Advances in Sustainable Energy: Policy, Materials and Devices [1 ed.] 3030744051, 9783030744052

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Yong-jun Gao Weixin Song Jingbo Louise Liu Sajid Bashir   Editors

Advances in Sustainable Energy Policy, Materials and Devices

Advances in Sustainable Energy

Yong-jun Gao • Weixin Song • Jingbo Louise Liu • Sajid Bashir Editors

Advances in Sustainable Energy Policy, Materials and Devices

Editors Yong-jun Gao Hangzhou Branch of the Zhejiang Tsinghua Yangtze River Delta Research Institute Clean Energy and Energy Conservation and Environmental Protection Centre Hangzhou City, China Jingbo Louise Liu Texas A&M University – Kingsville Kingsville, TX, USA

Weixin Song Department of Materials University of Oxford Oxford, UK Sajid Bashir Department of Chemistry Texas A&M University – Kingsville Kingsville, TX, USA

ISBN 978-3-030-74405-2 ISBN 978-3-030-74406-9 https://doi.org/10.1007/978-3-030-74406-9

(eBook)

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

Preface

At the time of starting this book with Springer Nature, I lost my doctorate advisor (‘academic father’), and at the start of writing for the actual monograph, I lost my biological father from severe acute respiratory syndrome (SARS), a ribonucleic acid encoded virus with a corona geometry. At present SARS is probably the most wellknown phrase and has affected everyone in all sectors of society. “To lose one parent, Mr. Worthing, may be regarded as a misfortune; to lose both looks like carelessness.”―Oscar Wilde, The Importance of Being Earnest. Like, Mr. Wilde, I reflect on, where we are as a society. During this reflection period, when most of us are working remotely, teleconferencing, or are limited in how many people can work in any laboratory, we have to ask is Sustainable Energy Necessary? and if yes, what is the likely energy production roadmap over the next five decades? This monograph aims to address this central thematic question and is split into several themes, relating to energy generation from coal, hydrogen, thermal heat, light, electrons, and associated storage to ensure a continuous supply of electricity as demanded by industry, residual, home and businesses. The authors cover themes related to the above mission and in this preface as editors will summarize and expand the scope (1-Bashir) contained within the pages of this monograph. The first three chapters deal with photocatalysis concerning solar energy (2- Li/Kong), hybrid nature-derived mimetics (3-Bai/Ran), and the first principle-based density level theory calculations (4-Tang). The next four chapters deal with the production of hydrogen or oxygen or their utilization, with clean hydrogen production technologies (5-Abdullah), from metal hydride (6-Corgnale) and oxygen reduction using ruthenium electrocatalysts for direct methanol fuel cell devices (7-Uribe-Godínez) and direct conversion of carbon dioxide to oxygenates (8-Spivey). We return to the theme of solar energy is respect to solar thermal fluids thermodynamics (9-Kalita), modeling (10-Fox/Bashir), and electrolysis (11-Haik).

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Preface

The theme of energy storage is dealt with in the next two chapters with a focus on transition metal-based electrochemical storage (12-Huda) and an emphasis on mixed vanadium oxides (13-Li). We then explore an application of certain metal oxides in the form of metal-organic frameworks for water adsorption (14-Wang), advanced materials for supercapacitors (15-He), or metal layered two-dimensional MXenes for electrochemical storage as a departure from traditional metal oxides, addressed earlier in the monograph (16-Luo). The role of solid-electrolyte in high-density lithium-ion batteries is discussed next (17-Qiao) and the modern characterization approach to better understand energy materials is further explored as a standalone topic for X-ray energy (18-Tan/Shearing) or advanced transmission electron microscopy(19-Robertson). We then break the next segment to discuss clean coal (20-Meyer), coal and biomass gasification (21-Guan), energy from lignocellulosic materials (22-El-Azazy/Shibl), and how these materials can be used for sustainable development. Two examples are used to illustrate this approach. The role of Municipal solid waste incineration (23-Quina) and carbon sugars as fuels for microbial fuel cells, both theory and actual practice is extensively reviewed (24-Liu). The last segment policy, energy materials and electrification in china in two perspectives, one from the application of artificial intelligence in the monitoring of smart devices and applications such as electric vehicles (25-Gao) and the application of blockchain, smart devices and photovoltaic/wind/hydrogen stored energy in a delocalized grid as a framework for renewables energy (26-Gao) and the monograph ends with a general review of themes covered by one of our editors (27-Song) with the back matter subject index. All authors justify the role of energy storage or application of sustainable energy to lessen the influence of carbon dioxide that is derived from the combustion of coal to generate electricity. Fossil fuels such as coal and natural gas are the mainstay for electricity generation both within the United States and globally as observed in figure 1, however at the cost of reliable and relatively cheap electricity from coal is the expense of emissions from soluble oxides of carbon, nitrogen, and sulfur, as well as particulate matter. This has an environmental cost associated with it and will contribute towards respiratory distress, which in turn will enhance the susceptibility of a slice of the population towards SARS. Lastly, should carbon be used to generate thermal heat or used to manufacture materials, drugs, plastics, paints and other useful resources is a question that is being asked to generate a circular economy where waste is minimized and each atom is utilized, a form of conversation of matter?. The authors within this monograph highlight a few of the approaches that are currently under development at the technical level of attainment. The review and case studies are broad because the challenges are diverse, such as the development of clean biofuels for aviation, transportation and heating and cooling, waste to energy conversion technologies, a value-added material transformation from waste and life cycle assessment. On the proceedings page, we as the editorial team will further develop some areas that we consider emerging and pivotal towards the development of strategies to engineering a future with zero carbon emissions and cheaper energy resources that are sustainable.

Preface

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Any project, be it a greeting card, a book chapter, a review article, or a grant proposal is an intense endeavor and almost always a team effort; this monograph is no different, which we acknowledge here. Charlotte Hollingworth, for the day-today management, pre-print layout of the various manuscripts and Solomon, Cynthia, and Sharon also at Springer (Springer Science+Business Media, LLC) for taking an idea based on our work within the American Chemistry Society (ACS) division of energy and fuels (ENFL, Liu and Bashir), and support from Texas A&M UniversityKingsville. Lastly, we pay our respects to Peter J Derrick and Mohammed Bashir who died recently. A short biography can be found here for Dr. Derrick at https://doi.org/ 10.1177/1469066717739174. Thank you, Peter, and Dad your legacy will endure. Kingsville, TX, USA

Sajid Bashir

Kingsville, TX, USA

Sai Chava

Oxford, UK Hangzhou City, China Kingsville, TX, USA

Weixin Song Yong-jun Gao Jingbo Louise Liu

Contents

1

2

3

4

Promising Clean Energy Development: Practice, Challenges, and Policy Implications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Sajid Bashir, Sai Chava, Weixin Song, Yong-jun Gao, and Jingbo Louise Liu

1

Applications and Fundamentals of Photocatalysis with Solar Energy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Qiuyang Huang, Dan Kong, and Yongdan Li

27

Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells: Design, Fabrication, and Performance . . . . . . . . . . . . . . . . . . . . . . Dan Bai, Huhu Feng, Xingchen Yu, Chenxin Ran, and Wei Huang

67

First-Principles Calculations for the Interfaces of Perovskite Solar Cells . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Jun-Peng An, Ying Tian, Hong-Tao Xue, Jun-Chen Li, Jun-Qiang Ren, Xue-Feng Lu, and Fu-Ling Tang

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5

Clean Hydrogen Production Technologies . . . . . . . . . . . . . . . . . . . . 159 Mohammad Yusuf, Mohamad Sahban Alnarabiji, and Bawadi Abdullah

6

Thermal Hydrogen Compression Based on Metal Hydride Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 171 Claudio Corgnale, Robert C. BowmanJr., and Theodore Motyka

7

Oxygen Reduction Reaction Performed by Ru-Based Catalysts . . . . 193 J. Uribe-Godínez and A. Altamirano-Gutiérrez

8

Direct Catalytic Low-Temperature Conversion of CO2 and Methane to Oxygenates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 227 Ashraf Abedin and James J. Spivey

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Contents

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Heat Transfer Analysis in Solar Thermal Collectors . . . . . . . . . . . . 251 Pankaj Kalita, Dudul Das, Samar Das, Rabindra Kangsha Banik, and Urbashi Bordoloi

10

Heat Transfer Fluids in Concentrating Solar Power Systems: Principle and Practice . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 279 Elise B. Fox, Sai Raghuveer Chava, Jingbo Louise Liu, and Sajid Bashir

11

Electrocatalysis for the Water Splitting: Recent Strategies for Improving the Performance of Electrocatalyst . . . . . . . . . . . . . . 315 Tanveer ul Haq and Yousef Haik

12

Future of Electrochemical Energy Storage and Its Impact on the Transition Metals . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 341 Nazmul Huda and Shahjadi Hisan Farjana

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based) Oxides . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 359 Yue Wang, Fan Li, Sajid Bashir, and Jingbo Louise Liu

14

Water-Stable Metal-Organic Frameworks for Water Adsorption . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 387 Xuan Wang and Charles Lee

15

Supercapacitors: History, Theory, Emerging Technologies, and Applications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 417 Yiyang Liu, Paul R. Shearing, Guanjie He, and Dan J. L. Brett

16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy Storage . . . . . . . . . . . . . . . . . . . . . . . . 451 Jianmin Luo

17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 479 Rajesh Pathak, Yue Zhou, and Qiquan Qiao

18

3D X-Ray Characterization of Energy Storage and Conversion Devices . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 513 Chun Tan, Andrew S. Leach, Thomas M. M. Heenan, Rhodri Jervis, Dan J. L. Brett, and Paul R. Shearing

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In Situ Transmission Electron Microscopy for Studying Lithium-Ion Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 545 Chen Gong, Shengda Pu, and Alex W. Robertson

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Clean Coal Conversion Processes–The Present and Future Challenges . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 571 E. Gerald Meyer, Sai Raghuveer Chava, Jingbo Louise Liu, and Sajid Bashir

Contents

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21

Coal Gasification with Exergy Recuperation and CO2 Recovery . . . 593 Zhongkai Zhao, Yohanes Andre Situmorang, Atsushi Tsutsumi, Xiaogang Hao, Abuliti Abudula, and Guoqing Guan

22

Lignin and Lignocellulosic Materials: A Glance on the Current Opportunities for Energy and Sustainability . . . . . . . . . . . . . . . . . . 621 Marwa El-Azazy, Sajid Bashir, Jingbo Louise Liu, and Mohamed F. Shibl

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Municipal Solid Waste Incineration and Sustainable Development . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 653 Beatriz Sales Bandarra and Margarida J. Quina

24

Microbial Fuel Cells: Design and Evaluation of Catalysts and Device . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 681 Sajid Bashir, Shawn P. Mulvaney, William Houf, Luis Villanueva, Zhaohui Wang, Gregory Buck, and Jingbo Louise Liu

25

Observation on Comprehensive Energy Trend . . . . . . . . . . . . . . . . 765 Rui Guan, Yunchuan Wang, Sai Raghuveer Chava, Jingbo Louise Liu, Sajid Bashir, and Yong-jun Gao

26

Smart Energy Trend Observation . . . . . . . . . . . . . . . . . . . . . . . . . . 797 Ran Wei, Yong-jun Gao, Zhihua Wu, Sai Raghuveer Chava, Jingbo Louise Liu, and Sajid Bashir

27

Postface: Conclusion on Renewable Energy Strategies for a Sustainable Future: Part A: Role of Energy Storage . . . . . . . . 839 Weixin Song, Yong-jun Gao, Sajid Bashir, and Jingbo Louise Liu

Index . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 847

About the Editors

Yong-jun Gao, a Postgraduate trained and is at present the Deputy Director of Clean Energy and Energy Conservation and Environmental Protection Center, Zhejiang Yangtze River Delta. Since 2001, he has been doing research and work in the fields of energy and environmental protection. He has won the second academic prize of China Nuclear Energy Association, the third academic prize of China Electricity Council, and the second prize of the science and technology management innovation of State Power Investment Corporation Limited. He has published many articles as the first author or co-author in publications such as China Nuclear Power, Nuclear Power Engineering, China Power Enterprise Management, East China Electric Power, and Nanostructured Materials for nextgeneration Energy Storage and Conversion. He was awarded the model worker and the May Day medal by a Chinese city.

Weixin Song is an MPLS Enterprise Fellow and Postdoctoral Research Fellow at the University of Oxford. He received his Ph.D. in Material Electrochemistry from Imperial College London in 2019 funded by the President’s Ph.D. Scholarship. He has been a member of the Centre for Doctoral Training (CDT) of Advanced Characterization of Materials and the London Centre for Nanotechnology (LCN). He completed his BSc in 2012 and MSc in 2015 from Central South University, China. He has research interests in materials electrochemistry and characterization with materials xiii

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About the Editors

application in energy storage, electrocatalysis, and photo electrocatalysis. He is particularly interested in the materials and interface studies using electron microscopy and spectroscopy. He has 22 first-author publications and 18 licensed patents. He has contributed chapters on one book published by Springer. He received the highest student honor of Central South University (2014), STFC Future Early Career Award (2016), Armourers & Brasiers Rolls-Royce Scheme Grant (2018), National award for outstanding selffinanced students abroad (2018), MPLS Enterprise Fellowship (2019) and Constance Fligg Tipper Centenary Memorial prize (2020).

