Green Energy Materials Handbook [1 ed.] 9781138605916, 9780429466281, 9780429881169, 9780429881152, 9780429881176

Green Energy Materials Handbook gives a systematic review of the development of reliable, low-cost, and high-performance

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Green Energy Materials Handbook [1 ed.]
 9781138605916, 9780429466281, 9780429881169, 9780429881152, 9780429881176

Table of contents :

Introduction


Molecular effects of functional polymer binders on Li+ transport on the cathode surface within lithium ion battery


2.1 Introduction


2.2 Molecular dynamics simulation details


2.3 Results and discussion


2.4 Summary and future perspectives



Essential properties of Li/Li+ graphite intercalation compounds


3.1 Introduction


3.2 The theoretical model


3.3 Rich geometric structures of graphites and graphite intercalation compounds


3.4 Unusual band structures of graphite-related systems


3.5 van Hove singularities in density of states


3.6 Chemical bondings and charge distributions


3.7 Summary



Defective and amorphous graphene as anode materials for Li-ion batteries: a first-principles study


4.1 Introduction


4.2 Computational methods


4.3 Results and discussions


4.4 Conclusion



Rich Essential Properties of Si-Doped Graphene


5.1 Introduction


5.2 Computational methods


5.3 Geometric structures of Si-adsorbed and Si-substituted graphene


5.4 Rich electronic structures


5.5 Spatial charge densities


5.6 The diverse density of states


5.7 Summary



Diversified essential properties in transition metals adsorbed Graphene


6.1 Introduction


6.2 The theoretical model


6.3 Results and discussions


6.4 Summary



Combining neural network with first-principles calculations for computational screening of electrolyte additives in lithium ion batteries


7.1 Introduction


7.2 Materials and methods


7.3 Results and disscussions


7.4 Conclusion



Metal oxide-reduced graphene oxide (MO-RGO) nanocomposite as high performance anode materials in Lithium ion batteries


8.1 Introduction


8.2 Potential binary metal oxides asanode materials in LIBs


8.3 Complex metal oxides as anode materials in LIBs


8.4 Metal oxide-graphene/reduced graphene oxide nanocomposite as anode materials in LIBs


8.5 Our research contribution toward LIB


8.6 Conclusions



In-situ X-ray and Neutron Analysis Techniques on Lithium/Sodium ion batteries


9.1 Introduction


9.2 Methodology for in-situ X-ray and neutron scattering experiments


9.3 In-situ X-ray analysis on synergistic effects of Si anode materials


9.4 In-operando X-ray diffraction - a quantitative analysis on Si-graphite negative electrode


9.5 In-situ X-ray diffraction analysis of lithiation-induced crystal restructuring of Sn/TiO2 nanocrystallites


9.6 In-operando neutron diffraction analysis on low temperature lithium diffusion behaviors in 18650 Li-ion battery


9.7 In-operando neutron diffraction Studies on P2-Na2/3Fe1/3Mn2/3O2 cathode in a sodium ion battery


9.8 Summary



Micro-Phase Separated poly(VdF-co-HFP)/Ionic Liquid/Carbonate as Gel Polymer Electrolytes for Lithium-Ion Batteries


10.1 Introduction


10.2 Experimental


10.3 Results and discussion


10.4 Conclusion



Gel and solid electrolytes for Lithium ion batteries


11.1 Introduction


11.2 Solid-state electrolytes (SSEs)


11.3 Gel Polymer Electrolytes (GPEs)


11.4 Summary



Silicon-Nanowire Based Hybrid Solar Cells


12.1 Introduction


12.2 Silicon nanowires fabrication


12.3 PEDOT: PSS polymer as the p-type layer of hybrid solar cell application


12.4 Silicon Nanowire based Hybrid Solar Cells


12.5 Conclusion



Characterization and Performance of Li-ZnO Nanofiber and Nanoforest Photoanodes for Dye-sensitized Solar Cell


13.1 Introduction


13.2 Experimental


13.3 Results and discussion


13.4 Conclusion



Review of monolithic dye-sensitized solar cells and perovskite solar cells


14.1 Introduction


14.2 Monolithic dye-sensitized solar cells





Mesoporous electrode for monolithic perovskite solar cells



Conclusion

15. High-Performance Quasi-Solid-State Polymer Electrolytes for Dye-Sensitized Solar Cell Applications


16. Concluding Remarks


17. Perspective on Battery Research


Index

Citation preview

Green Energy Materials Handbook

Green Energy Materials Handbook

Edited by

Ming-Fa Lin and Wen-Dung Hsu

MATLAB® and Simulink® are trademarks of the MathWorks, Inc. and are used with permission. The MathWorks does not warrant the accuracy of the text or exercises in this book. This book’s use or discussion of MATLAB® and Simulink® software or related products does not constitute endorsement or sponsorship by the MathWorks of a particular pedagogical approach or particular use of the MATLAB® and Simulink® software.

CRC Press Taylor & Francis Group 6000 Broken Sound Parkway NW, Suite 300 Boca Raton, FL 33487-2742 © 2019 by Taylor & Francis Group, LLC CRC Press is an imprint of Taylor & Francis Group, an Informa business No claim to original U.S. Government works Printed on acid-free paper International Standard Book Number-13: 978-1-138-60591-6 (Hardback) This book contains information obtained from authentic and highly regarded sources. Reasonable efforts have been made to publish reliable data and information, but the author and publisher cannot assume responsibility for the validity of all materials or the consequences of their use. The authors and publishers have attempted to trace the copyright holders of all material reproduced in this publication and apologize to copyright holders if permission to publish in this form has not been obtained. If any copyright material has not been acknowledged, please write and let us know so we may rectify in any future reprint. Except as permitted under U.S. Copyright Law, no part of this book may be reprinted, reproduced, transmitted, or utilized in any form by any electronic, mechanical, or other means, now known or hereafter invented, including photocopying, microfilming, and recording, or in any information storage or retrieval system, without written permission from the publishers. For permission to photocopy or use material electronically from this work, please access www.copyright.com (http://www.copyright.com/) or contact the Copyright Clearance Center, Inc. (CCC), 222 Rosewood Drive, Danvers, MA 01923, 978-750-8400. CCC is a not-for-profit organization that provides licenses and registration for a variety of users. For organizations that have been granted a photocopy license by the CCC, a separate system of payment has been arranged. Trademark Notice: Product or corporate names may be trademarks or registered trademarks, and are used only for identification and explanation without intent to infringe. Library of Congress Cataloging‑in‑Publication Data Names: Lin, Ming-Fa, author. | Hsu, Wen-Dung, author. Title: Green energy materials handbook / Ming-Fa Lin and Wen-Dung Hsu. Description: Boca Raton : Taylor & Francis, a CRC title, part of the Taylor & Francis imprint, a member of the Taylor & Francis Group, the academic division of T&F Informa, plc, [2019] | Includes bibliographical references and index. Identifiers: LCCN 2019002151| ISBN 9781138605916 (hardback : acid-free paper) | ISBN 9780429466281 (ebook) Subjects: LCSH: Electric batteries--Materials. | Solar cells--Materials. | Green chemistry. Classification: LCC TK2896 .L56 2019 | DDC 621.3028/6--dc23 LC record available at https://lccn.loc.gov/2019002151 Visit the Taylor & Francis Web site at http://www.taylorandfrancis.com and the CRC Press Web site at http://www.crcpress.com

Contents Preface......................................................................................................................vii Acknowledgments......................................................................................................ix Editors........................................................................................................................xi Contributors............................................................................................................ xiii Chapter 1 Introduction...........................................................................................1 Jow-Lay Huang, Chi-Cheng Chiu, Shih-Yang Lin, Chin-Lung Kuo, Duy Khanh Nguyen, Ngoc Thanh Thuy Tran, Wen-Dung Hsu, Chia-Chin Chang, Jeng-Shiung Jan, Hsisheng Teng, ChiaYun Chen, I-Ming Hung, Peter Chen, Yuh-Lang Lee, and Ming-Fa Lin Chapter 2 Molecular Effects of Functional Polymer Binders on Li+ Transport on the Cathode Surface within Lithium-Ion Batteries....... 19 Kun-You Chen and Chi-Cheng Chiu Chapter 3 Essential Properties of Li/Li+ Graphite-Intercalation Compounds..... 37 Shih-Yang Lin, Wei-Bang Li, Ngoc Thanh Thuy Tran, WenDung Hsu, Hsin-Yi Liu, and Ming-Fa Lin Chapter 4 Defective and Amorphous Graphene as Anode Materials for Li-Ion Batteries: A First-Principles Study........................................... 65 Yu-Jen Tsai and Chin-Lung Kuo Chapter 5 Rich Essential Properties of Si-Doped Graphene............................... 83 Duy Khanh Nguyen, Shih-Yang Lin, Ngoc Thanh Thuy Tran, Hsin-Yi Liu, and Ming-Fa Lin Chapter 6 Diversified Essential Properties in Transition Metal–Adsorbed Graphene........................................................................................... 109 Ngoc Thanh Thuy Tran, Chun-Hsien Kuo, Hai Duong Pham, Shih-Kang Lin, and Ming-Fa Lin Chapter 7 Combining Neural Network with First-Principles Calculations for Computational Screening of Electrolyte Additives in Lithium-Ion Batteries........................................................................ 125 Chia-Jung Lee, Ngoc Thanh Thuy Tran, Chih-Ao Liao, MingHsiu Wu, and Wen-Dung Hsu v

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Contents

Chapter 8 Metal Oxide–Reduced Graphene Oxide (MO–RGO) Nanocomposites as High-Performance Anode Materials in Lithium-Ion Batteries........................................................................ 145 Sanjaya Brahma, Shao-Chieh Weng, and Jow-Lay Huang Chapter 9 In-Situ X-Ray and Neutron Analysis Techniques on Lithium/ Sodium-Ion Batteries......................................................................... 165 Lakshmanan Saravanan, Tsan-Yao Chen, Chih-Wei Hu, Hung-Yuan Chen, Yu-Fan Su, Jow-Lay Huang, Chun-Ming Wu, Ping-Ching Wu, Yung-Der Juang, and Chia-Chin Chang Chapter 10 Micro-Phase Separated Poly(VdF-co-HFP)/Ionic Liquid/ Carbonate as Gel Polymer Electrolytes for Lithium-Ion Batteries......... 197 Jeng-Shiung Jan, Guan-Ying Fu, and Yu-Chao Tseng Chapter 11 Gel and Solid-State Electrolytes for Lithium-Ion Batteries.............. 215 Ramesh Subramani and Hsisheng Teng Chapter 12 Silicon-Nanowire-Based Hybrid Solar Cells.................................... 235 Ilham Ramadhan Putra, Pawan Kumar Singh, and Chia-Yun Chen Chapter 13 Characterization and Performance of Li-ZnO Nanofiber and Nanoforest Photoanodes for Dye-Sensitized Solar Cells.................. 253 I-Ming Hung, Jing-Ru Chen, and Yi-Hung Wang Chapter 14 Monolithic Dye-Sensitized and Perovskite Solar Cells.................... 269 Ming-Hsien Li, Kuan-Yu Lin, and Peter Chen Chapter 15 High-Performance Quasi-Solid-State Polymer Electrolytes for Dye-Sensitized Solar Cell Applications............................................ 281 Shanmuganathan Venketasan and Yuh-Lang Lee Chapter 16 Concluding Remarks......................................................................... 331 Jow-Lay Huang,Chi-Cheng Chiu, Shih-Yang Lin, Chin-Lung Kuo, Duy Khanh Nguyen, Ngoc Thanh Thuy Tran, Wen-Dung Hsu, Chia-Chin Chang, Jeng-Shiung Jan, Hsisheng Teng, ChiaYun Chen, I-Ming Hung, Peter Chen, Yuh-Lang Lee, and Ming-Fa Lin Chapter 17 Perspective on Battery Research....................................................... 341 Ralph Nicolai Nasara, Chai-Hao Tu, and Shih-Kang Lin Index....................................................................................................................... 357

Preface Environment and energy are always key issues related to the standard of living of human beings. A traditional energy resource, fossil fuel, is cheap but emits a lot of greenhouse gases, which pollute the environment. Thus, developing new technology on green energy is necessary more than ever. In order to compete with the cost of fossil fuels, the new materials need to have high efficiency, have a long cycle life, and be capable of being recycled and reactivated. To fulfill the goal, we need to have a deep understanding of how the materials degrade while operating. We also need to know how to recycle the device, separate it, and recover the materials. Both computational designs and experimental fabrications and characterizations are important to promote the materials into practical applications. This book aims to provide a systematic review of the development of reliable, low-cost, and high-performance green energy materials, covering both up-to-date mainstream computational and experimental studies. This work presents reliable and complete experimental measurements and computational results as well as potential applications. Among green technologies, electrochemical and energy storage technologies are considered the most practicable, environmentally friendly, and workable to make full use of renewable energy sources. On the other hand, computational materials science is capable of studying the thermodynamics and kinetics of the reactions that occur during operation and, thus, become an important tool for new materials design for green energy. It can be divided into four different methods: first-principle methods, atomic-level methods, meso-scale methods, and continuum methods. Details on how to utilize different computational approaches on materials design for new green energy materials will be discussed. The systematic studies proposed in this book can greatly promote the basic and applied sciences. They could attract much attention from researchers in the scientific community, not only for the study of green energy materials but also for the exploration of other emergent systems. This book contains 17 chapters. The content covers significant theoretical predictions and experimental works. For lithium-ion batteries (LIBs), the molecular dynamics simulations, first-principles method, and the direct combination of the neural network and first-principles calculations are, respectively, suitable for thoroughly exploring molecular effects of functional polymer binders on the cathode surface, such as the adatom chemisorptions, atom/ion intercalations, local defects and amorphous structures in graphene-based anode materials, and the essential electronic properties of electrolyte additives. The experimental examinations are conducted on the high-performance anode materials of metal O–reduced graphene oxide nanocomposites in LIBs, in situ X-ray and neutron inspections on LIBs, micro-phase separated poly(VdF-co-HFP)/ionic liquid/carbonate as gel polymer electrolytes for LIBs, gel and solid electrolytes for LIBs, silicon-nanowire-based hybrid solar cells, Li-ZnO nanofiber and nanoforest photoanodes for dye-sensitized solar cells,

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Preface

monolithic dye-sensitized solar cells and perovskite solar cells, and quasi-solid-state dye-sensitized solar cells. Certain significant issues are proposed to be worthy of near-future systematic studies. MATLAB® is a registered trademark of The MathWorks, Inc. For product information, please contact: The MathWorks, Inc. 3 Apple Hill Drive Natick, MA 01760-2098 USA Tel: 508 647 7000 Fax: 508-647-7001 E-mail: [email protected] Web: www.mathworks.com

Acknowledgments This work was financially supported by the Hierarchical Green-Energy Materials (Hi-GEM) Research Center, from the Featured Areas Research Center Program within the framework of the Higher Education Sprout Project by the Ministry of Education (MOE) and the Ministry of Science and Technology (MOST 107-3017-F006-003) in Taiwan.

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Editors Ming-Fa Lin is a distinguished professor in the Department of Physics, National Cheng Kung University, Taiwan. He received his PhD in physics in 1993 from the National Tsing Hua University, Taiwan. His main scientific interests focus on the essential properties of carbonrelated materials and low-dimensional systems. He is a member of the American Physical Society, the American Chemical Society, and the Physical Society of Republic of China (Taiwan).

Wen-Dung Hsu’s is currently an associate professor in the Department of Materials Science and Engineering, National Cheng Kung University. His expertise is utilizing computational materials-science methods, including first-principle calculations, molecular dynamics simulations, Monte Carlo methods, and finite-element methods to study materials issues. His research interests are mechanical properties of materials from atomic to macro scale, lithium-ion batteries, solid-oxide fuel cell, ferroelectrics, solid catalyst design for biodiesel, and processing design for single-crystal growth. He obtained his PhD from the Department of Materials Science and Engineering, the University of Florida in 2007. He then served as a postdoctoral researcher in the Department of Mechanical Engineering at the University of Michigan. He has been with National Cheng Kung University since 2008.

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Contributors Sanjaya Brahma Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan

Jing-Ru Chen Department of Chemical Engineering and Materials Science Yuan Ze University Chung-Li, Taiwan

Chia-Chin Chang Research and Development Center for Li-ion Battery National University of Tainan Tainan, Taiwan

Kun-You Chen Department of Chemical Engineering National Cheng Kung University Tainan, Taiwan

and Department of Greenergy National University of Tainan Tainan, Taiwan and Hierarchical Green-Energy Materials Research Center National Cheng Kung University Tainan, Taiwan Chia-Yun Chen Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan and Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan Hung-Yuan Chen Department of Greenergy National University of Tainan Tainan, Taiwan

Peter Chen Department of Photonics National Cheng Kung University Tainan, Taiwan Center for Micro/Nano Science and Technology (CMNST) National Cheng Kung University Tainan, Taiwan and Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan Tsan-Yao Chen Department of Engineering and System Science National Tsing Hua University Hsinchu, Taiwan and Hierarchical Green-Energy Materials Research Center National Cheng Kung University Tainan, Taiwan

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Chi-Cheng Chiu Department of Chemical Engineering National Cheng Kung University Tainan, Taiwan and Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan Guan-Ying Fu Department of Chemical Engineering National Cheng Kung University Tainan, Taiwan Chih-Wei Hu Department of Engineering and System Science National Tsing Hua University Hsinchu, Taiwan Jow-Lay Huang Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan and Center for Micro/Nano Science and Technology National Cheng Kung University Tainan, Taiwan and Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan I-Ming Hung Department of Chemical Engineering and Materials Science Yuan Ze University Chung-Li, Taiwan

Contributors

and Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan Jeng-Shiung Jan Department of Chemical Engineering National Cheng Kung University Tainan, Taiwan Yung-Der Juang Department of Materials Science National University of Tainan Tainan, Taiwan Chin-Lung Kuo Department of Materials Science and Engineering National Taiwan University Taipei, Taiwan Chun-Hsien Kuo Department of Physics National Cheng Kung University Tainan, Taiwan Chia-Jung Lee Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan Yuh-Lang Lee Department of Chemical Engineering National Cheng Kung University Tainan, Taiwan and Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan

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Contributors

Ming-Hsien Li Department of Photonics National Cheng Kung University Tainan, Taiwan and

Ralph Nicolai Nasara Department of Materials Science and Engineering National Cheng Kung University Tainan City, Taiwan

Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan

Duy Khanh Nguyen Department of Physics National Cheng Kung University Tainan, Taiwan

Wei-Bang Li Department of Physics National Cheng Kung University Tainan, Taiwan

Hai Duong Pham Department of Physics National Cheng Kung University Tainan, Taiwan

Chih-Ao Liao Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan

Ilham Ramadhan Putra Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan

Kuan-Yu Lin Department of Photonics National Cheng Kung University Tainan, Taiwan

Lakshmanan Saravanan Research and Development Center for Li-ion Battery National University of Tainan Tainan, Taiwan

Shih-Kang Lin Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan and

Pawan Kumar Singh Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan

Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan

Yu-Fan Su Department of Greenergy National University of Tainan Tainan, Taiwan

Shih-Yang Lin Department of Physics National Chung Cheng University Chiayi, Taiwan

Ramesh Subramani Department of Chemical Engineering National Cheng Kung University Tainan, Taiwan

Hsin-Yi Liu Department of Physics National Cheng Kung University Tainan, Taiwan

Hsisheng Teng Department of Chemical Engineering National Cheng Kung University Tainan, Taiwan

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and Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan and Center of Applied Nanomedicine National Cheng Kung University Tainan, Taiwan Ngoc Thanh Thuy Tran Hierarchical Green-Energy Materials (Hi-GEM) Research Center National Cheng Kung University Tainan, Taiwan Yu-Jen Tsai Department of Materials Science and Engineering National Taiwan University Taipei, Taiwan Yu-Chao Tseng Department of Chemical Engineering National Cheng Kung University Tainan, Taiwan Chai-Hao Tu Hierarchical Green-Energy Materials (Hi-GEM) Materials Research Center National Cheng Kung University Tainan City, Taiwan

Contributors

Shanmuganathan Venketasan Department of Chemical Engineering National Cheng Kung University Tainan, Taiwan Yi-Hung Wang Department of Chemical Engineering and Materials Science Yuan Ze University Chung-Li, Taiwan Shao-Chieh Weng Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan Chun-Ming Wu National Synchrotron Radiation Research Center Hsinchu, Taiwan Ming-Hsiu Wu Department of Materials Science and Engineering National Cheng Kung University Tainan, Taiwan Ping-Ching Wu Department of Greenergy National University of Tainan Tainan, Taiwan

1

Introduction Jow-Lay Huang, Chi-Cheng Chiu, Shih-Yang Lin, Chin-Lung Kuo, Duy Khanh Nguyen, Ngoc Thanh Thuy Tran, Wen-Dung Hsu, Chia-Chin Chang, Jeng-Shiung Jan, Hsisheng Teng, Chia-Yun Chen, I-Ming Hung, Peter Chen, Yuh-Lang Lee, and Ming-Fa Lin

Energy, which is used in everyday living, can be saved in various forms, such as chemical batteries,1 solar electromagnetic fields,2 hydrogen,3 flowing water,4 blowing wind,5 radiative atoms,6 oil mines,6 oil gas,6 and coal mines.7 To greatly reduce the environmental impact, plenty of theoretical and experimental studies have been done for developing the various green energy materials.8–10 For example, the up-todate well-established potential applications cover the battery-driven cell phones11 and electric vehicles,12 the solar-cell factories,13 the hydrogen-based buses,14 watergenerated electric power,15 and the wind turbines.16 Specifically, this book is focused on lithium-ion batteries (LIBs),17 dye-sensitized solar cells,18 and perovskite solar cells.19 Furthermore, how to design and fabricate the electronic and optical devices with excellent performance, low cost, light weight, high safety, long lifetime, operating at a controllable temperature, and suitable voltage range are the main issues.20 The distinct theoretical models are proposed/developed to fully comprehend the diverse physical, chemical, and material phenomena. According to the previous studies, the molecular dynamics simulations,21 the first-principle calculations under the local charge density approximations,22 and the neutral network methods23 are available in thoroughly exploring the fundamental properties and solving the critical issues. As for LIBs, detailed analyses will be conducted on the anode materials accompanied with the chemical modified electrolytes and the significant additives,24 and the different functional polymer binders.25 On the experimental side, the successful syntheses of the emergent materials;26 the high-resolution measurements on geometric, electronic, optical, and transport properties;27 and the delicate examinations on the battery performance and the photon-to-electron conversion efficiency28 will be finished under a series of systematic studies. Detailed comparisons between the experimental measurements and the theoretical predictions are also made. Part of the inconsistency arising from them become new and open issues proposed in the contents. Developing functional polymer binders for LIB cathodes and anodes has drawn much attention for improving LIB capacity, due to their low overall content yet critical role at the electrode interface. One of the widely applied commercial binders for LIBs is poly(vinylidene difluoride) (PVDF) for its good electrochemical stability 1

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Green Energy Materials Handbook

and adhesive properties. Yet PVDF binder is an inert conductor of lithium ions (Li+), leading to high polarization resistance near the LIB electrode at a high charging/discharging rate. Hence, introducing ion-conducting polymers such as PEO and PAN into binder development is one of the common strategies to enhance the overall performance of LIBs. Gong et al.38 applied PAN as binders for various anode materials, including graphite, Li4Ti5O12, and Si/C, and showed improved adhesion, reduced charge transfer resistance, and enhanced capacity endurance. More recently, Tsao et al. utilized PEO-b-PAN copolymer as LiFePO4 cathode binder for LIB and this showed an improved charge transfer resistance and high-capacity retention under a high C-rate. Another study by the same group developed a water-borne binder of fluorinated copolymer functionalized with PEO featuring small impedance during charging and discharging.29 A novel PEDOT:PSS developed by Das et al. as LiFePO4 cathode binders showed improved LIB capacities after a cycling test. Note that current studies have utilized various complex formulas such as polymer blends or copolymers, making it difficult to identify the molecular effects of each functional polymer. Also, the detailed mechanisms of the aforementioned novel functional polymers on affecting Li+ transports at the electrode interface may differ from polymer electrolyte and thus remain elusive. Graphite, in which graphene layers are periodically arranged along the z-direction, has been extensively utilized in everyday living for a long time. Three kinds of typical stacking configurations have been successfully identified from the experimental measurements. There exist AAA, ABA, and ABC stackings, namely, the simple hexagonal, Bernal, and rhombohedral graphites. The second one dominates in natural graphite, while the third one only corresponds to the partial system. Apparently, the theoretical and experimental studies on them show a lot of unusual fundamental properties (e.g., electronic properties,30 magnetic quantization,31 optical absorption, and reflectance spectra32,33) and transport properties.34 All pristine graphites belong to the semimetals,35 mainly owing to the interlayer atomic interactions of C-2pz orbitals. Their interlayer attractive forces mainly originate from the van der Waals interactions. They are weak but significant; therefore, many different guest atoms/ions are easily intercalated into the graphitic spacings. On the other side, alkali atoms can create active chemical environments to form the critical interactions with other atoms or molecules, since each of them possesses an s-state electron in the outmost orbital. They are suitable for serving as guest atoms intercalated into the layered graphite, leading to a very high electrical conductivity.36 Up to now, the stage-n alkali graphite-intercalation compounds have been successfully synthesized except for Na guest atoms. Apparently, there are important differences between Li and other alkali atoms (M = K, Rb, Cs). For example, the stage-1 systems are, respectively, LiC6 and MC8 with the distinct unit cell. In particular, the stacking configuration of the neighboring graphitic layers is AAA or ABA, being sensitive to the type and concentration of alkali atoms.37 As to the intercalation and deintercalation of Li+ ions in graphite, such actions might appear frequently in the charging and discharging processes of lithium-based batteries. In addition, whether the Li+ ions become neutral Li atoms is an open issue since it is very difficult to identify the orbital configurations of the extra components from the current experimental measurements. Both kinds of graphite-intercalation compounds are predicted to present the similar

Introduction

3

geometric structures.8 It is worthy of a systematic study on the essential properties in pristine graphites, as well as Li-atom and Li+-ion graphite-intercalation compounds by using the first-principles method. The existence of defects in graphene-based anode materials has been believed to enhance the specific capacity of LIBs, but the mechanisms remain unclear to date. Our results based on first-principles calculations indicated that the local geometry of vacancy defects strongly influences the Li capacity instead of the number of missing atoms. In addition to vacancy defects, not only the disordered region generated by rotating carbon bonds in pristine graphene but also a large number of structural disorders in amorphous graphene can lead to greater Li capacity. According to the electronic structure analysis, the enhanced Li capacity could be attributed to the increased amount of states just above the Fermi level (p-type doping effect) induced by structural disorders in defective or amorphous graphene. Up to now, the graphite or silicon–graphite composite system frequently serves as an efficient anode material in the Li+-ion batteries,38–40 where the latter might have the better performance.41 The former consists of periodical graphene layers with honeycomb lattices due to very strong C atom sp2 bondings, while the silicon atoms in the latter might dramatically change the critical orbital hybridizations and thus greatly modify the fundamental properties. Chemical modifications on layered graphenes by silicon-guest-atom chemisorptions and substitutions (dopings) will become one of the mainstream topics. Pristine graphene systems, with various layer numbers and stacking configurations, have stirred a lot of theoretical42,43 and experimental44 studies in terms of basic science,45 engineering,46 and potential applications.47 They are expected to exhibit the diverse physical, chemical, and material phenomena,48–50 as verified from the high-resolution experimental measurements51—for example, the unusual electronic structures,52 magnetic-field-dependent Landau-level spectra,53 quantum Hall effects,54 optical absorption spectra,55 superhigh mechanical properties,56 Coulomb excitations57 and decays,58 and adatom- and molecule-diversified essential properties.59–65 Previous research has shown that surface adsorptions by distinct atoms and functional groups can create diverse phenomena, such as the opening of the energy gap,66 the semiconductor-metal transition,67 the destruction of the Dirac-cone band structure,68 and richer van Hove singularities in the density of states.69 The similar or different behaviors are revealed in the doping cases.70,71 Up to now, systematic studies on silicon-related chemical modifications have been absent; therefore, they will be finished under first-principles calculations, accompanied by the development of a theoretical framework. The calculated results cover the Si-enriched geometric structures, the Si- and C-dominated/C-codominated valence and conduction bands, the spatial charge densities and their variations after chemical modifications, the charge transfers of Bader analyses, and the atom- and orbital-decomposed densities of states. Such physical quantities could provide full information in determining the critical multiorbital hybridizations of Si–C bonds for the surface adsorptions and substituted dopings. The electronic properties of adatom-doped graphenes are one of the important topics in physics, chemistry, and materials. They are greatly diversified by various adatom adsorptions. The adatom-doped graphenes have attracted a lot of theoretical72–76 and experimental77–79 studies. More recently, the metal-doped graphenes have

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attracted great attention. Based on first-principles calculations, the adsorption of metal atoms (e.g., aluminum, and iron/cobalt/nickel) could lead the gapless semiconducting behavior of pristine graphene to become n-type dopings. After adsorption on graphene surfaces, they are expected to induce more complicated multiorbital hybridizations in the significant carbon–metal bonds. The theoretical predictions could highly promote further understanding of experimental measurements. For example, aluminum-based batteries have been developed quickly to greatly enhance the charging and discharging reactions and to reduce the cost of the metallic anode.80 The Al-adsorbed graphene is predicted to create high numbers of free carriers, while the highest concentration is limited up to 25%.81 Recently, transition metal-absorbed graphenes have been shown to exhibit magnetic properties and other interesting properties that are promising for spintronic devices and batteries.82–84 Four different stable structures have been observed on Fe-adsorbed graphene, and the corresponding first-principles calculations show that these structures all possess magnetic moments.82 On the other hand, the adsorption site, magnetic configuration, and anisotropy of single Co adatoms on graphene are determined by experiments.85 The different geometric structures, the free carrier density, the electronic properties, and the magnetic configurations of transition metal (Fe, Co, Ni)-adsorbed graphene will be thoroughly discussed. Recently, a great deal of research regarding energy density and voltage of LIBs, which are being developed for serving grid support in many applications, has gained increasing interest.86–91 However, high-voltage application in LIBs makes it unstable and lowers performance crucially.92,93 Using various solvent formulas and functional additives can address the issue with stabilizing the electrolytes from decomposition. The selection of the best additive among phosphides, sulfonate esters and lactam, and so on, which perfectly fits the electrodes, can accelerate and improve the developments in LIBs.94 Several parameters, including the energy of lowest unoccupied molecular orbital (LUMO), highest occupied molecular orbital (HOMO), oxidation potentials (OPs), and reduction potentials (RPs), can be used for first screening of additive selection. A good additive can help enhance the cycle life of Li-ion, especially at higher temperatures, and keeps the internal resistance low with use and age. However, due to the limited calculation resource, it would not be affordable to obtain more molecule properties in the solvent state. An artificial neural network (ANN) is a computational intelligence tool comprising many nonlinear processing elements connected to each other. The combination of an ANN with density functional theory calculations is considered a good choice to investigate the effects of molecular structure on the electrochemical parameters of molecules and find out potential candidates for electrolyte additives. LIBs are considered future energy storage devices that have the potential to meet the energy demands from all possible consumer electronic devices ranging from portable toys to high-end electric hybrid vehicles. The physical (structure, morphology) and chemical properties of the electrodes (anode/cathode) play a very important role in the overall enhancement of the electrochemical properties of LIBs. Graphite is usually used as the anode in LIBs because of excellent cyclic stability and rate capability, but the low theoretical capacity (372 mAh/g) obstructs its applications in future devices that need high capacity. Therefore, continuous

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research is ongoing to develop alternate anode materials that can deliver high capacity for a large number of cycles at a high current rate. Metal oxides (e.g., SnO2, TiO2, MnO2, Mn3O4, Fe2O3, Co3O4 etc.) with high theoretical capacity (>750 mAh/g) have been investigated for their applications in LIBs. However, the issue still lies with the low electrical conductivity of the host oxide and the severe volume expansion (>300%) that destroy the electrode after a few charge/discharge cycles and undermine overall device performance. A variety of different strategies have been undertaken to overcome these issues: (a) reduce the particle size of the host oxide materials to nanosize to accommodate the volume expansion as well as strain generated during the charge/discharge cycles; (b) modify the morphology to sheet, rod, tube, and hollow spheres and create pores within the materials for more lithium-ion storage; (c) employ surface engineering; and (d) incorporate carbon-based materials (activated carbon, carbon nanotubes, graphene) that can act as a buffer for the growth of oxide nanomaterials and increase the conductivity of the overall composite. In this chapter, we review the LIB performance of some potential binary/ternary metal oxides and highlight the enhancement in the electrochemical performance of the metal oxides after incorporation in graphene/ reduced graphene oxide. LIBs have been widely used in electric vehicles, portable devices, grid energy storage, and so on, especially during the past decades because of their high specific energy densities and stable cycling performance.95,96 Although LIBs have been widely used as a rechargeable power source, there are still many hidden physical and chemical phenomena to be uncovered and understood for performance improvement. Considering the variety and adaptability of energy conversion and its applications, the abundance and cost of metal ions are key factors in mass energy storage systems. In these considerations, the sodium-ion battery (SIB) has also attracted much attention in the academic and industrial sectors.97 A SIB comprises electrochemical and geometric configurations identical to those of a LIB, so such a device could be produced without significant modification of production facilities. Investigations of the phenomena at the atomic scale are essential for full understanding of the processes during battery operation.98 On the basis of elemental abundance, NaxFeO2 (NFO) is a proper cathode candidate for the commercial production of SIBs. Fe sites are replaced with Mn dopant to improve the stability and capacity of NFO (e.g., Fe1/2Mn1/2O2), which shows a reversible capacity of 170–180 mAh g−1 as compatible with that of the cathode in a LIB suggests to be a potential for the commercialization of SIBs.99,100 Batteries operated in both high- and low-temperature conditions and the temperature change during charging/discharging affect the performance and safety limits. Graphite has been used as a standard anode material in commercial LIBs because of its low cost, low lithium intercalation potential, and good cycling stability. Li-plating on graphite anodes severely degrades the performance of the battery, including capacity loss, impedance rise, activity slow down, and aging rate increase. It is now known that intercalation into graphite and plating onto graphite surface can both occur when Li-ions return to the graphite anode upon charging.101 On the other hand, Li-plating can also lead to loss of active lithium and capacity fading. Plated lithium may react with the electrolyte (adding to solid−electrolyte interphase [SEI]

