnetism, 6e7 First-order reversal curve (FORC), 170e172, 171f Fluorine-doped tin oxide (FTO), 122e123

Index

G Galvanostatic measurements, 78e80 Gas sensors, 338e340 Gasochromic material, 7e8 Gasochromic sensors, 340 Gel electrolytes, electrochromic devices with NiO film, 144 PEGMA, 142e144 PMMA, 144e146 sandwich structure, 142e144 Glancing angle deposition (GLAD) method, 57e58, 259

H Halochromic, material, 7e8 Heterostructure-based films individual films, 281 In-doped CdS film, 279e281 MoO3/CdS double layer, 274e276 optical density spectra, 278e279, 279f Smakula equation, 274e276 Urbach’s rule, 278 Hydrogen gas, 8 Hygro/hydrochromics, 7e8

I Indium tin oxide (ITO), 74e75, 262 e263, 310 Infrared laser wave-front mapping, 337 Interference security image structures (ISISs), 341 Intervalence charge transfer (IVCT), 29 e30, 248 Ion storage (IS) layer, 75, 77 Ionochromic material, 7e8

K Kirchhoff’s law, 223

L Liquid electrolytes, electrochromic devices with molecular monolayer, 130e131 PEC devices, 132e133 TiO2 and WO3 film, 130e131 Lithium, 343

351

Lithium borate (LiBO2), 136 Lithium niobate (LiNbO3), 82 Lithium perchlorate (LiClO4), 144 Low-pressure CVD (LPCVD), 64e65

M Magnéli phases, 26e27 Magnesium fluoride (MgF2), 58e60 Magnetochromic materials, 6e7 Maxwell Garnett theory, 207 Maxwell’s equations, 293e294 Mechanochromic. See Piezochromic materials Memristive devices, 337e338 Mesoporous thin films, 93e94 Metal-to-insulator transition (MIT), 19e20, 154e155, 337e338 Mirrors, 334e335 Molybdenum oxide (MoO3), 340 Molybdenum trioxide (MoO3), 24e25, 146e147 nanostructured TMO thin films, electrochromic properties of AFM images, 118, 119f angle-deposited, 120e122 dry lithiation method, 120e122 electrodeposited, evolution of, 123e124 GLAD method, 120e122 mixed-phase, 124 MoOxCy, 126e127 orthorhombic phase, 118 photochromics DABS values, 263 AFM images, 259 Eg value, 259e260 electric field, 263 ellipsometric studies, 259 GLAD method, 259e260 metal film/particles, 261e262 metal nanoclusters, 273e274 Mo-doped, 282e287 physisorbed water, 258e259 Schottky barrier, 263, 265f semiconductor bandgap of, 258e259 visible light photons, 266

352

Index

Molybdenum trioxide (MoO3) (Continued) TMO chromogenic properties corner-sharing oxygen atoms, 46 crystal structure, 46 octahedral units, 46 polymorphic structures, 43e46 Monoxides (MO), 20 Mott transitions, 16, 19e20 MotteHubbard splitting, 19e20

N Nanostructured (NS), 8e9 Nanothermochromics, 205e207 Nickel oxide (NiO) layer, 18, 127e128 Niobium (Nb), 181e184 Normal film, 47e49

O Oblique angle deposition (OAD), 57e58 Octahedral unit, 20 MoO6, 46 WO3 electrochromism, 26e27 Optical characterization atmospheric absorption bands, 84e86 optical ellipsometry, 87 photopic transmittance, 86e87 Planck’s radiation law, 84 solar radiation, 84e86 spectral transmittance, 86e87 transparent windows, 84 Organo-molybdenum oxides (MoOxCy), 126e127 Orthorhombic structure, 23e24, 23f Oxidizing gases, 339e340

P PEG methyl ether methacrylate (PEGMA), 142e144 Peierls model, 40 Perovskite structure, 25e26 Peroxoisopolymolybdenum acid (Mo-IPA), 283e284 Peroxoisopolytungstic acid (W-IPA), 283e284 Phonon density of states (PDOS), 160e161