Jingbo Louise Liu received her Ph.D. in Materials Science and Engineering from the University of Science and Technology Beijing (China) in 2001. She completed her postgraduate training at the University of Calgary in Alberta (Canada) in electrochemistry analysis of nanomaterials. She is at present a Full Professor at Texas A&M University-Kingsville (TAMUK, USA) and focused on materials preparation, characterization, and applications. She is a Fellow of the Royal Society of Chemistry, a Vebleo Fellow, and a Fellow of the International Association of Advanced Materials Society. She also holds Chartered Chemist status and separately Chartered Scientists status from the Science Council, UK. She also holds DEBI Faculty Fellow status at the US Air Force Research Laboratory and is a past JSPS IF (Japan) and FFSI in Israel awardee recipient respectively. She has authored and co-authored, books, book chapters, and peer-reviewed journal articles (> 100). During her 12.5-year services at TAMUK, in chemistry, she has taught > 8700 students and trained about 150 students and scholars to conduct leading-edge research. She directed and /or participated in the projects (> 40) supported by the NSF (US and China), NSERC (CANADA), ACS Petroleum Research Funds, and Department of Education as PI, Co-PI, and senior personnel. She was recently elected as the Division Program Chair, Chair-Elect, and Chair of Energy and Fuels of the American Chemical Society.

About the Editors

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Sajid Bashir received his Ph.D. training in matrixassisted laser desorption/ionization time-of-flight mass spectrometry from the University of Warwick (UK) in 2001 and previously graduate training in Fourier transform ion cyclotron resonance mass spectrometry from the University of New York at Buffalo (USA). He was a postgraduate research associate at Cornell University (USA) in the field of plant proteomics. Currently, he is a Full Professor at Texas A&M University-Kingsville (TAMUK) and a past Faculty Fellow at the US Air Force. He has directed and participated in more than 20 projects supported by the Welch Foundation, TAMUK, Texas Workforce Commission, and US National Institute of Health. He has co-authored > 80 book chapters and peer-reviewed journal articles. He is a fellow of the Royal Society of Chemistry, and also holds Chartered Chemist and separately Chartered Scientists from the Science Council (UK). He is also the American Chemical Society, Energy and Fuels Division Technical Secretary (2018–2022). During his service at TAMUK, he trained more than 3000 students on both undergraduate and graduate levels. He created online courses and established safety training protocols in conjunction with Risk Management. Currently, he collaborated with local law enforcement as a consultant in forensic chemistry.

Abbreviations

δ α λ Φ θ Φfb ΔT φ (A)ZIB abs AC ACN AFC AOR ASU ATR BFB BP CAES CB CBM CCS CCS CCT CE CFB CFS CH CHP CLG CNF

Delta Elevation Angle Wavelength Polarization Potential Theta Flat Band Potential Charging/ Discharging Cycle Time Latitude Of The Collector (Aqueous) Zinc Ion Batteries Absorption Activated Carbon Acetonitrile Alkaline Fuel Cell Alcohol Oxidation Reaction Air Separation Unit Auto-thermal Reforming of Methane Bubbling Fluidized Bed Black Phosphorus Compressed Air Energy Storage Conduction Band Conduction Band Minimum Carbon Capture and Storage CO2 Capture and Storage Clean Coal Technology Coulombic Efficiency Circulating Fluidized Bed Capacitive Faradic Storage Helmholtz layer capacitance Combined Heat, and Power Chemical Looping Gasifier Carbon Nanofibers xvii

xviii

CNT CNT CO CO2 COF COS CP CQD CRM CS2 CSC CSC CSP CT CTMA CV CVD DBCFB DEPG des DFT dI DL DMC DMF DMFC DOS dx E Ec EC EDC EDLC EF EMC EPRI ESPW ESR ETC EU EUV Ev FEC FOM

Abbreviations

Carbon Nanotubes Carbon Nanotubes Carbon Monoxide Carbon Dioxide Covalent Organic Frameworks Carbonyl Sulfide Conductive Polymers Carbon Quantum Dots Critical Raw Materials Carbon Disulfide Charge layer capacitance China Scholarship Council Concentrated Solar Power Computed Tomography Cetyltrimethylammonium Chloride Cyclic Voltammetry Chemical Vapor Deposition Dual-Bed Circulating Fluidized Bed Dimethyl Ether of Polyethylene Glycol Desorption Density Functional Theory Decrement of the Light Intensity Double Layer Dimethyl Carbonate N, N-Dimethylacetamide Direct Methanol Fuel Cell Density Of States Film Thickness Occupied by the electrons CB Energy Ethylene Carbonate Ethylene Dicarbonate Electric Double-Layer Capacitors Fermi Level Enterprise Management Console Electric Power Research Institute Electrochemical Stable Potential Window Equivalent Series Resistance Evacuated tube collectors European Union Extreme Ultraviolet VB Energy Fluoroethylene Carbonate Figure Of Merit

Abbreviations

FPC FST GBL G-C-S GHG GO GPL GTE H 2O H 2S HCN HDFC HER HF HHV HOMO HOR HRSG HTFT HTMM HyPr-RING hν I IBA ICT IEA IFA IGCC IGFC IHP IoT IRENA ISE k KIB KPI LAB LCI LDH LIB LMA LMB LSV LUMO

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Flat Plate Collectors Foundation for Science and Technology γ-ButyroLactone Gouy-Chapman-Stern Greenhouse Gases Graphene Oxides Gel Polymer Electrolyte Gross Thermal Efficiencies Water Hydrogen Sulfide Hydrogen Cyanide Housing Development Finance Corporation Hydrogen Evolution Reaction Heating Fluid Higher Heating Value Highest Occupied Molecular Orbital Hydrogen Oxidation Reaction Heat Recovery Steam Generator High-Temperature Fischer-Tropsch Synthesis High-Temperature Mixing Method Hydrogen Production by Reaction-Integrated Novel Gasification Photon Energy Light Intensity Bottom Ash Information and Communications Technology International Energy Agency Fly Ash Integrated Coal Gasification Combined Cycle Integrated Coal Gasification Fuel Cell Inner Helmholtz Plane Internet of Things International Renewable Energy Agency Inorganic Solid Electrolytes Boltzmann Constant Potassium Ion Batteries Key Performance Indicators Lead-Acid Batteries Life Cycle Impact Layered Double Hydroxides Lithium-Ion Batteries Lithium Metal Anode Lithium Metal Batteries Linear Sweep Voltammetry Lowest Unoccupied Molecular Orbital

xx

MCFC MEA MeOH METI MEXT MH MIEC MO MOF MSW Nc NEA NFCS NMP Nv OER OHP OOH OOR ORR PAFC PAH PANI PC PC PCB PCBM PCDD PCDF PCE PCM PCP PDDA PEC PEMFC PHES PM PPy PSC PTFE PVA QFL RDE

Abbreviations

Molten Carbonate Fuel Cell Membrane-Electrode Assemblies Methanol Ministry of Economy, Trade, and Industry Ministry of Education, Culture, Sport, Science, and Technology of Japan Metal Hydride Mixed Ionic/electronic Conductors Methyl Orange Metal-Organic Frameworks Municipal Solid Waste Effective Densities of CB National Energy Administration Non-Faradaic Capacitive Storage N-Methyl-2-Pyrrolidone Effective Densities of VB Oxygen Evolution Reaction Outer Helmholtz Plane Adsorbed (per) Hydroxide species Oxidation of Organic Reactant Oxygen Reduction Reaction Phosphoric Acid Fuel Cell Polycyclic Aromatic Hydrocarbons Porous Nano flower Polyaniline Electrode Propylene Carbonate Propylene Carbonate Polychlorinated Biphenyls Phenyl-C61-Butyric Acid Methyl Ester Polychlorinated Dibenzo-p-Dioxins Polychlorinated Dibenzofurans Power Conversion Efficiency Phase Change Material Porous Coordination Polymers Poly Diallyl Dimethylammonium Chloride Photoelectrochemical Proton Exchange Membrane Fuel Cell Pumped Hydro Energy Storage Particulate Matter Polypyrrole Perovskite Solar Cells Ethylene Polytetrafluoride Polyvinyl Alcohol Quasi-Fermi Level Rotating Disk Electrode

Abbreviations

RDS ROS SARS SC SDC SEI SGSP SHE SIB SNCR SOFC SOHIO SPE SRM SSE STEM STP TBCFB TES TFE TGA TiO2 TMB TMO TOC TXM UPD VB VBM VOC

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Rate-Determining Step Reactive Oxygen Species Severe Acute Respiratory Syndrome Supercapacitors Static Dielectric Constant Solid Electrolyte Interphase Salt Gradient Solar Pond Standard Hydrogen Electrode Sodium-Ion Batteries Selective Noncatalytic Reduction Solid Oxide Fuel Cell Standard Oil of Ohio Solid Polymer Electrolyte Steam Reforming of Methane Solid State Electrolyte Scanning Transmission Electron Microscopy Standard Temperature and Pressure Triple-Bed Combined Circulating Fluidized Bed Total Energy Supply TetraFluoroethylene Copolymer Thermo Gravimetric Analysis Titanium Dioxide Transition-Metal Based Transition Metal Oxides Total Organic Carbon Transmission X-ray Microscopy Under-Potential Deposition Valence Band Valence Band Minimum Volatile Organic Compounds

Chapter 1

Promising Clean Energy Development: Practice, Challenges, and Policy Implications Sajid Bashir, Sai Chava, Weixin Song, Yong-jun Gao, and Jingbo Louise Liu

Dedicated to Dr. Peter J Derrick and Mr. Mohammed Bashir, Rest in Peace.

1.1

Fossil Fuels: A Prolog

To achieve this conversation of mass principle where atoms are converted from one form to another, waste heat and energy should be recaptured and retasked, and converted to value-added materials. As is known in freshmen chemistry, one mole of carbon generates one mole of carbon dioxide, using methane as an example. Coal

S. Bashir (*) Department of Chemistry, Texas A&M University-Kingsville, Kingsville, TX, USA e-mail: [email protected] S. Chava Department of Chemistry, Texas A&M University-Kingsville, Kingsville, TX, USA Emergent Biosolutions, San Diego, CA, USA e-mail: [email protected] W. Song Department of Materials, University of Oxford, Oxford, UK e-mail: [email protected] Y.-j. Gao Center for Clean Energy and Energy Conservation and Environmental Protection, Zhejiang Yangtze River Delta, Hangzhou, People’s Republic of China Sunshine Times Law Firm, Hangzhou, People’s Republic of China e-mail: [email protected] J. L. Liu Department of Chemistry, Texas A&M University-Kingsville, Kingsville, TX, USA Texas A&M Energy Institute, College Station, TX, USA e-mail: [email protected]; [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_1

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being a multimeric component generates multiple moles of carbon dioxide, as well as nitric oxides and oxides of sulfur: 4 C85 H5 O7 N2 S2 þ 351 O2 ! 340 CO2 þ 8 NO2 þ 8 SO3 þ 10 H2 O Combustion of bituminous coal will generate approximately 85 moles of CO2 per mole of coal or 24 megajoules per kilogram or 1000 g of CO2eq for each kWh generated, while natural gas is approximately at 500 g CO2eq per kWh. Therefore, technologies that improved combined heat and power such as recapture heat from waste incineration will result in lower consumption of coal and natural gas. In general, the electricity demand has increased annually globally by 3% although, with the more advanced infrastructure of the United States, the increase is less, there is an increased demand in energy generation from a mix of resources, principally coal and Natural gas, followed by hydro and nuclear [1]. Since 2010, there has been a renewed emphasis on a transition to renewable energy (Fig. 1.1), where the first transition was going from wood/peat moss locally to coal enabling the early societies to become more mobile. In the modern preindustrial era, coal becomes the engine for change, which enables the Romans, Greeks, and Chinese to use coal for heat and forging of weapons. The British industrial revolution was powered by coal to provide heat for steam engines, buildings, and generate electricity and manufacture steel. The second transition is in the early 1970s as the economy transitioned to more natural Gas due to the Oil

Fig. 1.1 Summary of energy generation by nuclear, petroleum distillates, and renewable resources (left-axis) and fossil fuels (Coal, and Natural Gas, right axis; [2–6])

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

3

Crisis and to lower carbon dioxide emissions. We are in the third type of transition to a more sustainable form of generating energy [7]. A levelized cost of electricity from different energy resources by the Institute of Energy Research suggests that retrofit of current coal and gas power plants to extend their useless operating lifespans by 15 years could yield near-zero carbon dioxide emissions as the most reliable and cheapest form of energy production. Under such a scenario, soluble oxide of carbon, nitrogen, and sulfur are captured and carbon dioxide is stored in geological sites suited for long-term storage [8]. The current general-purpose analyses reported in the literature fail to connect the levelized cost of energy to the reduction of particulate matter, solid oxides of carbon, nitrogen, and sulfur, and second the cost to implement these technologies and their lifecycle energy and emission [9]. This we believe is necessary before the first step before an environmental roadmap can be fully developed and deployed [10] (Fig. 1.2). In the area of energy capture and recycling, that waste to biofuels will take an increasing role in the next 50 years as a necessary aid to transition away from the utilization of coal for the generation of electricity. There are two broad reasons: One, these resources currently exist, do not need to be mined or extracted but are generated as a result of industrial, agricultural, and process chemistries, and are essential ‘lost’ to the economy. If these resources can be integrated, waste is minimized, and the extracted energy can be applied to the grid [13]. Second, the ‘lost energy’ is currently ‘replaced’ by coal and natural gas and this will enable lesser dependence on coal. Examples of atom-conversion are the generation of biofuels