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growth) or become disconnected from the graphite, forming a reservoir of inactive metallic lithium.102,103 Unfortunately, studying Li-plating ex-situ is difficult, since the deposition of metallic lithium is partly reversible. Also, it is difficult and potentially dangerous to open a cell in the charged state because of possible short-circuiting. The high amount of Li inside a battery in combination with its high reactivity makes the battery very sensitive to the conditions of operation, and its improper use was already reported to be the reason for serious hazards. The high chemical activity of the battery constituents results in a significant materials interaction, and the control of possible risks requires a systematic investigation of the Li-ion battery under real operation conditions “in operando.” An experimental investigation using a nondestructive probe that is capable of revealing the in-situ Li+ migration rate during discharging–charging can thus help develop optimal operating conditions to avoid Li-plating. The formation of a SEI over electrodes during charging/discharging, composed of electrolytes that have been reduced at reactive sites on the electrode surface, can cause significant initial irreversible capacity loss and reduce the battery life.104 Despite the importance of the SEI to battery performance, it remains a formidable challenge to observe the growth of this layer, and the mechanism of SEI formation in carbon anodes remains a key unsolved problem in battery research.105 Several researchers have been worked on understanding the complex mechanism of the SEI layer formation and improving the performance in terms of cycle life of a battery.106 The SEI on a graphite anode is a highly passivating film that protects the graphite particles from the cointercalation of solvent molecules, prevents the exfoliation of graphene layers, and is permeable only to Li+ ions. The details of the SEI layer formation with the electrochemical cycling in the batteries have been elusive, and it is necessary to examine the evolution of the SEI by using in-situ techniques. Other methods such as Fourier-transform infrared spectroscopy (FTIR), X-ray absorption spectroscopy (XAS), optical observation, and atomic force microscopy (AFM), have also been adapted to in-situ studies of rechargeable batteries,107,108 but these methods do not provide the combination of microstructural and compositional information compared to in situ neutron scattering. The in-situ diagnostic techniques will play a critical role in materials and process innovation, system optimization, safety analysis, failure diagnosis, and lifetime prediction in developing next-generation LIBs. Recent reviews109–111 provide detailed descriptions on designing the in-situ electrochemical cells for neutron diffraction, the experiments on neutron scattering, and the in-situ neutron powder diffraction to probe the relationship between electrochemical performance and structural properties within the electrode materials. Most importantly, the electrochemical performance of the in-situ cell should be similar to that of the real (commercial) battery, so that the structural information extracted can be directly compared to that found within the technological applications. The use of in-situ methods to probe atomic and molecular-scale battery function during operation is becoming increasingly common as the information gained in such experiments can direct methods to improve existing battery materials (e.g., by identifying possible failure mechanisms112,113 and by revealing crystal structures that could be considered for the next generation of materials).114 In the case of in-situ X-ray powder diffraction studies,

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Reimers et al.115 and Morcrette et al.116 showed that LixCoO2 undergoes multiple phase transitions on the removal of Li+. Recently, in-situ X-ray and neutron diffraction has been successfully used to reveal changes in the crystal structures of electrode materials during battery operation.117−119 This chapter describes some of the experimental investigations of employing the in-situ neutron and X-ray scattering analysis on the structural disorder and phase transition of different anode and cathode materials in the Li- and Na-ion batteries. The experiments are the in-situ neutron diffraction study on an 18650-type (NMC/ MGP) Li-ion battery in discharging-charging operations, carried out at and below the ambient temperature (−20˚C and −5˚C), to investigate the lithium diffusion into the anode and to study the phases observed during the intercalation/deintercalation; the in-situ X-ray diffraction study on submicro-sized Si nanostructure prepared by the high-energy mechanical milling (HEMM) and wet milling techniques to investigate strategies for alleviating volume changes and high conductivity to make practical silicon-based anodes; in-operando X-ray diffraction (XRD) to study the changes observed in the crystal structure of artificial conducting graphite (FSN) as a function of applied potential to analyze the effects of high-energy milling on silicide formation and lithiation/delithiation affinity of Si materials; and the in-operando neutron diffraction studies on local structure fading induced irreversibility in a 18650 cell with P2-Na2/3Fe1/3Mn2/3O2 (NFMO) crystal as a cathode in a SIB. LIBs have been regarded as a developing power source for portable mobile devices, energy storage systems, and hybrid/electric vehicles of the next generation owing to their advantages, such as high capacity, high-energy density, cyclability, and adjustable appearance.120–123 Electrolyte systems, which transport ions from the anode to the cathode, are the key component of LIBs. The current commercial electrolyte systems used for LIBs are conventionally composed of mixtures of organic carbonates and lithium salt. However, the volatility and flammability of the above liquid electrolytes remain problems, causing LIBs to carry the risks of leakage and related fire hazards.124–126 The safety concerns posted by these commercial electrolytes have deterred the practical application of LIBs at a large scale and driven the development of alternative electrolytes that are electrochemically stable, nonflammable, and nonvolatile. Ionic liquids (ILs) have recently emerged as a new class of electrolyte and appeared to be the suitable replacement of organic carbonate due to their wide electrochemical window, negligible flammability, and vapor pressure,127–131 which are especially in need of stable and safe electrolytes.132–134 However, the relatively poor ionic conductivity of IL-based electrolytes compared with that of organic carbonatebased electrolytes allowed the simple replacement of carbonates with ILs result in poor cycling performance for LIBs.135,136 Accordingly, various strategies have been pursued to improve the performance of IL-based electrolytes, one of which is mixing the IL with carbonate.137–142 This approach can allow better ion transport with respect to pure IL-based electrolytes and also improve the interface compatibility between electrodes and electrolytes. Unfortunately, the ratio of the carbonate solvent in the mixtures is usually quite high, most of which are above 20 wt%. Although the flammability can be suppressed, issues such as the accumulation of volatile carbonates inside the batteries or electrolyte leakage accompanied by carbonate evaporation are still unresolved.

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Hence, a quaternary system composed of a polymer, IL, organic carbonate, and lithium salt may have the potential to optimize the electrolyte system. We reported the preparation of novel polymer electrolytes based on poly(VdF-co-HFP), EMIMTFSI, EC, and LiTFSI by a solvent casting method, and their use as the electrolyte system for Li/LiFePO4 cells was researched. Additionally, the thermal properties and electrical performances, such as ionic conductivity, electrochemical stability, interfacial behaviors, and the effect of morphology on the Li/LiFePO4 cell performance for these electrolytes, were also discussed in detail. Since the commercialization of LIBs in 1991, LIBs have been the most attractive battery technology in all kinds of portable electronics due to their advantages and electrochemical properties.143–146 Despite their establishment in energy storage technology, the improvement in practical energy density and advancement in technology were sluggish. The energy density of current LIBs (theoretically ~390 W h kg−1 and practically 100–200 W h kg−1) cannot meet the demands of energy storage.147–149 To fulfill consumers’ growing demand for advanced future technologies, the development of batteries with safer electrolytes, design, high-energy density, and enhanced battery life is urgently needed. The use of metallic Li as an anode is a key design for developing high-energy density LIBs due to the high specific capacity (3860 mAh g−1) and low electrochemical potential (–3.04 V vs SHE) of metallic Li.148–150 Usage of Li metal as anodes in current LIBs is not feasible due to the instability of liquid electrolytes against Li metal. This leads to Li dendrite growth at the Li metal–liquid electrolyte interface, and the short circuit of dendrites results in a fire hazard.143,149,151 Moreover, the developments of high-energy density metal batteries with safer electrolytes are required for large-scale applications such as electric vehicles and grid energy storage systems. Gel/solid-state electrolytes (SSEs) have the potential to overcome the safety issues of conventional organic liquid electrolytes. Gel polymer electrolytes (GPEs) comprise a polymer framework, an organic solvent/plasticizer, and an electrolytic salt.152 An adequate design of the functional groups enables the polymer framework to effectively entrap the solvent molecules and facilitate ion transport. On the other hand, SSEs are considered next-generation electrolytes due to many advantages such as no leakage of electrolytes, chemical stability, thermal stability, and wide electrochemical window.153–156 However, understanding of their ion-conducting mechanism and interface compatibility against electrodes is still in the early stage. Hybrid solar cells are composed of organic and inorganic materials for constructing the p–n junction that can efficiently convert the light to the electric power. In this aspect, the involved nanostructures, the electrical conductivity of organic components, and structural designs will be introduced. A dye-sensitized solar cell (DSSC) is a photoelectrochemical system, also known as the Gratzel cell, based on the photosensitized anode, electrolyte, and cathode.157–160 The most attractive feature of the DSSC is the simple manufacture process based on roll-printing techniques,161 which provides a variety of uses applicable not only to glass-based systems but also to the polymer-based substrate, and the most of the materials used are low cost. One important factor affecting the conversion efficiency of DSSC is the surface area and morphology of anode material.162–168 In this chapter, nanostructures such as nanofibers and nanoforests offer a large surface area for dye adsorption and/or a direct pathway for electron transport, which improve the conversion efficiency of DSSC.

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Emerging solar cells, such as DSSCs and perovskite solar cells (PSCs), have become promising alternatives to conventional silicon-based solar cells due to their ease of fabrication and use of cost-effective materials.169–171 Conventional DSSCs based on the sandwich architecture mainly consist of two fluorine-doped tin oxide (FTO)–coated glass substrates with the electrolyte sandwiched between them. One FTO substrate is coated with mesoporous TiO2 as a working electrode for electron collection, onto which the dyes are anchored for light harvesting. The other one is platinized to form an ohmic contact with the electrolyte to serve as a cathode. The high cost and heavy weight of FTO glass hinder future application. For the aim of low cost and light weight, elimination of one FTO substrate for DSSCs was proposed to fabricate a monolithic device structure. The organic–inorganic hybrid halide perovskite materials have attracted tremendous attention due to their promising electrical and optical properties and versatile fabrication process, resulting in high power conversion efficiency (PCE) for the PSCs.172–175 The merits of strong absorption over the visible range, promising ambipolar transport property with a long carrier diffusion length, high carrier mobility,172,176–179 and facile solution-processed fabrication for perovskite, make it as a promising light absorber for cost-effective solar cells to compete with the current existing photovoltaic technologies. The conventional perovskite solar cells are mainly composed of two selective contact electrodes, and a perovskite light absorber is sandwiched by them. Because perovskite is vulnerable to high-temperature processes and polar solvents, low-temperature processed soft chemistry is favorable for perovskite solar cells. As inspired by the monolithic dye-sensitized solar cells,180 mesoporous counter electrodes have recently been employed for PSCs, where the perovskite precursor is ultimately infiltrated in the final step for device fabrication. Thus, the deposition of perovskite in the last step of device fabrication offers several advantages. First, it allows us to apply high temperatures or harsh chemical processes for making selective electrodes and provides more selections of charge transport materials as well as their fabrication process. As a result, perovskite deposition as the last step in device fabrication prevents an unnecessary reaction from degrading the PSCs. Second, the demand for a large area smooth morphology is alleviated as the contact is a mesoscopic junction. DSSCs are a suitable alternative for classical silicon solar cells in terms of low fabrication cost, easy preparation procedures, and high-energy conversion efficiency.181 However, liquid-state DSSCs shows poor long-term performance because volatile liquid electrolytes are utilized. An easy method for improving the long-term performance as well as the stability of the cells is by substituting liquid electrolytes with quasi-solid-state polymer gel electrolytes.182 In this chapter, the introduction describes the overall development of DSSC electrolytes such as liquid, polymer, and nanocomposite electrolytes. The general aspects of DSSC electrolytes partly describe the basic structure, principles of DSSCs, and importance and drawbacks of liquid, polymer gel, and nanocomposite electrolytes. The fabrication of the DSSCs describes the fabrication of laboratory and submodule cells. The characterization of DSSCs describes various methods involved in the characterization of DSSCs. The chapter183–185 describes the performance of the poly(acrylonitrile-co-vinyl acetate), poly(vinylidene fluoride-hexafluoropropylene), poly(ethylene oxide), and poly(methyl methacrylate)-based gel electrolytes in DSSCs.

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The organization of this book is as follows. Chapter 1 discusses the theoretical and experimental progress. The details of the theoretical modeling are included in Chapters 2–7. For Chapter 2, the molecular dynamics simulations are available in systematically examining the molecular effects due to five kinds of functional polymer binders on the LiFePO4 cathode surface of LIBs, and a detailed comparison with the commercially applied PVDF binder is thoroughly done. Chapter 3 covers the first-principles investigations on Li storage capacity in graphene-based anode materials, in which it is sensitive to different types of defect structures, various amorphous morphologies, vacancies, and Stone–Wales (SW) defects. The simulation method is suitable for fully exploring the chemical modifications on the fundamental properties of graphene-related systems, as clearly illustrated in Chapters 4–6. The geometric electronic and magnetic properties are conducted on the Li/Li+ stage-n graphite-intercalation compounds (Chapter 4), the silicon chemisorptions and substitutions on graphene surfaces (Chapter 5), and (Fe, Co, Ni)–adsorbed graphene systems (Chapter 6). The critical mechanisms, the single- or multiorbital hybridizations in chemical bonds, will be accurately identified from the atom-dominated energy bands, the spatial charge distributions, and the atom-, orbital-, and spin-projected density of states. Moreover, Chapter 7 focuses on the direct and efficient combination of the neutral network and first-principles calculations, being very useful in initial screening of electrolyte additives. On the experimental side, Chapter 8 investigates how to increase the Li+-transport performance of LIBs by using alternative anode materials of metal O-reduced graphene oxide nanocomposites, especially its relations with Li storage capacity and cycle number. A platform with proper design of the configuration of secondary ion battery components and protocol for in-situ X-ray and neutron analyses are introduced in Chapter 9. The gel polymer electrolytes of LIBs are synthesized and examined in Chapter 10, where they consist of micro-phase separated poly(VdF-co-HFP)/ ionic liquid/carbonate. A simple experimental method is identified to have excellent thermal stability and electrochemical/cycling performance. The different gel and solid electrolytes in LIBs are also discussed in Chapter 11 to greatly enhance ion transport. As for solar cells, the silicon-nanowire-based hybrid materials, as clearly indicated in Chapter 12, are explored for the great enhancement of the photovoltaic performance under the various methods, such as generation of nanostructures on n-type components, increase of the electrical conductivity in organic components, adopting interface engineering, and optimization of device structure. In Chapter 13, the Li-doped ZnO (Li-ZnO) DSSCs are successfully fabricated, and the ­typical geometric structures are accurately identified from scanning electron microscopy (SEM) and transmission electron microscopy (TEM) measurements. Whether such materials could serve as the highly efficient photoanode will be tested thoroughly. Chapter 14 discusses monolithic DSSCs and PSCs using porous electrodes. In Chapter  15, high-performance quasi-solid-state polymer electrolytes for dye-­ sensitized solar cell applications are presented. Most important, the advantages of light weight and low cost are discussed in detail. The concluding remarks for the specific chapters are made in Chapter 16. Finally, the focuses of Chapter 17 are nearfuture systematic investigations of open issues.

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REFERENCES 1. Tarascon, J.-M.; Armand, M. Materials for Sustainable Energy: A Collection of Peer-Reviewed Research and Review Articles from Nature Publishing Group; World Scientific, 2011; pp. 171–179. 2. Maier, S. A.; Atwater, H. A. Plasmonics: Localization and guiding of electromagnetic energy in metal/dielectric structures. Journal of Applied Physics 2005, 98, 10. 3. Dunn, S. Hydrogen futures: Toward a sustainable energy system. International Journal of Hydrogen Energy 2002, 27, 235–264. 4. O’Brien, E. Water-energy dynamics, climate, and prediction of woody plant species richness: An interim general model. Journal of Biogeography 1998, 25, 379–398. 5. Burton, T.; Jenkins, N.; Sharpe, D.; Bossanyi, E. Wind Energy Handbook; Chichester: John Wiley & Sons, 2011. 6. Hubbert, M. K. Nuclear energy and the fossil fuel. Drilling and Production Practice; Washington, D.C.: American Petroleum Institute, 1956. 7. Jevons, W. S. The Coal Question: An Inquiry Concerning the Progress of the Nation, and the Probable Exhaustion of Our Coal-Mines; London: Macmillan, 1906. 8. Green, M. A.; Emery, K.; Hishikawa, Y.; Warta, W. Solar cell efficiency tables (version 37). Progress in Photovoltaics: Research and Applications 2011, 19, 84–92. 9. Zhang, Q.; Sun, Y.; Xu, W.; Zhu, D. Organic thermoelectric materials: Emerging green energy materials converting heat to electricity directly and efficiently. Advanced Materials 2014, 26, 6829–6851. 10. Deng, D.; Kim, M. G.; Lee, J. Y.; Cho, J. Green energy storage materials: Nanostructured TiO 2 and Sn-based anodes for lithium-ion batteries. Energy & Environmental Science 2009, 2, 818–837. 11. Bruce, P. G.; Scrosati, B.; Tarascon, J.-M. Nanomaterials for rechargeable lithium batteries. Angewandte Chemie International Edition 2008, 47, 2930–2946. 12. Clement-Nyns, K.; Haesen, E.; Driesen, J. The impact of charging plug-in hybrid electric vehicles on a residential distribution grid. IEEE Transactions on Power Systems 2010, 25, 371–380. 13. Freundlich, A.; Alemu, A. Multi quantum well multijunction solar cell for space applications. Physica Status Solidi (C) 2005, 2, 2978–2981. 14. Zhao, J.; Melaina, M. W. Transition to hydrogen-based transportation in China: Lessons learned from alternative fuel vehicle programs in the United States and China. Energy Policy 2006, 34, 1299–1309. 15. Arvanitidits, N. V.; Rosing, J. Composite representation of a multireservoir hydroelectric power system. IEEE Transactions on Power Apparatus and Systems 1970, 2, 319–326. 16. Muller, S.; Deicke, M.; De Doncker, R. W. Doubly fed induction generator systems for wind turbines. IEEE Industry Applications Magazine 2002, 8, 26–33. 17. Etacheri, V.; Marom, R.; Elazari, R.; Salitra, G.; Aurbach, D. Challenges in the development of advanced Li-ion batteries: A review. Energy & Environmental Science 2011, 4, 3243–3262. 18. Hagfeldt, A.; Boschloo, G.; Sun, L.; Kloo, L.; Pettersson, H. Dye-sensitized solar cells. Chemical Reviews 2010, 110, 6595–6663. 19. Liu, M.; Johnston, M. B.; Snaith, H. J. Efficient planar heterojunction perovskite solar cells by vapour deposition. Nature 2013, 501, 395. 20. Lu, L.; Han, X.; Li, J.; Hua, J.; Ouyang, M. A review on the key issues for lithium-ion battery management in electric vehicles. Journal of Power Sources 2013, 226, 272–288. 21. Leung, K.; Budzien, J. L. Ab initio molecular dynamics simulations of the initial stages of solid–electrolyte interphase formation on lithium ion battery graphitic anodes. Physical Chemistry Chemical Physics 2010, 12, 6583–6586.

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22. Ouyang, C.; Shi, S.; Wang, Z.; Huang, X.; Chen, L. First-principles study of Li ion diff in LiFePO4. Physical Review B 2004, 69, 104303. 23. Charkhgard, M.; Farrokhi, M. State-of-charge estimation for lithium-ion batteries using neural networks and EKF. IEEE Transactions on Industrial Electronics 2010, 57, 4178–4187. 24. Goriparti, S.; Miele, E.; De Angelis, F.; Di Fabrizio, E.; Zaccaria, R. P.; Capiglia, C. Review on recent progress of nanostructured anode materials for Li-ion batteries. Journal of Power Sources 2014, 257, 421–443. 25. Koo, B.; Kim, H.; Cho, Y.; Lee, K. T.; Choi, N.-S.; Cho, J. A highly cross-linked polymeric binder for high-performance silicon negative electrodes in lithium ion batteries. Angewandte Chemie 2012, 124, 8892–8897. 26. Murtola, T.; Bunker, A.; Vattulainen, I.; Deserno, M.; Karttunen, M. Multiscale modeling of emergent materials: Biological and soft matter. Physical Chemistry Chemical Physics 2009, 11, 1869–1892. 27. others, et al. Boron nitride substrates for high-quality graphene electronics. Nature Nanotechnology 2010, 5, 722. 28. Snaith, H. J.; Ducati, C. SnO2-based dye-sensitized hybrid solar cells exhibiting near unity absorbed photon-to-electron conversion efficiency. Nano Letters 2010, 10, 1259–1265. 29. Tsao, C.-H.; Wu, E.-T.; Lee, W.-H.; Chiu, C.-C.; Kuo, P.-L. Fluorinated copolymer functionalized with ethylene oxide as novel water-borne binder for a high-power lithium ion battery: Synthesis, mechanism, and application. ACS Applied Energy Materials 2018, 1, 3999–4008. 30. Tatar, R.; Rabii, S. Electronic properties of graphite: A unified theoretical study. Physical Review B 1982, 25, 4126. 31. Kopelevich, Y.; Lemanov, V.; Moehlecke, S.; Torres, J. Landau level quantization and possible superconducting instabilities in highly oriented pyrolitic graphite. Physics of the Solid State 1999, 41, 1959–1962. 32. Taft, E.; Philipp, H. Optical properties of graphite. Physical Review 1965, 138, A197. 33. Kuzmenko, A.; Van Heumen, E.; Carbone, F.; Van Der Marel, D. Universal optical conductance of graphite. Physical Review Letters 2008, 100, 117401. 34. Spain, I. L. Electronic transport-properties of graphite, carbons, and related materials. Chemistry and Physics of Carbon 1981, 16, 119–304. 35. Brandt, N. B.; Chudinov, S. M.; Ponomarev, Y. G. Semimetals: 1. Graphite and Its Compounds; Amsterdam: Elsevier, 2012. 36. Matsumoto, R.; Arakawa, M.; Yoshida, H.; Akuzawa, N. Alkali-metal-graphite intercalation compounds prepared from flexible graphite sheets exhibiting high air stability and electrical conductivity. Synthetic Metals 2012, 162, 2149–2154. 37. Dresselhaus, M. S.; Dresselhaus, G. Intercalation compounds of graphite. Advances in Physics 2002, 51, 1–186. 38. Mao, C. et al. Selecting the best graphite for long-life, high-energy Li-ion batteries. Journal of the Electrochemical Society 2018, 165, A1837–A1845. 39. Badi, N. et al. Low-cost carbon-silicon nanocomposite anodes for lithium ion batteries. Nanoscale Research Letters 2014, 9, 360. 40. Yim, C. H.; Courtel, F. M.; Abu-Lebdeh, Y. A high capacity silicon-graphite composite as anode for lithium-ion batteries using low content amorphous silicon and compatible binders. Journal of Materials Chemistry A 2013, 1, 8234–8243. 41. Son, I. H. et al. Silicon carbide-free graphene growth on silicon for lithium-ion battery with high volumetric energy density. Nature Communications 2015, 6, ncomms8393. 42. Nilsson, J.; Neto, A. C.; Guinea, F.; Peres, N. M. R. Electronic properties of bilayer and multilayer graphene. Physical Review B 2008, 78, 045405. 43. Ho, Y. H.; Wu, J. Y.; Chiu, Y. H.; Wang, J.; Lin, M. F. Electronic and optical properties of monolayer and bilayer graphene. Philosophical transactions of the Royal Society of London A: Mathematical. Physical and Engineering Sciences 2010, 368, 5445–5458.

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89. Osiak, M. et al. Structuring materials for lithium-ion batteries: Advancements in nanomaterial structure, composition, and defined assembly on cell performance. Journal of Materials Chemistry A 2014, 2(25), 9433–9460. 90. Blomgren, G. E. The development and future of lithium ion batteries. Journal of the Electrochemical Society 2017, 164(1), A5019–A5025. 91. Shojan, J.; Chitturi, V. R; Soler, J.; Resto, O.; West, W. C.; Katiyar, R. S. High energy xLi2MnO3-(1− x) LiNi2/3Co1/6Mn1/6O2 composite cathode for advanced Li-ion batteries. Journal of Power Sources 2015, 274, 440–450. 92. Goodenough, J. B.; Kim, Y. Challenges for rechargeable Li batteries. Chemistry of Materials 2009, 22(3), 587–603. 93. Li, W.; Song, B.; Manthiram, A. High-voltage positive electrode materials for lithiumion batteries. Chemical Society Reviews 2017, 46, 3006–3059. 94. Haregewoin, A. M.; Wotango, A. S.; Hwang, B.-J. Electrolyte additives for lithium ion battery electrodes: Progress and perspectives. Energy Environmental Science 2016, 9(6), 1955–1988. 95. Choi, J. W.; Aurbach, D. Promise and reality of post-lithium-ion batteries with high energy densities. Nature Reviews Materials 2016, 1, 16013. 96. Sun, Y.; Liu, N.; Cui, Y. Promises and challenges of nanomaterials for lithium-based rechargeable batteries. Nature Energy 2016, 1, 16071. 97. Yabuuchi, N.; Kubota, K.; Dahbi, M.; Komaba, S. Research development on sodiumion batteries. Chemical Reviews 2014, 114(23), 11636–11682. 98. Senyshyn, A.; Mühlbauer, M. J.; Nikolowski, K.; Pirling, T.; Ehrenberg, H. Journal of Power Sources 2012, 203, 126–129. 99. Wang, H.; Yang, B.; Liao, X.-Z.; Xu, J.; Yang, D.; He, Y.-S.; Ma, Z.-F. Electrochemical properties of P2-Na2/3[Ni1/3Mn2/3]O2 cathode material for sodium ion batteries when cycled in different voltage ranges. Electrochimica Acta 2013, 113, 200–204. 100. Yabuuchi, N.; Kajiyama, M.; Iwatate, J.; Nishikawa, H.; Hitomi, S.; Okuyama, R.; Usui, R.; Yamada, Y.; Komaba, S. P2-type Nax[Fe1/2Mn1/2]O2 made from earth-abundant elements for rechargeable Na batteries. Nature Materials 2012, 11, 512–517. 101. Petzl, M.; Kasper, M.; Danze, M. A. Journal of Power Sources 2015, 275, 799–808. 102. Dubarry, M.; Truchot, C.; Liaw, B. Y. ; Gering, K.; Sazhin, S.; Jamison, D.; Michelbacher, C. Journal of the Electrochemical Society 2013, 160(1), A191–A199. 103. Li, Z.; Huang, J.; Yann Liaw, B.; Metzler, V.; Zhang, J. Journal of Power Sources 2014, 254 168–182. 104. Aurbach, D.; Eineli, Y.; Chusid, O.; Carmeli, Y.; Babai, M.; Yamin, H. J. Electrochemical Society 1994, 141, 603. 105. Xu, K. Chemical Reviews 2004, 104, 4303. 106. Zheng, M. S.; Dong, Q.-F.; Cai, H. Q.; Jin, M.-G. Formation and influence factors of solid electrolyte interphase film on the negative electrode surface in lithium ion batteries, fuel cells, and energy conversion. Journal of the Electrochemical Society 2005, 152(11), A2207–A2210. 107. Amalraj, S. F.; Aurbach, D. J. Solid State Electrochemistry 2011, 15, 877. 108. Hong, Y.-S.; Li, N.; Chen, H.; Wang, P.; Song, W.-L.; Fang, D. In operando observation of chemical and mechanical stability of Li and Na dendrites under quasi-zero electrochemical field. Energy Storage Materials 2018, 11, 118–126. 109. Harks, P. P. R. M. L.; Mulder, F. M.; Notten, P. H. L. In-situ methods for Li-ion battery research: A review of recent developments. Journal of Power Sources 2015, 288, 92. 110. Hansen, T. C.; Kohlmann, H. Chemical reactions followed by in-situ neutron powder diffraction. Zeitschrift für Anorganische und Allgemeine Chemie 2014, 640, 3044. 111. Dong, B.; Biendicho, J. J.; Hull, S.; Smith, R. I.; West, A. R. In-situ neutron studies of electrodes for Li-ion batteries using a deuterated electrolyte: LiCoO2 as a case study. Journal of the Electrochemical Society 2018, 165(5), A793–A801.

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112. Sharma, N. et al. Structural changes in a commercial lithium-ion battery during electrochemical cycling: An in situ neutron diffraction study. Journal Power Sources 2010, 195, 8258–8266. 113. Senyshyn, A.; Muehlbauer, M. J.; Nikolowski, K.; Pirling, T.; Ehrenberg, H. In-operando’ neutron scattering studies on Li-ion batteries. Journal of Power Sources 2012, 203, 126–129. 114. Pang, W. K.; Sharma, N.; Peterson, V. K.; Shiu, J. J.; Wu, S. H. In-situ neutron diffraction study of the simultaneous structural evolution of a LiNi0.5Mn1.5O4 cathode and a Li4Ti5O12 anode in a LiNi0.5Mn1.5O4 parallel to Li4Ti5O12 full cell. Journal of Power Sources 2014, 246, 464–472. 115. Reimers, J. N.; Dahn, J. R. Electrochemical and in-situ X-ray diffraction studies of lithium intercalation in LixCoO2. Journal of the Electrochemical Society 1992, 139, 2091. 116. Morcrette, M.; Chabre, Y.; Vaughan, G.; Amatucci, G.; Leriche, J. B.; Patoux, S.; Masquelier, C.; Tarascon, J. M. In situ X-ray diffraction techniques as a powerful tool to study battery electrode materials. Electrochimica Acta 2002, 47, 3137. 117. Senyshyn, A.; Mühlbauer, M. J.; Dolotko, O.; Hofmann, M.; Pirling, T.; Ehrenberg, H. Spatially resolved in operando neutron scattering studies on Li-ion batteries. Journal of Power Sources 2014, 245, 678–683. 118. Zinth, V.; von Lüders, C.; Hofmann, M.; Hattendorff, J.; Buchberger, I.; Erhard, S.; Rebelo-Kornmeier, J.; Jossen, A.; Gilles, R. Lithium plating in lithium-ion batteries at sub-ambient temperatures investigated by in situ neutron diffraction. Journal of Power Sources 2014, 271, 152–159. 119. Canas, N. A.; Einsiedel, P.; Freitag, O. T.; Heim, C.; Steinhauer, M.; Park, D.-W.; Friedrich, K. A. Operando x-ray diffraction during battery cycling at elevated temperatures: A quantitative analysis of lithium-graphite intercalation compounds. Carbon 2017, 116, 255–263. 120. Armand, M.; Tarascon, J. M. Nature 2008, 451, 652–657. 121. Kang, B.; Ceder, G. Nature 2009, 458, 190–193. 122. Sun, Y. K.; Myung, S. T.; Park, B. C.; Prakash, J.; Belharouak, I.; Amine, K. Nature Materials 2009, 8, 320–324. 123. Scrosati, B.; Garche, J. Journal of Power Sources 2010, 195, 2419–2430. 124. Balakrishnan, P. G.; Ramesh, R.; Prem Kumar, T. Journal of Power Sources 2006, 155, 401–414. 125. Goodenough, J. B.; Kim, Y. Chemistry of Materials 2010, 22, 587–603. 126. Lux, S. F.; Lucas, I. T.; Pollak, E.; Passerini, S.; Winter, M.; Kostecki, R. Electrochemistry Communications 2012, 14, 47–50. 127. Smiglak, M.; Reichert, W. M.; Holbrey, J. D.; Wilkes, J. S.; Sun, L.; Thrasher, J. S.; Kirichenko, K.; Singh, S.; Katritzky, A. R.; Rogers, R. D. Chemical Communications (Cambridge) 2006, 24, 2554–2556. 128. Earle, M. J.; Esperanca, J. M.; Gilea, M. A.; Lopes, J. N.; Rebelo, L. P.; Magee, J. W.; Seddon, K. R.; Widegren, J. A. Nature 2006, 439, 831–834. 129. Nakagawa, H.; Fujino, Y.; Kozono, S.; Katayama, Y.; Nukuda, T.; Sakaebe, H.; Matsumoto, H.; Tatsumi, K. Journal of Power Sources 2007, 174, 1021–1026. 130. Jin, Y.; Fang, S.; Chai, M.; Yang, L.; Tachibana, K.; Hirano, S.-I. Journal of Power Sources 2013, 226, 210–218. 131. Sun, X.-G.; Liao, C.; Shao, N.; Bell, J. R.; Guo, B.; Luo, H.; Jiang, D.-E.; Dai, S. Journal of Power Sources 2013, 237, 5–12. 132. Sakaebe, H.; Matsumoto, H.; Tatsumi, K. Electrochimica Acta 2007, 53, 1048–1054. 133. Fernicola, A.; Croce, F.; Scrosati, B.; Watanabe, T.; Ohno, H. Journal of Power Sources 2007, 174, 342–348. 134. Galiński, M.; Lewandowski, A.; Stępniak, I. Electrochimica Acta 2006, 51, 5567–5580.