Photochromics characterization methods chemical composition measurements, 254e256 electrical measurements, 256 optical measurements, 252e254 colorations, 2e4, 248, 250e251, 271e273 electrochromic coloration, 250e251, 271e273 electromagnetic radiation, 247e248 electron hole pairs, 251 Gaussian-shaped absorption curve, 248e250 heterostructure-based films individual films, 281 In-doped CdS film, 279e281 MoO3/CdS double layer, 274e276 optical density spectra, 278e279, 279f Smakula equation, 274e276 Urbach’s rule, 278 molybdenum trioxide DABS values, 263 AFM images, 259 Eg value, 259e260 electric field, 263 ellipsometric studies, 259 GLAD method, 259e260 metal film/particles, 261e262 metal nanoclusters, 273e274 Mo-doped, 282e287 physisorbed water, 258e259 Schottky barrier, 263, 265f semiconductor bandgap of, 258e259 visible light photons, 266 organic and inorganic materials, 247e248 photochromically tunable photonic crystals, 323e325 polaron absorption model, 251e252 Smakula’s equation, 248e250 tungsten trioxide DABS, 269 metal nanoclusters, 273e274 oxalic acid concentration, 266, 268 undoped WO3 composite, 283e287 visible light, 269e271

Index

Photoelectrochromic (PEC) devices, 132e133 Photoelectrochromic photonic crystal (PEC-PC) device, 324e325 Photonic bandgaps (PBGs), 60e61, 102e104, 342 Photonic crystals, 60e61, 65e66, 88 advantages, 291e292 band structure calculation, 293e294 Bloch functions, 293e294 Bragg mirror, 289 energies, electrons with, 292 general electromagnetic approach, 293e294 heterostructures, 294, 296e298 high- and low-refractive-index media, 289 Maxwell’s equations, 293e294 PBG, 289e291 Schrödinger’s equations, 293e294 self-assembly method, 294e295 transfer matrix methods, 294 tunable electrochromically tunable, 303e318 lattice constant (periodicity), 299 parameters of, 298e299 photochromically tunable, 323e325 refractive index (dielectric constant), 300e301 thermochromically tunable, 318e323 two- and three-dimensional, 295e296 Photovoltachromic cell (PVCC), 133e134 Photovoltaic (PV) devices, 128e130, 332 Physical vapor deposition (PVD) methods, 47, 100e102 Piezochromic materials, 5e6 Planck’s radiation law, 84, 223 Polaron model, 33e34, 33f Poly(ethylene terephthalate) (PET), 126e127 Poly(N-isopropylacrylamide) (PNIPAm), 216e217 Polycarbonate, 144 Polyethylene glycol (PEG), 97e100

353

Polymethylmethacrylate (PMMA), 144, 299 Polymorphic structures, 43e46 Polystyrene (PS), 306e307, 343 Potassium nitrate (KNO3), 269e271 Preisach model, 170e171

R Raman spectroscopy, 254e256 Rapid thermal annealing and cooling (RTAC), 176e177 Reversible optical change, 1 Rhenium cations, 24e25 Rock salt structure, 20, 21f Rutile form, 21e22, 21f

S Satellite windows, 335e337 Scanning electron microscopy (SEM) images, 313e315 Schottky barrier, 263, 265f Schrödinger’s equations, 293e294 Security, 341 Self-assembly method, 294e295 Silicon-based polymers, 4e5 Smart radiation device (SRD), 336e337 Smart windows, 190e191, 331e334 Sol gel method, 63 WO3 thin films, preparation of, 282e283 Solar energy applications AP-AACVD method, 191e193 AR coating, 195e198 electrochemical reactions, 201e202 five-layer structure, 199e201, 200fe201f hightemperature (metallic) and low-temperature (dielectric) states, 193e195 hybrid hydrogel, 219e220, 219t IR region, 190e191 laminate structures, 216e217 luminous and solar transmittances, 207e211 Maxwell Garnett theory, 207 Mg doping, 202e204 nanocomposite, 211, 212f, 213e216

354

Index

Solar energy applications (Continued) nanothermochromics, 205e207 PNIPAm, 216e219 semiconducting state, 211e213 smart window application, 190e191 spectral transmittances, 207e211 TOAB, 191e193 Solid-state electrochromic devices coloration and bleaching cycles, 138e140 DMPII, 133e134 dry lithiation method, 134e135 integrated transmittance values, 140 IR transmittance, 141e142 LiBO2, 136 MgF2 layer, 140 solar and photopic transmittance values, 136e138 Solvatochromic materials, 7e8 Spectrophotometry, 252e253 Spin-coating method, 63 Stoichiometry, 22e25 Surface plasmon resonance (SPR), 339e340