Fig. 1.2 Summary of emissions of common volatiles and particulate matter (left axis) and the price of natural gas and coal on a megawatt-hour basis (right axis; [11, 12])

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from nonfood feedstocks, such as nonedible seed oil, wood chips, grape marc, chicken/cow manure, fruit peel, rotten plant matter, or recycling of municipal solid waste by a thermal process, followed by aerobic respiration, fractionation using a biorefinery process to generate bioethanol and char. These could be supported using clay-based heterogeneous catalysis and a porous metal–organic framework to capture and store biogas [14–15]. Where the carbon source is lignocellulose biomass, the typical pretreatment is acid wash and digestion; alkaline hydrogen peroxide application may also achieve similar goals with lower energy balance and lesser useless by-products. Under such a scenario, lignocellulose biomass is pretreated with hydrogen peroxide [16]. As hydrogen peroxide has a greater oxidative potential than sulfuric acid, longer incubation periods may yield phenomenologically similar results in terms of delignation followed by hydrolysis described by Eq. 1.1: 

E : H2 O2 þ 2Hþ þ 2e ! 2H2 O, E ¼ 1:78 V vs:S2 O8 2 þ 2e 

! 2SO4 2 , E : 2:01 V

ð1:1Þ

The waste can also be heated after pretreatment or subjected to co-digestion through controlled pyrolysis to generate biochar. As biochar exhibits a high carbon content, its caloric values are like coal and could be used for heat and power generation or converted to use materials such as activated carbon or porous storage materials. In the area of biofuels, lignocellulosic organics from olive trees or organic manure could be treated with solid weak acids using heterogeneous catalysts and powered by concentrated solar power to generate biofuels, or if anaerobic respiration is used to generate lactic acid as a precursor for the synthesis of poly-lactic acid polymers from waste by-products, adding value back to the supply chain [17]. Like solid waste, wastewater treatment is a high-energy process particularly when secondary sludge is formed, the degree of microbial activity is still high. Wastewater microbial-rich sludge could be subjected to anaerobic respiration and converted to inorganic salts consisting of soluble oxides of nitrogen and phosphorus, which could be used as renewable fertilizers, the organic digestates as biogas.

1.1.1

Life-Cycle Assessment: Common Approaches

The life-cycle assessment should account for environmental, economic, and technical parameters in the application of any energy production platform, summarized in Fig. 1.3. Examples of environmental parameters include generation of global warming potential, acid rain and acidification potential, eutrophication potential, and cumulative energy demand. In the area of economy, the unit cost per megawatt-hour, operational life-cycle, infrastructure costs, and power plant operational costs

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

5

Fig. 1.3 Block diagram summary of factors that underpin life-cycle assessment cradle-to-grave type of analysis for energy generation (fossil fuels) and energy storage (sustainable energy resources, [18])

including the cost of decommissioning and the technical parameters are the cost of feedstocks, generating capacity, generating reliability, technological maturity of the proposed process, and transmission and storage potential. An example of the above approach is when an analysis of electrical storage is conducted on the ‘short-term’ (< 1 min), ‘medium-term’ (~ 270 min), or ‘long-term’ (> 130,000 min). By using the general equation to determine the levelized cost of electricity as described by Eqs. 1.2 and 1.3: ðCapC þ OpC  AFÞ AF  W

ð1:2Þ

1  ð1 þ itsÞ1 , its

ð1:3Þ

where AF ¼

where capital expenditures (CapC, $), operational expenditures (OpC, $) including the balance of plant cost, W is the annualized energy output from the storage system (kWh). The energy output is dependent upon the number of cycles (n)  discharge power (kW) and discharge per cycle time (ms), AF ¼ annuity factor (for l at fixed 5%), its ¼ is the cost of in-state discounts or tariffs, or subsidies ($) and l is the lifetime of the storage system (years) in increments of 1kWh and 8-216 GWh used from storage systems. The cycles for storage depend on a short–medium or long time scale, with 40 (thermal or isothermal energy storage) or 100 (adiabatic energy storage) cycles

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Fig. 1.4A Summary of energy storage by the form of storage (Mechanical, electrical, chemical, or thermal in nature, [20]) Performance Costs at 100 MW 3.5

LCOE (LOG Skwh)

3

Pump Hydro Medium-term Storage

2.5

Compressed Air Energy Medium-term Storage Pump Hydro long-term Storage

2

Battery Medium-term Storage

1.5

H2/CH4 to Power Long-term Storage H2/CH4 to Power Medium-term Storage

1

Battery Short-term Storage

0.5

Compressed Air Energy Short-term Storage Pump Hydro Short-term

0 0 –0.5

50

100

150

200

250

300

350

400

450

500

H2/CH4 to Power Short-term Storage

CO2 Emissions (g CO2e kWh)

Fig. 1.4B Summary of LCOE for different systems for fast, medium, and long-term energy retrieval on a log basis with area based on modal value. (Data from Table 3.4 of [11, 19])

per day for 1 year and cost factors from [19] (Data from Table 3.4 of [19 and Appendices of 11]) and results summarized in Fig. 1.4A by energy type, Fig. 1.4B by LCOE and Table 1.1 [51].

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7

Table 1.1 Summary of different storage systems in terms of power, specific energy, efficiency, discharge capacity, the response time (fast is seconds to minutes), lifetimes (years or discharge cycles), stability in self-discharge tendency, storage costs, power conversion system (PCS) cost, and balance of plant (BOP) associated costs [21]

Storage technology Pumped hydro storage Conventional compressed air energy storage Lead acetate (nickel cadmium, if different) Lithium ion Hydrogen Thermal Storage technology

Power (MW) 100–5000 5–3000

0–40

Specific energy (Wh/kg) 0.5–1.5 30–60

Efficiency (%) 65–80 41–75

Discharge at capacity (h) 1–24 1–24

Response time Fast Fast

70–80 (60– 80) 65–88 35–42 14–18 Storage cost ($/kWh) 5–136 2–141

1–5 (6–8) 0.017–2 12 1–24 PCS ($/kW)

¼ cycle

403–4644 432–1674

3–30 3–30

132–915 (609– 1210) 282–4104 2–15 0.059– 0.101

211–648 (281–355)

46–140 (76–130)

161–4320 540–4809 1–50 (kJ/kg/ K)

0–130 11–43 59–101

Pumped hydro storage Conventional compressed air energy storage Lead acetate (nickel cadmium, if different)

15 k 10 k

30–50 (45– 80) 75–200 400–1000 80–200 Self-discharge/day (%) Low Low

0.25–1.5 (1.5–3.0)

0.1–5 (0.2–0.3)

Lithium-ion Hydrogen Thermal

0.6–1.2 0.1–1.0 5– 15 years

1–5 Low 0.05–1

0.015–50 0–50 50–250 Lifetime (cycles, k)

¼ cycle ¼ cycle Fast BOP cost ($/kW)

Sources: For Pumped Hydro and Compressed Air Storage: [19, 20, 22–24]. For lead acetate storage: [25–29]. For lithium-ion storage: [24, 26, 27, 28, 30, 31, 32]. For hydrogen Storage: [20, 22, 24, 26, 33–35]. For thermal Storage: [24, 36–38]. For levelized costs of electricity and lifetime cost analysis: [22, 29, 39–50, 122]

The analysis using the LCOE equation demonstrates that Li-ion batteries and advanced adiabatic and isothermal compressed air energy storage as suitable energy storage platforms with low levelized costs and high reliability, whereas under a medium time scale, pumped hydro and advanced adiabatic compressed air or energy storage or thermal storage using molten salts or mineral oil and long-term power to gas to power or pumped hydro, where hydrogen or methane is used to generate electricity. The electricity would be generated using a combined cycle power plant or proton exchange membrane fuel cells [52, 53].

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Role of Energy Storage

Energy storage is required to enable constant power load and is driven by pumped hydro as the historically largest share form of energy storage, with lithium-ion batteries being a recent addition due to the rise of both solar and wind power. No lifetime assessment cycle reviewed encompasses all the required parameters, however, parameters that drive cost and applicability are a lifetime of storage technology, the efficiency of energy retrieval, and cycle length and depend on end-usage, with lithium-ion batteries gaining a significant share in all stationary applications due to greater research and development as a result of increased uptake of solar and wind as sustainable energy resources. For long-term storage and high-need applications, pumped hydro, compressed air, hydrogen and thermal energy storage are likely to increase due to stricter emissions standards for automobiles, fuel cells driven transport is now feasible, utilizing zeolite-based heterogeneous catalysts, for example, bismuth vanadate as an intermediate bandgap photocatalyst that could be coupled solar to photocatalysis through recombination of photoinduced electrons and holes to generate hydrogen from water or coal [54, 55].

1.2

The Electron Economy

While the last two decades have seen a dramatic rise in solar and wind energy generation resources, the future most likely will look toward a hybrid system where energy storage will play an increasingly prominent role coupled with clean coal and carbon dioxide capture. Figure 1.5 shows a conceptual workflow of how this could be accomplished. The energy generation resource such as a photovoltaic array, wind-powered turbine, solar concentrator to steam to turbine, or hydrogen feed gas will need to generate electricity, store it and then allow the system to utilize the stored energy. This will be done using inverters, and heat pumps, and a direct current (DC) bus,

Fig. 1.5 Workflow of hybrid energy generation and storage [56]

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9

using a DC-to-DC converter and DC bus connection to alternating current (AC) bus using inverters and the electricity or lead load connected to the AC bus. In this manner, energy is drawn with surplus energy used to recharge the secondary or flow batteries or electrolyzes to generate hydrogen from water, if the load is more than the immediate power, then the energy from the stored batteries can be tapped or stored hydrogen to generate electricity with the excess heat used to assist with steam generation. Currently, approximately 2=3 of the total annualized cost is due to components from the wind-powered turbine and the rest for batteries, inverters, and hydrogen storage systems, and fuel cells. The fuel cell catalyst is the next greatest expense and is likely to decrease as more research and development is geared toward the automobile and residential energy sectors, where the likelihood of supercapacitor, battery, flex-fuel, and fuel cells are used or some combination for automobile and stationary applications [57–59]. The rapid development in alkali-metal-based batteries is driven by chemical and physical characteristics of lithium and sodium and this is related to atomic radii, electron affinity, and charge density, from which energy density and kinetics of energy release are derived: e.g. 6.9 atomic mass unit, amu per mole, 0.76 Å, 180.5  C, 3.0 V, 3860 mAg.g1 and $5 k for lithium, Li vs. 23 amu, 1.06 Å, 97.7  C, 2.7 V 1160 mAg.g1 and $0.15 k for sodium, Na. The molar mass is in atomic mass units, cation radii in Angstroms, melting point in degrees Celsius, voltage relative to standard hydrogen electrode in volts, the energy density in milliampere hour per gram, and cost of carbonate salt on a ton basis in United States dollars, thousand increments, respectively. These energy parameters favor lithium-ion batteries, while environmental and economic factors favor sodium ion-based batteries [60]. A bifurcated role is likely with lithium-ion batteries being favored in automobiles where high energy density and fast kinetics are required and sodium-ion batteries for stationary energy storage where high density is not required, such as buildings or a configuration shown in Fig. 1.5 [61]. The near-future bottleneck to battery and fuel cell development is the electrocatalysts and electrode architecture, which will ultimately define high gravi- and volumetric energy densities at the solid electrolyte interphase (SEI) layer that limits spontaneous reactions at the expense of thermodynamic instability and lower cell lifetime, which may be accomplished through alloying, described in Eqs. 1.4 and 1.5 [62]: E þ yAþ þ xē ! Ay M,

ð1:4Þ

where A represents alkali metals Li or Na and E represent metalloids Si, Ge, Sn, Sb, and P, respectively, or conversion [63]: Ma Ob þ ðb:nÞA ! aM þ bAn O,

ð1:5Þ

where M.O are metal oxides and A is alkali metals (Li or Na) such as Silicon-based anodes with Li iron phosphate batteries, e.g. Li17Si4.

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1.2.1

Public Policy and Discourse

The above paragraphs imply an electron-driven economy, while this is true to a greater degree today, this is not a complete characterization of the likely environment-food-water-energy nexus in the next 50 years, because, within the United States, much policy is driven at the federal level due to the muscular resources available and at the time of writing this preface, president-elect Biden has 48 senate votes most likely confirming his policies and the best-case scenario is of a clean sweep of the State of Georgia Senate races, tying the US senate at fifty senators from the Grand Old Party (Republican GOP) and the democratic party that as a slim majority within the U.S house of representatives and either a 48–52, 49–51 minority or a tie broken by the Vice-President and confirming Democratic policies on behalf of the presidency of the United States. This political reality means major policy initiates will require bipartisan support to be enacted into law. This means the Biden Green initiative on the environment that pledges great reduction in carbon dioxide emissions could be met through not an electron economy but a carbon majority and hydrogen minority-driven economy, whereas natural gas and coal-powered generation of electricity are accomplished using plant upgrades to capture carbon dioxide and conversion of biomass, coal and natural resources to hydrogen that is used to generate energy and promote fuel cell-driven transport. Here, the U.S Department of energy could utilize its public-partnership research fund to promote companies to conduct subsidized research in fuel cells-driven transport, similar to a model employed to assist E. Musk in his development of a sports electric vehicle in the early zero–zero decades. There was bipartisan support to elicit more private sector funding for technologies that have yet to reach maturity and widespread deployment but were further in the development cycle than emerging technologies, much as like the National Institute of Health (NIH) Model in support vaccine research in developing world countries related to the human immunodeficiency viruses/acquired immunodeficiency syndrome and tropical diseases that do not reach the profit margin for many US Pharma to invest in heavily. This NIH initiative has also received support from prominent foundations such as the Gates Foundation and could be a scenario whereby, the politics is removed from the policy.