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Molecular Effects of Functional Polymer Binders on Li+ Transport on the Cathode Surface within LithiumIon Batteries Kun-You Chen and Chi-Cheng Chiu

CONTENTS 2.1 Introduction...................................................................................................... 19 2.2 Molecular Dynamics Simulation Details......................................................... 21 2.3 Results and Discussion.................................................................................... 23 2.3.1 Characterization of Li+ Affinity of Functional Polymers..................... 23 2.3.2 Thermodynamic and Kinetic Effects of Functional Polymers.............25 2.3.3 Li+ Mobility under E-field...................................................................28 2.4 Summary and Future Perspectives................................................................... 31 Acknowledgments..................................................................................................... 32 References................................................................................................................. 32

2.1 INTRODUCTION Due to high power and energy densities, lithium-ion batteries (LIBs) have become common energy storage devices for electronic systems and mobile digital devices. Over the past few years, there has been growing interest in the applications of secondary lithium batteries in electric vehicles and energy storage systems in a smart grid.1–3 Current research has been focusing on the new material developments to increase the capacity, operating voltage, energy density, and heat stability to improve the overall performance and efficiencies of LIBs.4–6 Particularly, due to their low overall content yet critical role at the electrode interface, developing functional binders for LIB cathodes and anodes has drawn much attention for improving LIB capacity.7–10 One of the widely applied commercial binders for LIBs is poly(vinylidene difluoride) (PVDF), due to its good electrochemical stability and adhesion.11,12 Yet the slurry 19

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fabrication of PVDF requires the N-methyl-2-pyrrolidone (NMP) solvent, leading to various issues, including increased manufacturing costs and inherent environmental, safety, and health (ESH) problems. Thus, developing new hydrophilic polymer binders (e.g., carboxymethyl cellulose [CMC], styrene-butadiene rubber [SBR], and polyacrylic acid [PAA]) has attracted much interest in an aqueous-based fabrication process.13–16 Furthermore, the conventional PVDF binder can cause a high polarization resistance at a high charging/discharging rate, due to its inert lithium-ion (Li+) conductivity. Consequently, introducing Li+-conducting polymers into binder development has become one of the strategies to enhance the overall performance of LIBs.17–21 Lithium-ion-conducting polymers, for example, poly(ethylene oxide) (PEO) and polyacrylonitrile (PAN), have been applied to develop various polymer electrolytes.22–25 PEO is a common polymer electrolyte known for its high Li+ conductivity, electrochemical stability, and interfacial stability between electrode and electrolytes. Also, hydrophilic PEO can act as surfactants to stably disperse hydrophobic polymer moieties in aqueous solution.26–30 Three Li+ transport mechanisms of PEO have been proposed: (1) intrachain Li+ conduction along the polymer chain, (2) cooperative motion of Li+ coordinated with the PEO segments, and (3) intersegmental hopping of Li+ between PEO chains.30–33 Other than PEO, PAN has also been applied in the gel polymer electrolyte system since first demonstrated by Watanabe et al., due to its dimensional stability, thermal stability, and liquid electrolyte absorbance.22,34 The Li+ transport mechanism of PAN-based gel polymer electrolytes, as first proposed by Huang, is via the Li+ hopping among the PAN polar nitrile groups and the carbonyl groups of ethylene carbonate (EC), one component of the absorbed liquid electrolyte.22,35 Wang et al. further showed that the competition between EC and PAN to interact with Li+ can assist the formation of isolate Li+, leading to enhanced Li+ conductivity.36 Recently, various ionic polymers have also been applied for polymer electrolyte development. The single-ion polymer electrolyte based on polyanionic triblock polymers P(STFSILi)–PEO-P(STFSILi) developed by Bouchet et al. exhibited high lithium-ion transport number, due to the immobilized anionic groups.24,37 These single-ion polymer electrolytes also have excellent mechanical strength, a stable electrochemical window up to 5 V, and good ionic conductivity. Applying the concepts of polymer electrolytes, introducing ion-conducting polymers into binders to improve LIB performance has been reported in various studies.7 Gong et al. applied PAN as binders for various anode materials, such as graphite, Li4Ti5O12, and Si/C, and showed improved adhesion, reduced charge transfer resistance, and improved LIB capacity endurance.38 More recently, Tsao et al. utilized PEO-b-PAN copolymer as a LiFePO4 cathode binder for LIB and showed an improved charge transfer resistance and high-capacity retention under high C-rate.18 Another study by Tsao et al. demonstrated a water-borne SF binder of fluorinated copolymer functionalized with PEO, which features small impedance after charging and discharging.39 Das et al. utilized PEDOT:PSS as LiFePO4 cathode binders to replace conventional carbon black and PVDF during the cathode slurry process.40 The resulting cathode modified with PEDOT:PSS demonstrated improved capacities after the cycling test. Although there has been great progress on developing new polymer binders, the detailed mechanisms of these novel functional polymers on

Molecular Effect of Functional Polymer Binders on Li+ Transport

21

affecting Li+ transports at the electrode interface still remain elusive. Furthermore, current studies have utilized various complex formulas such as polymer blends or copolymers, making it difficult to identify the molecular effects of each functional polymer. In this work, we applied molecular dynamics (MD) simulations to investigate the molecular effects of various functional polymer binders on the LiFePO4 (LFP) cathode surface. We considered five different polymer systems of PVDF, PAN, PEO, poly(styrene sulfonate) (PSS), and poly(N-vinylformamide) (PNVF).7,41 Using equilibrium MD, we characterized the effects of functional polymers on the Li+ distributions at the LFP/electrolyte interface, the Li+ adsorption free energy, and the Li+ diffusion on the LFP surface. Through nonequilibrium MD of applying an external electric field, we further investigated the molecular mechanisms of different functional binders on modulating lithium-ion intercalation and deintercalation at the cathode surface. The provided molecular insights can serve as steppingstones for the future binder designs for high-power LIBs.

2.2 MOLECULAR DYNAMICS SIMULATION DETAILS Figure 2.1 illustrates the chemical structures for the five different polymer binders focused in this work, including PVDF, PAN, PEO, PSS, and PNVF. The electrolyte was composed of 1 M LiPF6 in ethylene carbonate (EC) and diethyl carbonate (DEC) with 1:1 volume ratio. The all-atom optimized potentials for liquid simulations (OPLS-AA) force field was applied to describe all molecules, including polymers, organic solvents, and lithium salts.42 The OPLS force field has been widely used in MD studies of ionic liquids, polymers, and solid polymer electrolyte systems.43–46 The parameters for PF− were taken from updated OPLS potentials developed by Lopes and Pauda.45 The LiFePO4 (LFP) cathode contained 3 × 8 × 6 unit crystal cell with the [010] surface normal to the z-axis.47 The van der Waals interactions of LFP atoms were described with the OPLS force field, whereas the atomic charges were assigned based on the molecular model by Smith et al.47 As illustrated in Figure 2.2, the initial configuration for each binder system consisted of LFP cathode solvated with 532 EC molecules, 292 DEC molecules, and 78 LiPF6. Each binder systems contained eight polymer chains of 16 monomers. Due to

FIGURE 2.1  Chemical structures for all the functional polymer binders, including (a) poly(vinylidene fluoride) PVDF, (b) polyacrylonitrile PAN, (c) poly(ethylene oxide) PEO, (d) poly(styrene sulfonate) PSS, and (e) poly(N-vinylformamide) PNVF.

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FIGURE 2.2  The representative conformation for a binder system used in MD simulations. Each system contained an LFP cathode colored in pink, LiPF6 with Li+ emphasized with orange spheres, and EC/DEC solvent represented with colored lines. Each cathode/electrolyte interface is covered with four polymer chains represented with colored spheres.

the three-dimensional periodic boundary condition, the system contained two cathode/ electrolyte interfaces, each covered with four polymer chains. The initial configurations were built using the PACKMOL software with the initial box size of 3.1 × 4.2 × 11.5 nm3.48 Note that the simulation box size is much smaller than the LFP particle size reported experimentally. Hence, the cathode/electrolyte interface in the MD simulation was considered a microscopic representation of the LFP particle surface. All MD simulations of binder systems were conducted via LAMMPS.49 Periodic boundary conditions were applied in all three dimensions. Van der Waals and shortrange electrostatic interactions were calculated with a 1.2-nm cutoff, while longrange electrostatic force was evaluated using the Ewald method.50 Temperature and pressure were controlled at 298 K and 1 bar by Nose–Hoover and Parrinello–Rahman algorithms, respectively.51–55 Motion equations were evaluated with an integration time step of 1 fs. Each initial configuration was first energy minimized through the steepest descent minimization algorithm, followed by 10-ps NVT (constant temperature and constant volume) simulations with a 1-fs timestep to eliminate high-energy configurations. Next, the system was equilibrated under the NPzT ensemble with the fixed x- and y-dimensions within 100 ps. Each simulation was then equilibrated for 50 ns under the NVT ensemble with the system coordinates saved every 5 ps. The last 20-ns trajectories were taken for the analyses of structural and dynamic

23

Molecular Effect of Functional Polymer Binders on Li+ Transport

properties. The recent simulation studies by Ponce et al. utilized an external electric field to guide the motion of Li+ and mimic the charging and discharging processes within a Li+ nanobattery.56,57 Inspired by the work, we also applied an external electric field (E-field) to the binder systems. With the E-field in the +Z direction, Li+ ions were driven to move in the +Z direction. This resulted in the deintercalation at the  +Z LFP interface and the intercalation at the −Z LFP interface. Hence, it allowed us to study the discharging and charging processes at the cathode interface simultaneously in one simulation. In this work, we applied an E-field of 0.5 V/Å in the +Z direction to the equilibrated binder system for 20 ps to examine the effects of functional binders on the Li+ deintercalation and the intercalation. The system was further re-equilibrated for 20 ps, followed by an E-field of 0.5 V/Å applied in the −Z direction for 20 ps and reequilibrated for another 20 ps to complete one cycle. Each system was simulated with 20 E-field cycles to examine the binder durability.

2.3 RESULTS AND DISCUSSION 2.3.1 Characterization of Li+ Affinity of Functional Polymers One of the disadvantages of conventional LIB binder PVDF is lacking Li+ affinity and the ability to transport the Li+ ion. To characterize the Li+ affinity for the five different polymers, we first calculated the radial distribution functions (RDFs) between the Li+ and the functional groups for each polymer type (i.e., the fluorine for PVDF, the nitrile group for PAN, the oxide oxygen for PEO, the sulfonate oxygen for PSS, and the formamide oxygen for PNVF). Compared with the conventional PVDF, all the tested functional polymers show higher intensities for the first peak near 0.2 nm as illustrated in Figure 2.3, indicating greater affinities toward Li+. PEO can wrap around and coordinate with Li+ to form a stable coordinated complex, leading to Li - Polymer Radial Distribution Function

Distribution

150 PVDF PAN PEO PSS PNVF

100

50

0

0

0.1

0.2

0.3 r (nm)

0.4

0.5

0.6

FIGURE 2.3  The radial distribution functions between Li+ and the functional groups of the five tested polymers.

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the highest RDF peak intensity at ∼0.18 nm among all tested polymers. For PSS, the negatively charged sulfonate can also attract Li+ through coulombic interaction, resulting in a strong RDF peak at ∼0.18 nm with the double intensity of the PVDF system. For the PAN systems, the dipole of the PAN nitrile group can also interact with Li+, leading to a strong RDF peak at ∼0.2 nm. Indeed, earlier work by Wang et al. showed that PAN can form a coordination complex with Li+, which is one of the critical effects for the PAN-based polymer electrolyte.22,35,36 The first RDF peak of PNVF has similar intensity to that of PAN. This is due to the strong dipole of the PNVF formamide group, leading to a similar dipole-charge interaction as in the PAN system. Figure 2.4 shows the density profiles of Li+ and polymers along the interface normal (i.e., Z direction), where the cathode interface is located at 1.8 nm. Among all the functional polymers, PEO has the most similar distribution as the PVDF systems. This is due to the oxide group of PEO along the polymer main chains. Hence, the PEO polymer can be oriented on the cathode surface as PVDF. Note that Li+ Distribution 0.03 0.025

PVDF PAN PEO PSS PNVF

0.02 0.015

Number Density (1/ Å3)

0.01 0.005 0

1

2

3

4

5

4

5

Polymer Distribution 0.12 0.1 0.08 0.06 0.04 0.02 0

1

2

3 Z Position (nm)

FIGURE 2.4  The Li+ distribution profiles (top panel) and the polymer distribution profiles (bottom panel) for the five tested polymer systems along the interface normal (i.e., the Z direction). The cathode interface located at 1.8 nm is labeled with the shaded line.

Molecular Effect of Functional Polymer Binders on Li+ Transport

25

PEO can also wrap around Li+ to form stable complexes, leading to a wider distribution at Z > 2 nm. The strong affinity toward Li+ of PEO also can lead to a higher Li+ distribution near the cathode surface at Z  PEO. This suggests the lowest adhesive ability of PEO as a binder.

2.4 SUMMARY AND FUTURE PERSPECTIVES In this chapter, we utilized molecular dynamic simulations to examine the molecular effects of functional polymers on the Li+ transport at the LFP cathode/electrolyte interface. We systematically characterized and compared the thermodynamics and kinetic properties of Li+ on LFP surfaces covered with PVDF, PAN, PEO, PSS, and PNVF. Compared with conventional PVDF binders, all other four functional polymers have higher Li+ affinity, according to the RDF analyses. PAN and PNVF interact with Li+ via the polar side chains of PAN nitrile and PNVF formamide, respectively; PEO forms a stable coordination complex with Li+ using the backbone oxide, and PSS attracts Li+ through its negatively charged sulfonate. These functional polymer binders can alter the electrolyte structure at the cathode/electrolyte interface, leading to various changes in surface electrostatic potentials, Li+ adsorption free energy, and Li+ mobility. In summary, PEO increases surface Li+ density via its stable coordination with Li+, negatively charged PSS can significantly decrease the surface voltage, PAN has the most effect on reducing the Li+ adsorption free energy on the LFP surface, and PNVF greatly enhances the surface Li+ diffusivity. Furthermore, we found that PSS greatly enhances average Li+ drift velocity under an external E-field for both Li+ intercalation and deintercalation at the cathode interface. In contrast, PEO forms a stable complex with Li+, leading to decreased Li+ mobilities at both charging and discharging interfaces. From the cycling E-field test, we found PSS and PNVF exhibit the highest adhesive durability among the tested polymers other than conventional PVDF. Comparing simulation and experimental results, we can correlate the molecular effects of functional binders with the key characteristics of LIB performance. Several experiments have shown that a novel binder using PEO and PAN functional groups can significantly reduce the cell impedance.18,38,39 The interfacial impedance reduction can thus be related to the high Li+ affinity of the polymers and the decreased Li+ adsorption free energy to the cathode surface observed in simulation. The strong Li+ coordination of PEO can be further related to the improved dispersibility of

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active materials.29 A recent study on PEDOT:PSS binder demonstrated an improved cell capacity under high C-rate,40 which can be correlated with the enhanced Li+ drift velocity induced by PSS shown by MD results. And a recent study on PNVF binder has shown that it has a stronger adhesion, a lower internal resistance, and an improved coulombic resistance under high C-rate, which can be related to the simulation results of high adhesive durability, reduced Li+ adsorption free energy, and enhanced Li+ diffusivity for PNVF, respectively. Note, however, experiments often utilize various strategies to incorporate different characteristics from various functional polymers, such as blending mixtures, composite materials, and copolymers. The presented MD study characterized detailed molecular mechanisms of each type of functional polymers, which provides novel insights for new binder design. Future work will focus on exploring the synergistic effects among different polymer types via blending or copolymerization.

ACKNOWLEDGMENTS This work was financially supported by the Hierarchical Green-Energy Materials (Hi-GEM) Research Center, from the Featured Areas Research Center Program within the framework of the Higher Education Sprout Project by the Ministry of Education (MOE) in Taiwan. The authors also thank the Ministry of Science and Technology in Taiwan for supporting this research under grants MOST 106–2923E-006–007 and 107–3113-E-006–006.

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11. Zhang, Z.; Zeng, T.; Lai, Y.; Jia, M.; Li, J. A comparative study of different binders and their effects on electrochemical properties of LiMn 2 O 4 cathode in lithium ion batteries. J. Power Sources 2014, 247, 1–8. 12. Maleki, H.; Deng, G.; Haller, I. K.; Anani, A.; Howard, J. N. Thermal stability studies of binder materials in anodes for lithium-ion batteries. J. Electrochem. Soc. 2000, 147, 4470–4475. 13. Lee, J.-H.; Paik, U.; Hackley, V. A.; Choi, Y.-M. Effect of carboxymethyl cellulose on aqueous processing of natural graphite negative electrodes and their electrochemical performance for lithium batteries. J. Electrochem. Soc. 2005, 152, A1763–A1769. 14. Li, J.; Lewis, R. B.; Dahn, J. R. Sodium carboxymethyl cellulose: A potential binder for Si negative electrodes for Li-ion batteries. Electrochem. Solid-State Lett. 2007, 10, A17–A20. 15. Lux, S. F.; Schappacher, F.; Balducci, A.; Passerini, S.; Winter, M. Low cost, environmentally benign binders for lithium-ion batteries. J. Electrochem. Soc. 2010, 157, A320–A325. 16. Li, J.; Klöpsch, R.; Nowak, S.; Kunze, M.; Winter, M.; Passerini, S. Investigations on cellulose-based high voltage composite cathodes for lithium ion batteries. J. Power Sources 2011, 196, 7687–7691. 17. Li, J.; Le, D.-B.; Ferguson, P. P.; Dahn, J. R. Lithium polyacrylate as a binder for tin– cobalt–carbon negative electrodes in lithium-ion batteries. Electrochim. Acta 2010, 55, 2991–2995. 18. Tsao, C.-H.; Hsu, C.-H.; Kuo, P.-L. Ionic conducting and surface active binder of poly (ethylene oxide)-block-poly(acrylonitrile) for high power lithium-ion battery. Electrochimica Acta 2016, 196, 41–47. 19. Wei, Z.; Xue, L.; Nie, F.; Sheng, J.; Shi, Q.; Zhao, X. Study of sulfonated polyether ether ketone with pendant lithiated fluorinated sulfonic groups as ion conductive binder in lithium-ion batteries. J. Power Sources 2014, 256, 28–31. 20. Shi, Q.; Xue, L.; Wei, Z.; Liu, F.; Du, X.; DesMarteau, D. D. Improvement in LiFePO4– Li battery performance via poly(perfluoroalkylsulfonyl)imide (PFSI) based ionene composite binder. J. Mater. Chem. A 2013, 1, 15016–15021. 21. Chiu, K.-F.; Su, S. H.; Leu, H.-J.; Chen, Y. S. Application of lithiated perfluorosulfonate ionomer binders to enhance high rate capability in LiMn2O4 cathodes for lithium ion batteries. Electrochim. Acta 2014, 117, 134–138. 22. Mindemark, J.; Lacey, M. J.; Bowden, T.; Brandell, D. Beyond PEO—Alternative host materials for Li+-conducting solid polymer electrolytes. Prog. Polym. Sci. 2018, 81, 114–143. 23. Aziz, S. B.; Woo, T. J.; Kadir, M. F. Z.; Ahmed, H. M. A conceptual review on polymer electrolytes and ion transport models. J. Sci.: Adv. Mater. Devices 2018, 3, 1–17. 24. Yue, L.; Ma, J.; Zhang, J.; Zhao, J.; Dong, S.; Liu, Z.; Cui, G.; Chen, L. All solid-state polymer electrolytes for high-performance lithium ion batteries. Energy Storage Mater. 2016, 5, 139–164. 25. Sengodu, P.; Deshmukh, A. D. Conducting polymers and their inorganic composites for advanced Li-ion batteries: A review. RSC Adv. 2015, 5, 42109–42130. 26. Tran, B.; Oladeji, I. O.; Wang, Z.; Calderon, J.; Chai, G.; Atherton, D.; Zhai, L. Thick LiCoO 2/nickel foam cathode prepared by an adhesive and water-soluble PEG-based copolymer binder. J. Electrochem. Soc. 2012, 159, A1928–A1933. 27. Huang, C.-W.; Wu, C.-A.; Hou, S.-S.; Kuo, P.-L.; Hsieh, C.-T.; Teng, H. Gel electrolyte derived from poly(ethylene glycol) blending poly(acrylonitrile) applicable to roll-to-roll assembly of electric double layer capacitors. Adv. Funct. Mater.2012, 22, 4677–4685. 28. Kuo, P.-L.; Wu, C.-A.; Lu, C.-Y.; Tsao, C.-H.; Hsu, C.-H.; Hou, S.-S. High performance of transferring lithium ion for polyacrylonitrile-interpenetrating crosslinked polyoxyethylene network as gel polymer electrolyte. ACS Appl. Mater. Interfaces 2014, 6, 3156–3162.

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29. Tsao, C.-H.; Hsiao, Y.-H.; Hsu, C.-H.; Kuo, P.-L. Stable lithium deposition generated from ceramic-cross-linked gel polymer electrolytes for lithium anode. ACS Appl. Mater. Interfaces 2016, 8, 15216–15224. 30. Mogurampelly, S.; Borodin, O.; Ganesan, V. Computer simulations of ion transport in polymer electrolyte membranes. Annu. Rev. Chem. Biomol. Eng. 2016, 7, 349–371. 31. Maitra, A.; Heuer, A. Cation transport in polymer electrolytes: A microscopic approach. Phys. Rev. Lett. 2007, 98, 227802. 32. Webb, M. A.; Jung, Y.; Pesko, D. M.; Savoie, B. M.; Yamamoto, U.; Coates, G. W.; Balsara, N. P.; Wang, Z.-G.; Miller III, T. F. Systematic computational and experimental investigation of lithium-ion transport mechanisms in polyester-based polymer electrolytes. ACS Cent. Sci. 2015, 1, 198–205. 33. Borodin, O.; Smith, G. D. Mechanism of ion transport in amorphous poly(ethylene ox- ide)/LiTFSI from molecular dynamics simulations. Macromolecules 2006, 39, 1620–1629. 34. Watanabe, M.; Kanba, M.; Nagaoka, K.; Shinohara, I. Ionic conductivity of hybrid films composed of polyacrylonitrile, ethylene carbonate, and LiClO4. J. Polym. Sci. Polym. Phys. Ed. 1983, 21, 939–948. 35. Huang, B. The mechanism of lithium ion transport in polyacrylonitrile-based polymer electrolytes. Solid State Ionics 1996, 91, 279–284. 36. Wang, Z. Spectroscopic investigation of interactions among components and ion transport mechanism in polyacrylonitrile based electrolytes. Solid State Ionics 1999, 121, 141–156. 37. Bouchet, R.; Maria, S.; Meziane, R.; Aboulaich, A.; Lienafa, L.; Bonnet, J.-P.; Phan, T. N. T.; Bertin, D.; Gigmes, D.; Devaux, D.; Denoyel, R.; Armand, M. Single-ion BAB triblock copolymers as highly efficient electrolytes for lithium-metal batteries. Nat. Mater. 2013, 12, 452–457. 38. Gong, L.; Nguyen, M. H. T.; Oh, E.-S. High polar polyacrylonitrile as a potential binder for negative electrodes in lithium ion batteries. Electrochem. Commun. 2013, 29, 45–47. 39. Tsao, C.-H.; Wu, E.-T.; Lee, W.-H.; Chiu, C.-C.; Kuo, P.-L. Fluorinated copolymer functionalized with ethylene oxide as novel water-borne binder for a high-power lithium ion battery: Synthesis, mechanism, and application. ACS Appl. Energy Mater. 2018, 1, 3999–4008. 40. Das, P. R.; Komsiyska, L.; Osters, O.; Wittstock, G. PEDOT: PSS as a functional binder for cathodes in lithium ion batteries. J. Electrochem. Soc. 2015, 162, A674–A678. 41. Okada, H.; Fujikawa, D.; Nodono, M.; Momose, F.; Shimonaka, A.; Itou, M.; Ishii, A.; Momose, H.; Mitsubishi Rayon Co Ltd, Binder for electrode of electrochemical element, composition for electrode of electrochemical element, electrode of electrochemical element and electrochemical element. US Patent 2014O193709A1. 42. Jorgensen, W. L.; Maxwell, D. S.; Tirado-Rives, J. Development and testing of the OPLS all-atom force field on conformational energetics and properties of organic liquids. J. Am. Chem. Soc. 1996, 118, 11225–11236. 43. Sambasivarao, S. V.; Acevedo, O. Development of OPLS-AA force field parameters for 68 unique ionic liquids. J. Chem. Theory Comput. 2009, 5, 1038–1050. 44. Canongia Lopes, J. N.; Deschamps, J.; Pádua, A. A. H. Modeling ionic liquids using a systematic all-atom force field. J. Phys. Chem. B 2004, 108, 2038–2047. 45. Canongia Lopes, J. N.; Pádua, A. A. H. Molecular force field for ionic liquids composed of triflate or bistriflylimide anions. J. Phys. Chem. B 2004, 108, 16893–16898. 46. Doherty, B.; Zhong, X.; Gathiaka, S.; Li, B.; Acevedo, O. Revisiting OPLS force field parameters for ionic liquid simulations. J. Chem. Theory Comput. 2017, 13, 6131–6145. 47. Smith, G. D.; Borodin, O.; Russo, S. P.; Rees, R. J.; Hollenkamp, A. F. A molecular dynamics simulation study of LiFePO4/electrolyte interfaces: Structure and Li+ transport in carbonate and ionic liquid electrolytes. Phys. Chem. Chem. Phys. 2009, 11, 9884.

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48. Martínez, L.; Andrade, R.; Birgin, E. G.; Martínez, J. M. PACKMOL: A package for building initial configurations for molecular dynamics simulations. J. Comput. Chem. 2009, 30, 2157–2164. 49. Plimpton, S. Fast parallel algorithms for short-range molecular dynamics. J. Comput. Phys. 1995, 117, 1–19. 50. Essmann, U.; Perera, L.; Berkowitz, M. L.; Darden, T.; Lee, H.; Pedersen, L. G. A smooth particle mesh Ewald method. J. Chem. Phys. 1995, 103, 8577–8596. 51. Nose, S. A unified formulation of the constant temperature molecular-dynamics methods. J. Chem. Phys. 1984, 81, 511–519. 52. Nose, S. A molecular-dynamics method for simulations in the canonical ensemble. Mol. Phys. 1984, 52, 255–268. 53. Hoover, W. G. Canonical dynamics: Equilibrium phase-space distributions. Phys. Rev. A 1985, 31, 1695–1697. 54. Martyna, G. J.; Klein, M. L.; Tuckerman, M. Nosé–Hoover chains: The canonical ensemble via continuous dynamics. J. Chem. Phys. 1992, 97, 2635–2643. 55. Parrinello, M.; Rahman, A. Polymorphic transitions in single crystals: A new molecular dynamics. J. Appl. Phys. 1981, 52, 7182–7190. 56. Ponce, V.; Galvez-Aranda, D. E.; Seminario, J. M. Analysis of a Li-ion nanobattery with graphite anode using molecular dynamics simulations. J. Phys. Chem. C 2017, 121, 12959–12971. 57. Galvez-Aranda, D. E.; Ponce, V.; Seminario, J. M. Molecular dynamics simulations of the first charge of a Li-ion—Si-anode nanobattery. J. Mol. Model. 2017, 23, 4424. 58. Zhou, F.; Schulten, K. Molecular dynamics study of a membrane-water interface. J. Phys. Chem. 1995, 99, 2194–2207.

3

Essential Properties of Li/ Li+ Graphite-Intercalation Compounds Shih-Yang Lin, Wei-Bang Li, Ngoc Thanh Thuy Tran, Wen-Dung Hsu, Hsin-Yi Liu, and Ming-Fa Lin

CONTENTS 3.1 Introduction...................................................................................................... 37 3.2 The Theoretical Model..................................................................................... 38 3.3 Rich Geometric Structures of Graphites and Graphite-Intercalation Compounds......................................................................................................40 3.4 Unusual Band Structures of Graphite-Related Systems..................................44 3.5 Van Hove Singularities in Density of states..................................................... 53 3.6 Chemical Bondings and Charge Distributions................................................. 56 3.7 Summary.......................................................................................................... 57 References................................................................................................................. 59

3.1 INTRODUCTION It is well known that graphite is one of the most investigated materials both theoretically and experimentally. Up to now, it has served as the best anode in the Li+-band battery. This layered system is purely composed of the hexagonal-symmetry carbon layers, in which the weak but significant van der Walls interactions would greatly modify the low-lying π-electronic structure and thus dominate the essential physical properties. Monolayer graphene is identified to be a zero-gap semiconductor, while a 3D graphite belongs to a semimetal. The electronic properties strongly depend on the way the graphitic planes are stacked on each other. In general, there are three kinds of stacking configurations: AAA (simple hexagonal), ABAB (Bernal), and ABCABC (rhombohedral). The total free carrier density is predicted to be 3.5 × 1020 e/cm3 in simple hexagonal graphite and ∼1019 in Bernal graphite at room temperature. When various atoms and molecules are further intercalated into the AB-stacked graphite, many graphite-intercalation compounds are formed. When many free conductions (holes) are induced after the intercalation, such systems exhibit the donor-type (acceptor-type) behaviors. Among these compounds, only the stage n lithium intercalation systems display the AAA stacking configuration, as confirmed from the X-ray diffraction patterns. Here, n clearly indicates the number of graphitic sheets between two periodical guest-atom layers, in which n = 1, 2, 3, and 4 (Figure 3.8c–f), being 37

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arranged from the highest concentration to the lower one, will be studied thoroughly in terms of the essential properties. It is also noticed that the other alkali intercalation compounds present the MC8 structure in the stage-1 configuration (M = K, Rb; Cs). Obviously, both LiC6 and MC8 should have very different band structures and fundamental properties. Up to now, there has been a lot of theoretical and experimental research on alkaliatom graphite-intercalation compounds but few studies about alkali-ion ones. As to the former, the first-principle method1–4 has been utilized to investigate the total ground-state energies, optimal guest-atom distributions, interlayer distances, stacking configuration between neighboring graphitic layers, intrinsic and atom-doped electronic energy spectra, and density of states, transport properties, and phonon spectra. Such systems cover LiCx (x = 6, 12, 18; 24) and MCx (x = 8, 16, 24; 32). Specifically, the thermodynamic and kinetic properties of lithium atoms in graphite-intercalation compounds are thoroughly explored by the first-principle calculations1–3 and Monte Carlo method.2 The Li-atom differences in graphite are classified into three types: (I) vacancy, (II) interstitial, and (III) interstitial mechanisms. The similar numerical investigations are applied to the Li+-ion difference on the graphene layers,5,6 in which their results cover the optimal height and site of guest atoms, the absorption energy, and energy barrier along the specific transport path. Moreover, the tight-bind model and the superlattice model are frequently utilized to study the fundamental properties for pristine graphites and the atom/molecule-intercalated graphite compounds (e.g., optical absorption spectra, coulomb excitations, and magnetic properties). As to the various experimental measurements on graphite-related systems, they include the X-ray/TEM/STM/LEED patterns, angle-resolved photoemission spectra, optical reflectance/absorption/transmission spectra, electrical conductivities, magnetic properties, phonon energy spectra, electron energy loss spectra, and femtosecond excited carrier dynamics. Specifically, the planar Li-atom distribution configurations of stage-1, stage-2, and stage-3 compounds, as well as their periodical interlayer distances, are examined by the X-ray diffraction spectra. The intercalation-induced high conduction-electron densities are also observed in the optical and transport measurements. The rich and unique phenomena in alkali-atom graphite-intercalation compounds are expected to have the significant differences under a systematic comparison with those in alkali-ion cases.

3.2 THE THEORETICAL MODEL A typical condensed-matter system is made of a periodic crystal potential, where each atom would contribute several valence electrons around an ionic core. Obviously, each one exhibits the complex composite effects mainly arising in the electron– electron Coulomb interactions and the electron–ion crystal potential, especially for the many-body effects. This becomes a high barrier in solving the many-particle Schrodinger equation. The difficulty of numerical calculations is greatly enhanced when the various chemical environments need to be taken into consideration (e.g., the very weak van der Waals interactions between two neighboring graphitic layers and the chemisorptions on graphite surface). Some approximate methods have been proposed to achieve reliable geometric structures and electronic properties. Up

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to now, the first-principle calculations have been frequently utilized to obtain the quantum states of periodic systems. Such numerical calculations are very efficient for fully exploring the fundamental physical properties. Specifically, the Vienna ab initio simulation package (VASP)7 evaluates an approximate solution within the density functional theory by solving the so-called Kohn–Sham equations.8 The charge distribution can determine all the intrinsic interactions in condensed-matter systems; that is, the carrier density is responsible for the ground-state energy and the essential properties. The spatial charge density could be solved by the numerical self-consistent scheme, as clearly shown in a flowchart of detailed evaluations [Figure 3.1]. Compared with the first-principle method, the tight-binding model, with the hopping integrals (the parameters), cannot study the optimal geometric structures and thus identify the complicated orbital hybridizations in various chemical bonds. By a detailed comparison of these two modes in the low-lying energy bands, the reliable parameters of the latter are thus obtained (e.g., the vertical and nonvertical atomic interactions in AA-, AB-, and ABC-stacked graphites). And then, they

FIGURE 3.1  The flowchart of VASP calculations.