T Tantalum pentoxide (Ta2O5), 141e142 Tetraacetylammonium bromide (TOAB), 191e193 Tetrahedral unit, 22 Thermochromic materials, 4e5 Thermochromic smart radiator devices (SRDs) design considerations, 222e225 atomic oxygen, 224e225 IR emittance, 223e224 Kirchhoff’s law, 223 Planck’s law, 223 Surface ESD, 224e225 Wien’s law, 223 emissivity, 221e222 experimental results dielectric SiO2 layer, 230e231 emissivity values, 227e230 emittance spectrum, 230e231 Gaussian curve, 230e231 optical constants (n, k), 227e230

SEM micrograph, 232e233, 232f thin silver film, 232e233 static surfaces, 220e221 sun-facing position, 220e221 sun-obscuring (earth shade) positions, 220e221 terrestrial and extraterrestrial conditions, 220e221 thin film-based SRD technologies, 221e222 Thermochromic (TC) thermochromically tunable photonic crystals Bragg reflection peak, 321e322 differential reflectance, 320e321 PBG, 318 reflectance spectra, 320e321, 320f semiconductor-to-metal transition, 318 silica spheres, 319e320 transition temperature, 320e321 VO2 doping of, 181e189 free electron density, 156 hysteresis curves, 166e167 MIT, 154e156, 160e161 monoclinic M1, 155e156 monoclinic M2, 155e156 Mott model, 156e157 nano- and microcrystalline nature of, 163 nanostructured vo2 thin films and applications, 189e238 PDOS, 160e161 percolation transition, 158e159 phase transition, historical development of, 154e155, 155t phonon contribution, 161e163 resistance hysteresis, 166e167, 167f stoichiometric, behavior of, 164 tetragonal phase, 155e156 thermal hysteresis in, 168e181 thermochromic characteristics of, 163 transmittance and reflectance, 165e166 XRD patterns, 167e168

Index

Thin film nanostructure advanced materials, 46e47 applications of, 47 bulk (3-D) material, 47e49, 48f chemical methods, 47 CVD, 64e65 1-D size effect, 49 ECD EC layer, 127 gel electrolytes, 142 half cells, 128e130 IC layer, 128e130 IS layer, 127e128 liquid electrolytes, 130e133 solid-state, 133e142 TC layers, 128e130 high-pressure evaporation/sublimation and condensation higher substrate temperature, 53 interparticle collisions, 53e55 mean free path, 53e55 sourceesubstrate distance, 55e57 ultrahigh-vacuum conditions, 55 light interaction, 50e51 oblique/glancing angle deposition columnar films, 57e58 film deposition, 58e60, 60f nanosculpted, 58e60 nonperiodic and periodic nanostructures, 62 OAD, 58e60 photonic crystals, 60e61 tangent rule, 58 physical methods, 47 PVD methods, 47 solegel method, 63 thickest copper film, 51e52 Titanium dioxide (TiO2), 8e9 Transition elements, 14e15 Transition metal oxides (TMOs), 8e9, 301e303, 331 chromogenic properties molybdenum trioxide, 43e46 tungsten trioxide, 26e35 vanadium dioxide, 35e42 vanadium pentoxide, 42e43

355

crystal structure dioxides, 21e22 M2O5, 23e24 MO3, 24e25 monoxides, 20 perovskite structure, 25e26 sesquioxidesecorundum type, 22 spinelseM3O4, 22e23 EC. See Electrochromic (EC) electronic structure charge transfer insulator, 17e18, 17f charge transfer regime, 17e18 chemical classification scheme, 14e15 configuration, 15e16 Coulomb repulsion, 16 covalent bond formation, 14e15 Hubbard bands, 19e20 metallic behavior, 18 Mott transitions, 16, 19e20 MotteHubbard insulator, 17e18, 17f MotteHubbard splitting, 19e20 oxygen atom, 14e16 properties, 14 transition elements, 14e15 vanadium atom (VO2), 15e16 photochromic effect. See Photochromics photonic crystals. See Photonic crystals TC. See Thermochromic (TC) thin film nanostructuring. See Thin film nanostructure Transparent conducting (TC) layer, 78 Tunable photonic crystals chromogenic effects, 341e344 EC CE spectra, 315e316, 315f dry lithiation process, 305 electrical and optical property, 316e318 high-efficiency coloration, 303 LiClO4 solution, 310 measurements, types of, 310 PBG, 304e305 polystyrene gel composites, 303e304