1.2.2

The Hydrogen Economy

How could such an economy work in principle? Where would the hydrogen be generated? Currently, hydrogen production, approximately 48% from natural gas, 30% from Naphtha and Oils, and 18% from coal [28] or biomass. When fossil fuels are the feedstock, hydrogen production can be accomplished by hydrocarbon reforming, pyrolysis. Reforming from hydrocarbons can be accomplished by steam reforming, partial oxidation, and autothermal reforming. When generated from biomass, this may be by bio photolysis, dark or photo fermentation, if thermochemical, by pyrolysis, gasification, combustion or liquefaction, and from water splitting as indicated in Fig. 1.5. Here, electrolysis, thermolysis, and photolysis

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Table 1.2 Summary of the energy density of common fuels Fuel Hydrogen Methane Ethane Butane Methanol Ethanol Gasoline (Cetane # 8– 14) Diesel (Cetane # 51) Coal Wood (C6H10O5)n Sugar (glucose)

Phase at STP Gas Gas Gas Gas Liquid Liquid Liquid

Specific energy (MJ/ Kg) 141.9 55.5 51.9 49.5 20 29.7 45.4

Energy Density (MJ/ kg) 119.9 50.0 47.8 29.7 17.1 23.3 42.6

Liquid Solid Solid Solid

46.0 15–27 15.0 15.6

43.0 12.0–36.5 9.0 23.9

Sources and Notes’: Standard Temperature and Pressure. [64–68]

could be utilized. The advantages of hydrogen are in its high energy density when compared with carbon-based fuels, as shown in Table 1.2 and workflow in Fig. 1.6. Hydrogen production by steam reformation is nickel catalytic conversion of hydrocarbons and steam to form hydrogen and carbon oxide through a series of sequential steps, which result in the synthesis of gas (syngas), water-gas shift, and methanation or gas purification. The feed is light hydrocarbons or naphtha; however, the materials need to be carbon and hydrogen to avoid poisoning of the catalyst. If the coal or feed contains sulfur, a desulfurization step is performed [69] at 3.5 MPa and 900  C temperature using a 3:5 ratio of steam to carbon [70]. The heat is recovered and the channel promotes the water gas shift reaction, where carbon monoxide reacts with steam to produce hydrogen gas and is purified through a pressure swing adsorption [71] at the expense of carbon dioxide emission, which could be captured using supercritical conditions and the liquefied carbon dioxide stored in geological reservoirs. Alternately, the carbon dioxide could be reacted with hydrogen to generate wet methane [72]. This methane is re-fed back to the reformer to generate additional hydrogen. The different gases could be separated using Pd-based Inorganic membranes with high selectivity and at lower operational temperatures (550  C) lowering the cost of hydrogen production to $2.0 kg with Carbon dioxide capture and storage [73] described by the steam reformation, watergas shift reaction, or methanation described in Eqs. 1.6, 1.7 and 1.8: Cn Hm ðgÞ þ nH2 O ðgÞ ! nCO ðgÞ þ ðn þ ½mÞH2 ðgÞ

ð1:6Þ

WGS : CO ðgÞ þ H2 O ðgÞ ! CO2 ðgÞ þ H2 ðgÞ

ð1:7Þ

Methanator : CO ðgÞ þ 3H2 ðgÞ ! CH4 ðgÞ þ H2 O ðlÞ

ð1:8Þ

Partial oxidation is a similar method but using oxygen instead of air and occurs at around 1315  C for the non-catalytic process using coal, oils, or methane [71]. The

Fig. 1.6 Workflow process to produce hydrogen using ten (〇!〇) distinct processes. Each unique process and the pathway are indicated using a numbered sequence. For example, hydrogen production by Steam reforming is shown by the process 〇, membrane separation by the process ❷, coal gasification by 〇, biomass-based processes are shown by sequences 〇,〇 and〇. Thermo-photo- or electrolysis processes by sequences 〇,〇,〇, and 〇. CRA represents carbonrich acids, CBP Carbon By-products; Pe the electrical energy, HeTF Heat Transfer Fluids, HiTHeE High-temperature Heat Exchanger [68]

12 S. Bashir et al.

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13

feed gas or materials are oxidized directly, or sulfur is removed, and oxygen is used to partially oxidize hydrocarbons and syngas in a similar workflow to the steam reformation for the catalytic process conducted at lower operating temperatures summarized in Eqs. 1.9, 1.10, 1.11 and 1.12 [74]: Catalytic Reformer : Cn Hm ðgÞ þ ½nO2 ðgÞ ! nCO ðgÞ þ ½mH2 ðgÞ

ð1:9Þ

Non  catalytic : Cn Hm ðg=lÞ þ nH2 O ðg=lÞ ! nCO ðgÞ þ ðn þ ½mÞH2 ðgÞ

ð1:10Þ

WGS : CO ðgÞ þ H2 O ðg=lÞ ! CO2 ðgÞ þ H2 ðgÞ

ð1:11Þ

Methanator : CO ðgÞ þ 3H2 ðgÞ ! CH4 ðgÞ þ H2 O ðlÞ

ð1:12Þ

Because coal has lower hydrogen to carbon ratio than methane, higher pressure and temperature are required (6 MPa and 1315  C), generating non-reactant ash and more water [75]. The above two processes are endothermal and require additional energy input, for autothermal reformation; the initial oxidation is exothermal, and the thermal energy is used to drive the endothermal steam reforming process to drive hydrogen production. The oxidizer is steam, oxygen, or air as the inlet gas to the reformer to drive the reformation and oxidation reactions to occur [76], summarized in Eq. 1.13: Cn Hm ðlÞ þ ½nH2 O ðlÞ þ ¼nO2 ðgÞ ! nCO ðgÞ þ ð½n þ ½mÞH2 ðgÞ

ð1:13Þ

Because part of the thermal energy is utilized, the overall energy demand is lesser than the before processes, lowering the cost of production. Methane conversion to hydrogen could be separated using a Pd membrane and could be accomplished at lower temperatures than the non-catalytic reaction when the air was used, approximately 900  C, that has to be tuned to the operating parameters of the inorganic membrane, as a higher temperature may lead to membrane deformation and loss of selectivity of hydrogen [77]. Hydrocarbon pyrolysis is the breakdown of hydrocarbons to hydrogen via a thermocatalytic breakdown to carbon and hydrogen at the boiling point of the liquid hydrocarbons. When heavy oil hydrocarbons are utilized, the decomposition temperature is higher around 350  C to promote gasification of the hydrocarbon followed by cracking or the methane to hydrogen gas as a form of decarbonization at under 1000  C and 1 atm pressure. Pyrolysis is a gasification process and does not involve a water gas shift or carbon dioxide absorption step. The hydrogen gas is separated using Pd-Ag alloyed membrane at lower temperatures to avoid damage to the membrane, the rate of separation is dependent upon the partial pressure of the incoming gas, noting that the gas needs to be devoid of sulfur to avoid catalytic poisoning [78]. Hydrogen gas can also be formed using biomass, where the cellulose lignin-based materials can be converted to methanol, sugars that were generated from carbon dioxide, and sunlight known as photosynthesis. The carbon biomass is then

14

S. Bashir et al.

thermochemically broken via pyrolysis to generate a gasified-rich distillate in methane and carbon monoxide that can be processed to generate hydrogen. Alternately, the biomass can be liquified and combusted to ultimately generate hydrogen at 0.5 MPa and 530  C anaerobically to generate small neutral gases, hydrocarbons, and char [79] described by Eqs. 1.14, 1.15, 1.16 and 1.17: Cn Hm ðs=lÞ þ nH2 O ðlÞ ! nCO ðgÞ þ ðn þ ½mÞH2 ðgÞ

ð1:14Þ

CO ðgÞ þ H2 O ðlÞ ! CO2 ðgÞ þ H2 ðgÞ

ð1:15Þ

In biomass gasification, the biomass is converted into syngas using air, oxygen, or steam at 1400  C and 32 atm: Biomass ðs=lÞ þ Air ðgÞ ! H2 ðgÞ þ CO2 ðgÞ þ CO ðgÞ þ N2 ðgÞ þ CH4 ðgÞ þ Cn Hm ðlÞ þ H2 O ðlÞ þ char ðsÞ ð1:16Þ Biomass ðs=lÞ þ Steam ðg=lÞ ! H2 ðgÞ þ CO2 ðgÞ þ CO ðgÞ þ CH4 ðgÞ þ Cn Hm ðlÞ þ H2 O ðlÞ þ char ðsÞ þ tar ðsÞ ð1:17Þ Photosynthesis from plants or fermentation by cyanobacteria, algae, and fungi can also convert biomass to sugar-rich feedstocks that can be converted to carbonrich acids (CRA) and then to hydrogen gas using light energy or in the dark using two electron reduction by a hydrogenase or nitrogenase. Some elect microbes can directly convert water to hydrogen in a process known as photolysis, summarized in Eqs. 1.18, 1.19, 1.20, 1.21 and 1.22 [80]: 2H2 O ðlÞ þ hv ! 2H2 ðgÞ þ O2 ðgÞ or 12H2 O ðlÞ þ 6CO2 ðgÞ þ hv ! C6 H12 O6 ðsÞ þ 6O2 ðgÞ and

ð1:18Þ ð1:19Þ

C6 H12 O6 ðsÞ þ 12H2 O ðlÞ þ hv ! 12Hþ ðaqÞ

þ6CO2 ðgÞ or sugars or carbon rich acids ðCRAÞ ð1:20Þ C6 H12 O6 ðsÞ þ 2H2 O ðlÞ ! 2CH3 COOH ðCRAÞ ðlÞ þ 4H2 ðgÞ þ 2CO2 ðgÞ ðfermentationÞ

ð1:21Þ

C6 H12 O6 ðsÞ þ 2H2 O ðlÞ ! CH3 ðCH2 Þ2 COOH ðCRAÞðlÞ þ 2H2 ðgÞ þ 2CO2 ðgÞ ðfermentationÞ

ð1:22Þ

Cellulose/Starch are the common bio feed stocks and fermentation is carried out at mild acidic conditions in the dark. When sunlight is used, nitrogenase can convert CRA (e.g. acetic acid) into hydrogen and carbon dioxide, summarized in Eq. 1.23 [81]:

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

CH3 COOH ðlÞ þ 2H2 O ðlÞ þ hv ! 4H2 ðgÞ þ 2CO2 ðgÞ

15

ð1:23Þ

These reactions could be accomplished using a bio-reactor, where sugars are digested by bacteria anaerobically in the dark to generate carbon-rich acids, that could be used by photosynthetic bacteria using light and water described by Eqs. 1.24 and 1.25 [82]: C6 H12 O6 ðsÞ þ 2H2 O ðlÞ ! 2CH3 COOH ðlÞ þ 2CO2 ðgÞ þ 4H2 ðgÞ

ð1:24Þ

2CH3 COOH ðsÞ þ 4H2 O ðlÞ þ hv ! 8H2 ðgÞ þ 4CO2 ðgÞ

ð1:25Þ

Using glucose as the biomass, approximately 7 mol of hydrogen could be generated per mol of glucose [83] under mild acidic conditions. The last method to generate hydrogen is direct water splitting using sunlight [84] and redox of water. As this is an endothermic reaction, energy from sunlight is used using a fuel cell/ electrolyzer design consisting of a cathode and anode and electrolyte with an external current. Here, water is electrolytically broken to hydrogen and oxygen, shown in Eq. 1.26 [85]: 2H2 O ðlÞ ! 2H2 ðgÞ þ O2 ðgÞ

ð1:26Þ

The electrolysis can be accomplished using fuel cells such as proton exchange membrane fuel cells, solid oxide electrolysis cells, where they generate proton migrates across a membrane to the cathode to generate hydrogen, leaving water and oxygen [86, 87]. In the solid oxide electrolyzer operated under alkaline conditions, water is broken down to H+ and OH, with the hydroxide traveling through the aqueous electrolyte to the anode to form oxygen [88], with the driving energy being thermal instead of electrical in a fuel cell, with hydrogen in the steam feed and using fuel cells described by Eqs. 1.27, 1.28, 1.29 and 1.30 [89]: Alkaline Solid Oxygen Electrolyzer (Anode). OH ðaqÞ ! O2 ðgÞ þ 2H2 O ðlÞ þ 4ē

ð1:27Þ

2H2 O ðlÞ þ 2ē ! 2OH  ðaqÞ þ H2 ðgÞðcathodeÞ

ð1:28Þ

Proton exchange membrane 2H2 O ðlÞ ! O2 ðgÞ þ 4Hþ ðaqÞ þ 4ē ðanodeÞ þ

4H ðaqÞ þ 4ē ! 2H2 ðgÞðcathodeÞ

ð1:29Þ ð1:30Þ

The efficiency and relative merits of hydrogen production summarized in Eqs. (1.4, 1.5, 1.6, 1.7, 1.8, 1.9, 1.10, 1.11, 1.12, 1.13, 1.14, 1.15, 1.16, 1.17, 1.18, 1.19, 1.20, 1.21, 1.22, 1.23, 1.24, 1.25, 1.26, 1.27, 1.28, 1.29 and 1.30) are summarized in Table 1.3A. The cost of solar thermal is due to the electrolyzer unit and for solar, it is peak and off-peak costs plus the cost of the electrolyzer unit. For

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S. Bashir et al.