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are very useful in understanding other essential properties (e.g., optical properties and coulomb excitations in three kinds of graphites). In addition, the generalized tight-binding model, but not the first-principles method, is suitable for studying the essential properties under the external fields (e.g., the quantized Landau levels in uniform magnetic fields). In this chapter, the optimal geometric structures and electronic properties and magnetic configurations are thoroughly studied for 3D graphite-related systems by utilizing VASP.7 Such systems cover monolayer graphene, Bernal graphite, simple hexagonal graphite, and stage-1 to stage-4 Li/Li+ graphite-intercalation compounds. The Perdew–Burke–Ernzerhof function within the generalized gradient approximation can deal with the many-particle coulomb interactions, the exchange and correlation energies of valence, and conduction electrons.9 The projector-augmented wave pseudo-potentials characterize electron–ion interactions.10 It should be noticed that to correctly describe the weak but significant atomic interactions between the neighboring graphitic layers, the van der Waals forces must be included in the calculations by the semiempirical DFT-D2 correction of Grimme.11 When one solves the many-body Schrodinger equation, plane waves, with a maximum energy cutoff of 400 eV, consist of a complete set in building the Bloch wave functions. For charged system calculations, monopole, dipole, and quadrupole corrections in all directions are calculated. The 3D periodic boundary condition is along , in which a primitive unit cell depends on the geometric distribution of host and guest atoms. The Brillouin zone is sampled by 30 × 30 × 30 and 70 × 70 × 70 k point meshes within the Monkhorst–Pack scheme, respectively, corresponding to the numerical calculations of geometric optimizations and electronic structures. The energy convergence is set to be 10 −5 eV for two neighboring simulation steps; furthermore, the maximum Hellmann–Feynman force acting on each atom is less than 0.01 eV/Å during the process of ionic relaxations. The details of the optimal calculation processes are revealed in Figure 3.1.

3.3 RICH GEOMETRIC STRUCTURES OF GRAPHITES AND GRAPHITE-INTERCALATION COMPOUNDS Apparently, the pristine graphites and Li/Li+ graphite-intercalation compounds display the rich geometric structures by the VASP calculations on the total ground-state energies. The periodical graphitic layers in the AA stacking are bound together by the weak van der Waals interactions due to the perpendicular 2pz orbitals on two neighboring planes, so it needs to have the delicate numerical evaluations on the geometry-dependent total ground-state energies. These intrinsic interactions can create the vertical and nonvertical hopping integrals in the tight-binding model and thus dominate the low-energy electronic properties. Of course, the very strong σ bondings also exist on each graphitic layer. They are only slightly changed during variation from a 2D graphene to a 3D graphite, as confirmed by the C–C bond lengths of them (1.420 Å and 1.423 Å, respectively). By the detailed calculations in Figure 3.2, simple hexagonal graphite is identified to present an optimal interlayer distance of Iz = 3.489 Å, with the lowest ground-state energy of −0.238 eV. This system possesses the largest interlayer distance and the highest ground-state energy,

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FIGURE 3.2  The dependence of the total ground-state energy on the interlayer distance for (a) a simple graphite, (b) LiC6, and (c) Li+C6.

compared with those ([3.35 Å, −0.6 eV] and [3.40 Å, −0.4 eV]) of Bernal graphite and rhombohedral graphite. Such significant differences are responsible for the fact that a natural graphite is made up of AB and ABC stackings, especially for the former configuration. The distinct stacking symmetries would dominate the interlayer atomic interactions and greatly diversify the fundamental chemical and physical properties. The geometric symmetries are greatly diversified by the chemical intercalation. The Li/Li+ could be easily intercalated into graphitic layers so that such guest atoms/ ions create a periodical planar structure. According to their concentrations, there exist the stage-dependent graphite-intercalation compounds, being clearly identified in the experimental syntheses.12–16 Three types of absorption positions, hollow, top, and bridge ones, are frequently observed in the engineering of chemical modifications.17 Among them, the hollow-site positions are examined to have the lowest ground states. That is to say, such positions present the strongest chemical bondings/ interactions. Similar results are revealed in other alkali-atom/ion intercalations.18,19

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By detailed calculations, the height-dependent (interlayer-distance-dependent) ground-state energies show the optimal interlayer distance with the lowest one. For example, the stage-1 LiC6/Li+C6 has the interlayer distance of 3.728 Å/3.059 Å under the ground-state energy of −58.689 eV/−65.417 eV. Obviously, the ground-state energy decreases as the stage number of n grows, in which the declining concentration of Li–C/Li+–C bonds is the main reason. Compared with the original interlayer distance of two neighboring graphitic layers 3.489 Å in Figure 3.2a, those with and without the Li/Li+ intercalations are, respectively, enhanced and reduced, except for the Li+C6 case (e.g., ∼3.728−3.768 Å/3.059−3.539 Å and ∼3.323−3.364 Å/2.918−3.348 Å under the stage-1 to stage-4 configurations). A simple relation between n and the interlayer distances might be absent. However, a detailed comparison between the Li and Li+ intercalation cases clearly shows that the latter always have the shorter interlayer distances corresponding to the lower total ground-state energies. Apparently, such significant results suggest the stronger intrinsic interactions in Li+ intercalation compounds and more stable ion status; that is, they might account for the rapid charging and discharging closely related to the intercalation and deintercalation processes, respectively. As to the C–C bond lengths, they are only slightly changed by the Li-atom and Li+-ion intercalations. The stage-1, stage-2, stage-3, and stage-4 Li (Li+) compounds, respectively, have 1.443 Å, 1.429 Å, 1.428 Å, and 1.426 Å (1.414 Å, 1.414 Å, 1.416 Å, and 1.418 Å). Their enhancement and decrease in the atomic and ionic cases suggest charge transfers of Li → C and C → Li+. Charge transfer (fs) stacking configurations (AA, AB, ABC) are worthy of a closer examination. By the Bader analyses in the VASP calculations, the stage-1, stage-2, stage-3, and stage-4 Li-intercalated (Li+-based) graphite compounds, respectively, exhibit the charge transfers of 0.855, 0.857, 0.873, and 0.876 (0.183, 0.138, 0.117, and 0.132), being almost independent of the number of stage n. The 2s orbitals of Li atoms are mostly transferred to the neighboring C atoms, while they do not present the 100% transfer. The 3D free electron density is deduced to be inversely proportional to n; that is, f could be utilized to evaluate the conduction electron density. Also, it could be estimated from the area covered by the curve of density-of-state versus energy (discussed later). The 2s orbitals of Li atoms are mostly transferred to the neighboring C atoms, while they do not present the 100% transfer. To easily compare all the calculated results, the stacking configuration is assumed to be the sequence of AA. However, the very low guest-atom/ion concentration or the pristine case could belong to the AB stacking, as observed in natural graphite. How to transform from AA to AB stackings during the decline of guest atoms or ions could be an interesting focus of future studies. Scanning tunneling microscopy (STM) is the most powerful experimental technique for resolving surface structure, being able to characterize surface topographies in real space with both lateral and vertical atomic resolution (e.g., the nanoscaled bond lengths, crystal orientations, planar/nonplanar structures, step edges, local vacancies, dislocations, atomic or molecular adsorptions, and nanoclusters). The upto-date STM measurements on the graphite-related systems have confirmed the complex relations among the hexagonal lattice, the finite-size confinement, the flexible feature, and the diverse chemical bondings of carbon atoms. For example, graphene nanoribbons exhibit the nanoscale-width planar honeycomb lattice, accompanied

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with the achiral (armchair and zigzag) or chiral edge structures.20–22 Furthermore, they could also be formed in the curved,21 folded, 23 scrolled,24 and stacked lattice structures.25 Carbon nanotubes possess achiral or chiral arrangements of hexagons on a cylindrical surface.26,27 The atomic-scale measurements directly identify the AB,28,29 ABC,30,31 and AAB32 stacking configurations in few-layer graphene systems, the corrugated substrate and buff graphene layer,33,34 the rippled structures of graphene islands,35–37 and the adatom distributions on graphene surfaces.38,39 Layered graphite could exhibit 2D networks of local defects on the surface: the pyridinicnitrogen and graphitic-N structures.40 Apparently, the high-resolution experimental measurements of STM could be utilized to verify the unique surface structures of Li-atom and Li+-ion graphite-intercalation compounds (e.g., the C–C bonds affected by the guest-atom or guest-ion intercalation, the optimal adsorption site, and the planar periodical distribution of host atoms and guest atoms/ions). Transmission electron microscopy (TEM) is a microscopy technique in which an electron beam, with a uniform current density, is transmitted through an ultrathin specimen to create an image, as a result of the interactions between incident charges and sample. TEM is a very important experimental technique in directly visualizing the crystal structure, locating and identifying the type of defects, and studying structural phase transitions. This measurement has a rather big atomic scattering factor, being ∼10,000 times of that from the X-ray diffraction. TEM provides electron diffraction an advantage to observe even the weakest diffracted spot. However, its resolution is seriously limited by spherical and chromatic aberrations of the lenses. More delicate techniques for improving the diffraction resolution become indispensable. By applying a monochromator and a Cs corrector into TEM, which is called the high-resolution TEM (HR-TEM), the structural resolution can reach less than 0.5 Å. HR-TEM has been successfully and extensively utilized to analyze crystal structures and lattice imperfections in various nanomaterials. The TEM/HR-TEM measurements on graphene-related systems are very suitable in identifying the sp2-bonding-enriched nanoscale structures, such as the multiwalled cylindrical structures of carbon nanotubes41,42; the curved,42–44 folded,45,46 and scrolled profiles24,47,48 of graphene nanoribbons; and the stacking configurations and the interlayer distances of few-layer graphene systems.49–51 Obviously, the high-resolution TEM measurements are suitable in examining the rich geometric properties of pristine (AA and ABC) and Li- and Li+-intercalated graphites, such as the periodical distance along the z-direction, the distances of layered graphene in the high-n systems, the stacking configuration of graphitic layers, and the optimal distribution configuration and position of guest atoms or ions. It is well known that X-ray diffraction techniques could provide the most powerful tools in exploring the crystal symmetries, especially for the 3D condensed systems.12–16 Specifically, such experimental measurements are very suitable for the stage-n structures in Li-/Li+-graphite-intercalation compounds (e.g., the lattice constants, interlayer distances, and order/disorder transformations). Up to now, X-ray diffraction patterns have been utilized to confirm the existence of stage-1,12–14,16 stage-2,12,14,16 and stage-315 Li-graphite-intercalation compounds. The optimal planar structure belongs to √3 × √3 R30° with a lattice constant ∼4.305 Å, as observed from both LiC6 and LiC12.12,14,16 The periodical interlayer distances in stage-1 and stage-2 are, respectively,

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identified to be ∼3.705−3.706 Å and ∼7.025−7.065 Å. Furthermore, the graphitic layers of them present the AA stacking along the z-direction. However, the stacking ordering of ABAαABAα... is examined to survive in the stage-3 system (LiC18).15 The highresolution X-ray measurements are required to thoroughly examine the stacking configurations of the larger-n Li-graphite-intercalation compounds. The similarities and important differences between the stage-n Li-atom and Li+-ion systems in the optimal geometric structures are worthy of the detailed examinations using the X-ray diffraction spectra, especially for the experimental measurements on the latter. They could reveal the important information about the available electronic configurations associated with the neutral atoms and full ions in the anode of Li+-based battery.

3.4 UNUSUAL BAND STRUCTURES OF GRAPHITE-RELATED SYSTEMS Pristine graphite possesses an unusual electronic structure with a large or small overlap of valence and conduction bands, being sensitive to the stacking configuration, for example, the AA, AB, and ABC stackings associated with Figure 3.3a–c, respectively. It is well known that monolayer graphene is a zero-gap semiconductor.52 There exist the linearly intersecting valence and conduction bands at the Dirac point (not shown), so the density of states at the Fermi level (the band overlap) is vanishing. Only the finite temperature can induce some free carrier density.53 The low-energy electronic states are initiated from two equivalent valleys. However, the significant interlayer atomic interactions lead to the drastic changes in the low-lying energy bands. All the 3D graphitic systems and the occupied valence bands are asymmetric to the unoccupied conduction bands about the Fermi level, mainly owing to the significant interlayer hopping integrals. Among three kinds of graphites, AA-stacked graphite has a strong vertical interlayer hopping integral.54 The Dirac-point energies, corresponding to the corner states in the first Brillouin zone (Figure 3.2c), strongly depends on the kz-components of 3D wave vectors. As a result, an obvious band overlap is revealed in the kz-dependent energy dispersion along ΓA. The periodical arrangement of graphitic layers in the z-direction can create band width as wide as ∼1 eV. Furthermore, on the (k x, k y) plane, an electron (hole) pocket comes to exist near the A (Γ) point. That is to say, the free electrons and holes, which appear in the density of states (Figure 3.4a), are, respectively, located in the range of 0 ≤ E ≤ 0.5 eV and −0.5 eV≤ E ≤ 0. The abovementioned low-lying energy bands, with |Ec,v | ≤ 3 eV, mainly come from the π bondings of the C–2pz orbitals. Such electronic states are responsible for most of the essential physical properties. Specifically, the parabolic energy dispersions belong to the saddle points at the middle points of the first Brillouin zone, for example, the electronic energy spectra at the M and L points [2.5 eV and 1 eV] and [−2 eV and −3 eV], as shown in Figure 3.3a,d. They possess a very high density of states (discussed later in Figures 3.5–3.7), so their features are useful in distinguishing the diversified intercalation effects. As to the deeper-energy electronic states (Ev ≤ −3 eV), part of them originates the σ bondings of the (2px, 2py , 2s) orbitals, especially for those near the Γ and A points (∼−3 eV). In addition, the initial σ-electronic states just have the same energy with the π-electronic saddle point of L. They will exhibit splitting behavior after intercalation.

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FIGURE 3.3  Electronic structures for the (a) AA-stacked graphites, (b) AB-stacked pristine graphites, and (c) ABC-stacked graphites. Also shown in (d) is that of the first system under the enlarged unit cell with six carbon atoms, (e) corresponding to the original and reduced first Brillouin zones.

There are certain important differences among typical kinds of graphites. The linear Dirac-cone structures on the (k x, k y) planes are clearly revealed in the AA-stacked (Figure 3.8a) and ABC-stacked graphites (Figure 3.8c). For simple hexagonal graphite, such unusual energy bands are just initiated from the corners of the original/reduced first Brillouin zone. Furthermore, rhombohedral graphite exhibits a novel spiral configuration along the kz-direction (a slight deviation from the planar corner), mainly owing to the specific stacking symmetry.55 In general, the energy of Dirac points is very sensitive to the kz component of the wave vector. On the other side, a Bernal graphite presents the bilayer- and monolayer-like energy dispersions

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FIGURE 3.4  The spatial charge distributions before/after Li intercalation: (a) a simple hexagonal graphite, (b, c) LiC6, (d, e) LiC12, (f, g) LiC18, and (h, i) LiC24.

(the parabolic and linear ones), as indicated in Figure 3.8b. This result is consistent with the previous theoretical calculations and the experimental ARPES measurements. In terms of the kz-dependent π-electronic band widths related to the 2pz orbitals (Figure 3.8a–c), they are, respectively, about 1.0 eV, 0.2 eV, and 0.03 eV for simple hexagonal, Bernal, and rhombohedral graphites, such as those near the Fermi level along KH and at the middle energy along ML (±2 eV). Obviously, the first and third systems have the largest and smallest valence and conduction overlaps, respectively, and so do the free electron/hole density purely due to the interlayer atomic interactions. The abovementioned results suggest that the interlayer hopping integrals of the tight-binding model are most complicated for ABC-stacked graphite when the critical parameters are from the first-principle calculations. In short, the AA-, AB-, and ABC-stacked graphites belong to semimetals, while the latter two only possess very few 3D free carrier densities at low temperatures. These systems are thus

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FIGURE 3.5  Similar plots as Figure 3.4b–i but indicated for (a, b) Li+C6, (c, d) Li+C12, (e, f) Li+C18, and (g, h) Li+C24.

expected to display diverse physical and chemical properties (e.g., unusual magnetic quantization phenomena, magneto-optical properties, and many-particle coulomb excitations). Electronic structures exhibit drastic changes after the intercalation of Li atoms, as clearly illustrated in Figure 3.9a–d for stage-1 to stage-4 graphite-intercalation compounds, respectively. For stage-1 LiC6, the asymmetry of electron and hole bands, being about the Fermi level, becomes more serious, as identified from a detailed comparison with the pristine case (Figures 3.3d and 3.9a). Apparently, the main mechanisms originate from more orbital hybridizations due to the Li–C bonds (Figure 3.8c). Apparently, the Fermi level presents the blue-shift phenomenon, in which it is situated at certain conduction bands and does not intersect with any valence bands. The free conduction electrons are purely induced by the Li-atom intercalation; furthermore, the free holes related to the unoccupied valence states thoroughly disappear. That is, two kinds of free carriers arising from the interlayer atomic interactions are fully replaced by the conduction ones. It should be noticed that the energy bands just below the Fermi level are split along the ΓA direction, and the originate double

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Green Energy Materials Handbook 1.2 Total 2s 2px 2py 2pz

(a) AA

1 0.8 0.6 0.4 0.2 0 –4

–3

–2

–1

0

1

2

–1

0

1

2

–1

0

1

2

1

DOS (states/eV.cell)

(b) AB 0.8 0.6 0.4 0.2 0 –4

–3

–2

1.2 (c) ABC

1 0.8 0.6 0.4 0.2 0 –4

–3

–2

E (eV)

FIGURE 3.6  The carbon-orbital-projected DOSs: (a) the AA-, (b) AB-, and (c) ABCstacked graphites.

degeneracy, as shown in Figure 3.3d, is only created by the zone-folding effect with an enlarged unit cell with six carbon atoms. The upper and lower energy bands, respectively, represent the conduction and valence states. The maximum and minimum values of the valence band (indicated by the red dashed curves) are −0.6 eV and −2.01 eV (Γ and A points). The latter, the maximum energy range between the Fermi

49

Properties of Li/Li+ Graphite-Intercalation Compounds 1.2 Total 2s 2px 2py 2pz Li-2s

1 0.8 0.6

1.8 1.6 1.4 1.2 1 0.8 0.6 0.4 0.2 0

(a) LiC6

DOS (states/eV.cell)

0.4 0.2 0 –4

–3

–2

–1

0

2.4

1

2

(c) LiC18

1.6 1.2 0.8 0.4 0 –3

–2

–1

E (eV)

0

1

–4

–3

–2

–1

0

4 3.5 3 2.5 2 1.5 1 0.5 0

2

–4

(b) LiC12

2

1

2

(d) LiC24

–4

–3

–2

–1

0

1

2

E (eV)

FIGURE 3.7  The C- and Li-orbital-decomposed DOSs for (a) LiC6, (b) LiC12, (c) LiC18, and (d) LiC24.

level and the top of valence state, plays a critical role in determining the blue-shift value of the Fermi level (i.e., it will determine the Li-intercalation-induced 3D conduction electron density). Moreover, the low-lying energy bands do not present the crossing and anticrossing behaviors between any two energy subbands. Specifically, the π-electronic energy spectra at the saddle points of L and M are not identical to the initial σ-electronic ones along ΓA. This will lead to the separation of van Hove singularities. Band structures are getting more complex as the number of stage n grows. Concerning stage-2 to stage-4 compounds (Figure 3.9b–d), there are more energy subbands as a result of the enlarged unit cell. However, the scale of the first Brillouin zone is greatly reduced. The crossing and anticrossing behaviors come to exist frequently. Obviously, the effective blue shift of the Fermi level, the maximum energy range of EF, and the top valence state at the A point, which is investigated later by a more accurate method for density of states (Figure 3.6a–d), present the declining behaviors. The stage-2 LiC12 has two valence subbands slightly crossing the Fermi level near the Γ point (Figure 3.9b). The similar phenomenon becomes obvious and even appears at the A point for the LiC18 and LiC24 compounds (Figure 3.9c,d). Apparently, the 3D conduction electron density decreases in the increment of n. Moreover, the π-electronic parabolic dispersions, being near the M and L, exhibit the obvious splitting behaviors. There exist (one, two, three, four) valence and conduction saddle points, respectively, in the stage-1, stage-2, stage-3, and stage-4 compounds. They represent the most important feature in distinguishing the diverse intercalation effects. The n-enriched electronic structures lie in the drastic changes of the guest-atom concentration and their distribution configuration.

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FIGURE 3.8  Geometric structures: (a) a simple hexagonal graphite, (b) top views of a pristine system and LiC6, (c) LiC6/Li+C6 (stage-1 lithium graphite-intercalation compound), (d) LiC12/Li+C12 (stage-2), (e) LiC18/Li+C18 (stage-3), and (f) LiC24/Li+C24 (stage-4).

Li+-ion intercalation can greatly diversify electronic structures, as shown by a detailed comparison with the pristine and guest-atom cases. For example, Li + C6, LiC6, and AA-stacked graphite (Figures 3.10a, 3.9a, and 3.3d) are rather different from one another in the electronic properties, especially for the low-lying valence and conduction bands. The stage-1 Li+C6 exhibits two splitting energy subbands along ΓA near the Fermi level. Furthermore, the conduction and valence ones are partially occupied and unoccupied, respectively, clearly illustrated by the energy dispersions along AH and ΓM. However, the valence bands are fully occupied in the stage-1 LiC6 (Figure 3.9a). Obviously, the splitting behavior indicates the separation and distortion of the linear Dirac-cone structure (Figures 3.10a and 3.3d), for example, the energy spacing of ∼0.45−0.9 eV in the separated valence and conduction bands. The linearly gapless Dirac cones, being revealed in simple hexagonal graphite, might be vanishing through the Li+ intercalation. According to various band structures near the Fermi level (Figure 3.10a–d), the stage-n graphite Li+-ion intercalation compounds are predicted to have the same free electron and hole densities,

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FIGURE 3.9  Valence and conduction bands of (a) LiC6, (b) LiC12, (c) LiC18, and (d) LiC24.

being similar to those in the AA-stacked graphite. That is, they do not possess the high-density conduction electrons, or the charge transfer almost disappears during the intercalation/deintercalation of guest Li+ ions. The Fermi level almost remains at the middle between the unoccupied valence states and the occupied conduction ones. Moreover, energy bands become more complicated with the increasing stage n. On the experimental side, ARPES is best for exploring the quasi-particle energies of the occupied electronic states within the Brillouin zone,56 and the measured dispersion relations can directly examine those evaluated from the first-principle method tight-binding model. The ARPES chamber is associated with the instrument of sample synthesis to identify the in situ band structures. Up to now, the measured results have confirmed the feature-rich energy bands in the carbon-related systems,

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FIGURE 3.10  Similar plots as Figure 3.9a–d, but shown for (a) Li+C6, (b) Li+C12, (c) Li+C18, and (d) Li+C24.

as verified under the distinct dimensions,20,57–61 layer numbers,57,62–64 stacking configurations,65 substrates,57,58,62 and adatom/molecule chemisorptions.66–68 Graphene nanoribbons are identified to exhibit 1D parabolic energy subbands centered at the high-symmetry point, accompanied by an energy gap and different subband spacings.20 Recently, plenty of ARPES measurements have been done for few-layer graphenes, obviously showing the linear Dirac-cone structure of the monolayer system,57–59 two pairs of parabolic bands in bilayer AB stacking,57,65 the coexistent linear and parabolic dispersions in symmetry-broken bilayer system,63 the monolayer- and bilayer-like energy bands in trilayer ABA stacking,57,62 the linear and partially flat and sombrero-shaped energy bands of trilayer ABC stacking,62 the substrate-induced

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large energy spacing between the π and π* bands in bilayer AB stacking,57,62 the substrate-created oscillatory bands in few-layer ABC stacking,57 and the metal-semiconductor transition and the tunable low-lying energy bands after the adatom/molecule chemisorptions on the graphene surface.66–68 It should be noticed that the high-technique ARPES is required to examine the 3D band structure of the AB-stacked natural graphite,60,61 since the conservation of the transferred momentum is destroyed along the kz-direction. The experimental measurements clearly illustrate that the 3D energy bands the bilayer- and monolayer-like energy dispersions at kz = 0 and 1 (the K and H in the first Brillouin zone), respectively; furthermore, there exists the strong trigonal warping effect around the KH axis and a weakly dispersive band near the Fermi energy due to the stepped surface with a zigzag edge. The high-resolution ARPES is required to examine the theoretical predictions on the occupied valence and conduction states of graphite-related systems, covering electronic energy spectra of simple hexagonal graphite (Figure 3.3a), rhombohedral graphite (Figure 3.3c), and stage-n Li-atom (Figure 3.9a–d) and Li+-ion graphite intercalation (Figure 3.10a–d). Such experimental measurements could provide the full information on the important effects due to the stacking configuration, the significant Li–C chemical bonding (charge transfer), and the zone folding (the lower symmetry of Li+-ion distribution). They would identify the critical electronic structures below the Fermi level. The detailed emanations, which are conducted on the kz-dependent band widths and linear Dirac-cone structures, are sufficient in distinguishing the AA and ABC stacking configurations. The experimental identifications on the large blue shift of the Fermi level are obvious evidence about the Li-atom intercalation effect (the strong charge transfer from Li to C). A declining relation between EF and stage n is reflected in the red shift of the lower conduction subbands (the higher valence ones). However, the Fermi level is almost identical for the pristine and Li+-intercalation cases. The latter exhibit the splitting of energy subbands along the ΓA direction, in which this behavior becomes more serious in the increment of n because of the reduced distribution symmetry. Another important focus of Li/Li+ intercalation effects covers the n-split π-electronic parabolic dispersions near the saddle points of M and L and their state energies separated from the initial σ-electronic ones.

3.5 VAN HOVE SINGULARITIES IN DENSITY OF STATES The van Hove singularities in 3D graphite-related systems are revealed as the special structures, being quite different from those in layered graphene systems. The atomand orbital-decomposed density of states, as shown in Figures 3.6, 3.7, and 3.11, could provide the full information and be very useful in examining the atom/ion intercalation effects. Their structures principally originate from the critical points in the energy-wave-vector space, in which such points might form a 1D space with a continued kz-dependence. For simple hexagonal graphite (Figure 3.6a), density of states is asymmetric about the Fermi level because of the interlayer atomic interactions. It belongs to a semimetal, since density of states (DOS), a finite but minimum value at E = 0, being situated in the ranges of 0  2GLi). Khurana et al. proved that a high-modulus SPE is not the only factor for the suppression of Li dendrite growth.60 Zeng et al. developed an interpenetrating network of poly(ether acrylate) (inp-PEA) through the photopolymerization of PEO and branched acrylate.61 This inp-PEA SPE demonstrated both a high ionic conductivity of 0.22 mS cm−1 at RT and an exceptional mechanical strength of 12 GPa. Because of its high mechanical strength, this network could effectively suppress Li dendrite growth. Pan et al. developed a cross-linked inorganic polyhedral oligomeric silsesquioxane (POSS) and PEO to inhibit lithium dendrite growth.62 The small particle size of POSS resulted in the uniform formation of SPE, and POSS provided the mechanical strength required to resist the lithium dendrite growth of SPE. In addition to PEO-based SPEs, non-PEO polymers, including poly(acrylonitrile) (PAN), poly(vinylidene fluoride) (PVDF), PVDF-HFP, poly(propylene oxide), and poly(methyl methacrylate) (PMMA), have received attention for use in LIBs. 11.2.3.2 SPEs with Inorganic Fillers The addition of inorganic fillers into the polymer network results in high ionic conductivity, high mechanical strength, a high transference number, and an enlarged electrochemical window. As mentioned, crystallization is the primary concern for low ionic conductivity in PEO-based SPEs. Adding fillers to an SPE may transform the local crystalline structure of the polymer to an amorphous state and increase the chain mobility of polymers. Thus, SPEs with fillers exhibit higher ionic conductivity than SPEs without fillers. Inorganic fillers with a small particle size demonstrate high ionic conductivity. Inorganic fillers are classified into two types: passive fillers and active fillers. Passive fillers (Al2O3, SiO2, TiO2, and ZrO2) are nonconductive and not involved in ion conduction, whereas active fillers (Li6.4La3Zr1.4Ta0.6O12 and Li1 + xAlxTi2–x(PO4)3) are conductive and are involved in Li ion conduction. Cui and coworkers synthesized PEO SPE with monodispersed nonconductive ultrasmall SiO2 nanospheres, which exhibit a high ionic conductivity of 4.4 × 10 −5 S cm−1 at 30°C and 1.2 × 10 −3 S cm−1 at 60°C as well as a wide electrochemical window (5.5 V vs Li+/Li).63 The ionic conductivity of SPEs with active fillers is one to two orders of magnitude higher than that of SPEs with passive fillers. Thus, active fillers attract more attention in SPEs for Li metal-based batteries. Lee et al. produced a composite SPE

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with tetragonal Li7La3Zr2O12 incorporated into the PEO:LiClO4 matrix that exhibited an ionic conductivity of 4.4 × 10 −4 S cm−1 at 55°C.64 Li et al. produced a composite SPE with Li6.4La3Zr1.4Ta0.6O12 (LLZTO) particles and a Li salt–free PEO matrix, revealing a high ionic conductivity of 2.1 × 10 −4 S cm−1 at 30°C and 5.6 × 10 −4 S cm−1 at 60°C.65 The electrochemical window of composite SPE is greater than 4.7 V, and the Li transference number is 0.46. In contrast to isolated nanoparticles, such as nanofillers, nanowire or nanofiber can offer extended Li ion transport pathways in the polymer matrix. Chan et al. developed composite SPEs with LLZO nanowires as passive fillers in a PAN-based polymer matrix with ionic conductivity in the order of 10 −4 S cm−1.66 Yang et al. produced composite SPEs with vertically aligned and connected NASICON-type Li1 + xAlxTi2−x(PO4)3 (LATP) nanoparticles (NPs) in the PEO matrix.67 This vertical alignment of composite SPEs achieved an ionic conductivity 0.52 × 10 −4 S cm−1, which is 3.6 times higher than that of the SPE with randomly dispersed LATP NPs. However, the agglomeration of individual isolated fillers is the main problem that arises when using isolated NPs as fillers in composite SPEs. To avoid this problem, Hu et al. created 3D garnet nanofiber networks to establish continuous Li ion transfer channels in PEO-based composite electrolytes.68 This flexible 3D composite electrolyte exhibited an ionic conductivity of 2.5 × 10 −4 S cm−1 at RT. 3D garnet structure consists of interconnected nanofiber networks to provide continuous ion-conducting pathways. A garnet LLZO nanofiber network was prepared through the electrospinning method and then sintered at 800°C. The PEO–Li salt mixture then infiltrated the 3D garnet nanofiber networks. This composite electrolyte demonstrated excellent lithium stripping and plating performance at a current density of 0.5 mA cm−2 over the period of 300 h. 11.2.3.3 Single-Ion Conductors Single-ion conductors differ from dual-ion conductors, in which anions are fixed to the polymer backbone through covalent bonding with neutral molecules while cations are transported freely through a hopping process because the delocalization of anions results in weak electrostatic force. The immobilized anions cause the cation transference number of the single-ion conductor SPE to approach unity, similar to the case with ISEs. Polymer electrolytes usually have a low lithium transference number as a result of dual-ion movement. The single-ion conductor, which can approach a transference number of one, has become the focus of research for developing future lithium metalbased batteries. Bouchet et al. developed single-ion BAB triblock copolymers as a new type of SPE for lithium metal batteries.69 In single-ion BAB triblock copolymers, the B block is a polyelectrolyte based on poly(styrene trifluoromethane-sulphonylimide of lithium) P (STFSILi) and connected to the central A block of the linear PEO (Figure 11.5). This structure ensures the delocalization of negative charge; thus, the Li ion enables a high dissociation level as a result of a weak interaction with the anionic structure. The single-ion BAB triblock SPE exhibits an excellent cation transport number (>0.85) that approaches unity, with reasonable conductivity (1.3 × 10−5 S cm−1 at 60°C). Furthermore, this single-ion-conducting SPE has excellent mechanical properties (10 MPa at 40°C) and an enlarged electrochemical window of 5 V compared with PEO. A battery assembled using a single-ion BAB triblock copolymer as SPE, LiFePO4 as the cathode, and Li metal as the anode exhibited excellent cyclic performance with

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FIGURE 11.5  Chemical structure of the single-ion conductor triblock copolymer P (STFSILi)-b-PEO-b-P (STFSILi). (Reproduced with permission from ref. [69]. Copyright (2013), Springer Nature.)

capacity retention of more than 85% at the C/2 rate at temperatures between 60°C and 80°C. Combining nanoparticles with anion-fixed single-ion-conducting SPEs is an effective method for increasing ionic conductivity. Lago et al. combined Al2O3 with PEO/polyethylene glycol–dimethyl ether (PEGDME) and reported superior ionic conductivity (4 × 10 −4 S cm−1 at 70°C) in comparison with other results.70 However, singleion-conducting SPEs require further improvements in terms of ionic conductivity for application in LIBs.