356

Index

Tunable photonic crystals (Continued) PS crystal template, 306e307, 313e315 scanning electron microscopy images, 303e304 WO3, band structure calculations, 307e310, 309f X-ray diffraction analysis, 304e305 photochromics, 323e325 TC Bragg reflection peak, 321e322 differential reflectance, 320e321 PBG, 318 reflectance spectra, 320e321, 320f semiconductor-to-metal transition, 318 silica spheres, 319e320 transition temperature, 320e321 Tungsten trioxide (WO3), 146e147 nanostructured TMO thin films, electrochromic properties of amorphous WO3 films, 89e90 annealing temperature, 100 CE, 100e102 coarse-grained, 93e94, 102e104 colloidal solution approach, 97e100 crystallineWO3 films, 89e90 EC-related work, 88e89 electrodeposition method, 93e95, 100 electronic conductivity, 107e108 GLAD method, 102e104 large-area tungsten oxide dihydrate sheets, 97 mesoporous, 97e100 mixed conductivity, 107e108 PBG, 102e104 PEG, 97e100 polycrystalline, 90e91 porosity of, 100e102 PVD methods, 102e104 scanning electron micrograph, 93, 94f sol gel method, 97 solar and photopic transmittance values, 90e91, 93t source substrate distance, 104e107 transmittance spectra, 97

photochromics DABS, 269 metal nanoclusters, 273e274 oxalic acid concentration, 266, 268 undoped WO3 composite, 283e287 visible light, 269e271 TMO chromogenic properties amorphous, 27e29, 32e35 crystalline, 30e32 electrochromic coloration, 29e30 Magnéli phases, 26e27 octahedral WO6, network of, 26e27 valence and conduction bands, 27e29

U UV irradiation, 254e256, 273e274

V Vanadium atom, 15e16 Vanadium dioxide (VO2), 4e5 doping of Mo doping, 181e184, 183f Nb-doped, 181e184, 182f tungsten doping, 184e186, 184t free electron density, 156 hysteresis curves, 166e167 MIT, 154e156, 160e161 monoclinic structure M1 phase, 37e40, 38f, 155e156 M2 phase, 39, 39f, 155e156 Mott model, 156e157 nano- and microcrystalline nature of, 163 nanostructured VO2 thin films and applications Ge-coated Kapton sheet, 236e238 Hadamard masks, 236e238 high-resolution image, 236 mid-IR, 233e235 near-IR, 233e235 solar energy applications, 190e220 TC behavior of, 189e190 thermochromic smart radiator devices, 220e233 threshold power, 236

Index

PDOS, 160e161 percolation transition, 158e159 phase transition electroneelectron interaction, 40 historical development of, 154e155, 155t MotteHubbard model, 40 phonon contribution, 161e163 resistance hysteresis, 166e167, 167f stoichiometric, behavior of, 164 tetragonal phase, 155e156 structure, 36e37 thermal hysteresis in AFM image, 175e176, 175fe176f aspects of, 168e169 cooling cycle, 172 dielectric (semiconductor) domains, 170 EMA, 171e172 ferromagnetic materials, 169e170 FORC diagram method, 170e172, 171f interdomain interaction, 172 magnetic inertia, 169e170 metallic domains, 170 metal-to-insulator/semiconductor transition of, 176e177 nanocrystalline, scanning electron micrograph of, 173e174, 173f

357

optical and electrical properties, 181 polycrystalline films, 177 quantitative analysis of, 169e170 R M2 and M2 M1 transitions, 180 RTAC method, 176e177 thermochromic characteristics of, 163 transmittance and reflectance, 165e166 XRD patterns, 167e168 Vanadium pentoxide (V2O5), 23e24 nanostructured TMO thin films, electrochromic properties of anodically coloring, 110, 115e117 cathodically coloring, 110, 115e117 CV cycles, 113e115 deposition conditions, 111e113 lithium intercalation, 111e115 optical and structural properties, 115e117 phase irreversibility, 110e111 Raman profile of, 115e117, 116f TMO chromogenic properties, 42e43

W Wien’s law, 223

X X-ray diffraction (XRD), 167e168 X-ray photoelectron spectroscopy (XPS), 254e256