Table 1.3A Hydrogen production by energy source, rate, capacity factor, and approximate cost in the year of study Electricity platform Nuclear Solar thermal Solar photovoltaic Wind BioSystem Photolysis Fermentation (Dark) Fermentation (Light)

H2 production rate (kg/day) 1000 1000 1350 35,000 Synthesis rate (mmol H2/h)

Capacity factor (%) 95 40 30 60

0.05–0.5 9.5

Hydrogen cost ($/kg) 4.15 7.00 10.50 5.25 Bioreactor volume (m3) 340–1710 8.50

0.15

750

Sources: [68, 86, 90–96]

wind, the unit can be used in the cogeneration of hydrogen and electricity or hydrogen varying the cost. The values chosen are to the nearest-semi-decade or multiples of five and are based on a standard fuel stack of 5 kW [97]. The synthesis rate for biological processes is highly dependent on the source of carbon, whether the photolysis is by direct or indirect means and the availability of oxygen. For this reason, the modal value by semi-decade is used. As the above is an energy-intensive process, currently the energy is offset from renewable resources such as wind, solar, or concentrated solar thermal. At temperature above 2500  C, water splitting is energetically favorable and could be accomplished catalytically at lower temperatures using binuclear catalysts based on copper, summarized in Eqs. 1.31, 1.32, 1.33 and 1.34 [98, 99]): 

2H2 O ðlÞ ! 2H2 ðgÞ þ O2 ðgÞ, Temp > 2500 C

ð1:31Þ

~ C CuCl2 ðsÞ þ H2 O ðgÞ ! CuO*CuCl2 ðsÞ þ 2HCl ðgÞ, T400 ~ C CuO*CuCl2 ! 2CuCl ðlÞ þ ½O2 ðgÞ, Temp500

ð1:33Þ

~ C SnO ðsÞ þ H2 O ðgÞ ! SnO2 ðsÞ þ H2 ðgÞ, Temp550

ð1:34Þ



1.2.3

ð1:32Þ

The Photon Economy

Concentrating Solar Power is one attractive vehicle by which the thermolysis of water can be accomplished. Lastly, photolysis where sunlight is absorbed by a semiconductor catalyst where photon energy above the semiconductor bandgap at the anode will result in the generation of an electron-hole pair which migrates toward

1 Promising Clean Energy Development: Practice, Challenges, and Policy. . .

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Table 1.3B Hydrogen production by thermal source, rate, catalyst, and approximate cost in the year of study Thermal source Nuclear Nuclear Solar Solar

H2 production rate (kg/day) 7000 800,000 6000 6000

Solar

6000

Catalyst CuCl SI ZnO/Zn Fe3O4/ FeO Fe2O3/ Fe3O4

Operating temperature ( C) 550 850 1730 1630

Hydrogen cost ($/kg) 2.15 2.55 8.00 8.40

1330

8.40

Sources: [68, 86, 90–96, 104–110]

the appropriate biased electrodes and the electrolyte, with the holes migrating at the anode, where water is split into protons that migrate through the electrolyte to the cathode and oxygen that remains with water, while electron flow is through an external circuit to the cathode, where protons combine to form hydrogen gas, summarized in Eqs. 1.35, 1.36 and 1.37 [100–103]: þ

2 p þH2 O ðlÞ ! ½O2 ðgÞ þ 2Hþ ðaqÞ ðanodeÞ

ð1:35Þ

2ē þ 2H þ ðaqÞ ! H2 ðgÞðcathodeÞ

ð1:36Þ

H2 O ðlÞ ! H2 ðgÞ þ ½O2 ðgÞ

ð1:37Þ

The efficiency and relative merits of hydrogen production summarized in Eqs. (1.31, 1.32, 1.33 and 1.34) is summarized in Table 1.3B: This reaction would occur at a bandgap of 1.23 eV without the application of an external bias, with the most common semiconductor photocatalysts being SiC and TiO2 [111].

1.2.4

The Circular Economy

Steam methane reforming is the most common with biomass being used more as a carbon-neutral mechanism to generate hydrogen gas by thermochemical processes where hydrogen gas is generated from pyrolysis and gasification, whereas biochemical processes utilize fermentation using bioreactors. The energy costs for these processes can be aligned using concentrated solar power to drive coal gasification and water splitting and storage that is achieved using high pressure or liquefaction using cryogenic temperatures and cooling limiting the total capacity and are summarized in Table 1.4 for relative costs. Solid-state storage using metal–organic frameworks at lower temperatures may be achieved at lower operating pressures,

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Table 1.4 Summary of relative merits of each hydrogen generation process. The efficiency is the highest practical efficiency and the cost is the modal costs where the process is highly variable such as wind, solar, or biochemical by semi-decade. Nuclear thermal is always cheaper than solar thermal due to the greater energy density of nuclear fission and greater maturity of the technology Process Steam reformation Partial oxidation Autothermal reforming Pyrolysis Coal gasification Photolysis

Efficiency (%) 85

Energy source Hydrocarbons

Feedstock Methane

Hydrogen cost ($/kg) 2.30

75 75

Hydrocarbons Hydrocarbons

Methane Methane

1.65 1.50

50 30 10

Steam Hydrocarbons Solar

Biomass Coal Water and Algae/ Cyanobacteria CRA

2.05 1.60 2.15

Water

6.50

Water Water

2.60 8.50

Dark fermentation Electrolysis

80

STP (dark)

60

Thermolysis Thermolysis

45 45

Wind\Solar\ Nuclear Nuclear Solar

2.60

Notes and Sources: CRA Carbon-rich acids; Hydrocarbon are fossil fuels including natural gas and/or coal; STP Standard Temperature and Pressure; [68, 96, 104–110] Table 1.5 Summary of current storage methods for hydrogen

Process Liquid hydrogen Compressed hydrogen Activated carbon Interstitial metal hydride Salt-like metal hydride

Storage density (mat wt %)/[volumetric density, by wt%] 100/[14]

Volume density (kg/m3)/[energy release, MJ/kg H2] 70/[120]

Temperature ( C) 253

Pressure (MPa) 0.1

100/[3]

40/[120]

25

77

10.5/[6.5]

4/[3.5]

200

10

2/[3]

150/[15]

0

1

7.5/[18]

150/[37]

330

1

Sources: [112–120]

or chemical conversion to generate metal hydrides, which can be decomposed at close to standard temperature and pressure (STP) and are summarized in Table 1.5. The energy demand is unlikely to diminish in the near and intermediate future and the environmental stress from soluble oxides of carbon, nitrogen, and sulfur are unlikely to diminish, including the long-term effects of SARS [121]. As this

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pathology in part is due to respiratory distress, then energy production without particulate matter and carbon dioxide would greatly reduce the risk of proclivities to being susceptible to the virus. The development of integrated fossil fuel and renewable energy production, recycling, waste management, and lifecycle assessment will need to be developed and deployed. Hydrogen gas can be generated by several abiotic and biotic methods with steam methane reformation the most common or coal gasification using thermal pyrolysis as possible new approaches using concentrated solar thermal or nuclear thermal and separation by a membrane-based reactor with combined cycle gas turbine exhausts recycling heat and exhaust gas driven by concentrated solar power. Alternately water-splitting methods may be employed and hydrogen stored using porous materials or through compression with new safety codes for its usage and a newer transportation pipeline that can potentially improve the infrastructure, create jobs, and transition to a hydrogen economy for power and heating as well as fuel cell-based transportation, with fossil fuel as feedstocks and nuclear or thermal as energy resources to drive the synthesis of hydrogen, which we believe will garnish bipartisan support at the US Senate which will need to sign off any major environmental or energy policy initiates regardless of the occupancy of the Whitehouse. Acknowledgments This work is supported by the Petroleum Research Fund of the American Chemical Society (53827-UR10) and the Robert Welch Foundation (Departmental Grant, AC-0006). We thank the program chair and Dr. E. Gerald Meyer of the ENFL America Chemical Society for the opportunity to run or moderate symposia at the technical sessions. Lastly, The leadership at Texas A&M University-Kingsville, Department, College, and University level, as well as Springer Science+Business Media, LLC technical staff for their assistance in copy editing this and other book chapters. Author Contributions S. Bashir completed the initial draft including equations relating to thermodynamics. J.L. Liu, data for figures and Sai. Chava the data for hydrogen production and figure 6, and W. Song and Y. Gao who reviewed the final draft which was submitted by S. Chava.

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Sajid Bashir received his Ph.D. in Analytical Chemistry from the University of Warwick, England, in 2001. He was a postgraduate research associate at Cornell University to research in the field of plant genetics. Currently, he is a Full Professor at Texas A&M University-Kingsville (TAMUK) and a Faculty Fellow at the US Air Force. He has directed and participated in more than 20 projects supported by the Welch Foundation, TAMUK, Texas Workforce Commission, and US National Institute of Health. He has coauthored more than 80 book chapters and peer-reviewed journal articles. He is a fellow of the Royal Society of Chemistry and Chartered Chemist & Chartered Scientists of Science Council. During his service in TAMUK, he trained more than 3000 students on both undergraduate and graduate levels. He created online courses and established safety training protocols in conjunction with risk management. Currently, he collaborated with local law enforcement as a consultant. Sai Raghuveer Chava received his M.S. in Chemistry from Texas A&M University-Kingsville in 2010. He was a Graduate Research Assistant at NNTRC, Kingsville who researched snake venoms. Currently, he is working as a Scientist on Vaccine research focusing on the analytical method development, formulation, and process-related problems using the analytical data. He has also worked on ophthalmic formulations, process characterization, method developments, and method validations. He has coauthored five journal articles and a book chapter and peerreviewed journal articles. He is an elected member of the Royal Society of Chemistry and Sigma Xi.

Dr. Weixin Song is an MPLS Enterprise Fellow and Postdoctoral Research Fellow at the University of Oxford. He received his Ph. D. in Material Electrochemistry from Imperial College London in 2019 funded by the President’s Ph.D. Scholarship. He has been a member of the Centre for Doctoral Training (CDT) of Advanced Characterization of Materials and the London Centre for Nanotechnology (LCN). He completed his B.Sc. in 2012 and M.Sc. in 2015 from Central South University, China. He has research interests in materials electrochemistry and characterization with materials application in energy storage, electrocatalysis, and photo electrocatalysis. He is particularly interested in the materials and interface studies using electron microscopy and spectroscopy. He

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S. Bashir et al. has 22 first-author publications and 18 licensed patents. He has contributed chapters on one book published by Springer. He received the highest student honor of Central South University (2014), STFC Future Early Career Award (2016), Armourers & Brasiers Rolls-Royce Scheme Grant (2018), National award for outstanding self-financed students abroad (2018), MPLS Enterprise Fellowship (2019), and Constance Fligg Tipper Centenary Memorial prize (2020). Yong-jun Gao Postgraduate, Deputy Director of Clean Energy and Energy Conservation and Environmental Protection Center, Zhejiang Yangtze River Delta. Since 2001, he has been doing research and work in the fields of energy and environmental protection. He has won the second academic prize of China Nuclear Energy Association, the third academic prize of China Electricity Council, and the second prize of the science and technology management innovation of State Power Investment Corporation Limited. He has published many articles as the first author or co-author in publications such as China Nuclear Power, Nuclear Power Engineering, China Power Enterprise Management, East China Electric Power, and Nanostructured Materials for Next-Generation Energy Storage and Conversion. He was awarded the model worker and the May Day medal by a Chinese city. Jingbo Louise Liu received her Ph.D. in Materials Science and Engineering from the University of Science and Technology Beijing in 2001. She is a Full Professor at Texas A&M UniversityKingsville (TAMUK) and focused on materials preparation, characterization, and applications. She is a Fellow of the Royal Society of Chemistry and a DEBI Faculty Fellow at the US Air Force Research Laboratory. She has authored and co-authored books, book chapters, and peer-reviewed journal articles (> 100). During her 12.5-year services in TAMUK, she taught more than 8700 students and trained about 150 students and scholars to conduct leading-edge research. She directed and/or participated in the projects (> 40) supported by the NSF (US and China), NSERC (CANADA), ACS Petroleum Research Funds, and Department of Education as PI, Co-PI, and senior personnel. She was recently elected as the Division Chair-Elect of Energy and Fuels, American Chemical Society.

Chapter 2

Applications and Fundamentals of Photocatalysis with Solar Energy Qiuyang Huang, Dan Kong, and Yongdan Li

2.1

Introduction

This section will give a brief introduction to photocatalysis, mainly on the historical background, applications, and reaction systems of photocatalysis.