11.3 GPEs GPEs consist of a liquid plasticizer/organic solvent incorporated into a polymer–salt system. The concept of a GPE was first proposed in 1975 by Feuillade et al., who incorporated liquid plasticizer propylene carbonate (PC) into a copolymer matrix of vinylidene fluoride and hexafluoropropylene with NH4ClO4 salt.71 GPEs are an intermediate between

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SPEs and LEs because of their physical and chemical properties. The ion transport mechanism of GPEs differs from that of SPEs, in which lithium ions are transported in a liquid or swollen-gel phase rather than through the segmental motion of polymer chains caused by the introduction of a liquid plasticizer/organic solvent constituent.72 Because of their unique ion transport mechanism, most GPEs can achieve high ionic conductivity (10−3 S cm−1). GPEs are a promising candidate because they overcome the low mechanical strength of conventional LEs and poor ionic conductivity of SPEs. GPEs should have high mechanical strength, high ionic conductivity, and excellent electrochemical stability that facilitate practical use in LIBs. Because of the advantages of GPEs, many types of GPEs with different polymer matrices have been developed, including PEO,73–75 PAN,76–79 PMMA,80,81 PVDF,82–84 and PVDF-HFP.85,86 Because each polymer matrix has its own drawbacks, blended polymers, cross-linked polymers, and copolymer matrices have been developed to fulfill the requirements of LIBs. Plasticizers and lithium salts are the two constituents of GPEs other than polymer matrices. Generally, plasticizers are used to increase the amount of the amorphous phase in polymer electrolytes and promote the dissociation of Li salts to prompt an increase in ionic conductivity.75,87 Commonly used plasticizers include low-molarmass organics, organic solvents, and ionic liquids.72 PEG has been widely used as a low-molar-mass plasticizer in PEO–salt matrices.88,89 The ionic conductivity of GPEs is greatly influenced by both the content and molecular weight of the PEG.89 Other low-molar-mass organic compounds that have been used as plasticizers in polymer–salt systems include PEGDME,74,75 borate esters,73,90 phthalates,91–93 and succinonitrile.94,95 Organic solvents, such as ethylene carbonate, PC, diethyl carbonate, dimethyl carbonate, and γ-butyrolactone (γ-BL), are commonly used as plasticizers and facilitate the solvation of Li ions and transportation. Combining solvents with combined viscosity and dielectric permittivity helps to produce ionic conductivity greater than that of a single solvent.96 Ionic liquids have also been researched as plasticizer alternatives to organic solvents because of the advantages they pose in terms of high chemical and thermal stability, nonflammability, and high electrochemical stability.97–99 Other than plasticizers, lithium salt is another key constituent in GPEs; LiPF6 and LiTFSI are the most commonly used lithium salts in LIBs.

11.3.1 Fillers in GPEs GPEs have low mechanical strength because of the presence of the plasticizer. The addition of ceramic fillers, such as SiO2, TiO2, Al2O3, BaTiO3, and MgO into GPEs, is an effective method for solving this problem. The inclusion of ceramic fillers not only increases the mechanical strength of GPEs but also improves their ionic transport and electrochemical properties. A thin space-charge layer forms at the boundaries of ceramic grains dispersed in LEs, and overlapping of these space-charge layers leads to the formation of favorable lithium ion transport pathways.100,101 Zalewska et al. produced GPEs with submicron-sized SiO2 fillers incorporated into a PVDF-HFP polymer matrix modified with methacryloxy and vinyl groups.102 These GPEs exhibited a high ionic conductivity near 10 −2 S cm−1. Li et al. reported that Al2O3 NPs in PVDFHFP reduced the degree of crystallization of the polymer matrix as a result of the solid plasticization effect.103 With the optimum 10 wt% of Al2O3, GPEs exhibit an ionic

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conductivity of 1.95 × 10 −3 S cm−1 and lithium ion transference number of 0.73. Teng et al. developed a GPE with TiO2 NP–decorated poly(acrylonitrile-co-vinyl acetate)– PMMA (PAVM):TiO2 polymer matrix.104 They reported that PAN segments and TiO2 demonstrated exceptional absorbance of PF6 anions and created 3D percolative spacecharge pathways for Li ion transport (Figure 11.6a). This GPE exhibited a high ionic conductivity of 4.5 × 10 −3 S cm−1 and a transference number of 0.7. The battery fabricated using GPE–PAVM:TiO2, a graphite anode, and a LiFePO4 cathode presented a discharge capacity of 152 and 84 mAh g−1 at a 0.1- and 20-C rates, respectively. In 2017, Teng and coworkers reported a positive–intrinsic–negative (PIN)–diode-like GPE for LIBs.105 The diode-like GPE was composed of an intrinsic poly(acrylonitrile-co-methylacrylate) polymer matrix, and the positive and negative layers in the diode-like structure were achieved by doping TiO2 and SiO2 (which respectively exhibited positive and negative zeta potentials) into the intrinsic polymer host. As indicated in Figure 11.6b, the positive layer made contact with the LiFePO4, at which point the TiO2 nanoparticles adsorbed the PF6 anions to prevent their accumulation on the electrode surface. The neighboring layer with SiO2 nanoparticles attracted Li ions to the graphite anode and adsorbed desolvated Li ions on the SiO2 surface to prevent the cointercalation of solvent molecules into the graphite. Thus, the developed PIN-configuration GPE demonstrated excellent electrochemical performance when converted into a full cell. In 2018, Teng et al. used graphene oxide quantum dots (GOQDs) as fillers in a PAVM polymer matrix.106 The addition of GOQDs into the polymer matrix increased the ionic conductivity and Li transference number (0.77) by suppressing the formation of ion–solvent clusters and immobilizing the PF6 anions. The assembled Li/ electrolyte/LiFePO4 battery exhibited excellent electrochemical stability with 100% retention after 500 charge–discharge cycles at 5 C rate. Although GPEs exhibit increased mechanical strength and ionic conductivity with the addition of fillers, they cannot achieve the mechanical strength required for safe use in large-scale applications. Further developments for GPEs are required to establish their widespread use in commercial LIBs.

FIGURE 11.6  (a) Schematic of GPE-PAVM: TiO2 nanoparticles connect the space charge layers of the nitrile groups to form a 3D percolation Li-ion pathway. (Reproduced with permission from ref. [104]. Copyright (2016), American Chemical Society.) (b) Conceptual illustration of the positive–intrinsic–negative (PIN) diode-like asymmetric GPE. (Reproduced with permission from ref. [105]. Copyright (2017), The Royal Society of Chemistry.)

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11.4 SUMMARY In this chapter, we provide a detailed discussion of the various types and properties of SSEs, including ISEs and SPEs. In addition to SSEs, GPEs for LIBs are also briefly discussed. GPEs comprise the incorporation of a liquid plasticizer/organic solvent into a polymer–salt system. Liquid electrolyte components remain in the GPE system, indicating that safety concerns cannot be entirely resolved. Although ISEs boast various advantages over LEs in terms of mechanical strength, safety, battery life, and electrochemical window, they have the drawbacks of high interfacial resistance and low ionic conductivity. Efforts have been made to overcome problems related to interfacial resistance; specifically, the wettability of lithium in garnet-based solid electrolytes was improved through the coating of metal/metal oxide layers on the garnet surface. Another effective method for decreasing interfacial resistance is combining polymer electrolytes with ISEs for the interlayer. The hybrid electrolyte system with ceramic and polymer electrolytes is a promising candidate for Li metal batteries. SPEs have certain advantages over ISEs, such as flexibility, easy processing, and low cost. Despite these advantages, low mechanical strength and poor ionic conductivity are major challenges that hinder their wide application. The addition of inorganic fillers can improve the ionic conductivity and mechanical strength of SPEs. Although great efforts and progress have been made regarding SSEs, a comprehensive understanding of their charge-transport mechanism, interfacial reactions, cell architecture, and electrochemical performance has yet to be achieved. Further development of SSEs in terms of reducing the cost of materials and improving ionic conductivity through an understanding of their transport mechanism is necessary for their wide application in Li metal batteries in the future.

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71. Feuillade, G. and P. Perche, Ion-conductive macromolecular gels and membranes for solid lithium cells. Journal of Applied Electrochemistry, 1975. 5(1): pp. 63–69. 72. Long, L., et al., Polymer electrolytes for lithium polymer batteries. Journal of Materials Chemistry A, 2016. 4(26): pp. 10038–10069. 73. Chakrabarti, A., R. Filler, and B.K. Mandal, Synthesis and properties of a new class of fluorine-containing dilithium salts for lithium-ion batteries. Solid State Ionics, 2010. 180(40): pp. 1640–1645. 74. Kim, H., B. Oh, and Y. Kang, Preparation and electrochemical properties of nonwoven reinforced solid polymer electrolytes. Polymer Bulletin, 2000. 44(5–6): pp. 509–515. 75. Kim, Y.-T. and E.S. Smotkin, The effect of plasticizers on transport and electrochemical properties of PEO-based electrolytes for lithium rechargeable batteries. Solid State Ionics, 2002. 149(1–2): pp. 29–37. 76. Carol, P., et al., Preparation and characterization of electrospun poly (acrylonitrile) fibrous membrane based gel polymer electrolytes for lithium-ion batteries. Journal of Power Sources, 2011. 196(23): pp. 10156–10162. 77. Min, H.-S., J.-M. Ko, and D.-W. Kim, Preparation and characterization of porous polyacrylonitrile membranes for lithium-ion polymer batteries. Journal of Power Sources, 2003. 119: pp. 469–472. 78. Choi, S., et al., Electrochemical and spectroscopic properties of electrospun PAN-based fibrous polymer electrolytes. Journal of the Electrochemical Society, 2005. 152(5): pp. A989–A995. 79. Wang, S.-H., et al., Design of poly (acrylonitrile)-based gel electrolytes for high-performance lithium ion batteries. ACS Applied Materials & Interfaces, 2014. 6(21): pp. 19360–19370. 80. Meneghetti, P., S. Qutubuddin, and A. Webber, Synthesis of polymer gel electrolyte with high molecular weight poly (methyl methacrylate)–clay nanocomposite. Electrochimica Acta, 2004. 49(27): pp. 4923–4931. 81. Ramesh, S. and G. Ang, Impedance and FTIR studies on plasticized PMMA–LiN (CF 3 SO 2) 2 nanocomposite polymer electrolytes. Ionics, 2010. 16(5): pp. 465–473. 82. Choi, S.W., et al., An electrospun poly (vinylidene fluoride) nanofibrous membrane and its battery applications. Advanced Materials, 2003. 15(23): pp. 2027–2032. 83. Wang, X., et al., Gelled microporous polymer electrolyte with low liquid leakage for lithium-ion batteries. Journal of Membrane Science, 2014. 454: pp. 298–304. 84. Kim, J.R., et al., Electrospun PVdF-based fibrous polymer electrolytes for lithium ion polymer batteries. Electrochimica Acta, 2004. 50(1): pp. 69–75. 85. Yan, P., et al., Composite-porous polymer membrane with reduced crystalline for lithium–ion battery via non-solvent evaporate method. Ionics, 2015. 21(2): pp. 593–599. 86. Li, Z., et al., A foaming process to prepare porous polymer membrane for lithium ion batteries. Electrochimica Acta, 2009. 54(18): pp. 4403–4407. 87. Chaurasia, S., et al., Role of ionic liquid [BMIMPF 6] in modifying the crystallization kinetics behavior of the polymer electrolyte PEO-LiClO 4. RSC Advances, 2015. 5(11): pp. 8263–8277. 88. Ito, Y., et al., Ionic conductivity of electrolytes formed from PEO-LiCF3SO3 complex low molecular weight poly (ethylene glycol). Journal of Materials Science, 1987. 22(5): pp. 1845–1849. 89. Yang, L., et al., Effects of plasticizers on properties of poly (ethylene oxide) complex electrolytes. Solid State Ionics, 1990. 40: pp. 616–619. 90. Karatas, Y., et al., Synthesis and modeling of polysiloxane-based salt-in-polymer electrolytes with various additives. The Journal of Physical Chemistry B, 2009. 113(47): pp. 15473–15484. 91. Johan, M.R. and L.B. Fen, Combined effect of CuO nanofillers and DBP plasticizer on ionic conductivity enhancement in the solid polymer electrolyte PEO–LiCF3SO3. Ionics, 2010. 16(4): pp. 335–338.

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92. Rajendran, S., M. Sivakumar, and R. Subadevi, Li-ion conduction of plasticized PVA solid polymer electrolytes complexed with various lithium salts. Solid State Ionics, 2004. 167(3–4): pp. 335–339. 93. Klongkan, S. and J. Pumchusak, Effects of nano alumina and plasticizers on morphology, ionic conductivity, thermal and mechanical properties of PEO-LiCF3SO3 solid polymer electrolyte. Electrochimica Acta, 2015. 161: pp. 171–176. 94. Wu, X.-L., et al., Enhanced Li+ conductivity in PEO–LiBOB polymer electrolytes by using succinonitrile as a plasticizer. Solid State Ionics, 2011. 186(1): pp. 1–6. 95. Echeverri, M., N. Kim, and T. Kyu, Ionic conductivity in relation to ternary phase diagram of poly (ethylene oxide), succinonitrile, and lithium bis (trifluoromethane) sulfonimide blends. Macromolecules, 2012. 45(15): pp. 6068–6077. 96. Bandara, L., M. Dissanayake, and B.-E. Mellander, Ionic conductivity of plasticized (PEO)-LiCF3SO3 electrolytes. Electrochimica Acta, 1998. 43(10–11): pp. 1447–1451. 97. Wetjen, M., et al., Temperature dependence of electrochemical properties of crosslinked poly (ethylene oxide)–lithium bis (trifluoromethanesulfonyl) imide–N-butyl-Nmethylpyrrolidinium bis (trifluoromethanesulfonyl) imide solid polymer electrolytes for lithium batteries. Electrochimica Acta, 2013. 87: pp. 779–787. 98. Kim, G.-T., et al., UV cross-linked, lithium-conducting ternary polymer electrolytes containing ionic liquids. Journal of Power Sources, 2010. 195(18): pp. 6130–6137. 99. Rupp, B., et al., Polymer electrolyte for lithium batteries based on photochemically crosslinked poly (ethylene oxide) and ionic liquid. European Polymer Journal, 2008. 44(9): pp. 2986–2990. 100. Kumar, B., Heterogeneous electrolytes: Variables for and uncertainty in conductivity measurements. Journal of Power Sources, 2008. 179(1): pp. 401–406. 101. Osińska, M., et al., Study of the role of ceramic filler in composite gel electrolytes based on microporous polymer membranes. Journal of Membrane Science, 2009. 326(2): pp. 582–588. 102. Zalewska, A., et al., Study of the interfacial stability of PVdF/HFP gel electrolytes with sub-micro-and nano-sized surface-modified silicas. Electrochimica Acta, 2010. 55(4): pp. 1308–1313. 103. Li, Z., et al., Effect of Al2O3 nanoparticles on the electrochemical characteristics of P (VDF-HFP)-based polymer electrolyte. Electrochimica Acta, 2004. 49(26): pp. 4633–4639. 104. Wang, S.-H., et al., Immobilization of anions on polymer matrices for gel electrolytes with high conductivity and stability in lithium ion batteries. ACS Applied Materials & Interfaces, 2016. 8(23): pp. 14776–14787. 105. Lin, Y.-Y., et al., Diode-like gel polymer electrolytes for full-cell lithium ion batteries. Journal of Materials Chemistry A, 2017. 5(33): pp. 17476–17481. 106. Chen, Y.M., et al., Minimization of ion–solvent clusters in gel electrolytes containing graphene oxide quantum dots for lithium‐ion batteries. Small, 2018. 14(12): p. 1703571.

12

Silicon-Nanowire-Based Hybrid Solar Cells Ilham Ramadhan Putra, Pawan Kumar Singh, and Chia-Yun Chen

CONTENTS 12.1 Introduction.................................................................................................. 235 12.2 Silicon-Nanowire Fabrication...................................................................... 236 12.3 PEDOT:PSS Polymer as the P-Type Layer of Hybrid Solar Cell Application��������������������������������������������������������������������������������������������������243 12.4 Silicon-Nanowire-Based Hybrid Solar Cells............................................... 245 12.5 Conclusion................................................................................................... 250 References............................................................................................................... 251

12.1 INTRODUCTION Alternative energy as a renewable resource has been found to maintain energy stability in the world since the energy crisis has become a very urgent problem. A lot of studies proved that the greenhouse effect has an impact on life through several phenomena such as climate changes, the increase in temperature of Earth’s atmosphere, and organic pollution that have caused the use of current energy sources. In order to decrease the use of current energy sources, it is very important to find a potential solution with renewable energy that can be used in human daily life such as by utilizing alternative energy consumption via solar, wind, geothermal, and other energy resources. Among them, solar energy is an alternative energy source, where continuous solar light irradiations that can be converted into electricity through photovoltaic mechanisms. Regarding this, solar-cell devices offer superior conversion among other devices, and they are still developing to achieve the proper mass-need fulfillment. Many types of solar cells have been developed in recent years, and the most common type that is commercially used are silicon-based solar cells. However, the production of these kinds of solar cells needs more fabrication cost and energy consumption within fabrications. A hybrid solar cell is one type of energy conversion device that can convert the photon from sunlight into the charges regarding the production of electrical energy, and this has become a concern since the development of its efficiency is getting closer to commercial solar cells today. Furthermore, these kinds of solar cells have a low-cost fabrication process, and the limited energy needed in fabrication results in better efficiency and performance. One interesting type of hybrid solar cells is by combining a p-type organic polymer such as PEDOT:PSS with an n-type silicon as 235

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the active layers to provide the p–n junction. By a simple fabrication process to promote the photovoltaic mechanism in the device, it is a promising feature to explore new ideas in producing high-power conversion efficiency solar cells instead of the silicon inorganic-based solar cells. In the development of PEDOT:PSS/Si-based solar cells, many engineering processes were originated in achieving better power conversion efficiency (PCE). Nanostructuring of silicon wafer, treatment of organic polymers, interface engineering, and other methods have been developed in regards to achieving better performance of the solar cells. In this chapter, several discussions related to the engineering mechanism of silicon nanowire-based hybrid solar cells are performed.

12.2 SILICON-NANOWIRE FABRICATION One-dimensional silicon nanostructures such as nanowires have been realized as promising materials for a lot of applications in the semiconductor fields and renewable energy systems. Several approaches in silicon-nanowire fabrication were proposed in many studies. The fabrication processes, including top-bottom or bottom-top mechanisms, have been developed. The vapor-liquid-solid (VLS) growth mechanism first proposed by Wagner and Ellis [1] around the 1960s became the key technique in silicon-nanowire growth. Over time, other examples for siliconwire growth techniques, including chemical vapor deposition (CVD), annealing in the reactive atmosphere, SiO evaporation, molecular beam epitaxy (MBE), laser ablation, and solution-based techniques, such as supercritical-fluid-liquid solid (SFLS) and solution-liquid-solid (SLS) methods, had been proposed to develop the growth of silicon nanowires in several applications. Regarding the obtainment of several properties and characteristics of silicon nanostructures that influence the application of the device, such as crystalline quality and orientation, relative orientation to the substrate, size, and strain, further numerous methods were developed. Several methods using etching mechanisms have potential since they provide a superior yield of silicon nanowires with a lower cost. The examples of this kind of process are the reactive ion etching (RIE) and metal-assisted chemical etching (MACE) processes. MACE is the preferred process for gaining better fabrication of silicon nanostructures such as nanowires. It offers a sound structure of nanowire specification such as orientation, length, diameter, uniformity, morphology, and porosity. It was demonstrated for the first time in 1997 in producing porous Si by etching the aluminum thin film-covered silicon substrates using H2O, HF, and HNO3 mixed solution. The silicon substrate was degraded to form porous silicon in a short incubation time due to the existence of aluminum thin films. In recent years, several notable metals such as Cu, Au, Ag, Pd, Pt, and others have been studied as an effective catalyst for the MACE of a silicon substrate. MACE has been developed into two main successful steps: one-step and two-step MACE. One-step MACE emerges when the process of metal catalyst nucleation and anisotropic etching takes place in etching solutions containing hydrofluoric acid (HF) and metal salts. In two-step MACE, meanwhile, the metal catalyst is preliminarily deposited on the wafer surface, followed by anisotropic etching using HF/oxidant. The oxidant for this two-step etching that is

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commonly used is H2O2 solution. Several demonstrations of the formation kinetics in silicon nanowires by single-step MACE process were studied by several groups. The first demonstration of silicon-nanowire fabrication was provided by the production of Si nanowires using a single-step metal-assisted chemical etching process on lithographically defined areas. The p-type (100) silicon wafers were used as the substrates with a resistivity in the range of 1–2 Ω cm. Single-step MACE process was carried out using HF/AgNO3 aqueous solution by immersing the silicon substrates for a preferred time to determine the silicon nanowires’ length. The mechanism of silicon-nanowire fabrication involves two different procedures that occur simultaneously: (a) Ag nanoparticle deposition on the silicon substrate and (b) catalytic etching mechanism of silicon at the section of deposited Ag nanoparticles. Regarding that, the silicon nanowires (SiNWs) are fabricated by the reaction that exists as a galvanic displacement. The reduction potential of Ag+/Ag couple is known to be more positive compared to the flat-band potential of the silicon substrate. At the silicon surface, Ag+ ion reduction (by the holes injection mechanism) into the silicon causes Ag metal deposition based on cathodic reactions since Ag+ is of the electron-consuming type. In another way, oxidation of silicon caused by injected holes (anodic reaction) occurs since it is the hole-consuming type. During this mechanism, electrons of the atoms at the silicon surface move to Ag+ ions in the HF solution [2]. HF solution dissolves the oxidized silicon and then brings about the etching process and produces the nanowires. The reduction of Ag+ and oxidation of Si mechanisms contribute to the formation of Ag atoms on the silicon surface and, at the beginning, initiate the Ag nuclei production, which has a higher electronegativity than Si that induces the electrons of Si to become negatively charged, thus bringing further Ag+ reduction with a catalytic surface. It leads to the formation of a quasi-Schottky Ag/Si interface by a lower energy barrier for the holes that provide a charge transfer by hole injection of Ag nuclei to Ag+ ions in the solution. The etching process takes place vertically to a silicon surface with (100) since the highest rate of reaction is known to be much higher in [100] orientations compared to other crystallographic orientations. It is made when the etching process is applied to the silicon substrate; the etched side takes place on not only the front surface but also the back side of the wafer that has crystallographic orientations equivalent to [100]. A study of the silicon nanowires in two-step Ag-assisted chemical etching also provides more understanding about the MACE process and can be an alternative way in silicon-nanowire fabrication. This kind of etching process contains two steps. The first step etching process was conducted by immersing the substrates for 30 seconds into an aqueous solution that contained 0.02 M AgNO3 and 4.5 M HF at 25°C. After the rinsing process, the substrates were soaked in HNO3 (6%) solution for 10 minutes. Then, the substrates were immersed into an aqueous solution of 0.3 M H2O2 and 4.5 M HF at 50°C for 40 seconds to fabricate separated nanoparticle arrays. Residual Ag nanoparticles were removed using HNO3 (65%) and then second step etching was done by dipping into a mixture of 0.02 M AgNO3 and 4.5 HF at 50°C for 1 minute to form the dense aligned nanowire arrays, and finally residual Ag NPs were again removed by concentrated HNO3. The first etching step is described by the dissolution and local oxidation of silicon atoms in the galvanic displacement approach resulting

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in Ag nanoparticles. These nanoparticles have a surface density on the order of 1010 cm−2. HNO3 reactant promotes the partial separation and dissolution of preliminary Ag nanoparticles, which decrease the surface density to 108 cm−2. The wide separated distributions of the Ag nanoparticles have a role as the cathodes in promoting the nanopores’ configuration. The dissolution process of Si atoms that takes place underneath the Ag+/Si interface can be described by the following reaction equation:

Si +

n 4−n H 2O + 6HF → nH 2O + H 2SiF6 + H 2 (12.1) 2 2

Furthermore, the residual Ag nanoparticles produced during the etching process were removed completely by HNO3 reactant with a concentration of 65% after the etching mechanism was done. The second step is fabricating the nanowire arrays directly on the nanoparticle array by second etching by accomplishing this equation:

Si + 4Ag + + 6HF → H 2SiF6 + 4Ag + 4H + (12.2)

The reduction of Ag+ ions on the Si surface initiated by an electrochemical reaction and it is clearly understood that the potential of Ag+/Ag is more positive than the Si Fermi level. The nanowire configuration was achieved by the formation of Ag seeds on the silicon and injecting holes silicon. The dissolution of Si takes place at the interface of Ag seed/Si and is maintained by reduction of Ag ions, resulting in the aligned nanowire arrays. Later on, the concentrated HNO3 reactant has a role to dissolve the metal nanoparticles that cover Si surfaces. The optical reflectivity for different thicknesses of the two-step nanowires has been observed as the investigation of antireflection properties of SiNWs in an application such as a photovoltaic device layer. The results visualized that the reflectivity gradually decreased by the increasing nanowire lengths [3]. The other study provides the fabrication of silicon nanowires by a two-step MACE process with the combination of an advanced fabrication process by utilizing electron beam lithography (EBL). The ordered silicon nanowires with the formation of nanogaps by applying MACE within actual dimension control were developed. The control silicon nanowire arrays were established by conducting an etching process using MACE on the silicon substrates patterned by the EBL process. It was found that those nanogaps at the edge of each silicon nanowire contribute to the reactant diffusion pathways in the Si dissolution process and cause a difference in etching rates, which depends on the nanogap spacing. A uniform deposition of 3-nm-thick adhesive Ti and 20-nm-thick Au layers was done in resulting patterns, and the diameter and pitch size of the remnant resists was 85 nm and 140 nm. The deviation size of these patterns was less than 4 nm, and the square shape at each dot was produced. A preferred lift-off process was conducted on the samples regarding the design pattern transfer process onto the Si substrates. It produced tuned dimensions of 50 µm × 1000 µm for each structure of sample size. The Ti/Au nanopatterns facilitated the MACE process of Si substrates in producing defined features of Si nanostructures. The etching process was developed by immersing the patterned Si

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substrates in the H2O2/HF mixed solutions for a different time at room temperature. The electrochemical reactions can be referred to as:

Cathodic : H 2O2 + 2H + → 2H 2O + 2h + → E 0 = +1.76 V (12.3)



Anodic : Si + 6F − → SiF62 − + 4e − → E 0 = −1.24 V (12.4)

Since the H2O2 oxidants had a higher reduction potential than the valence band of Si, the hole injections were initiated and produced catalytic etching by the Ti/Au metal assistance. These metal meshes had a role as the cathodes that facilitated the oxidation of Si substrate under them. The oxidized Si was dissolved by the HF solution, which allowed the unoxidized exposed Si to have contact with the metal and assist the next cycle of reaction series. This etching kinetics mechanism on a patterned Si substrate was described clearly in companion of the effective contact among H2O2/ HF reactants with the Au/Si interface, resulting in the charge transport mechanism. Si nanowires with excellent dimensions and uniformity were obtained by the combination of both EBL and MACE processes. The pitch size of nanopatterns from 140, 170, and 230 nm was observed and modulated regarding the analysis of etching kinetics in this process on the geometrical spacing of ordered nanopatterns. Intrinsically, those three different pitch sizes of EBL patterns were predefined on a crystalline Si (100) substrate with the distance between each designed pattern of 100 µm. This basic fabrication strategy was proposed to make sure the consistency of etching conditions could be achieved in different ways, including reaction time, temperature, and the concentrations of reactants. The morphologies of silicon nanowires resulted in similar diameters (85 nm) for all three pitch sizes. Interestingly, the reduction of pitch size can fabricate longer nanowires. A quantitative observation of the pitch sizes after an etching time of more than 8 minutes was conducted and notably, nanowire average spacing for different designed patterns of 228, 168, and 136 nm results in size deviation of less than 5 nm. These aspects provided the explanation that the combination of EBL patterning and Au-assisted etching process shows credibility in the fabrication of silicon nanowires with dimensions in the sub-100-nm level by such an etching technique. Furthermore, the controllable lengths of nanowires can be achieved with a linear correlation with formation rates where the increase of the lengths can directly result from the increase of etching time. A higher formation rate of nanowires can be carried out on the patterned arrays of a 140-nm pitch size compared to a 170-nm and 230-nm pitch, respectively. The nanogaps also exist at the edge of every nanowire when the width approaches 4 to 7 nm. This occurred in every conjunction of the nanowire, and the consistent size obtained with the widths and geometries was unchanged regardless of the time of the etching process (1 minute to 5 minutes). These findings, speculation of the formation of these periodic gaps, were proposed as these nanogaps have a role as an effective pathway in allowing the reactants to have contact with the metal catalyst. As the Au films have thick continuous layers that cover the Si substrates, they provided direct penetration of the reactants to metal catalyst-covered Si. H2O2 as the oxidant generated the holes, and it has important roles in the result of uniformly vertical movement to create

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the well-array nanowires. The H2O2 oxidants produced holes that in a special way depleted at the metal/Si interfaces regarding the lower Schottky barrier of hole injections toward Si. The concentration gradients of holes were obtained and, furthermore, motivated the H2O2 reactants to diffuse under the metal catalyst rather than take place at the edge of metal catalyst film entirely until the removal of oxidized Si at interfaces was completed. The homogeneous distribution of the holes in the Si layer beneath the Au films was obtained, resulting in a uniformly vertical movement of meshes sinking into Si. Regarding the measurement for etching rates on patterned Si, a model utilizing that etching mechanism was proposed. The electrochemical oxidation of Si, chemical dissolution of oxidized Si, and additional effect of mass transfer of reactants before reaching the metal/Si interfaces are the essential parts for determining the kinetic mechanism of the whole etching process by expressing this into kinetic resistance in a series noted as REC, RD, and RM, respectively. By stating that all the metal patterns take on with similar etching environment, the R EC and RD are notably the same in all cases, whereas R M can be considered the dominant factor that varies the etching kinetics. For example, in this etching process, the pitch size of metal meshes that have a larger spatial size increased the R M since it has the decreasing numbers nanogap arrangement, providing pathways for reactants to diffuse reaching the Si, so that suppression of Si dissolution is obtained. Simultaneously, increasing the coverage area of metal meshes, in particular, influences the active reactants to move farther regarding removal of Si at the interface completely. The combination of those effects slows down the entire dynamic process, and the extent of this deceleration is easily modulated using the metal coverage area. Regarding the analysis of the universal influence of metal coverage on Si etching, a modulation by using pitch size (P) and nanopatterns diameter (D) was offered. The following equation provides the ratio of Ti/Au coverage with the thickness aforementioned:

Area of metal coverage    (Area of metal coverage + Area of exposed Si)  ×100% (12.5)  

Based on this modulation, two representatives concerning the geometries of nanopatterns and influences on metal coverage were investigated. The first representative is sparse nanowires with pitch size (P) of 170 nm, diameter (D) of 60 nm, and metal coverage of 93%. By comparison, the dense nanowires (P = 85 nm, D = 85 nm, and metal coverage = 75%) was also fabricated. It has been demonstrated that the sparse nanowires possess a shorter nanowire lengths compared with the dense one under the similar etching condition, which can be explained by the diffusion of reactants via such etching reaction [4]. Another study of a proposed etching process by using isotropic etching, followed by the assistance of silver, expanded understanding of the MACE process. Furthermore, a simpler etching process to fabricate more desirable silicon nanowires was developed. Shape-diversified silicon nanowires were fabricated by utilizing a hydrofluoric acid (HF)/nitric acid (HNO3) solution with the facilitation of a catalyst using silver (Ag) metal. A novel strategy in etching based on HF/HNO3 and silver nanoseed support provides a contribution in preparing varied silicon nanostructure

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arrays with a notably smooth surface. Three different p-type silicon substrates were used, including (100), (110), and (111) orientation with resistivity values of 1 to 10 Ω cm. Before etching, the silicon substrates were cleaned in the ultrasonic machine using acetone, isopropyl alcohol, and DI water several times. The cleaned substrate was later immersed into diluted 0.1 M HF for one minute at room temperature. After finishing the dipping process, the substrates were rinsed into DI water and dried by N2 gas. The etching process was done by HF/HNO3 and the Ag nanoseed catalyst by fulfilling two basic steps: (1) the dipping of substrates into AgNO3 (0.005 M) and HF (4.5 M) mixed solutions for a few seconds, which can provide the silver (Ag) loading onto the silicon substrate mechanism, and (2) after rinsing the as-prepared samples into DI water and drying by N2 gas, the Ag-loaded silicon substrates were immersed in a mixture solution of HNO3 (0.6–2.4 M) and 4.8 M HF. The mechanism can be described by electrochemical reactions listed below: Cathodic reaction:

HNO3 + 3H + → NO + 2H 2O + 3h + (12.6)



HNO3 → HNO2 → NO → HNO → H 2 N 2O2 → N 2O (12.7)

Anodic reaction:

Si + 6F − → SiF62 − + 4e − (12.8)

After the complete removal of Ag nanoparticles, the as-prepared samples were rinsed with DI water several times and later dipped into a concentrated HNO3 (65%) solution for 2 minutes. In addition, the comparison for etching rates of three different Ag+ ions containing oxidants was proposed by subsequently dissolving 0.02 M of each oxidant into 4.5 M HF solution followed by magnetic stirring for 20 minutes. The dissolution reactions are presented following these reactions:

AgNO3 → Ag + + NO3− (12.9)



Ag2O + 2H + → 2Ag + + H 2O (12.10)



CH 3CO2 Ag + H + → Ag + + CH 3CO2H (12.11)

The involved anisotropic etching of HF/HNO3 using the assistance of the Ag mechanism can be divided into steps. In the first step, the galvanic displacement caused the deposition of Ag nanoparticles electroless onto the Si wafers. The distilled water (DI water) used to rinse the silicon substrates several times was then dried by N2 gas flow with gentle blowing. Later, in the dipping process of Ag-loaded Si samples that mixed aqueous with various concentrations of HNO3 (0.6–2.4 M) and 4.8 M HF at room temperature, Ag nanoparticles performed as the microscopic cathodes in initiating the reduction-oxidation process of HNO3 dissociation regarding hole injection into Si. As a result of the injected holes, the Si atoms that took place under the Ag seeds underwent the local oxidation reaction. In addition, since the HF etchant was used, the oxidized Si dissolved, leaving

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the oxide-free area exposed to the next recycling reaction process. It is notable that the volume ratio of HNO3 to HF for this process was in the range of 0.14 to 0.60, which is much lower than the HNA etching process and represents a highly efficient carrier transport pathway. Then, the application of HF/HNO3 etching assistance by Ag essentially was similar to the two-step MACE by utilizing the H2O2 oxidant solution. However, these different characteristics were conceptualized to distinguish this current process from the H2O2-based MACE process. From an electrochemical point of view, it is known that the reduction-oxidation (redox) potential of H2O2 reactant (E0 = +1.76 V) is larger than both Ag+/Ag (E0 = +0.80 V) and NO3− (+0.96 V) regarding oxidation reactions. Original Ag seeds experienced the decomposition and redisposition thermodynamically for the time of Si dissolution, thus resulting in difficulty in controlling the surface features of the etched structures. More notably, regarding the dissociation of preference, Ag is outstandingly diminished by the Ag-assisted HF/HNO3 etching process in comparison to the reduction potentials of Ag and HNO3. The shape ­prolongation of Ag seeds in a long-term catalytic transport within Si facilitates the stable and uniform etching of Si that obtains controllability of nanowire fabrication. A remarkably smooth nanowire array was obtained for different concentrations of HNO3, noted as 0.6 M and 1.2 M. The addition of HNO3 concentration did not influence the consistency of the smooth top surface utilization; further, the nanowire lengths were observed to increase. This smoother surface is different from H2O2-based MACE, which results in rough nanowires. This discovery can be described since the lower reduction potential of NO3− provides the oxidation and dissolution of Si. It promotes the supply of holes in a stable condition and consequently maintains the uniformity of Ag seeds sinking on the Si, which leads to smoother surface nanowires. When the HNO3 concentration was increased to 1.8 M, the top surfaces of wires became rough as expected because of the rigorous reaction of concentrated HNO3 from notable injections of holes into the Si. The smooth features of the silicon-nanowire surface benefit from the undesired scattering of light that can be reduced efficiently and enhances the light-trapping mechanism of Si nanowires. In addition, the surface wetting characteristics have been observed by the Cassie–Baxter model,

cos θCB = rΦ f cos θ − rΦ (1 − f ) (12.12)

where θCB is the contact angle (CA) of nanostructured Si, θ is the CA of polished Si surfaces, rΦ is the wetted area roughness, and f is the fraction area of solid-liquid interfaces on the nanostructured Si surface. The results show that the current process samples exhibit CA from 150°–163° (+0.4°) by increasing the HNO3 reactant concentration from 0.6–2.4 M compared to the 1.2-M H2O2-based MACE process, which has a CA of 170° (+0.2°), and the increasing CA is associated with the surface roughness, implying a smoother surface of current process nanowires. From the kinetic study, the results of nanowire lengths denote a linear relationship with etching time for every verified condition. It also provides understanding that the increasing nanowire lengths are affected by faster etching rates while the concentrated HNO3 is utilized in the etching process. Furthermore, the process temperature has an impact on reaction kinetics and becomes another critical factor to the

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correlation of nanowire length versus etching time. The activation energy involved in every catalytic etching mechanism can be explained by the Arrhenius equation:

R = A exp( − Ea / RT ) (12.13)

where R and A are the etching rate of Si and the preexponential factor, respectively. Ea, R, and T note the activation energy, gas constant, and reaction temperature, respectively. By analyzing the activation energy approach, it is established that the etching rate at each temperature is rationally confirmed with the extracted Ea values. In consort with the investigation of the time factors, it is proved that the involvement of the HF/HNO3 etching process catalyzed by Ag metals is predominantly determined by the oxidation reaction instead of the diffusion of reactive electrolytes. Moreover, it is known that the MACE process also is related to the current density (j) involved in the reduction-oxidation process that is described by the equation:

j = − zekc nsCredox exp( −U / kBT ) (12.14)

where z is the number of transferred electrons and e, kc, and ns are the charge of e­ lectrons, the rate constant of the redox process, and the electron density at the Si/ electrolyte interface, respectively. Credox corresponds to the concentration of Ag + ions. U and k B are the activation energy and Boltzmann constant, respectively [5].