2.2

Historical Background

The photoinduced redox reactions on semiconductors can be traced back to the 1930s when scientists reported that TiO2 could react with organic dyes [1]. Limited to the development of social needs, such phenomena did not attract much attention at Q. Huang Collaborative Innovation Center of Chemical Science and Engineering (Tianjin), State Key Laboratory of Chemical Engineering, Tianjin Key Laboratory of Applied Catalysis Science and Technology, School of Chemical Engineering and Technology, Tianjin University, Tianjin, China D. Kong (*) Department of Chemical and Metallurgical Engineering, School of Chemical Engineering, Aalto University, Aalto, Finland Department of Materials Science and Physics, China University of Mining and Technology, Xuzhou, China e-mail: [email protected] Y. Li Collaborative Innovation Center of Chemical Science and Engineering (Tianjin), State Key Laboratory of Chemical Engineering, Tianjin Key Laboratory of Applied Catalysis Science and Technology, School of Chemical Engineering and Technology, Tianjin University, Tianjin, China Department of Chemical and Metallurgical Engineering, School of Chemical Engineering, Aalto University, Aalto, Finland © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_2

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that time. With the development of human society, energy demand has been increasing dramatically. The negative effects caused by the large-scale consumption of fossil fuels also become a huge challenge such as climate change and air pollution. The chemical influence of air pollution affects the hydrologic cycle, thus contributing to water pollution and food security [2]. In the early 1970s, the outbreak of the energy crisis made people realize that it was urgent to find alternative sustainable resources to replace fossil fuels [3]. In 1972, Fujishima and Honda reported that TiO2 decomposed water to produce hydrogen under ultraviolet light irradiation [4], which immediately became a research hotspot because it provided a novel solution to directly convert solar energy to hydrogen, which is clean, renewable, and abundant on the earth. The research on the decomposition of toxic inorganic pollutants in water by photocatalysis was reported by Frank and Bard in 1977, where cyanide was decomposed in the presence of aqueous TiO2 suspensions [5]. Later on, it was discovered that TiO2 and some other semiconductors could catalyze the oxidation of organic compounds into CO2 and H2O under illumination, and similar effect on toxic inorganic compounds, such as chromium species, nitrites, and NOx [6]. In summary, photocatalysis technology is a promising alternative to protect the environment and relieve the energy crisis. In recent years, great progress has been made on semiconductor photocatalysis. Furthermore, advanced equipment and research methods have been developed and applied to the field of photocatalysis, which promotes the rapid development of both theoretical research and industrial technology.

2.3

Applications

The photocatalysis is widely applied in a variety of fields and products in the environmental and energy fields, including air and water purification, hydrogen evolution, and carbon dioxide conversion into gaseous hydrocarbons, sterilization, and disinfection.

2.3.1

Environmental Protection

Humankind has been facing up a situation of severe environmental pollution. Many kinds of industrial and domestic wastewater and gases containing toxic pollutants have been discharged to the environment. Due to the complex composition, low concentration, and poor degradation efficiency of many compounds, conventional removal methods have not been able to achieve the desired results. Photocatalytic technology provides a new way to deal with environmental pollution. It is an ideal solution because it utilizes solar energy and brings little secondary pollution. Furthermore, the photocatalysts are inexpensive and easy to produce. In 1976, Carey et al. [7] proposed that TiO2 could decompose chlorinated

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biphenyl under ultraviolet light. Since then, more and more photocatalysts have been developed for the degradation of organic/inorganic pollutants to protect the environment. The photocatalysts such as TiO2 can oxidize hydrocarbons [8, 9], halogenated hydrocarbons [10, 11], organic acids [12–14], organic dyes [15–17], and nitrogenous organic compounds [18, 19] into CO2 and H2O. Inorganic pollutants [20–22] such as Cr6+ can also be reduced by the photocatalytic reaction. The developed countries have tried to apply photocatalytic technology to solve environmental pollution problems [23, 24]. Spain’s PSA Center built the first industrial-scale wastewater treatment reactor to deal with the resin plant wastewater containing various organic pollutants such as phenol, ethylene glycol, and styrene. The photocatalytic technology is also applied to air cleaners and purification of nuclear-contaminated soil in America. Similarly, in Japan, the photocatalyst has been deposited onto the floodlights of the tunnels to decompose NOx and CO in exhausted gases emitted by the vehicles into N2 and CO2. Besides, the photocatalysts have been used as indoor formaldehyde cleaner.

2.3.2

Clean Energy Production: Water Splitting and Reduction of Carbon Dioxide

Hydrogen is considered to be one of the alternatives to fossil energy, due to its high calorific value, pollution-free combustion, and recyclability. At present, hydrogen can be produced by fossil fuel reforming, electrolysis of water, microbial fermentation [25]. However, fossil fuel reforming can bring pollution to the environment and its efficiency is low. The cost of water electrolysis is a high and similar problem with the microbial fermentation method. Using solar energy to split water is regarded as a promising method owing to its low cost, nonpollution, and sustainable utilization. Fujishima and Honda’s work in 1972 is the milestone of photoelectrochemical (PEC) water splitting. Many factors affect the efficiency of PEC water splitting including the recombination of photogenerated carriers, sluggish surface reaction kinetics, and poor light absorption [26]. Carbon dioxide emissions have increased dramatically due to the use of fossil fuels, which is mainly responsible for the greenhouse effect [27]. Therefore, reducing CO2 emissions is one of the most challenging tasks. In recent years, many scientists are working on techniques to reuse CO2 also called carbon capture and storage (CCS) [28], such as the CO2 geological storage, wet scrubbing with aqueous amine solutions to remove CO2 after combustion. The main problem is how to convert CO2 into useful chemicals and avoid the regeneration of CO2 with low energy consumption. CO2 can act as a carbon source for many organic products. However, CO2 is very stable under normal circumstances and thus large extra energy is needed to reduce it. The photocatalytic reduction of CO2 into chemicals has to attract increasing attention around the global [29]. The low light conversion efficiency is still a challenge.

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Sterilization and Disinfection

Bacteria in air or water are harmful to human health, so photocatalytic disinfection has also attracted people’s attention. The interaction between the photocatalysts and the bacteria produces active superoxide radicals and hydroxyl radicals, which can penetrate the cell wall of the bacteria and destroy the cell membranes to prevent the transport of the film-forming substances and block the respiratory system [30]. For example, the tiles decorated with TiO2 photocatalyst have the functions of bactericidal action, especially aimed at Escherichia coli, Pseudomonas aeruginosa, Staphylococcus aureus, Enterococcus faecium, etc. [31]. Such tiles could be widely used in sterilization in the daily environment (hospitals, pharmaceutical workshops, etc.). Many antibacterial tiles with photocatalyst TiO2 developed by several companies such as Toto (Japan) have been put into the market [32].

2.4

Reaction System of Photocatalysis

Up to now, there are two systems in the photocatalytic reaction: the semiconductor powder suspension system and the photoelectrochemical (PEC) cell [33]. For the former, the photocatalysts are directly dispersed in the solution, which is convenient and simple to operate. However, the photocatalysts are easy to aggregate and hard to be recovered. The latter is to fix the photocatalyst onto a conductive substrate to construct an electrode so that the oxidation and reduction reactions occur at the anode and the cathode, respectively, which facilitates the separation of the products. Meanwhile, the recombination of photogenerated carriers can be suppressed by the applied bias voltage.

2.4.1

Basic Principles of Photocatalysis

Photocatalysis belongs to the heterogeneous reaction that involves a solid phase of semiconductor and a liquid phase of electrolyte. The intrinsic photophysical properties of semiconductors play an important part in photocatalysis. This section will discuss the basic theory of photocatalysis in terms of the photophysical properties of semiconductors and the structure of the semiconductor-electrolyte interface.

2.4.2

Photophysical Properties of the Semiconductor

Semiconductors refer to materials having conductivity between conductors and insulators, which is closely related to the band structure. Illumination usually can

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enhance the conductivity of the semiconductor. This section will introduce the photophysical properties of semiconductors and the general laws of semiconductor-light interactions.

2.4.3

Band Structure of Semiconductor

The photocatalytic reaction involves the migration and reaction of photoinduced carriers in the semiconductors. The energy band structure of the semiconductor is of great significance. The energy band structure of the semiconductors is composed of the valence band (VB), conduction band (CB), and the forbidden band (the gap between VB and CB) [34]. Figure 2.1 shows the energy band of a semiconductor at a certain temperature. The actual semiconductors always have complex band structures such as impurity levels (Fig. 2.2). They can bind electrons and generate localized electron states and thus introducing impurity levels in the forbidden band. According to the valence difference, the impurity level can be divided into donor level and acceptor level; the location is generally close to the bottom of CB for the former and to the top of VB for the latter. The relevant semiconductors are called n-type and p-type semiconductors, respectively. The excitation of electrons from the donor level to CB is much easier than that from VB, that is, the electrical conduction often depends on the electrons excited by the donors to CB. It is also easier for the electron to be excited from VB to the acceptor level than that to CB. The donor and acceptor levels are also possible traps for the photogenerated charge carriers, which can be divided into shallow and deep traps. The shallow trap level approaches either the bottom of CB (for n-type) or the top of VB (for p-type), whereas the deep trap level nearly at the center of the Fig. 2.1 The energy band of a semiconductor

Fig. 2.2 Band structure of n-type and p-type semiconductors. The Ed is referred to as the donor level in the n-type semiconductor and the Ea is referred to as the acceptor level in the p-type semiconductor

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forbidden band. The deep trap level can become the recombination center, reducing the lifetime of the photogenerated charge carrier. The surface state energy levels also locate in the forbidden band and are also the recombination center.

2.4.4

Optical Absorption Spectrum of Semiconductor

When a semiconductor is illuminated by the light, the electron in the VB absorbs enough photon energy, passes over the forbidden band, and enters CB, becoming a free electron and simultaneously leaving a hole in the valence band; an electron-hole pair is generated. This process is called intrinsic absorption, which occurs only when the photon energy (hν) is greater than or equal to the forbidden bandgap (Eg) [35]. The corresponding characteristic wavelength (λ0) is the absorption edge. hν0 ¼

hc ¼ Eg λ0

ð2:1Þ

1240 Eg

ð2:2Þ

λ0 ¼

The electron transition related to intrinsic absorption can be divided into direct and indirect transitions, respectively. For direct transition, the lowest energy point of the potential energy surface of CB is vertically at the highest point of that of VB. When the photon energy absorbed by the electron is greater than the energy gap, a vertical transition occurs from VB to CB. For the indirect transition, the potential energy surface of CB drifts relative to that of VB. The electron not only absorbs photons but also exchanges a certain amount of vibration energy with the crystal lattice, that is, release or absorb a phonon according to the law of conservation of momentum. Since the energy of those extra transitions is small, the energy of indirect transition is still close to the energy gap. Figure 2.3 shows direct and indirect transitions. The electrons transfer from a low energy state to a high energy state due to the absorbed photons. The absorption coefficient (α) is used to describe the absorption intensity. Assuming that the light passes through a film with a thickness of dx, the decrement of the light intensity, dI, is proportional to the film thickness dx and the light intensity I: dI ¼ αIdx

ð2:3Þ

I ¼ I 0 exp ðαlÞ

ð2:4Þ

that is

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Fig. 2.3 (a) Direct transitions and (b) indirect transitions. Point A is the top of VB and point B(B0 ) is the bottom of CB. The hν  hν’, neglecting the phonon energy, but the wave vectors of electrons are different between point B and B0 after photoexcitation.

The light intensity I decay as the light propagates through the semiconductor medium. In addition to intrinsic absorption, exciton absorption, free carrier absorption, impurity absorption, etc. also exist and their absorption wavelengths appear in the long-wavelength region of the intrinsic absorption band. However, those absorptions are weak, so the absorption spectrum of the semiconductor is still dominated by the intrinsic absorption.

2.4.5

Emission Spectrum of the Semiconductor: Recombination

Light emission is the opposite of light absorption. For light absorption, electrons in the low energy level absorb a suitable photon and jump to the high energy level; while electrons in the high energy level can go back to the low energy level with a photon emitting during the light emission. In a balanced semiconductor, light absorption and emission occur alternately. When there are excess electrons and holes, additional radiation would be generated resulting from their recombination, and thus, the radioluminescence happens. The recombination of photogenerated electron-hole pairs in semiconductors can be classified into direct and indirect recombination. The direct recombination means the electrons jump directly back to VB from CB and recombine with the photogenerated hole. The indirect recombination means the photogenerated electron is captured by the trap level first and then leaps to VB from the trap level to recombine with the holes. The photocatalytic effect stems from the photogenerated electron-hole pairs. After the electron-hole pairs are generated, they separate and migrate to the surface of the semiconductor to carry out the redox reactions. The

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Fig. 2.4 Processes involved during photocatalysis. (i) Photoexcitation to generate charge carriers; (ii) charge carrier diffusion; (iii) bulk recombination; (iv) oxidation or reduction; (v) surface recombination

charge carrier recombination can occur during this process. Figure 2.4 shows the excitation and recombination process. After the charge carriers reach the surface of the semiconductors, whether they take part in the redox reactions depends on the positions of CB and VB of the semiconductor and the redox potential of the chemicals adsorbed on the surface of the semiconductor. Taking the PEC water-splitting reaction as an example, the conduction band potential of the semiconductor should be more negative than the hydrogen electrode potential, and the valence band potential to be more positive than the oxygen electrode potential. Many semiconductors (e.g., TiO2, ZnO, Fe2O3, CdS) can be used as photocatalysts for water splitting due to their suitable band structure. Figure 2.5 shows the band positions of several semiconductors in contact with aqueous electrolyte at pH ¼ 13.

2.5

Potential and Charge Distribution Across the Semiconductor–Electrolyte Interface

Photocatalysis is a solid-liquid heterogeneous reaction, and the interface properties will change after contacting the two phases. This section will describe the related properties of charges and potential at the interface.

2.5.1

Fermi Level

In solid-state physics, the Fermi level (EF) is the smallest possible increment of ground state energy caused by the addition of a particle in a system of noninteracting

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Fig. 2.5 Band positions of several semiconductors in contact with aqueous electrolyte at pH ¼ 1

fermions, that is, the chemical potential of the fermion system in the ground state at the absolute zero temperature, or the maximum energy of a single fermion in the ground state, or the highest energy level occupied by the fermion [36]. Based on the Fermi-Dirac statistics, the probability that an electronic state with an energy state of E is occupied by the electrons is: f ðE Þ ¼

1 exp

EE F kT




þ1

ð2:5Þ

where k is the Boltzmann constant. The Fermi level plays an important role in determining the energy of the entire system, as well as the distribution of the charge carriers. According to the statistical theory, the Fermi level is the chemical potential of the system. The Fermi level of an intrinsic semiconductor can be calculated according to Eq. (2.6): 1 1 N E F ¼ ðE v þ Ec Þ  kTln c 2 2 Nv

ð2:6Þ

where Ec, Ev is referred to CB and VB energy, respectively; Nc, Nv referred to the effective densities of CB and VB, respectively. For the intrinsic semiconductor, its valence band is filled with valence electrons, while the conduction band is empty. Correspondingly, the Fermi level is in the

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middle of the bandgap. For the n-type semiconductor, the more donors giving electrons to the conduction band are incorporated, the closer the Fermi level is to the bottom of CB. For p-type semiconductors, the higher the acceptor concentration is, the closer the Fermi level is to the top of VB. When the doping concentration is high enough to some extent, the Fermi level may enter CB or VB. The Fermi level of the solution can be obtained analogized to the solid Fermi level. When a semiconductor is immersed into a solution containing a redox couple, the oxidation state can accept electrons to become a reduction state, as the following equation shows. Ox þ ne Ð Red

ð2:7Þ

The Fermi level of the solution is equal to the electronic potential. The equilibrium of the solid-liquid system is reached when the Fermi levels or the electrochemical potentials of the solid and the solution are equal. The electronic potential (μe) is equal to the difference between the electrochemical potentials of Red (μ red) and Ox (μ ox), which can be described as: μox þ nμe ¼ μred

ð2:8Þ

The equilibrium redox potential of the standard hydrogen scale is given by the Nernst equation: φ ¼ φo 

RT C ln ox nF C red

ð2:9Þ

Combining eqs. (2.8) and (2.9): 1 RT C ln ox μe ¼ ðμred  μox Þ ¼ φo  n nF C red

E e ¼ eμe ¼ eφo 

kT C ln ox n Cred

ð2:10Þ

ð2:11Þ

Thus, the Fermi level of the solution can be described as follows: ) F redox ¼ F o 

kT C ln ox n C red

ð2:12Þ

where Fredox is the Fermi level of the solution. Cox and Cred are the concentration of the oxidation state and reduction state, respectively. Fo is the energy when both the oxidation state and reduction state are under equilibrium.

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2.5.2

37

Space Charge Layer and Band Bending

If a semiconductor and a solution containing a redox couple have different Fermi levels, the electrons will migrate at the interface, resulting in different concentration of charge carriers at the surface and in the bulk. When the Fermi levels of the semiconductor and the solution are the same, the system reaches equilibrium. Meanwhile, the semiconductor band bending occurs at the surface, that is, the positions of the energy bands are different from those in the bulk. The region of the bent surface band is called the space charge layer. An n-type semiconductor will be used to illustrate the formation of a space charge layer, while the case of a p-type semiconductor is just the opposite. Figure 2.6 illustrates the formation of three types of space charge layers of an ntype and a p-type semiconductor. If the Fermi levels of the semiconductor and the solution are equal, the electrons would not flow at the interface, thus no space charge layer forms at the semiconductor surface. If the Fermi level of the n-type semiconductor is lower than that of the solution, the electrons will migrate from the solution to the surface of the semiconductor until the Fermi levels are equal. Extra electrons gather on the surface of the semiconductor, causing the band to bend downward relative to the bulk. An accumulation layer is thus formed. On the contrary, if the Fermi level of the n-type semiconductor is higher, then the electrons transfer to the solution. The concentration of electrons at the surface is lower than that in the bulk. The surface energy band of the semiconductor will bend upward relative to the bulk phase to form a depletion layer. Based on the depletion layer, electrons continue to transfer from the semiconductor to the solution so that the concentration of majority carriers (electron) on the surface of the semiconductor is lower than the minority carriers (hole). In this case, an inversion layer is formed. For p-type semiconductor, if the Fermi layer of the semiconductor is higher than that of the solution, an accumulation layer is formed. Otherwise, an depletion layer is formed.

Fig. 2.6 Space charge layer models for n- and p-type semiconductors in contact with solutions

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Q. Huang et al.

Flat-Band Potential

The flat band potential is an important feature of the semiconductor/electrolyte system. It determines the semiconductor energy band position. When a semiconductor is immersed in the electrolyte, the band bending at the surface of the semiconductor occurs due to the difference of the Fermi levels of the semiconductor and the solution. A space charge layer is formed on the side of the semiconductor and a Helmholtz layer on the side of the electrolyte, as a result of the band bending on the surface of the semiconductor. If a voltage bias is applied to the semiconductor for polarization, the Fermi level of the semiconductor will be changed. The voltage bias to flatten the energy band at the surface is called the flat band potential (φfb). The flat band potential can be obtained by measuring the differential capacitance of the semiconductor space charge layer. The capacitance of the system includes the space charge layer capacitance (CSC) of the semiconductor and the Helmholtz layer capacitance (CH) of the solution in tandem. Generally, CSC is much larger than CH, so the system capacitance is approximately equivalent to CSC. CSC can be tuned by changing the polarization potential (φ). According to the Mott–Schottky equation, the relationship between CSC and V can be described as:   1 2 kT ¼ ΔφSC  CSC qεo εeN D e

ð2:13aÞ

where ΔφSC is the potential drop of the band bending, that is, ΔφSC ¼ φ  φfb. ε, εo are the dielectric constant of the semiconductor and the vacuum, respectively. ND is the doping concentration. It indicates that the reciprocal of CSC is linearly related to the polarization potential. The flat potential and doping concentration can be calculated from the x intercept and slope of 1/Csc ~ ΔφSC, respectively.

2.5.4

Quasi-Fermi Level

The electrons in the semiconductor have the same Fermi level under the thermal equilibrium state. However, if a disruption of such illumination is added to the semiconductor, the equilibrium is broken to generate nonequilibrium carriers. Because the timescale of carrier transfer within a band is much smaller than that of the electron-hole recombination across the bandgap, it is assumed that the carriers in the valence band or the conduction band are still in the equilibrium state, while the valence band and the conduction band themselves are not. Here, a local Fermi level, quasi-Fermi level (QFL), is introduced. The nonequilibrium between the valence band and the conduction band is manifested by different QFLs. QFL of the conduction band is also called the electron quasi-Fermi level, EFn. Correspondingly, the quasi-Fermi level of the valence band is called the hole quasi-Fermi level, EFp. The QFLs for electrons and holes are defined as below:

2 Applications and Fundamentals of Photocatalysis with Solar Energy

n NC p ¼ Ev þ kTln NV

E Fn ¼ Ec þ kTln E Fp

39

ð2:13bÞ ð2:14Þ

where n and p are the concentration of photostationary-state electron and hole, respectively. In general, QFL would deviate from the equilibrated Fermi level with increased concentrations of carriers. Besides, the QFL of the majority carrier deviates much less than that of the minority carrier. Figure 2.7 shows the QFL for an n-type semiconductor under illumination. EFn is close to EF, while EFp deviates from EF seriously and draws near the VB. The position of EFp at the surface is determined by the illumination intensity and the kinetics of interfacial charge transfer.

2.5.5

Surface States

The semiconductor surface is not regular because the ideal periodic lattice is interrupted at the surface. The atoms on the surface have unpaired electrons, that is, dangling bonds, which are capable of exchanging electrons or holes with the bulk, as well as adsorbing external species. These electronic energy levels are referred to as surface state energy (Ess) and they play an important role in the semiconductorliquid interface. Taking an n-type semiconductor as an example, it is used as a photoanode of a PEC cell [37, 38]. As illustrated in Fig. 2.8, the photoexcited holes in the valence band should diffuse to the surface to oxidize Red to Ox in the solution. However, they can be trapped by the surface states and recombine with the electrons. To avoid the surface state trapping, the position of the Fermi level in the Fig. 2.7 Quasi-Fermi level of an n-type semiconductor under illumination

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Fig. 2.8 The downward shift of the Fermi level at the surface alters the electron occupation of surface states, and the corresponding change in surface charge leads to Fermi level pinning

semiconductor is adjusted by extra bias to regulate the occupation of Ess. Initially, the Fermi level of electrons (EFn) is located above the Ess so the surface states are full of electrons and no redox reaction happens. With more positive potential added to the n-type semiconductor, the EFn shifts downward until it is equal to Ess. Meanwhile, the Fermi level of holes (EFp) matches with the Fredox and the holes can initiate the oxidation reaction. Finally, the EFn is smaller than Ess when bias is positive enough, leading to empty surface states and increased photocurrent.

2.6

Strategies for Improving the Performance of Photocatalysts

There are many types of semiconductor materials, for example, elementary substances represented by Si, transition metal oxides represented by TiO2, compounds of IIB group elements and VIA group elements represented by CdS, and compounds of IIIA-VA group elements represented by GaAs. With the development of photocatalysis research, many complex compounds such as titanates (NaTaO3 [39, 40]), vanadates (BiVO4 [41, 42]), tantalites (SrTiO3 [43–45]), tungstates (Bi2WO6 [46, 47]), oxyhalides (BiOCl [48, 49]), and solid solutions (Bi2InNbO7 [50]) have been constructed. Although these semiconductors have been developed for PEC water splitting and water pollutant decomposing, they still suffer from low efficiencies, which originates from poor light absorption, severe photogenerated carrier recombination, slow surface reaction kinetics, etc. Therefore, it is of great importance to design the photocatalysts rationally. This section will focus on four strategies to modify the semiconductors in terms of doping, heterostructure construction, morphology control, and cocatalyst loading.

2.6.1

Doping

Doping is an effective way to change the physicochemical properties of the semiconductors. The addition of foreign elements to the semiconductor crystal lattice can

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cause lattice defects, thereby affecting the energy band structure and the movement and distribution of photogenerated carriers. Typically, doping can serve the purposes as below: (i) Doping can introduce some impurity levels in the forbidden band, which makes it possible for the semiconductors with a wide bandgap to absorb longer wavelength photons in the visible light. (ii) The dopants can participate in the capture and release of photogenerated electrons and holes, control the diffusion of carriers during the photocatalytic reaction process. (iii) Doping can increase the carrier density and improve the electrical conductivity of the semiconductors. (iv) Doping affects the shape-controlled growth of semiconductors. Many methods can be applied for doping, such as hydrothermal methods [51, 52], ion thermal diffusion from the substrate [53, 54], dip coating [55], impregnation, high-energy ion implantation [56, 57]. For wide bandgap semiconductors, such as TiO2 with a bandgap of ~3.2 eV, they usually suffer from limited light absorption. Dopants have been introduced to overcome this issue. Lee et al. [52] prepared transitions metal-doped TiO2 by the hydrothermal reaction. The UV–visible spectroscopy analysis showed that Fe3+-, Cr4+-, Ni2+-doped samples can significantly enhance the light absorption in the visible region. Accordingly, the doped TiO2 achieved higher benzene photodecomposition efficiency in the artificial sunlight, especially the Cr4+-modified sample. Nonmetallic elements can also tune the TiO2 energy band. Asahi et al. [58] reported an N-doped TiO2 whose absorption edge was shifted to the visible light region. Due to the narrowed bandgap, it achieved photodegradation of methylene blue and acetaldehyde under visible light. Other nonmetallic elements such as C [59], S [60, 61] have also been studied. It should be noted that the impurity level should be tuned properly in the bandgap, otherwise it will become the recombination center. This is closely related to the concentration, distribution, and d-electron configuration of the dopants. Choi et al. [62] studied the effects of 21 kinds of metal ions doped TiO2 used for chloroform oxidation and carbon tetrachloride reduction. The Fe3+, Mo4+, Ru3+, Os3+, V4+, and Rh3+ ions with a doping concentration of 0.1–0.5% can significantly increase the TiO2 photocatalytic activity, whereas the Li+, Mg2+, A13+, and Ga3+ ions have a negative influence. The ions with positive effect can act as both electron and hole traps and the trapped charge carriers have a sufficiently long lifetime to reach the surface by detrapping and electron tunneling. Hematite (Fe2O3) can respond to visible light due to its narrow bandgap (1.9–2.2 eV), which is an advantage relative to TiO2. However, Fe2O3 has a poor conductivity due to the highly localized small polarons, leading to serious carrier bulk recombination. To enhance electrical conductivity, many elements have been incorporated such as Sn [51], Ti [54], Al [63], Cr [64], Si [65], P [66]. Ling et al. prepared Sn-doped Fe2O3 by hydrothermal method and Sn was introduced by a facile dip coating procedure. The Mott–Schottky plots showed that the donor density