12.3 PEDOT:PSS POLYMER AS THE P-TYPE LAYER OF HYBRID SOLAR CELL APPLICATION Poly(3,4-ethylene dioxythiophene), abbreviated as PEDOT, is one type of conductive polymer that is electrochemically stable and can be doped to achieve its high electrical conductivity but keeps maintaining its transparency. This material is easily fabricated by a simple processing technique since its capability to be dispersed in water or several solvents promotes more advantages. It also has good thermal stability that advances its application in photovoltaic devices such as a hybrid solar cell. In its commercial product, PEDOT was fabricated in a stabilized form with PSS (polystyrene sulfonate) as the ion counter and charge balance, hereinafter referred to as the PEDOT:PSS solution. PEDOT:PSS has a core–shell structure. PEDOT has a structure with shorter chains attached to longer PSS chains through Coulombic attraction among them [6]. It forms a necklace-like structure in the water as its solvent since it consists of hydrophobic PEDOT and hydrophilic PSS of polyelectrolyte. In the pristine form, these materials have low conductivity since the insulating properties of PSS influence the energy barriers for the charge movement across the PEDOT. In common products, it also notable that PEDOT:PSS has different grades for highly conductive polymer application. As-prepared (pristine) PEDOT:PSS films from those different grades of polymers almost show identical conductivity in the range of 0.2–1 S cm−1, but different conductivity improvements when treated by using cosolvents have been observed. Several treatments already have been applied to enhance the conductivity of PEDOT:PSS since it has a low conductivity in the pristine form that is not capable in several applications, especially in the semiconductor field for photovoltaics

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application in the hybrid solar cell. Several treatments before making the whole cell should be applied to the PEDOT:PSS organic layer regarding conductivity and efficiency enhancement. Physical treatment by using thermal strongly modified the morphology and structures of the PEDOT:PSS obviously enhancing the conductivity. Another physical treatment is by using light, and it is found that UV irradiation could influence the conductivity enhancement and the work function of this polymer. Interestingly, the combination of physical (by applying an annealing process) and chemical treatments have already modified PEDOT:PSS into a high-conductive polymer in several applications. Chemical treatment developed in many ways to obtain improvement in conductivity by adding several solutions such as solvents, surfactants, ionic liquids, acids, and so on [6, 7]. The conductivity enhancement of PEDOT:PSS by adding several chemical solutions can be determined by secondary doping methods—for example, adding a polar organic solvent such as dimethyl sulfoxide (DMSO) into the polymer solution. The addition of this cosolvent can enhance the electrical conductivity at room temperature into two orders. Different explanations about the increasing of the electrical conductivity of PEDOT:PSS are still in the approaching paths since various groups have found a different mechanism to describe how PEDOT:PSS can be more conductive when added into several kinds of solutions. The first approach correlated to the temperature-dependent conductivity of PEDOT:PSS since the charges transport across the PEDOT chain by a hopping mechanism. Another approach confirms that when added by cosolvents such as ethylene glycol (EG), the conductivity of PEDOT:PSS can achieve 600–700 S cm−1 and the PSS section can be removed by the identification of X-ray photoelectron spectroscopy (XPS) that has the S2p band in the range of 163–167 eV, which can be attributed to the sulfur atom in PEDOT. By observation of this XPS result, the identification of PEDOT to PSS can be attributed to the change of both sections, which implies the removal of PSS as the insulator parts with the enhancement of electrical conductivity achieved. In addition, the probability of a screening effect that happened in the PEDOT:PSS solution after the polar solvent addition indicates the rise of the conductivity of the polymer. It can be remarked that the higher dielectric constant involved in polar solvents can induce a stronger screening effect between the counter ions and charge carriers, acquiring the reduction of Coulomb interaction between PEDOT and PSS. Although the conductivity enhancement mechanism is still argued, the key point is that for hybrid solar cell application, it is necessary to increase the organic material layer by applying several treatments. An observation of the EG and DMSO cosolvents’ addition into the PEDOT:PSS in a hybrid solar cell application already has prevailed. Different weight percentages of both solvents influence the conductivity of the PEDOT:PSS and enhance the performance of the solar cells that can be observed by the relationship between current density and voltage [8]. The highest PCE is achieved by 12%, which was made by the addition of 7 wt% of EG into the pristine PEDOT:PSS. Furthermore, the sheet resistance of this concentration shows that the lowest one compared to another percentage of cosolvents achieves the value of 1.9 × 102 Ω sq, providing a higher conductivity compared to other concentrations.

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12.4 SILICON-NANOWIRE-BASED HYBRID SOLAR CELLS A simple arrangement of a hybrid solar cell can be obtained by simple combination of PEDOT:PSS as the p-type layer and n-type silicon. These two kinds of materials can create a p–n junction and generate charge carriers to produce current. It was reported that there are two approaches in the junction models of a PEDOT:PSS–silicon solar cell. These two models briefly describe the junction between the p-type PEDOT:PSS layer and the n-type organic semiconductor that is silicon. The first model is based on the Schottky junction theory by explaining the interface of the semiconductor to metal. The second one describes the abrupt junction that happened between the highly doped p-type and the moderately doped n-type semiconducting region. Regarding both cases, the simplest form to describe the current density-voltage characteristic for those types of photovoltaic junctions is by the ideal diode equation under the illumination:

  qV   J = J 0  exp  − 1 − J sc (12.15)  kT   

By rewriting the equation above at the condition of an open circuit (J = 0), it is shown that the open-circuit voltage (VOC) of the cell primarily is dependent on the dark saturation current density (JO) and also the short-circuit current density (JSC) and can be written by:

Voc ≈

kT  J sc  ln (12.16) q  J 0 

The dominating transport mechanism of a Schottky junction between the metal and high-mobility semiconductor is the thermionic emission of the majority carriers over the potential barrier (ΦB) that the formation is taking place at the interface.

æ qF ö J 0 = A**T 2exp ç - B ÷ (12.17) è kT ø

The equation above describes the J0 for the Schottky junction with A** noted as the reduced effective Richardson constant that comprised the effects of tunneling and scattering the majority carriers at phonons, as well as the correction factor for a number contribution of majority carriers diffusion for relatively doped silicon. A reasonably small applied field of A** at the value of about 110 A/(cm·K)2 is the applied value for room temperature application. In particular, for a Schottky junction of J0, it is dependent on the Schottky barrier height ΦB. However, in a junction between the p-type and n-type semiconductors, the diffusion of minority carriers becomes the primary reason in the transport process. The assumption that the profile of doping between the two types of semiconductors changes suddenly at the junction of the diode equation can be specified by following Shockley [9]. Supposing that the doping in the p-type semiconductor is essentially larger than the n-type one can be expressed by:

J0 =

ni2µ p kT (12.18) LP N D

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In this, μp is noted as the mobility, Lp is the diffusion length of the minority carriers (holes for n-Si), and ni denotes the intrinsic carrier concentration. J0 is simply dependent with the properties of the moderately doped n-type semiconductor and it is inverse to the doping concentration (ND) so that it can be recognized from the following equation. The VOC in a p–n junction has to increase by the increase of ND. The consideration that should be noticed in the Shockley equation is the only diffusion that is limited by bulk recombination mechanisms out of the space-charge region. It is sufficient for the high intrinsic carrier density semiconductor. For a low intrinsic charge carrier semiconductor like silicon, the more suitable approach is by considering generation and recombination at the traps inside of the space-charge region. With specifying the area of specific parallel (RP) and series resistance (RS) of the device, J–V characteristics of an uneven p–n junction solar cell can be visualized by the twodiode model following the equation:



   q(V − RSJ )    q(V − RSJ )   J = J 01  exp  − 1 + J 02  exp    − 1     kT 2kT    (V − RSJ ) + − J SC RP

(12.19)

For the application of Si nanowires and the PEDOT:PSS polymer, a fabrication of a hybrid solar cell with 8.40% of efficiency was achieved by designing a simple solar cell [9]. The main active layers consisted of two different materials to make a p–n junction. The n-type silicon nanowires and p-type PEDOT:PSS were proposed as this active design. To promote charge collection, two types of electrode were used: indium tin oxide (ITO) at the top and Ti/Ag at the bottom of the cell. The fabrication of SiNW/PEDOT:PSS solar cells was proposed by fabricating different lengths of nanowires. N-type silicon (100) wafers with phosphor-doped and resistivity of 1–10 Ω cm were used. The silicon nanowire array was fabricated on the polished single-side of the wafer by a MACE process with Ag+ ions (0.023 M) and HF (5.6 M). The mechanism that happened during the immersion process of the n-Si wafer was described as reduction-oxidation (redox) by the following explanation. Ag+ ions were reduced to form Ag nanoparticles and deposited on the n-Si surface. On the other hand, the silicon near the silver nanoparticles was oxidized, forming silicon dioxide (SiO2) between the Si and Ag. The silicon dioxide that formed later was etched by HF. Since the electronegativity of Ag (1.9) is higher than Si (1.8), the electrons of Si that located near the Ag nanoparticles were distributed through Ag. Regarding this, the attraction of Ag+ by partially negative charges of the Ag surface took place. In addition, the formation of Ag nanoparticles was supported by the silver ions that gathered and attracted the electrons that took place in the etching solution. At the same time, Si under the Ag formed the SiO2 and was etched by HF immediately. This mechanism can be explained by the following chemical reactions:

Ag + + e − → Ag (12.20)

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Si + 2H 2O → SiO2 + 4H + + 4e − (12.21)



SiO2 + 4HF → H 2SiF6 + 2H 2O (12.22)

The rate of reactions discovered in the [100] direction resulting arrays at the surface formed normal to the (100) crystallography. As known, the pretreatment process was applied to the wafer before being performed by MACE in a particular condition to produce SiNWs with a large density. In regards to producing different lengths of silicon nanowires, the etching time was applied over several minutes and controlled at 3, 10, and 20 minutes, resulting in 0.37, 2.15, and 5.59 µm of SiNWs, respectively. In the end fabrication of SiNWs, the samples were soaked by concentrated nitric acid and diluted HF in the goal of eliminating the residual Ag and SiO2. The fabrication of the hybrid solar cell continued to deposit the electrode such as Ti or Ag as the back side to gain charges and the cathode contact by using an e-gun evaporator. Before fabricating the p–n junction, the silicon-nanowire samples were saved in 60% humidity at room temperature for 5 hours to grow the oxide to enhance the hydrophilicity of the Si surface, resulting in better contact between the SiNWs and PEDOT:PSS polymer. Moreover, the quantum tunneling mechanism of the carriers can still happen in this p–i–n structure as long as the oxide layer is very thin. This structure can reduce bulk recombination that happens in the p–n junction structure. The PEDOT:PSS later was spin-coated on the ITO/glass electrode and stuck the SiNWs into PEDOT:PSS. Since the SiNWs have good hydrophilicity, the polymer solution can penetrate and fill the hole of the nanowires easily, resulting in better contact and junction of the cells. The final process is the annealing process by baking the cells on the hot plate at 140°C for 10 minutes in the nitrogen ambient. The device’s performance was investigated using the solar simulator in the dark and 1 sun AMG 1.5 G 100 mW cm−2 illumination [9]. The results showed that the performance of the device for SiNWs is superior compared with the flat-silicon-based devices. It is noticed that SiNWs can reduce reflection and enhance the absorption of the light. The JSC of solar cells with the incorporation of silicon nanowires with a length of 0.37 µm is two times (24.24 mAcm−2) higher than that of the flat one (11.14 mAcm−2). Moreover, the SiNWs enlarge the junction area from 7 to 9.35 times larger than that of the flat sample. SiNWs were also found to reduce the RS from 25.2 Ωcm2 to 2.95 Ωcm2 using nanowires 0.37 µm in length. Furthermore, the VOC of the cells for both planar or nanostructured samples with a 0.37-µm length of nanowires in the latter does not show big differences. The correlation between VOC and RSH can be referred to as:

VOC =

nkT  J PH   VOC   ln   −  + 1 (12.23) q  J 0   J 0 RSH  



J PH = qg(x )(LP + LN + W ) (12.24)



éæ 1 ö Dn æ 1 ö Dp ù J 0 = qni2 êç +ç ú (12.25) ÷ ÷ êëè N a ø tn 0 è N d ø tp 0 úû

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n is noted as the ideality factor, k is the Boltzmann constant, and T is the operation temperature, q is the electron charge, JPH is the photogenerated current density, and J0 is the reverse saturation current density. Furthermore, g(x), Lp, Ln, W, and ni are denoted as the generation rate of electron-hole pairs, minority carrier diffusion lengths in the n-type region, minority carrier diffusion lengths in the p-type region, carrier diffusion length in the depletion region, and intrinsic concentration, respectively. Na and Nd are dopant concentrations of the p-type and n-type region, respectively. Dp and Dn are denoted as the diffusion coefficient of holes and electrons, respectively. Τn0 and τp0 are the minority carrier lifetime of electrons and holes, respectively. Regarding this, VOC has a variation not only with Rsh but also with n, Jph, and J0. In addition, the correlation of FF in the finite condition of RS and Rsh can be found based on the following equation:

 (v + 0.7) FF0 (1 − rs )  FF = FF0 (1 − rs ) 1 − oc  (12.26) voc rsh  



 voc − ln ( voc + 0.72 )  FF0 =   (12.27) (voc + 1)  

(

)



voc =



rs =



rsh =

(Voc ) nkT  q 

(12.28)

( Rs ) (12.29) Voc J sc

(

)

( Rsh ) (12.30) Voc J sc

(

)

From those equations above, n is denoted as the ideality factor, k is the Boltzmann constant, T is noted as operation temperature, and q denotes the one electron charge. It is can be inferred that the FF is not only determined by Rs but also determined by n, Voc, Jsc, and Rsh. Therefore, the variation of FF with the wire length is not the same as Rs [10]. The improvement of efficiency from the simple fabrication process of hybrid solar cells was achieved by using several approaches. The etching process of n-type silicon substrates for the fabrication of nanowires was conducted by applying MACE by following these equations: Cathode reaction:

H 2 O 2 + 2H + → 2H 2 O + 2 h +

E0 = +1.76 V (12.31)

Anode reaction:

Si + 6F − → SiF62 − + 4e −

E0 = −1.24 V (12.32)

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The fabrication process was developed by employing several treatments for the materials. After preparing the SiNWs, the deposition of aluminum (Al) as the back electrode with the thickness of 200 nm followed by growing native oxide. The growth of native oxide has a role to increase adhesion properties of SiNWs when contacted with PEDOT:PSS and also as the barrier for bulk recombination. The PEDOT:PSS polymer was treated by solvents (EG) and surfactants (FS) regarding the increase of conductivity, performance enhancement, and adhesion improvement and was spin coated on the top of SiNWs. After the annealing process at 120°C, the ITO glass was integrated and had a role as the top electrode. The effects of the surfactants (fluorosurfactants) used were observed regarding the performance of SiNW-based hybrid solar cells. The optimum amount of surfactants for reaching the best performance of the solar cells was found to be 0.1 wt%. The value of Jsc was increased due to the decrease of internal resistance that occurred at the interface of PEDOT:PSS and Si surfaces, as presented in Figure 12.1. For further investigations, the contact-angle measurement was proposed and resulted in the decrease of CA from 132.84° to near 0° with the addition of EG as the solvent. This feature clearly improved the uniformity of interfaces between active layers, making charge separation of photogenerated carriers more efficient. The effects of different solvents are also observed in optimizing the addition of solvents for enhancing cell performance, as presented in Table 12.1.

FIGURE 12.1  Investigations of cell parameters on adding amounts of surfactants on polymer preparations with ethylene glycol as solvents: (a) fill factor, (b) Jsc, (c) Voc, and (d) PCE.

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TABLE 12.1 Photovoltaic Performances of Hybrid Solar Cells and Effects of Solvent on Photovoltaic Performance FF (%)

Voc (V)

Jsc (mA/m2)

1.0

62.04

0.3780

27.01

6.34

2.0 4.0 6.0 2.0 4.0 6.0 8.0 1.0 2.0 4.0

69.18 54.24 51.15 70.79 69.48 71.98 62.04 54.49 62.58 44.20

0.4039 0.3884 0.3960 0.3350 0.3580 0.3910 0.3780 0.2788 0.3020 0.3760

27.07 29.06 30.08 26.80 32.80 36.10 32.50 26.70 40.11 20.60

7.56 6.12 6.09 6.36 8.15 10.16 7.63 4.06 7.58 3.43

6.0

31.00

0.4159

14.30

1.43

Solvents

wt%

DMSO

EG

PEG

PCE (%)

Compared with DMSO and polyethylene glycol (PEG), the EG-based solar cells showed better performance since this solvent greatly decreased the sheet resistance attributed to the good arrangement of the structures in PEDOT within the PSS matrix. The nanowire lengths in the hybrid solar cells also influenced the power conversion efficiency (PCE), which is enhanced compared to the planar cells since the nanowires can increase the light-trapping phenomenon and in advance promote the large-area production for undergoing the formation of heterojunction in the assistance of separating photogenerated carriers.

12.5 CONCLUSION The MACE process is a promising way to produce silicon nanowires with controllable dimensions, surface roughness, and desirable uniformity that can be applied for several applications with a low-cost process. The electrochemical reactions and kinetic study of silicon-nanowire fabrication by this method were discussed, and several parameters in fabricating SiNWs, such as reactant concentrations and etching rates, were also discussed to give more point of views regarding the nanofabrication of silicon. Moreover, silicon nanowires exhibit outstanding properties such as low reflectivity and desirable surface properties so that they can be better approached to meet the requirements of potential photovoltaic applications. On the other hand, the PEDOT:PSS as the p-type conductive polymer has an important role in hybrid solar cells since it has a good conductivity after applying physical and chemical treatments. The conductivity of this polymer is importantly related to carrier mobility in the active layer of the hybrid solar cells, so it is critical to conduct the treatment process before applying this polymetric material in photovoltaic applications. The

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p–n junction of the hybrid solar cells by using PEDOT:PSS/Si as the active layer has been discussed in two models through the Schottky junction and abrupt junction that happened between highly doped p-type and moderately doped n-type semiconducting regions. It can be anticipated that silicon-nanowire-based hybrid solar cells could be the promising design strategy for the development of next-generation solar cells.

REFERENCES 1. Wagner, R. S. & Ellis, W. C. Vapor-liquid-solid mechanism of single crystal growth. Appl. Phys. Lett. 4, 89 (1964). 2. Peng, K., Lu, A., Zhang, R. & Lee, S. T. Motility of metal nanoparticles in silicon and induced anisotropic silicon etching. Adv. Funct. Mater. 18, 3026–3035 (2008). 3. Chen, C. Y., Li, W. J. & Chen, H. H. Tailoring broadband antireflection on a silicon surface through two-step silver-assisted chemical etching. ChemPhysChem 13, 1415–1420 (2012). 4. Chen, C. Y. & Liu, Y. R. Exploring the kinetics of ordered silicon nanowires with the formation of nanogaps using metal-assisted chemical etching. Phys. Chem. Chem. Phys. 16, 26711–26714 (2014). 5. Chen, C. Y. & Wong, C. P. Unveiling the shape-diversified silicon nanowires made by HF/HNO3 isotropic etching with the assistance of silver. Nanoscale 7, 1216–1223 (2015). 6. Ouyang, J. ‘Secondary doping’ methods to significantly enhance the conductivity of PEDOT:PSS for its application as transparent electrode of optoelectronic devices. Displays 34, 423–436 (2013). 7. Shi, H., Liu, C., Jiang, Q. & Xu, J. Effective approaches to improve the electrical conductivity of PEDOT:PSS: A review. Adv. Electron. Mater. 1, 1–16 (2015). 8. Thomas, J. P., Zhao, L., McGillivray, D. & Leung, K. T. High-efficiency hybrid solar cells by nanostructural modification in PEDOT:PSS with co-solvent addition. J. Mater. Chem. A 2, 2383 (2014). 9. Jäckle, S. et al. Junction formation and current transport mechanisms in hybrid n-Si/ PEDOT:PSS solar cells. Sci. Rep. 5, 1–12 (2015). 10. Syu, H. J., Shiu, S. C. & Lin, C. F. Silicon nanowire/organic hybrid solar cell with efficiency of 8.40%. Sol. Energy Mater. Sol. Cells 98, 267–272 (2012).

13

Characterization and Performance of Li-ZnO Nanofiber and Nanoforest Photoanodes for DyeSensitized Solar Cells I-Ming Hung, Jing-Ru Chen, and Yi-Hung Wang

CONTENTS 13.1 Introduction.................................................................................................. 253 13.2 Experimental................................................................................................ 255 13.3 Results and Discussion................................................................................ 256 13.4 Conclusion...................................................................................................266 Acknowledgments...................................................................................................266 References...............................................................................................................266

13.1 INTRODUCTION The world’s energy demands will grow with increasing standards of living and economic growth. The energy crisis has become a serious issue with the consumption of oil reserves; however, most current energy consumption from oil reserves has caused serious environmental pollution problems. To reverse this crisis, we must develop renewable energy sources that are cost-effective as well as environmentally friendly. Biomass, wind, solar, geothermal, and hydroelectric power are well-known alternative energy sources. In particular, solar radiation may be a promising source of green energy. By an electrochemical process, solar radiation can be converted into electricity through photovoltaic devices. During the last few decades, thin films of crystalline silicon and semiconductor compounds have been developed for use in solar cells [1]. However, such devices have a high production cost. In 1991, Gratzel et al. discovered dye-sensitized solar cells (DSSCs), which are considered a potentially low-cost and highly efficient photovoltaic alternative to the traditional silicon-based solar cell [2]. DSSCs rely on high specific surface area and wide-band gap semiconductor oxides as a photosensitized anode for organic or metalorganic-complex 253

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dye molecules adsorption. They recently emerged as promising third-generation photovoltaic cells because they are flexible, inexpensive, and easier to manufacture than silicon-based solar cells [3]. In a DSSC, photons from sunlight strike the dye molecules, which excite electrons to elevate the dye to a higher energy level. The electrons are injected into the conduction band of the porous oxide photoanode. The oxidized dye molecule is regenerated by an electrolyte containing a redox system R/ R−, usually iodide/tri-iodide. Recently, a maximum conversion efficiency of approximately 11% DSSC was obtained using highly porous and nanocrystalline TiO2 films with ruthenium-based dye adsorption [4]. Numerous other semiconductors have also been used as photoanode materials, and a number of them, for example, ZnO [5], SnO2–ZnO composite [6], F-doped ZnO [7], SnO2–MgO core–shell nanoparticles [8], and Zn2SnO4 [9], have achieved outstanding results. In addition, a variety of nanostructures such as nanoparticles, nanowires/nanofibers, nanotubes, and nanoforests offer a large surface area for dye adsorption and/or a direct pathway for electron transport, which helps them achieve high conversion efficiency [5]. In addition to the desired structure of the photoelectrode film, the achievement of such high conversion efficiency for DSSCs is also attributed to the use of ruthenium-based dyes as the photosensitizer. These dyes, known as N3, N719, or black dye, are highly efficient in capturing photons within the visible wavelength region. More importantly, the photon-to-electron generation in these dyes has an extended excited-state lifetime, which facilitates its effective injection from the dye molecules to the semiconductor (approximately 100 fs) before radiative or nonradiative recombination occurs (approximately 15 ns) [10]. ZnO, one of the most common II–VI semiconductors, is considered a promising alternative to TiO2 in DSSC because it has a wide direct band gap (3.37 eV) and large exaction binding energy (60 meV) at room temperature [5, 11, 12]. DSSC with ZnO is thought to be advantageous because of the crystallization, electrical conduction, and molecular structure, which can be fabricated into various nanostructures. Furthermore, ZnO possesses an electron mobility of 115 to 155 cm2V−1s−1, which is seven orders of magnitude higher than approximately 10 −5 cm2V−1s−1 for TiO2 [13]. Numerous studies have focused on the application of ZnO in DSSC. Researchers have observed conversion efficiencies of 5% for ZnO nanoparticles [14], 3.3% to 3.9% for ZnO nanosheets [15, 16], 5.08% for nanoporous ZnO films [17], 0.83% for ZnO nanowires [18], 1.6% for ZnO nanotubes [19], 2.63% for ZnO nanoforests [20], and 6.1% for Li-doped ZnO [21]. The instability of ZnO in a ruthenium-based dye solution could cause poor photovoltaic performance in DSSCs. In addition, the formation of an inactive Zn2+/dye complex layer on the ZnO surface is responsible for lowering the injection efficiency of electrons from the dye molecules to the ZnO semiconductor [22]. A high-efficient DSSC based on nanoparticle photoanodes is often strongly dependent on its disordered geometrical structures and interfacial interference in electron transport, which affects the enhancement of the scattering of free electrons. The electron mobility is lessened because of the recombination of electrons with either oxidized species in the electrolyte or oxidized dye molecules before reaching the collection electrode. Single-crystalline ZnO nanowires have a unique structural feature that creates enhanced surface activities and rapidly and efficiently transports electrons to the collection electrode with the recombination suppressed [10].

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Another challenge of ZnO is to obtain nanofibers with smaller diameters (less than 100 nm), controllable density, and higher c-axis orientation. Well-aligned ZnO nanostructures can be prepared by electrochemical deposition (ED) [23], chemical vapor deposition (CVD) [24–27], vapor–liquid–solid (VLS) growth [28], and hydrothermal processing [29, 30]. The hydrothermal process is more convenient and economical for large-scale preparation of well-aligned ZnO nanowire or nanorod arrays. Recently, Greene et al. [31] and Zhou [32] successfully reduced the diameter of a ZnO nanofiber to approximately 40 nm by adding polyethyleneimine. Boyle et al. [33] and Ko et al. [20] developed a multistep approach that involves a precoating ZnO seed layer and a subsequent hydrothermal process to produce perpendicularly oriented ZnO nanorods and nanoforests. This study demonstrates that a well-crystallized, high-density, and high specific surface area of Li-ZnO nanoforest photoanodes can increase power conversion efficiency substantially. The materials’ characterization and the photovoltaic performance of the DSSC based on both Li-ZnO nanofiber and nanoforest photoanodes were investigated in detail.

13.2 EXPERIMENTAL Depending on the hydrothermal growth steps, two types of Li-ZnO photoanodes were prepared: nanofibers and nanoforest [20, 34]. The Li-ZnO seed precursor solution contained 0.75 M zinc acetate (Zn(CH3COO)2·2H2O, J. T. Baker, 99.8%), 2.2905 g diethanolamine (DEA, J. T. Baker 99%), and 0.075 M LiAc·2H2O (Aldrich, 98%) in 50 ml ethanol. This solution was heated to 60°C for 30 min and then cooled to room temperature. The FTO glass substrates (7Ω/□) were coated with seed precursor solution by spin coating five times and heating at 100°C for 30 min and at 300°C for 10 min. Subsequently, the FTO glass substrates coated with seed layers were placed in a 50-ml aqueous hydrothermal solution at 95°C for 1 to 4 h. The hydrothermal solution contained 0.05 M zinc nitrate (Zn(NO3)2·6H2O, 99.7%, J. T. Baker), 0.005 M lithium nitrate (LiNO3, 99.1%, J. T. Baker), 0.05 M hexamethylenetetramine (C6H12N4, 99%, Alfa Aesar), and 0.6 g polyethyleneimine (PEI, (–NHCH2CH2–)x(– N(CH2CH2NH2)CH2CH2–)y MW 1800, 99%, Alfa Aesar) in 50 ml deionized water. We adjusted the pH of the solution to 10.3 using NH4OH. After the hydrothermal process, the samples were washed in deionized water and dried at 100°C for 30 min, followed by heat treatment at 450°C for 2 h, creating Li-ZnO nanofibers. Li-ZnO nanoforests were grown by immersing the Li-ZnO nanofiber samples, which were coated with seed layers again and heated at 100°C for 30 min and 300°C for 10 min, followed by placement into a hydrothermal solution at 95°C for 5 to 8 h. After the hydrothermal process, the samples were washed with deionized water and dried at 100°C for 30 min before undergoing heat treatment at 450°C for 2 h. Before the solar cell assembly, the Li-ZnO films were dried at 100°C and then immersed in a ruthenium-complex dye (Esolar N719, C58H86O8N8S2Ru, Everlight Chemical Industrial Co.) with a concentration of 5 × 10 –4 M in acetonitrile in a dark bottle for 20 min. We then rinsed the samples with acetonitrile to remove excess surface dye and dried them at 70°C for 30 min. Sheets of weight paper with a thickness of 30 μm were used as spacers. A spacer was placed around the Li-ZnO photoanode,

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and the counter electrode, consisting of a Pt-coated ITO glass, was placed on the top. We then sandwiched the two electrodes together with sealing gel. We observed the morphology of the photoanodes using a field-emission scanning electron microscope (FE-SEM, JEOL JSM-6701F) and transmission electron microscope (TEM, JEOL JEM-2100). We used an atomic force microscope (AFM, Shimadzu SJ9500) to analyze the surface topology of the photoanodes and measure the current density versus voltage (I–V) characteristics of the cells under AM 1.5 100 mW/cm2 illumination from a solar simulator (Yamashita Denso YSS-50A) immediately after cell assembly. The incident photon-to-electron conversion efficiency (IPCE) value for DSSC was measured using an IPCE analyzer (Newport 96000).

13.3 RESULTS AND DISCUSSION Figure 13.1 shows the scanning electron microscopy (SEM) micrograph of a seed layer. The seed particles are closely packed to form a film. The particle size distribution is narrow, and the size is approximately 7 nm. Numerous cracks can be observed on the seed layer, which we attributed to shrinkage during the drying and heating process. Figure 13.2 shows the transmission electron microscopy (TEM) images of seed crystallites. The crystallite size distribution ranges from 2 to 14 nm (Figure 13.2c), and the percentage of crystallite in the range of 6 to 9 nm is over 80%. This means that the crystallites’ size distribution is narrow. This result is consistent with the SEM observation shown in Figure 13.1. The structure of the Li-ZnO seed is a wurtzite hexagonal structure, confirmed by the selected-area electron diffraction (SAED) pattern, which is shown in the upper-right corner of Figure 13.2a, in the presence of the (002), (100), (101), (102), (110), (103), and (200) diffraction rings, which consist of the JCPDS 89-0510 of ZnO. In the typical high resolution transmission electron microscopy (HRTEM) image (Figure 13.2b) of crystallite, the lattice fringes are clearly observed. The

FIGURE 13.1  SEM micrograph of seed layer. The inside is the high-magnification micrograph.

Li-ZnO Nanofiber and Nanoforest Photoanodes for DSSCs (a)

257

(200) (102) (101) (100)

(002) (110) (103)

10 nm (b)

1 nm (c)

25

15

10

Percentage (%)

Frequency (%)

20

5 Diameter (nm)

0 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 Diameter (nm)

FIGURE 13.2  (a) TEM image of seed layer (the inset is the corresponding selectivearea electron diffraction pattern), (b) HR-TEM image, and (c) crystallite size distribution histogram.

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distance between the adjacent lattice planes is 0.25 nm, which corresponds to the d-spacing of the (101) plane of the wurtzite ZnO (Figure 13.3). We investigated the surface topography of the FTO substrate and Li-ZnO seed using an АFМ, which is shown in Figure 13.4. The investigated surface topography is on the 15-μm × 15-μm scanned surface area. The surface altitude difference (Rmax) and root-mean-square roughness (R rms) values are summarized in Table 13.1. The seed layer film has a small surface irregularity and an agglomerated structure of Rmax and R rms of 24 nm and 180 nm, respectively, which is lower than the FTO substrate.

FIGURE 13.3  The AFM images of (a) FTO substrate and (b) seed layer.

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FIGURE 13.4  SEM micrographs of the Li-ZnO nanofibers grown on seed layer with different growing periods of (a) 1 h, (b) 2 h, (c) 3 h, and (d) 4 h. The inset is the high-magnification micrograph. (e) Schematic diagram of Li-ZnO nanofibers grown from seed layer in the hydrothermal process. (f) Diameter size distribution histogram of Li-ZnO nanofibers with a different growing period.

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TABLE 13.1 The Surface Altitude Difference (Rmax) and Root-Mean-Square Roughness (Rrms) of FTO Substrate and Seed Layer Rmax (nm)

Rrms (nm)

FTO

30

227

seed layer

24

180

Sample

The smooth seed layer is due to the nanosized and narrow particle-size distribution seed particle and is uniformly coated on the surface of the FTO glass. Figure 13.4 shows the typical SEM micrographs of the Li-ZnO nanofibers grown on the seed layer in a hydrothermal process from 1 to 4 h. These SEM micrographs show that the morphology of all samples is homogeneous and remarkably similar. The high-magnification micrograph shows that the upper portions of the nanofiber were closed to each other because of the van der Waals forces. The Li-ZnO nanofibers grow along the C-axis direction and are perpendicular to the FTO substrate, which is illustrated by the schematic diagram in Figure 13.4e. The average diameter of the Li-ZnO nanofibers does not strongly depend on the growth period. When the growth period increased from 1 to 4 h, the average diameter of these nanofibers was approximately 28 nm, with a narrow size distribution. Figure 13.5 shows the TEM images and SAED of the Li-ZnO nanofiber. The diameter of the Li-ZnO nanofiber was found to be approximately 32 nm. The corresponding SAED shows that the nanofiber is a wurtzite-structured single crystal with excellent crystallization and grows in the [0001] direction. The lattice finger is clearly observed with the HR-TEM image (Figure 13.5c). The diameter, length, and aspect ratio are listed in Table 13.2. A Li-ZnO nanofiber can be produced of 826 to 3358 nm, with an approximate diameter of 26 nm. The aspect ratio of nanofibers with a growth time of 1 h is 31.8, which rises to 125 as the hydrothermal growth period increases to 4 h. Positively charged PEI molecules are well known to adsorb on the lateral facets of ZnO nanorods because of their electrostatic affinity [32, 35]. Thus, the lateral growth of the nanofibers is largely limited but allows for the axial growth of the ZnO nanofibers in the solution, thus yielding high aspect ratio nanofibers [31, 34]. The hierarchically branched Li-ZnO nanoforest thin films could easily be grown in a large area using a low-cost, multistep hydrothermal process. The quality of the Li-ZnO nanoforests has been characterized by SEM and TEM. Figure 13.6 shows the typical SEM micrographs of the Li-ZnO nanoforests in a hydrothermal solution for 5 to 8 h. These SEM micrographs show that the hierarchically branched Li-ZnO nanofibers grow perpendicular to the vertical orientation of the first-generation backbone Li-ZnO nanofibers’ surface. The size of the second-generation nanofibers is approximately 25 nm, and they are closed to each other because of the van der Waals forces. The second-generation nanofibers grow along the C-axis perpendicular direction of large-size first-generation nanofibers. The average

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FIGURE 13.5  (a) TEM image, (b) the corresponding selective-area electron diffraction pattern, and (c) HR-TEM image of Li-ZnO nanofibers grown for 4 h.

TABLE 13.2 Diameter, Length, and Aspect Ratio of the Li-ZnO Nanofibers Grown with Different Growing Periods Growing Period (h) 1 Diameter (nm)

2

3

4

26 ± 3

26 ± 5

26 ± 8

27 ± 8

Length (nm)

826 ± 90

1106 ± 95

2660 ± 109

3358 ± 378

c/a ratio

31.8 ± 0.2

42.5 ± 6

102.3 ± 29

125 ± 25

diameter of second-generation nanofibers does not strongly depend on the growth period. When the growth period increases from 5 to 8 h, the average diameter of these nanofibers is approximately 25 nm with a narrow size distribution, which is similar to that of the first-generation nanofibers. Figure 13.7 shows the TEM images and SAED of the second-generation nanofibers. We found that this nanofiber is approximately 28 nm in diameter. The corresponding SAED shows that the second-generation nanofiber is a composite of nanocrystallite with a wurtzite

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FIGURE 13.6  SEM micrographs of the Li-ZnO nanoforests grown with different period of (a, b) 5 h, (c, d) 6 h, (e, f) 7 h, and (g, h) 8 h.

structure. The HR-TEM image (Figure 13.7c) shows that the nanocrystallite shows good crystallization. Figure 13.8 shows the I-V characteristics of the DSSC based on nanofibers and nanoforest photoanodes, and Table 13.3 summarizes the consequent characteristic data. The DSSC based on the nanoforest photoanode has a significantly higher short-circuit density (Jsc) of 8.92 mA/cm 2, FF value of 0.54, and efficiency

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FIGURE 13.7  (a) TEM image, (b) the corresponding selective-area electron diffraction pattern, and (c) lattice spacing image of Li-ZnO nanoforests grown for 8 h. 7 Nanofibers Nanoforests

Photocurrent density (mA/cm2)

6 5 4 3 2 1 0 0

100

200

300

400

500

600

700

Voltage (mV)

FIGURE 13.8  I–V curves of DSSCs based on Li-ZnO nanofibers and Li-ZnO nanoforests.

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TABLE 13.3 Jsc, Voc, FF, and η of DSSCs Based on Li-ZnO Nanofibers and Li-ZnO Nanoforests Jsc (mA/cm2)

Voc (mV)

FF

Efficiency (η, %)

Li-ZnO nanofibers

3.12

655

0.51

1.04

Li-ZnO nanoforests

8.92

570

0.54

2.74

Sample

60 Nanofibers Nanoforests 50

IPCE (%)

40

30

20

10

0 400

450

500

550 600 650 Wavelength (nm)

700

750

800

FIGURE 13.9  IPCE curves of DSSCs based on Li-ZnO nanofibers and Li-ZnO nanoforests.

of 2.74% than the DSSC based on the nanofiber photoanode of 3.12 mA/cm 2, FF value of 0.51, and efficiency of 1.04%. The short-circuit current density of the branched Li-ZnO nanoforests is almost three times higher than that of Li-ZnO nanofibers. The efficiency-respective enhancement is 163%. This is higher than the reported values for DSSCs based on ZnO nanowires [20]. Figure 13.9 shows the IPCE curves of these two DSSCs. We found that the photon-to-electron conversion efficiency was examined in the visible range from 400 nm to 700 nm. The highest IPCE value of 525 nm of the DSSC based on the nanoforest photoanode is approximately 53%, which is substantially higher than the DSSC based on the nanofiber photoanode (approximately 30%). Figure 13.10 shows the AC impedance spectra of these two DSSCs, and Table 13.4 lists the resistances and capacitor. The RS, Rsh, CPE, W1, and Rtotal are the resistance between the photoanode and FTO glass, charge transform resistance between photoanode/dye and electrolyte, double capacitor, diffusion resistance, and total resistance, respectively. We found that the

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Li-ZnO Nanofiber and Nanoforest Photoanodes for DSSCs 80

Nanofibers Nanoforests

70 60 Rs

Z’’ (ohm)

50

W1 s

Rsh CPE

40 30 20 10 0 0

10

20

30

40

50

60

70

80

90

100

Z’ (ohm)

FIGURE 13.10  Nyquist plots of the impedance data of DSSCs based on Li-ZnO nanofiber and Li-ZnO nanoforest.

TABLE 13.4 The Electron Transport Properties of Li-ZnO Photoanode Measured by Impedance Analysis Sample

Rs (Ω)

Rsh (Ω)

CPE (C)

W1 (Ω)

Rtotal (Ω)

Li-ZnO nanofibers

6.64

47.49

0.88

0.089

54.13

Li-ZnO nanoforests

6.39

33.06

0.83

0.085

39.45

*Rtotal = Rs + Rsh

major resistance starts from the charge transform resistance between the photoanode/dye and the electrolyte. The Rsh and Rtotal of the DSSC based on nanoforest photoanodes are 33.06 Ω and 39.45 Ω, respectively, which is substantially lower than the DSSC based on the nanofiber photoanode of 47.49 Ω and 54.13 Ω. We can explain this increase in efficiency by considering a combination of several effects. According to the IPCE result, we suggest that the significantly improved efficiency of the DSSC based on nanoforest photoanodes is due to the dye adsorbed amount, which apparently increases with the augmented specific surface area of nanoforest

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morphology and correspondingly to a large Jsc. In addition, light harvesting can be enhanced by scattering enhancement and trapping because of the random Li-ZnO nanotree multigeneration branches [5, 20]. Furthermore, because the AC impedance shows a lower total resistance, we suggest that nanoforest morphology provides more conduction paths for electron transport from the injection point to the FTO substrate. In this study, we found that the first-generation nanofiber is a single crystal; however, the second-generation nanofiber is a composite of nanocrystallites. We suggest that the photovoltaic performance of the DSSC based on nanoforest photoanodes could be further improved if the single-crystal structure of second-generation nanofibers can be prepared successfully.

13.4 CONCLUSION In this study, we successfully prepared well-crystallized Li-ZnO nanofiber and nanoforest photoanodes using a hydrothermal process that was applied to DSSC. We prepared single-crystal first-generation nanofibers with a diameter of approximately 32 nm and multicrystal second-generation nanofibers with a diameter of approximately 28 nm using approximately 7-nm Li-ZnO seeds. The high efficiency of 2.74% of the DSSC based on nanoforest photoanodes is due to the substantial amount of dye adsorbed on the high specific surface area and the lower resistance of the good crystallite Li-ZnO photoanode.

ACKNOWLEDGMENTS This work was financially supported by the Hierarchical Green-Energy Materials (Hi-GEM) Research Center, from the Featured Areas Research Center Program within the framework of the Higher Education Sprout Project by the Ministry of Education (MOE) and the Ministry of Science and Technology (MOST 107-3017-F006-003) in Taiwan.

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12. J.H. Choy, E.S. Jang, J.H. Won, J.H. Chung, D.J. Jang, Y.W. Kim, Adv. Mater. 15 (2003) 1911–1914. 13. E.M. Kaidashev, M. Lorenz, H.V. Wenckstern, A. Rahm, H.C. Semmelhack, K.H. Han, G. Benndorf, C. Bundesmann, H. Hochmuth, M. Grundmann, Appl. Phys. Lett. 82 (2003) 3901–3903. 14. K. Keis, E. Magnusson, H. Lindstrom, S.E. Lindquist, A. Hagfeldt, Sol. Energy Mater. Sol. Cells 73 (2002) 51–58. 15. E. Hosono, S. Fujihara, I. Honna, H.S. Zhou, Adv. Mater. 17 (2005) 2091–2094. 16. K. Kakiuchi, M. Saito, S. Fujihara, Thin Solid Films 516 (2008) 2026–2030. 17. Z. Chen, Y. Tang, L. Zhang, L. Luo, Electrochim. Acta 51 (2006) 5870–5875. 18. J.J. Wu, G.R. Chen, H.H. Yang, C.H. Ku, J.Y. Lai, Appl. Phy. Lett. 90 (2007) 213109–213111. 19. A.B.F. Martinson, J.W. Elam, J.T. Hupp, M.J. Pellin, Nano Lett. 7 (2007) 2183–2187. 20. S.H. Ko, D. Lee, H.W. Kang, K.H. Nam, J.Y. Yeo, S.J. Hong, C.P. Grigoropoulos, H.J. Sung, Nano Lett. 11 (2011) 666–671. 21. H. Horiuchi, R. Katoh, K. Hara, M. Yanagida, S. Murata, H. Arakawa, M. Tachiya, J. Phys. Chem. B 107 (2003) 2570–2574. 22. Q. Zhang, C.S. Dandeneau, S. Candelaria, D. Liu, B.B. Garcia, X. Zhou, Y.H. Jeong, G. Cao, Chem. Mater. 22 (2010) 2427–2433. 23. R. Liu, A.A. Vertegel, E.W. Bohannan, T.A. Sorenson, J.A. Switzer, Chem. Mater. 13 (2001) 508–512. 24. J.J. Wu, S.C. Liu, Adv. Mater. 14 (2002) 215–218. 25. S. Mitra, K. Sridharan, J. Unnam, K. Ghosh, Thin Solid Films 516 (2008) 798–802. 26. D.I. Suh, S.Y. Lee, T.H. Kim, J.M. Chun, E.K. Suh, O.B. Yang, S.K. Lee, Chem. Phys. Lett. 442 (2007) 348–353. 27. J.B. Baxter, E.S. Aydil, Sol. Energy Mater. Sol. Cells 90 (2006) 607–622. 28. M.H. Huang, S. Mao, H. Feick, H. Yan, Y. Wu, H. Kind, E. Weber, R. Russo, P. Yang, Science 292 (2001) 1897–1899. 29. S. Yamabi, H. Imai, J. Mater. Chem. 12 (2002) 3773–3778. 30. M.N.R. Ashfold, R.P. Doherty, N.G. Ndifor-Angwafor, D.J. Riley, Y. Sun, Thin Solid Films 515 (2007) 8679–8683. 31. L.E. Greene, B.D. Yuhas, M. Law, D. Zitoun, P. Yang, Inorg. Chem. 45 (2006) 7535–7543. 32. Y. Zhou, W. Wu, G. Hu, H. Wu, S. Cui, Mater. Res. Bull. 43 (2008) 2113–2118. 33. D.S. Boyle, G.K. Ovender, P. O’Brien, Chem. Commun. 1 (2002) 80–81. 34. M. Law, L.E. Greene, J.C. Johnson, R. Saykally, P. Yang, Nat. Mater. 4 (2005) 455–459. 35. W.J. Li, E.W. Shi, W.Z. Zhong, Z.W. Yin, J. Cryst. Growth 203 (1999) 186–196.

14

Monolithic Dye-Sensitized Perovskite Solar Cells Ming-Hsien Li, Kuan-Yu Lin, and Peter Chen

CONTENTS 14.1 Introduction.................................................................................................. 269 14.2 Monolithic Dye-Sensitized Solar Cells........................................................ 270 14.2.1 Carbon-Based Monolithic Dye-Sensitized Solar Cells.................. 270 14.2.2 Solid-State Monolithic Dye-Sensitized Solar Cells....................... 271 14.2.3 Metal-Based Monolithic Dye-Sensitized Solar Cells.................... 271 14.3 Mesoporous Electrode for Monolithic Perovskite Solar Cells.................... 272 14.3.1 Carbon-Based Mesoporous Electrode............................................ 272 14.3.2 Metal-Based Mesoporous Electrode............................................... 276 14.4 Conclusion................................................................................................... 279 References............................................................................................................... 279

14.1 INTRODUCTION Dye-sensitized solar cells (DSSCs) have become promising alternatives to conventional silicon-based solar cells due to their ease of fabrication and use of costeffective materials [1–3]. The conventional dye-sensitized liquid junction solar cells are based on the sandwich architecture that mainly employs two fluorine-doped tin oxide (FTO) glass substrates with the electrolyte sandwiched between them. One FTO substrate is coated with mesoporous TiO2, onto which the dyes are anchored, and serves as the working electrode for light harvesting and electron collection. The other one, acting as a cathode, is platinized to form ohmic contact with the electrolyte. It is noted that FTO glass is the major cost in the whole device [4]. Moreover, its heavy weight due to its thickness limits future application for lightweight wearable devices. For the aim of low cost and light weight, elimination of one FTO substrate for DSSCs was proposed to fabricate the monolithic device structure. The emerging perovskite solar cells (PSCs) have attracted tremendous attention due to their promising electrical and optical properties and versatile fabrication process, resulting in high power conversion efficiency (PCE) [5–8]. The hybrid organicinorganic perovskite possesses strong absorption over the visible range, promising 269

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ambipolar transport property with a long carrier diffusion length and high carrier mobility [5, 9–12]. Moreover, the facile solution-processed fabrication of perovskite makes it promising for cost-effective solar cells. These advantages make PSCs able to compete with the current existing photovoltaic technologies. The conventional perovskite solar cells are mainly composed of two selective contact electrodes, and a perovskite light absorber is sandwiched by them. Because perovskite is vulnerable to high-temperature processes and polar solvents, low-temperature processed soft chemistry is favorable for perovskite solar cells. For example, spin-coating of organic hole transport materials (HTMs) or electron acceptors such as [6,6]-phenyl-C61-butyric acid methyl (PC61BM) ester are performed after perovskite to avoid the degradation of the underneath perovskite film in a layer-by-layer n-i-p or p-i-n heterojunction structure. As inspired by the monolithic dye-sensitized solar cells,[13] mesoporous counter electrodes have recently been employed for PSCs, where the perovskite precursor is ultimately infiltrated in the final step for device fabrication. Thus, deposition of perovskite in the last step for device fabrication offers several merits. First, it allows us to apply high temperatures or harsh chemical processes for making selective electrodes and provides more selection of charge transport materials as well as their fabrication process. Some investigations have pointed out the use of organic HTMs with their corresponding additive and solvent can react with perovskite [14]. As a result, perovskite deposition as the last step in device fabrication prevents the unnecessary reaction to degrade the PSCs. Second, the demand for a large area smooth morphology is alleviated as the contact is a mesoscopic junction.

14.2 MONOLITHIC DYE-SENSITIZED SOLAR CELLS 14.2.1 Carbon-Based Monolithic Dye-Sensitized Solar Cells The pioneer work of monolithic DSSCs was demonstrated in 1996 by Prof. Micheal Grätzel, who applied SnO2-coated glass substrate as the transparent conductive oxide (TCO) electrode. On the top of TCO substrate, mesoporous anatase and rutile TiO2 layers were sequentially screen-printed to serve as working electrode and spacer, respectively. A mesoporous carbon electrode that worked as the counter electrode was screen-printed on the top of a mesoporous TiO2 spacer to replace the commonly used Pt-coated FTO substrate. Screen-printing electrodes allows us to lower the device cost and apply for a roll-to-roll fabrication process when the substrate is replaced by a flexible one. The device structure of monolithic devices is shown in Figure 14.1. The DSSCs based on the monolithic structure delivered a decent power conversion efficiency (PCE) of 6.67% [13]. In 2012, Seigo Ito’s group introduced mesoporous (mp) TiO2 as a working electrode, mesoporous ZrO2 as an insulator scaffold, and 10-μm-thick mesoporous carbon as a counter electrode for the monolithic DSSCs with a structure shown in Figure  14.2. Employing iodine-based electrodes for the monolithic device can prevent the increasing temperature damage on the cell when the device is illuminated with one sun. The resultant device achieved 2.06% PCE. The low-cost carbon applied as a mesoporous counter electrode can reduce the device cost and facilitate dye adsorption and electrolyte penetration due to the enhanced specific surface area [15].

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FIGURE 14.1  Schematic of monolithic DSSC device structure employing porous carbon counter electrode.

FIGURE 14.2  (a) Schematic of mesoporous-electrode-based monolithic DSSCs device. (b) The normalized efficiency vs temperature soaking test.

14.2.2 Solid-State Monolithic Dye-Sensitized Solar Cells In 1998, Udo Bach’s group applied solid-state hole transport material (HTM), SpicoOMeTAD, to replace the electrolyte and introduced gold as a counter electrode for the monolithic architecture with a structure shown in Figure 14.3. Although the device obtained a low efficiency of 0.73%, the all solid-state DSSCs attracted enormous attention due to the merits of application of a single FTO substrate and all solid structure. The all solid-state DSSCs can exclude the risk of electrolyte leakage. It is noted that this device structure is the prototype of the emerging perovskite solar cells (PSCs) [16].

14.2.3 Metal-Based Monolithic Dye-Sensitized Solar Cells Our group proposed a novel method to fabricate porous Ti/TiN/Ti composite thin film for the use of working electrodes in monolithic DSSCs as shown in Figure 14.4. TiN has been confirmed to show corrosion resistivity in contact with the I−/I3− redox couple electrolyte. Furthermore, nanostructured TiN composite films exhibit high electrical conductivity and superior catalytic activity compared to conventional Pt electrodes when it is in contact with the I−/I3− redox couple. However, owing to the low conductivity of TiN, a Ti thin film was first deposited onto the mp-TiO2 to

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FIGURE 14.3  (a) Schematic of solid-state monolithic DSSC device. 1: FTO transparent electrode; 2: compact TiO2 layer; 3: dye-sensitized heterojunction; 4: Au counter electrode. (b) J-V curves of solid-state monolithic DSSCs.

FIGURE 14.4  The fabrication process for the monolithic DSSCs with a porous Ti/TiN/Ti working electrode.

enhance the porous electrode conductivity. The addition of first Ti can also enhance the adhesion between mp-TiO2 and TiN. The third Ti was deposited to suppress the recombination loss between TiN and electrolyte. Before the sputter deposition of Ti/ TiN/Ti composite film, an 800-nm polystyrene (PS) ball was first spin-coated onto the mp-TiO2. The PS ball was further removed by immersing the template in dichloromethane with an ultrasonic bath to fabricate the porous Ti/TiN/Ti working electrode. The monolithic DSSCs with porous TiN/Ti/TiN electrode achieved a PCE of 3.16% and 2.92% for the device using Z-907 and MK-2 electrolyte, respectively [17].

14.3 MESOPOROUS ELECTRODE FOR MONOLITHIC PEROVSKITE SOLAR CELLS 14.3.1 Carbon-Based Mesoporous Electrode Inspired by the mesoporous carbon counter electrode applied for monolithic DSSCs [13], Han’s group first introduced mesoporous carbon electrode to replace the metal electrode in PSCs in 2013. The device architecture for PSCs is similar to that of DSSCs, which are mainly composed of compact (cp) TiO2 layer, mesoporous (mp) TiO2 layer, mesoporous ZrO2 insulator, and mesoporous carbon electrode (referred to Figure 14.5). It is noted that the monolithic structure can be fabricated using an all screen-printing method except for the deposition of the compact layer and perovskite. An inserted mp-ZrO2 layer between the mp-TiO2 layer and the carbon electrode can prevent direct contact between mp-TiO2 and mp-carbon and suppress the electron hole recombination. The monolithic PSCs obtained an efficiency of

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FIGURE 14.5  Schematic of mesoporous-carbon-based monolithic PSC device.

FIGURE 14.6  (a) Schematic of a mesoporous-carbon-based monolithic PSC device. (b) Energy level alignment of the device.

6.64%. Compared to the conventional sandwiched device, the monolithic structure eliminates the use of HTM and employs the low-cost carbon as the counter electrode. The application of a mesoporous carbon electrode further reduces the device cost due to the elimination of metal deposition, which usually requires a high-energy demand of a vacuum environment. However, to meet the conductive requirement, mesoporous carbon electrode usually reaches thickness over 10 μm [18]. Wang’s group employed the hetero p–n Schottky junction of mp-TiO2/mp-NiO for the mp-carbon-based mesoscopic device, which facilely separates electron and hole flow in the opposite direction. A high-conduction band level of −1.8 eV for mp-NiO can effectively block the electron. The MAPbI3 perovskite was filled from the pores of mesoscopic layers via sequential deposition and made into PSCs with the structure of FTO/cp-TiO2/mp-TiO2/mp-NiO/carbon shown in Figure 14.6. The monolithic device showed an efficiency of 11.4% [19]. Based on a similar architecture of mp-TiO2/mp-Al2O3/mp-C, Wang’s group introduced a single-walled carbon nanotube (SWCNT) in the mesoporous carbon counter electrode (referred to in Figure 14.7a) to improve the device performance. The SWCNT-doped carbon electrode modifies the surface wettability to facilitate the PbI2 penetration. Doped SWCNT in the carbon electrode further enhances the work function from 4.1 eV to 4.7 eV, which is compatible with the valance band of perovskite. The matching energy level facilities carrier transport with low energy loss. Compared to the undoped device, the SWCNT-doped device delivers an enhanced efficiency of 14.7% (Device C in Figure 14.7b) due to enhanced recombination resistance [20]. Liu et al. inserted an insulating layer of mp-ZrO2 between the mp-TiO2 and the NiO nanosheet to restrain the recombination occurring at the TiO2/NiO interface, resulting in an improvement of the carrier transport and resultant performance with

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14.2% PCE. Perovskite solar cells employing all inorganic metal oxide as framework were composed of FTO/cp-TiO2/mp-TiO2/mp-ZrO2/mp-NiO/C/sequential deposited MAPbI3 [21]. Analogous structure design was employed by Xu et al., who fabricated PSCs with the structure of mp-TiO2/mp-ZrO2/mp-NiO/C and delivered a noteworthy PCE of 14.9% [22]. The same architecture shown in Figure 14.8 was further demonstrated by Cao et al., who reported that the presence of mp-NiO increased the charge collection and reduced the charge recombination compared to the device without the NiO interlayer. A PCE of 15.03% was achieved by well-controlling the thickness of each layer. The mesoscopic cells maintained over 80% of their initial PCE after thermal aging over 1000 h at 60°C in dark and 82% of their initial PCE after light soaking for about 500 h. Such stability resulted from the over-thick carbon serving as a protection layer [23]. Two-step sequential solution deposition of perovskite loading is commonly used to penetrate the pores within the monolithic template. PbI2 is first filled in the template and reacts with MAI solution to transform MAPbI3 perovskite. The reaction rate between PbI2 and MAI has a significant impact on the MAPbI3 perovskite filling within the template. Since the MAI solution penetrates from the mesoporous counter electrode, initially crystallized MAPbI3 could hinder the MAI diffusion toward FTO substrate. As a result, the underneath PbI2 that is not totally converted into MAPbI3

FIGURE 14.7  (a) Schematic of a monolithic PSC device employed SWCNT-doped-carbon mesoporous electrode. (b) The J-V curves of SWCNT-doped device (Device C) compared with reference devices of Device A (undoped monolithic device) and Device E (conventional device with Spirco and Au).

FIGURE 14.8  Schematic of a mp-TiO2/mp-ZrO2/mp-NiO/C monolithic PSC device.

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hampers the light harvesting, carrier transport, and the resultant device performance. Prof. Eric Diau controlled the reaction rate by changing the solvent of MAPbI3, including dimethylformamide (DMF), dimethyl sulfoxide (DMSO), g-butyrolactone (GBL), and N-methyl-2-pyrrolidone (NMP), among which NMP effectively improves the MAPbI3 filling. NMP has low vapor pressure and volatility, which alleviate the MAPbI3 crystallization rate and facilitate the MAI penetration within the template. By well-controlled temperature and relative humidity, the NMP-assisted MAPbI3 perovskite crystallization completes with a growth period of 120 h. Such slow crystallization of MAPbI3 facilitates the growth of (004) facet with enhanced grain size. For the conventional thermal annealing process, the perceptional (220) facet of MAPbI3 perovskite leads to small grain size (referred to in Figure 14.9a). The NMP-assisted device delivers a superior PCE of 15% compared to the others owing to the effective pore filling of perovskite (referred to in Figure 14.9b) [24]. Although the over-thick mesoporous electron counter electrode prevents the perovskite from moisture penetration, the stability of perovskite solar cells can be further enhanced by using durable perovskite materials such as low-dimensional perovskite. Prof. Nazeeruddin doped a large-sized cation of aminovaleric acid iodide (AVAI) in the perovskite to synthesize 2D/3D hybrid perovskite and applied it to the carbon-based monolithic template. Compared to the 3D perovskite crystal structure, 2D/3D hybrid perovskite achieved enhanced stability and moisture resistance. The monolithic template incorporated with AVAI-doped perovskite delivered PCE of 12.9% (Figure 14.10a) with promising stability when the device was tested under one sun light soaking for 10,000 h (Figure 14.10c). The carbon-based monolithic template is feasible for a scale-up process with the screen-printing method. The asfabricated module delivered a decent efficiency of 10.1% (Figure 14.10b), revealing the potential for commercial application [25].

FIGURE 14.9  (a) X-ray diffraction (XRD) of NMP-assisted MAPbI3 crystallinity. (b) The J–V curves of monolithic device based on NMP-assisted, sequential, and DMF-assisted solution process of MAPbI3 perovskite.

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14.3.2 Metal-Based Mesoporous Electrode A thickness over 10 μm of the mesoporous carbon electrode is usually applied to improve electrical conductivity; however, a fragile over-thick carbon hinders the recycle application. Potential candidates for metal-based mesoporous electrodes, such as Au, Ni, and Au:NiOx, are further employed to replace mp-carbon electrodes. Yu et al. dealloyed the Ag/Au alloy leaf to form the nanoporous gold (NPG) electrode (Figure 14.11c) and transferred it onto the mesoporous substrate composed

FIGURE 14.10  (a) The J-V curve of a carbon-based monolithic device employing 2D/3D hybrid perovskite. (b) The perovskite solar module with 10 × 10 cm2. (c) The device stability test under one sun light soaking.

FIGURE 14.11  (a) The schematic of a nanoporous-gold-based monolithic device. (b) The energy level alignment of the device. The surface SEM image of (c) a dealloyed nanoporous gold electrode and (d) PbI2-coated NPG.

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of FTO/cp-TiO2/mp-TiO2/mp-Al2O3/NPG as presented in Figure 14.11a. The HTMfree device after the two-step sequential deposition of perovskite obtained a PCE of 7.99% [26]. Fan et al. employed a mp-Ni electrode to substitute the mp-carbon by reducing the mp-NiO into mp-Ni with an annealing process in the presence of hydrogen. Figure 14.12a shows the architecture of the mp-Ni-based template. After filling the MAPbI3 perovskite, the as-fabricated device achieved an efficiency of 13.6%. It is noted that the perovskite solar cells were degraded to 6.12% since the MAPbI3 perovskite was highly sensitive to the moisture. The author further removed the degraded perovskite and refilled perovskite within the template. The renewed device delivered a comparable efficiency of 11.63% (Figure 14.12b). The all-inorganic template integrated with porous metal electrode allows us to recycle the template [27]. However, conversions of metal into mesoporous metal electrodes, such as Ag elimination or NiO reduction into Ni, suffer from some redundant or unnecessary processes that increase the cost and energy required for fabrication. Our group demonstrated a monolithic structure composed of an all-inorganic template with a fabrication process shown in Figure 14.13a. The schematic of the device

FIGURE 14.12  (a) The schematic of a mp-Ni-based monolithic device. (b) The J–V curves of the as-fabricated device (Device A), degraded Device A, and rejuvenated device (Device A+).

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FIGURE 14.13  (a) The fabrication process of a np-Au:NiOx-based monolithic device. (b) The schematic and cross-sectional SEM of the fabricated device. (c) The J–V curves of the as-fabricated device (300/150) and recycled devices.

architecture along with the cross-sectional SEM of the as-fabricated device is presented in Figure 14.13b. Such devices allowed the deposition of the perovskite layer to be the last step during PSC fabrication and reuse of the expensive substrate. This enables the application of high temperature and ambient processes for the deposition of metal oxide within the template. Annealing the Ni/Au composite under ambient atmosphere allowed the gold to form an interconnected morphology and embed in the oxidized Ni [28]. After adjustment to the annealing conditions and the thickness of the mesoporous layer, a PCE of 10.25% was achieved. As-fabricated PSCs employed the template of mp-TiO2/mp-Al2O3/np-Au:NiOx. During the test of reuse, the original PCE of the PSC was 8.52%. In the test of template reuse, the device was first rinsed with organic solvent DMF to remove the perovskite. The template was followed by reloading the perovskite to exhibit a PCE of 8.17%. The deposition of perovskite was repeated for the third time and delivered a PCE of 7.72%. The photovoltaic performance of the recycling template is presented in Figure 14.13c. The decreased performance of recycled devices is presumably attributed to the passivated mp-TiO2 surface that might be resulting from the binding of residual lead iodine onto the mp-TiO2 surface after rinsing the MAPbI3 perovskite active layer [29]. After replacing with a new perovskite, the bound lead iodine could not convert into perovskite and form a passivation channel to hinder the carrier transport at the TiO2/perovskite interface. The lead iodine bound mp-TiO2 surface could also reduce the loading amount of perovskite because of its reduced specific surface area. With increasing recycle times, such an effect has an impact on the device performance in terms of reduced JSC and VOC simultaneously. However, the compatible efficiency of

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recycled devices, which indicates reusing the costly transparent-conducting-oxide coated glass (FTO substrate) along with all the other selective layers is promising for recycling applications [30].

14.4 CONCLUSION This chapter briefly reviews the monolithic dye-sensitized solar cells and perovskite solar cells, including carbon- and metal-based porous electrodes. Application of ultrathin glass substrate in the monolithic solar cells further reduces the weight of devices and provides flexibility. It is worth noting that mesoporous carbon electrodebased monolithic perovskite solar cells integrated with stable perovskite (such as 2D/3D hybrid perovskite) achieve a remarkable device stability as well as decent power conversion efficiency. However, the fragile carbon electrode is incompatible with the flexible substrate. The porous metal electrode-based monolithic perovskite solar cells employ stable perovskite, and ultrathin substrate is expected to serve as durable and flexible perovskite solar cells, which is beneficial for the internet of things, wearable devices, and flexible devices.

REFERENCES



1. O’Regan, B. and M. Grätzel, A low-cost, high-efficiency solar cell based on dye-sensitized colloidal TiO2 films. Nature, 1991. 353: p. 737. 2. Yella, A., et al., Porphyrin-sensitized solar cells with cobalt (II/III)–Based redox electrolyte exceed 12 percent efficiency. Science, 2011. 334(6056): pp. 629–634. 3. Mathew, S., et al., Dye-sensitized solar cells with 13% efficiency achieved through the molecular engineering of porphyrin sensitizers. Nature Chemistry, 2014. 6: p. 242. 4. Kroon, J.M., et al., Nanocrystalline dye-sensitized solar cells having maximum performance. Progress in Photovoltaics: Research and Applications, 2007. 15(1): pp. 1–18. 5. Park, N.-G., Perovskite solar cells: An emerging photovoltaic technology. Materials Today, 2015. 18(2): pp. 65–72. 6. Seo, J., J.H. Noh, and S.I. Seok, Rational strategies for efficient perovskite solar cells. Accounts of Chemical Research, 2016. 49(3): pp. 562–572. 7. Hsiao, Y.-C., et al., Fundamental physics behind high-efficiency organo-metal halide perovskite solar cells. Journal of Materials Chemistry A, 2015. 3(30): pp. 15372–15385. 8. Stoumpos, C.C. and M.G. Kanatzidis, Halide perovskites: Poor man’s high-performance semiconductors. Advanced Materials, 2016. 28(28): pp. 5778–5793. 9. Stranks, S.D., et al., Electron-hole diffusion lengths exceeding 1 micrometer in an organometal trihalide perovskite absorber. Science, 2013. 342(6156): pp. 341–344. 10. Li, Y., et al., Direct observation of long electron-hole diffusion distance in CH3NH3PbI3 perovskite thin film. Scientific Reports, 2015. 5: p. 14485. 11. Wehrenfennig, C., et al., High charge carrier mobilities and lifetimes in organolead trihalide perovskites. Advanced Materials, 2014. 26(10): pp. 1584–1589. 12. Ponseca, C.S., et al., Organometal halide perovskite solar cell materials rationalized: Ultrafast charge generation, high and microsecond-long balanced mobilities, and slow recombination. Journal of the American Chemical Society, 2014. 136(14): pp. 5189–5192. 13. Kay, A. and M. Grätzel, Low cost photovoltaic modules based on dye sensitized nanocrystalline titanium dioxide and carbon powder. Solar Energy Materials and Solar Cells, 1996. 44(1): pp. 99–117.

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14. Ito, S., S. Kanaya, H. Nishino, T. Umeyama, H. Imahori, Material exchange property of organo lead halide perovskite with hole-transporting materials. Photonics, 2015. 2(4): pp. 1043–1053. 15. Takahashi, S.I.a.K., Fabrication of monolithic dye-sensitized solar cell using ionic liquid electrolyte. International Journal of Photoenergy, 2012. 2012: p. 915352. 16. Bach, U., et al., Solid-state dye-sensitized mesoporous TiO2 solar cells with high photon-to-electron conversion efficiencies. Nature, 1998. 395: p. 583. 17. Shen, P.-S., et al., A novel porous Ti/TiN/Ti thin film as a working electrode for backcontact, monolithic and non-TCO dye-sensitized solar cells. Sustainable Energy & Fuels, 2017. 1(4): pp. 851–858. 18. Ku, Z., et al., Full printable processed mesoscopic CH3NH3PbI3/TiO2 heterojunction solar cells with carbon counter electrode. Scientific Reports, 2013. 3: p. 3132. 19. Liu, Z., et al., p-Type mesoscopic NiO as an active interfacial layer for carbon counter electrode based perovskite solar cells. Dalton Transactions, 2015. 44(9): pp. 3967–3973. 20. Li, H., et al., 14.7% efficient mesoscopic perovskite solar cells using single walled carbon nanotubes/carbon composite counter electrodes. Nanoscale, 2016. 8(12): pp. 6379–6385. 21. Liu, Z., et al., NiO nanosheets as efficient top hole transporters for carbon counter electrode based perovskite solar cells. Journal of Materials Chemistry A, 2015. 3(47): pp. 24121–24127. 22. Xu, X., et al., Hole selective NiO contact for efficient perovskite solar cells with carbon electrode. Nano Letters, 2015. 15(4): pp. 2402–2408. 23. Cao, K., et al., Efficient screen printed perovskite solar cells based on mesoscopic TiO2/Al2O3/NiO/carbon architecture. Nano Energy, 2015. 17: pp. 171–179. 24. Tsai, C.-M., et al., Control of preferred orientation with slow crystallization for carbonbased mesoscopic perovskite solar cells attaining efficiency 15%. Journal of Materials Chemistry A, 2017. 5(2): pp. 739–747. 25. Grancini, G., et al., One-year stable perovskite solar cells by 2D/3D interface engineering. Nature Communications, 2017. 8: p. 15684. 26. Zhou, X., et al., Hole-transport-material-free perovskite solar cells based on nanoporous gold back electrode. RSC Advances, 2015. 5(72): pp. 58543–58548. 27. Ku, Z., et al., A mesoporous nickel counter electrode for printable and reusable perovskite solar cells. Nanoscale, 2015. 7(32): pp. 13363–13368. 28. Lai, W.-C., et al., Oxidized Ni/Au transparent electrode in efficient CH3NH3PbI3 perovskite/fullerene planar heterojunction hybrid solar cells. Advanced Materials, 2016. 28(17): pp. 3290–3297. 29. Huang, L., et al., New films on old substrates: Toward green and sustainable energy production via recycling of functional components from degraded perovskite solar cells. ACS Sustainable Chemistry & Engineering, 2017. 5: pp. 3261–3269. 30. Li, M.-H., et al., Robust and recyclable substrate template with an ultrathin nanoporous counter electrode for organic-hole-conductor-free monolithic perovskite solar cells. ACS Applied Materials & Interfaces, 2017. 9(48): pp. 41845–41854.

15

High-Performance Quasi-Solid-State Polymer Electrolytes for Dye-Sensitized Solar Cell Applications Shanmuganathan Venketasan and Yuh-Lang Lee

CONTENTS 15.1 Introduction.................................................................................................. 281 15.2 General Aspects of DSSC Electrolytes........................................................ 283 15.3 Fabrication of DSSCs.................................................................................. 286 15.4 Characterization of DSSCs.......................................................................... 289 15.5 PAN-VA Polymer Gel Electrolytes.............................................................. 292 15.6 PVDF-HFP Polymer Gel Electrolytes......................................................... 301 15.7 PEO Polymer Gel Electrolytes..................................................................... 310 15.8 PMMA Polymer Gel Electrolytes................................................................ 317 15.9 Conclusion................................................................................................... 322 References............................................................................................................... 322

15.1 INTRODUCTION The large utilization of fossil fuels such as coal, oil, and natural gases, which are triggering environmental problems such as acid rain, air pollution, and global warming, has influenced the world to focus on a number of renewable energy sources for the future.1,2 Among the renewable energy sources, solar cells have gained significant attention as a solution to these issues.3 Solar cells are electrical devices that turn the energy of light into electricity by the photovoltaic effect.4 So far, crystalline silicon, gallium arsenide, cadmium telluride, and copper indium gallium selenide-based solar cells have shown energy conversion efficiency well over 20%.5,6 However, problems associated with high production cost and the complicated preparation procedures involved in preparing these kinds of solar cells should be eliminated.7 Dye-sensitized solar cells (DSSCs) have been considered a promising alternative to silicon and compound film solar cells because they are fabricated from cheap and benign materials and can be prepared in several attractive colors with transparency.8,9 Figure 15.1 represents the small DSSC and flexible DSSCs that prepared in our laboratory. 281

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FIGURE 15.1  The laboratory DSSC (a) and (b) flexible DSSC.

DSSCs using conventional ruthenium complex and zinc porphyrin dyes have obtained energy conversion efficiencies of 12% and 13%, respectively.10–12 Very recently, by using two metal-free organic dyes, DSSCs have obtained the best energy conversion efficiency of 14.3%.13 These high-efficiency DSSCs always require liquid electrolytes to achieve such high cell efficiencies. These liquid electrolytes have high ionic conductivity and good interfacial conduction to the photoelectrode. However, liquid electrolytes are not good for long-term use in DSSCs owing to possible leakage of electrolytes and evaporation of the volatile organic solvents, especially at elevated temperatures. These problems have led to an increased interest in the use of gelated liquid electrolytes, organic and inorganic hole transport materials, and room temperature ionic liquids (ILs).14–17 These materials have high stability at elevated temperatures that make them desirable for long-term use in DSSCs. However, the cells using these materials have lower energy conversion efficiencies as compared to the liquid-state DSSCs. This is due to the poor interfacial contact with photoelectrodes and low conductivity of these materials. Several research methods have been made to use polymer-based electrolytes instead of liquid electrolytes.18–20 Among them, polymer gel electrolytes (PGEs) are the most suitable electrolytes for DSSCs because of their relatively high thermal stability, high ionic conductivity, and high pore-filling ability inside the nanoporous TiO2 films. PGEs with in-situ gelation characteristics and electron-spun polymer membrane electrolytes have also been proposed to improve the pore-filling ability of highly viscous PGEs.21–26 Although PGEs have several advantages over solid- and liquid-state electrolytes, the gel network at high viscosity limits the ionic diffusion, ionic conductivity, and energy conversion efficiency. To increase cell efficiency further, many studies have been made to replace liquid electrolytes with nanocomposite electrolytes (NCEs).18,20 These NCEs are prepared by mixing nanosized inorganic materials or nanofillers (NFs) into the electrolytes. The use of NCEs instead of other electrolytes is an easy and effective approach to enhance the efficiency and stability of the quasi-solid-state dye-sensitized solar cells (QS-DSSCs). Owing to the several advantages of using NFs, many efforts have been devoted in recent years to develop printable electrolytes (PEs) based on polymer electrolytes and NFs for highly efficient laboratory cells and submodule DSSCs.27–29 These PEs can be utilized in the

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mass production of large area module cells and flexible DSSCs. The chapter reviews several important PGEs and PEs in view of their preparations and physical and electrochemical characteristics for application in QS-DSSCs. The effects and related mechanism of several types of NFs on the properties and performance of the PGEs in QS-DSSCs are described.

15.2 GENERAL ASPECTS OF DSSC ELECTROLYTES Since the best work of O’Regan and Gratzel in 1991, several approaches have been made to enhance the performance of DSSCs.9 A classical DSSC consists of a dyesensitized mesoporous TiO2 photoelectrode, an organic liquid electrolyte solution containing redox mediator, and a platinum counter electrode (Figure 15.2a). The solar energy conversion principle of the DSSCs involves photoexcitation of the dye, followed by electron injection into the conduction band of the TiO2 photoelectrode. The oxidized dye molecule is reduced by gaining electrons from the I− ions in the electrolyte, and the resulting I3− is reduced at the Pt-counter electrode (Steps 1–5 in Figure 15.1b).30,31 Among the several components of the DSSCs, the electrolytes play an important role in determining both cell efficiency and stability.18 Electrolytes not only accelerate charge transfer between the photoelectrode and counter electrode but also regenerate both photosensitizer and themselves while DSSCs are working. A very good electrolyte for application in DSSCs should be specified by (a) high ion diffusivity and conductivity; (b) high compatibility with photosensitizer and Pt films; (c) zero or low absorption in the range of visible light; (d) high electrochemical, chemical, optical, interfacial, and thermal stability to produce cells with high long-term performance at high temperatures; and (e) low cost and low toxicity. The liquid electrolytes consist of organic solvents (acetonitrile [ACN], 3-methoxy-propionitrile [MPN], valeronitrile [VN], γ-butyrolactone [GBL], ethylene carbonate [EC], propylene carbonate [PC], etc.), redox mediators or couples (iodide/triiodide [I−/I3−], cobalt2+/cobalt3+ [Co2+/Co3+]), and some additives (4-tert-butylpyridine [4-tBP], guanidinium thiocyanate [GUSCN], lithium ions [Li+ cation]).18,32 The organic solvent is an important component in liquid electrolytes, and it offers a medium for the

FIGURE 15.2  (a) Basic structure and components of the DSSCs and (b) working principle of the DSSCs.

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dissolution and diffusion of the redox mediator. The physical characteristics of solvents such as viscosity, dielectric constant, and donor number strongly influence the cell performance. The solvent viscosity is the main factor affecting the diffusion of ions in the electrolytes. The ionic liquids and water are also used as a substitute for organic solvents. The redox mediator is prepared by mixing parts of a chemical species existing in different oxidation states. The redox mediator acts as an oxidizing or reducing agent in an electrochemical reaction in DSSCs. The presence of additives in the electrolytes improves the performance of the DSSCs, although they are not involved in the fundamental photoelectrochemical process. They improve the TiO2 surface state, redox mediator potential, and shift of the conduction band edge and recombination kinetics, as well as photovoltaic parameters of DSSCs. PGEs comprise polymers and liquid electrolytes. The polymer generally is called a gelator. The polymer or copolymers frequently utilized are polyacrylonitrile (PAN),33–36 poly(acrylonitrileco-vinyl acetate) (PAN-VA),24 polyvinylidene fluoride (PVDF),37 poly(vinylidene fluoride-co-hexafluoropropylene) (PVDF-HFP),38 polyethylene glycol (PEG) or polyethylene oxide (PEO),39,40 and polymethylmethacrylate (PMMA)41,42 (Figure 15.3). These polymers involve various roles in PGEs such as support to the polymer gel, solidifying, absorbing, swelling, holding and contacting with liquid electrolyte.19,43 The solvent of DSSC electrolyte, often called the plasticizer, offers the space and medium for ionic salt movement. The solvent decreases the glass temperature and crystallization of PGEs. This is because the solvent inhibits the polymer-polymer chain interaction and increases the free volume and segmental motion of the PGEs. When blending the liquid electrolyte with the polymer, the system changes slowly from a sol state to gel state. In this process, due to weak interaction between the gelator and the solvents, the PGEs are attained by the gelation, adsorption, inflation, and entanglement network of the gelator in liquid electrolytes. PGEs always show the liquid electrolyte characteristics with high electrolyte conductivity. This is due to the three-dimensional networks in the polymer matrix that provide cages for liquid electrolyte storage and channels for quick ion mobility. Since liquid electrolytes are trapped into the polymer cages, the solvent evaporation is minimized, which improves the stability of the DSSCs. However, the polymer network inhibits the charge transport to a certain extent, triggering a negative effect on cell performance. The high viscous PGEs also have poor penetration with the mesoporous photoelectrode. To solve this problem, in-situ polymerization of preinjected liquid electrolytes into PGEs has been introduced.44,45 However, the polymerization of a monomer is difficult to carry out in a solution containing I3−/I− redox couple because iodine deactivates radical intermediates.45 In response, some kinds of PGEs with in-situ gelation characteristics are proposed to improve the penetration ability of PGEs in the nanostructured TiO2.22–24 For the in-situ gelation systems, the gelation process is carried out at room temperature, and therefore, the PGEs can be injected into the DSSCs at the liquid state and then perform gelation inside the DSSCs. Electrospun polymer nanofiber gel electrolytes also have high porefilling ability to the TiO2 film, which is due to the ultrathin diameter nanofiber containing liquid electrolytes within the diameter nanofiber.25,26 The main advantage of the electrospinning process is related to the high degree of control that is available over porosity, morphology, and composition using simple equipment.

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FIGURE 15.3  Structure of the important polymers used for the preparation of polymerbased electrolytes.

Nanosized inorganic materials or NFs such as carbonaceous materials (activated carbon [Ac], carbon black, carbon dots, carbon sphere, graphite [Gra], graphene oxide, graphene [Gr], single-wall carbon nanotube, multiwall carbon nanotube, carbon nanohorn, and graphene oxide sponge [GOS]), metal oxides (silicon dioxide [SiO2], aluminum oxide, titanium dioxide [TiO2], zinc oxide, tin oxide [SnO2], cobalt oxide, nickel oxide, and copper aluminum oxide), metal carbides (titanium carbide [TiC]), metal nitrides (aluminum nitride) and metal sulfides (cobalt sulfide [CoS]), and layered double hydroxide (LDH) are added to the electrolytes to form quasi-solid-state electrolytes.46 The electrolytes are often termed nanocomposite electrolyte (NCEs). These NFs have different physical, chemical, and electrochemical properties. They have their own effect on the performance of the electrolytes in DSSCs. These effects are not interrelated because they follow various kinds of mechanisms aimed to boost the performance of DSSCs. These mechanisms are mostly related to the surface functional groups of the NFs and the components of the electrolytes. Several effects have been proposed for NFs, including (a) increasing the amorphous nature of polymers, (b) creating a transfer channel for redox mediators, (c) enhancing ion diffusivity and electrolyte conductivity, (d) increasing the charge transfer resistance at the counter electrode/electrolyte interface (RPt), (e) decreasing the charge transfer resistance at the photoanode/electrolyte interface (Rct), and (f)

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increasing the stability of polymer and electrolytes at high temperatures. Although NCEs have several advantages, they are not suitable for the fabrication of large-area module cells. This is because of the difficulties in the injection of electrolytes into the large area cells. The use of PEs pastes in QS-DSSCs eliminates the PGE injection problems.27–29 The PEs easily apply on the surface of the photoelectrode through the printing process such as the doctor blade or screen-printing techniques.47,48 The viscosity of the electrolyte determines the performance of the printable process in QS-DSSCs. By altering the composition of polymers in the PGEs, the optimal viscosity of the electrolyte for the printable process is obtained.29,48

15.3 FABRICATION OF DSSCS The DSSCs are assembled using fluorine-doped tin oxide (FTO) as the conductive substrate for the photoelectrode and counter electrode after a thorough cleaning.49 First, the FTO glass is cleaned sequentially with a neutral cleaner, deionized water, acetone, and isopropyl alcohol. Second, a TiCl4 underlayer or blocking layer is formed on the FTO substrate to prepare the photoelectrode. Next, an 8-µm-thick TiO2 layer, containing 4-µm-thick main layers (PST-18NR) and a scattering layer (PST 400C), is coated on the FTO using a screen-printing method. The electrode is dried at 70°C for 10 min and then sintered at 500°C for 30 min in an air atmosphere. The TiO2 film is immersed in 0.5 mM ethanol-based N719 dye solution for 24 h. After rinsing with ethanol, dye-sensitized TiO2 photoelectrode is obtained. Third, the counter electrode is prepared by coating platinum (Pt) on the FTO substrate with holes using a sputter coater (108auto, Cressington Scientific Instruments), operating at 40 mA for 105 s. For the fabrication of the liquid-state and gel-state DSSCs, the dye-coated photoelectrode is assembled with the Pt-counter electrode containing holes using a 60-µm-thick spacer (Figure 15.4a). The electrolytes are injected into the cell via the holes using a vacuum pump (Figure 15.4b). After the DSSCs are filled with electrolytes, the holes are sealed with epoxy resin (Figure 15.4c). In contrast, for the assembly of the PE-based DSSCs, the bare FTO surface of the TiO2/FTO is first plastered with Surlyn film (30 µm). Then, the electrolyte paste is printed on the TiO2 layer using a doctor blade printing method (Figure 15.4d).47,48 After the Teflon tape is detached, a Pt-FTO and the photoelectrode with electrolyte paste are assembled using a 30-µm Surlyn film (Figure 15.4e,f). Figure 15.5a,b also explains the schematic of the photoelectrode and counter electrode process for the fabrication of DSSCs. The assembly of the four-strip (5.0 × 5.0 cm 2) and rectangular-shaped (3.0 × 6.5 cm 2) submodule cells is similar to that of the laboratory DSSCs (Figure 15.5 c).27,48 FTO glasses are cut into 5.0 × 5.0 cm 2 and 3.0 × 6.5 cm 2 sheets, cleaned with acetone and ethanol, and then used as substrates for the preparation of photoelectrodes and counter electrodes. For the preparation of the photoelectrode, the cleaned FTO glass is treated with TiCl4, and then TiO2 paste is printed on the TiCl4 -FTO using a screen-printing method. The 8 + 4-µm-thick TiO2 film is dried at 100°C for 10 min between each printing. The active area for the submodule cells is 11 cm 2. Silver lines are printed around the TiO2 layers and annealed at 450°C for 30 min. These silver lines are covered with both a Surlyn film (60 µm thick)

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FIGURE 15.4  Schematic representation of liquid and gel-state DSSC assembly using (a–c) injection process and (d–f) printable process.

FIGURE 15.5  Schematic of (a) photoelectrode process, (b) counter electrode process, and (c) submodule cell assembly process.

and a commercial polyester film (0.1 mm thick) via hot pressing. The four-strip and rectangular-shaped TiO2/FTO is immersed in N719 and Z907 dye solutions, respectively, and then rinsed completely in ethanol to remove any excess dye. After drying, the polyester film is removed, and PE electrolyte paste is printed on the dye-sensitized TiO2 electrodes using the doctor blade method. Electrolyte-printed photoelectrodes are obtained after removing the Surlyn film. Counter electrodes

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are prepared by depositing Pt thin films onto FTO glass embedded with silver lines using sputtering. The PE-coated photoelectrode and counter electrode are sealed using Surlyn film with a hot pressing method; afterward, complete submodule cells are obtained (Figure 15.6). In recent years, polymer nanofiber electrolytes are also utilized for DSSCs.50 By using an electrospinning technique, polymer nanofibers with diameters in the range from several micrometers down to tens of nanometers are prepared. The nanofiber is prepared from a solution of the polymer in a mixture of acetone and N′Ndimethylacetamide (7:3 wt.%) at 80°C with constant stirring to get a homogeneous polymer solution, which is then cooled to room temperature.25,26 A schematic representation of the experimental setup is shown in Figure 15.7. The polymer solution is supplied to the stainless steel needle using a syringe pump, and a high voltage of 12

FIGURE 15.6  (a) Four-strip and (b) rectangular shape submodule QS-DSSCs.48

FIGURE 15.7  Schematic representation of experimental setup for the preparation of the electrospun membrane.50

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kV is then deposited onto a grounded, polished stainless steel plate, where the tip ground distance is fixed at 12 cm. The membrane is vacuum-dried overnight at 80°C to remove any remaining solvent and then hot pressed to reduce the thickness from 30 to 15 µm. The electrospun polymer membrane is soaked in a liquid electrolyte to obtain the electrospun membrane electrolyte.25,26

15.4 CHARACTERIZATION OF DSSCS DSSCs are generally evaluated by energy conversion efficiency that is calculated from the current density–voltage (J–V) curve of the cell.51 This curve is measured using a source meter (6240 A, computer controlled, ADCMT, Japan) under simulated sunlight. The light is produced using a solar simulator (XES, A.M 1.5 G, AAA class, San-EI Electric, Japan). By illuminating the DSSC, scanning the voltage, and measuring the current using a source meter/potentiostat, one can plot the J–V curve of the cell. The typical J–V and power (P)–V curves are given in Figure 15.8. The photocurrent at 0 voltage and photovoltage at 0 current are called shortcircuit current (Isc) or short-circuit current density (Jsc) and open-circuit voltage (Voc). The maximum power output (Pmax) produced by a DSSC is reached when the product of current and voltage is maximal. The Pmax is calculated using the following equation:

Pmax = Vmp × I mp = Vmp × J mp × Acell

where Imp = the current at the maximum output power, Vmp = the voltage at the maximum output power, Jmp = the current density at the maximum power output, Acell = the illuminated cell area. The overall energy conversion efficiency (η) is described by the ratio of the Pmax to the solar energy input or incident power (Pin). The Pin is the irradiance illuminated

FIGURE 15.8  J–V and P–V curve of the DSSCs.

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on the DSSCs that is measured in W/m2. The efficiency is between 0% and 100%. The η of the cell is calculated using the following equation:

η = Pmax / Pin × 100 = J mp × Vmp / Pin × 100 (15.1)

The fill factor (FF) is also an important parameter that determines the quality of DSSCs. FF is a value between 0 and 1 that explains the shape of the J–V curve. A high FF value implies a more preferable rectangular shape explained by the following equation:

FF = Pmax / J sc ´ Voc = J mp ´ Vmp / J sc ´ Voc (15.2)

By substituting equation (15.2) into equation (15.1),

h = J sc ´ Voc ´ FF / Pin ´ 100

Electrochemical impedance spectrometry (EIS) is a powerful steady-state technique that is used to elucidate the charge-transport behavior in the DSSCs.51 EIS analyses are performed either under dark or sunlight using a potentiostat (PGSTAT 30, Autolab) equipped with an FRA module. The impedance spectra are recorded at frequencies ranging from 100 kHz to 10 mHz at an AC amplitude of 10 mV. The spectra obtained are analyzed using Z-view software with appropriate equivalent circuits. The conductivity of the DSSC electrolytes is also measured using EIS. For this purpose, dummy cells are prepared using electrolyte samples and two Pt/ FTO electrodes with the spacer. The Nyquist plots of the EIS spectra, including two and three semicircles for the dummy cell and DSSC, respectively, are presented in Figure 15.9. They correspond, from left to right, respectively, to series resistance (Rs), which includes the sheet resistance of the conductive glass and the contact resistance of the cell, and the RPt, Rct, and Nernst diffusion within the electrolyte (Rd). The RPt and Rct are obtained by fitting the EIS spectra with an equivalent circuit shown in the inset. In this circuit, CPE1 (Cct) and CPE2 (Cµ) describe double-layer capacitance at the counter electrode and chemical capacitance of the TiO2 film. In addition to R Pt, Rct, and Rd, electron density (ns), capacitance, and electron life time (τn) values are

FIGURE 15.9  EIS curves of the (a) dummy cell and (b) DSSCs using liquid electrolytes.

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extracted from the Nyquist plots of the EIS spectra. The shape of the Nyquist plots of DSSCs is dependent on the light illumination conditions and applied voltage. The conductivity of the electrolyte is obtained by the following equation:

s = 1 / Wd ´ L / A

where Wd = Warburg diffusion, L = is the thickness of the electrolyte, and A = area of the electrode. The diffusivity of ions (D) in the electrolyte is measured using cyclic voltammetry. For this purpose, the dummy cells are generally used. Figure 15.10 shows the cyclic voltammogram (CV) of dummy cells based on the liquid and polymer gel electrolytes that recorded at a scan rate of 5 mV/s in the potential range of −0.8 to 0.6 V. The D in the electrolytes is evaluated from the saturated current conditions in cyclic voltammograms, where the ion diffusion in the electrolyte is limiting. The D in the electrolyte is proportional to the saturated current density (J) according to Fick’s law:

J = 2n FCD / l

where n is the number of electrons transferred during the reduction of I3− ion (i.e., n = 2 for iodide electrolyte), F is the Faraday’s constant, C is the bulk concentrations of the I3− ions (for iodide electrolyte), D  = diffusion coefficient of the I3− ions, and l is the thickness of the electrolyte layer. Another measurement of the performance of DSSCs is the external quantum efficiency, which is called the incident photon to current conversion efficiency (IPCE).29

FIGURE 15.10  (a) Steady-state cyclic voltammograms of the liquid and PGEs and (b) IPCE spectra of the cells using liquid and PEs.29

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The IPCE spectra are obtained using a quantum efficiency measurement system (QR-R3011, DC mode, Enlitech, Taiwan). Figure 15.10b shows the IPCE spectra of the cells using liquid and PEs. The IPCE value relates to the Jsc produced in the external circuit under monochromatic illumination of the DSSCs divided by the photon flux that strike the DSSC.

15.5 PAN-VA POLYMER GEL ELECTROLYTES PAN, due to its high cation solvating capacity, has been generally employed as a gelator to prepare PGEs for lithium-ion batteries. For the use of PAN in DSSCs (Table 15.1), the first problem experienced is the insoluble nature of PAN in the solvents generally utilized for liquid cells, that is, ACN and MPN. Therefore, EC and/or PC were used as solvents of PAN. However, the introduction of EC and PC triggers a higher viscosity and lower mechanical properties of PGEs, which are not suitable for the performance of a PGE.33–36,52 The highest efficiency of gel-state DSSC obtained by PAN-based PGE was only 7.27%.34 To increase the performance of the PAN PGE in DSSC, the fumed SiO2 NFs and PAN composite were prepared by Mohan et al.53 The composite was amorphous due to the interaction between the fumed SiO2 and the PAN polymers, as confirmed using XRD analyses. The surface functional group of the silica had participated in the reduction in ion pairing. The author proposed that the transportation of mobile ions occurred between the SiO2 NFs through space charge layers in the composite electrolytes. These two factors contributed to the ionic conductivity of the composite electrolyte. The conductivity of the composite electrolytes was 1.32 × 10 −3 S cm−1. The cell using PAN and 12 wt.% SiO2 NF-based electrolyte had a high efficiency of 7.51%. In 2015, our group prepared the PANbased PEs for DSSCs.27 In this study, we observed that the suitable viscosity of the PE reduced both the ionic conductivity and energy conversion efficiency. This issue was resolved by the addition of TiO2 NFs into the PAN PEs. The cell based on the 4 wt.% TiO2 NFs and PAN electrolyte had an efficiency of 7.85%, which was similar to that of the cell based on liquid electrolyte. The QS-DSSC using 4 wt.% TiO2 NFs had a low RPt value of 0.98 Ω.cm2. The presence of TiO2 NFs caused a gradual desorption of dye in the stability test, which decreased the DSSC performance. This problem was solved by using TiC NFs. The ionic conductivity in the 4 wt.% TiC NFs/PAN PE was higher than in the pristine PAN PE. The cell based on the 4 wt.% TiC NFs/PAN had an efficiency of 7.68%, which was higher than that of the efficiency of the cell based on the pristine PE. The cell using TiC NFs had a high stability. The cell efficiency retained 96% of its initial value after a 600-h testing. Mohan et al. prepared Ac/PAN composite electrolyte film.53 The composite film had a thickness in the range of 25–30 µm. The film had some pores (~5 µm diameter) as evidenced by scanning electron microscope (SEM) analysis. In the composite electrolyte, the pores were filled with ionic liquid, which permitted the iodide ions to move easily from the photoelectrode to counter electrode. The ionic conductivity in the PAN/5 wt.% Ac was 8.67 × 10 −3 S cm−1. In the PGEs, the interaction between the carbon particles and the structure of the surrounding liquid electrolytes formed space charge layers. High ionic conductivity occurred through a special conduction pathway created by overlapping of the space

PAN (0.225 g)

PAN (0.025 g)

— PAN PAN (9 wt.%)

PAN PAN (10 wt.%)

— PAN-VA (7 wt.%)

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5

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TiO2 sphere (