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was much higher after ex situ doping [67]. Cesar et al. [68] prepared Si-doped Fe2O3 and found that, in addition to the increased conductivity, Si doping reduced the particle size which could reduce the diffusion length for carriers from the bulk to the surface, further decreasing the recombination. The distribution of dopants in lattices is critical for the photocatalytic activity. Luo et al. [66] proposed a P gradient-doped Fe2O3 nanoarray by controlling the annealing time after a facile dipping. From Fig. 2.9a, b, the Mott–Schottky plots indicated that whether gradient doping or homodoping was helpful to increase the carrier density. However, the graded distribution of the P element made the band bending widen and accelerated the separation of photogenerated carrier compared to that of homogeneous distribution so as to enhance photocurrent. Besides introducing foreign elements, native defects like vacancies or interstitials also influence the carrier densities of semiconductors. Oxygen vacancy is very common in metal oxides and useful in tuning the electrical conductivity and photocatalytic activity [69–72]. For example, high-temperature calcination is generally needed for the iron oxide materials prepared by the hydrothermal method. Xiao et al. found that rapid cooling caused a partial reduction of the surface of iron oxide (Fig. 2.10a), that is, oxygen vacancies were introduced, which contributes to increased carrier densities (Fig. 2.10b), and thus, better PEC water-splitting performance [73]. Similar to doping, the fabrication of a solid solution that incorporates a certain component into a crystal usually through a solid reaction, and the product retains the structure of the original one, is also an approach to regulate the crystal properties. There are many multicomponent semiconductors reported for photocatalysis, for example, GaN:ZnO [74–76], AgSrNbTiOx [77, 78], Cu3MS4 (M ¼ V, Nb, Ta) [79], ZnxCd1-xS [80, 81], CdxIn1-xS [82], In–Ni–Ta–O–N [83], GaP–ZnS [84]. The bandgap of multicomponent semiconductors can be adjusted by changing the stoichiometry of the different components. Li et al. [80] found that the absorption band

Fig. 2.9 (a) Mott–Schottky plots of bare Fe2O3, homo-P:Fe2O3, and grad-P:Fe2O3 photoanodes. (b) I–V characteristics of bare Fe2O3, homo-P:Fe2O3, grad-P:Fe2O3, homo-P:Fe2O3/Co-Pi, and grad-P:Fe2O3/Co-Pi photoanodes measured in 1 M KOH solution under AM 1.5G illumination. (Reprinted with permission from [65], copyright 2017 Royal Society of Chemistry)

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Fig. 2.10 (a) High-resolution XPS spectra of the Fe 2p. (b) Mott–Schottky plots. RC referred to rapid cooling, NC referred to natural cooling, SC referred to slowing cooling. (Reprinted with permission from [72], copyright 2017 Elsevier B.V.)

edge was red shifted continuously as the Cd content increased in the ZnxCd1-xS solid solution as shown in Fig. 2.11A, consistent with the color change from white to yellow. The sample with a Zn/Cd molar ratio of 1:1 exhibited the best H2 evolution performance. Another typical example is the GaN:ZnO solid solution. The bandgaps of ZnO and GaN are too wide to absorb visible light, while their solid solutions show a narrowed bandgap for the visible light absorption reflected from the UV–visible absorption spectrum in Fig. 2.11B [74]. Although separated InN or GaZnON shows no water-splitting activity, the Ni-InN/GaZnON solid solution exhibits overall water splitting under visible light and good photocatalytic stability [85].

2.6.2

Heterojunction/Homojunction

As a photocatalyst, semiconductors usually should have a suitable bandgap to ensure light absorption, efficient photogenerated carrier separation, and resistance to photocorrosion. Also, the energy levels of the conduction/valence band are proper to drive the redox reaction. However, no single semiconductor has met all of the above conditions so far. Correspondingly, the integration of two or more materials, that is, design heterojunction/homojunction, is a solution. A homojunction is from the interface of the same semiconductor materials but with different doping. On the contrary, a heterojunction is from the interface of different semiconductors (semiconductor-semiconductor heterojunction) or that of metals and semiconductors (metal-semiconductor heterojunction). Both can be built to improve the photocatalytic performance.

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Fig. 2.11 (A) UVvisible diffuse reflection spectra of the Zn1xCdxS (x ¼ 0, 0.1, 0.3, 0.5, 0.7, 0.9, and 1.0) samples. Reprinted with permission from [79], copyright 2013 Royal Society of Chemistry. (B) UV–visible diffuse reflectance spectra: (a) GaN, (b) GaN:ZnO (Zn 3.4 atom %), (c) GaN: ZnO (Zn 6.4 atom %), (d) GaN:ZnO (Zn 13.3 atom %), and (e) ZnO. Reprinted with permission from [73], copyright 2005 American Chemical Society

2.6.3

Semiconductor-Semiconductor Heterojunction

Regardless of the influence of surface states, the energy band of heterojunction is shown in Fig. 2.12. When two n-type semiconductors (SC 1 and SC2) with different energy band positions and Fermi level contact with each other, the electrons flow from the side with higher Fermi level (SC1) to the other side with lower Fermi level (SC 2). A built-in electric field is constructed at the interface, whose electric field direction is from SC1 to SC2. When the Fermi levels of the two semiconductors reach the same, the electrons will stop flowing and the heterojunction is formed. The charge layers are formed on both sides of the semiconductor contact surface. If both the semiconductors are n-type, a depletion layer will be formed on the surface of SC1 who loses electrons and an accumulation layer will be formed on the surface of SC2 who accepts excess electrons. The potential difference between the two sides is the contact potential difference which is equal to the difference between the two Fermi levels. Due to the redistribution of electrons at the interface, the carrier concentration of the semiconductor surface is different from that of the bulk, resulting in band bending. An ideal semiconductor-semiconductor composite system should satisfy the following conditions: (a) the energy band structures of the coupled semiconductors can be matched to each other so that the band bending at the interface can force photogenerated electron-hole pairs to separate; (b) the ohmic contact is necessary, so the charge can move smoothly at the interface; (c) after the charge is separated at the interface, the semiconductor as an electron acceptor should have excellent electron transport capability, and the other semiconductor as electron donor should be in favor of hole transport.

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Fig. 2.12 Heterojunctions between two semiconductors. (a) n-n heterojunction and (b) p-n heterojunction

Wide bandgap semiconductors have limited absorption of sunlight, but combining them with narrow bandgap semiconductors can solve this problem. TiO2 is a typical wide bandgap semiconductor and hardly responds to visible light. Many semiconductors (metal oxides: Bi2WO6 [86], WO3 [87, 88], Fe2O3 [89–91]; sulfides: CdS [92, 93], MoS2 [94]) have been integrated with TiO2 and successfully improved the visible light absorption. Aside from visible light activity, a reasonable heterojunction can also promote the separation of photoinduced carriers due to the built-in electric field, such as p-n junctions. Dai et al. [95] deposited BiOI on the TiO2 nanotube wall. A p-n heterojunction was formed at the interface. As shown in Fig. 2.13a, such an internal field facilitated the excited electrons on the conduction band of the p-type BiOI to transfer to that of n-type TiO2, while the holes remained in the BiOI valence band, suppressing the carrier recombination. Similarly, MoS2 nanosheet-coated TiO2 nanobelts could split water to produce hydrogen and decompose the organic dye Rhodamine B [96] due to the energy band matching of MoS2 and TiO2, which not only increases the light absorption, but also inhibits the photogenerated carrier recombination. Cong et al. [97] deposited iron oxide nanoparticles onto TiO2 nanotubes by electrodeposition. The modified TiO2 nanotubes could decompose phenol under visible light, and the decomposition efficiency was greatly improved owing to the formation of heterojunction. Graphitic carbon nitride (g-C3N4) is a metal-free semiconductor and has attracted great attention since it was applied in photocatalysis considering its suitable bandgap for visible light absorption, high stability in an aqueous solution, and nontoxicity [98–100]. It can also be used for heterojunction formation with other semiconductors like TiO2 [101–103], ZnO [104–107], CdS [108, 109], and Fe2O3 [110–112]. When g-C3N4 is loaded on a semiconductor, the light absorption and the photogenerated carrier separation efficiency can be significantly improved if band structures are matched. The absorbance intensity of the g-C3N4–ZnO composite photocatalysts

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Fig. 2.13 (a) Schematic illustration for the formation of the p-n junction between p-type BiOI and n-type TiO2 under visible light irradiation. Reprinted with permission from [94], copyright 2011 American Chemical Society. (b) UV–vis diffuse reflectance spectra of g-C3N4, ZnO, and g-C3N4– ZnO composites. Reprinted with permission from [103], copyright 2012 Royal Society of Chemistry

became stronger and the absorption edge shifted to the lower energy displayed in Fig. 2.13b. Accordingly, an enhanced photodegradation of methyl orange (MO) and p-nitrophenol was obtained [104].

2.6.4

Metal-Semiconductor (M-S) Heterojunction

The heterojunction can also be formed between metals and semiconductors, generally based on the n-type semiconductors. When the metal is in close contact with the n-type semiconductor, a heterojunction, also called Schottky junction, with an energy barrier can be created. Metals and n-type semiconductors have different Fermi levels. Usually, the work function of the metal (φm) is higher than that of the semiconductor (φs). Therefore, the electrons will continuously transfer from the semiconductor to the metal until a new equilibrium if the two are closely integrated. The metal acts as an electron collector, and the excess positive charge will accumulate on the surface of the n-type semiconductor so that the bands will be bent upward and a depletion layer is generated (Fig. 2.14). M-S heterojunction has been employed in photocatalysis, especially for Ag [113– 115], Au [116–118] nanoparticles with the property of localized surface plasmon resonance. Intrinsically they have strong absorption for visible light and can promote the charge transfer with semiconductors where plasmonic nanoparticles are loaded on. Wu et al. [119] produced Ag/TiO2 nanotubes. After the Ag nanoparticles were deposited, the visible light absorption was enhanced. Other noble metals like Pd [120] and Pt [121, 122] also appear in M-S heterojunction for photocatalysis. They can act as electron collectors to accelerate charge carrier separation at the surface of semiconductors [123, 124]. Owing to the high cost of noble metals, cheaper alternatives should be developed. Some earth abundant metals, such as Ni [125], Co [126], Cu [127] can play a similar role. Chen et al. [128] demonstrated that the H2

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Fig. 2.14 Schematic of Schottky junction formed between metals and n-type semiconductors

production rate over 1.2 wt% Ni-loaded CdS is nearly twice higher than that of 1.2 wt% Pt-loaded CdS under visible light.

2.6.5

Carbon-Semiconductor Heterojunction

In addition to the non-noble metals, carbon materials such as graphite, carbon nanotubes (CNTs), and carbon quantum dots (CQDs) are good alternatives to noble metals. Besides excellent conductivity, graphite and CNTs can supply a large surface area. Yu et al. [129] prepare mesoporous multiwalled CNTs/titanium dioxide. Only the chemically bonded composites could facilitate the separation of photogenerated electron-hole pairs and decrease their recombination rate and thus significantly enhanced photocatalytic activity for degrading acetone in the air under UV irradiation. Li et al. [130] synthesized graphene nanosheets decorated CdS clusters through a solvothermal method. These nanosized composites exhibited a high H2 production rate owing to the improved carrier transport by graphene. For CQDs, they stand out due to the size-dependent photoluminescence and the up-conversion photoluminescence, that is, CQDs can help harvest more photons. Yu et al. [131] reported that the absorption band edge of the CQDs-sensitized TiO2 composites was extended to the visible light region. The CQDs serve as electron reservoirs to trap the photoinduced electrons and enhance the separation of photoinduced electron-hole pairs. Accordingly, the photocatalytic activity of H2 production was improved. A variety of heterojunctions can be combined to modify the photocatalysts, which means the semiconductor-semiconductor heterojunctions and M-S heterojunctions can be produced in tandem. Zou’s group fabricated a ternary MoS2-graphene/ ZnIn2S4 composite, which showed a higher H2 photocatalytic evolution rate than Pt-loaded ZnIn2S4 photocatalyst [132]. A schematic was exhibited in Fig. 2.15a. It was believed that the graphene served as an electron transport bridge between the MoS2 and the ZnIn2S4, resulting in reduced charge recombination. Zhang et al.

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deposited Ag nanoparticles to the interface of nanoscale ZnO/CdS, obtaining a ZnO/Ag/CdS nanocomposite. The modification with Ag nanoparticles significantly enhances the light absorption and facilitates the separation of photogenerated carriers through the localized surface plasma resonance (Fig. 2.15b) [133]. Li’s group constructed a novel nanocomposite NiO/CdS@ZnO photoanode for water splitting. The well-matched band structures between the composite components significantly promoted the transfer of charge carriers as evidenced by the decreased photoluminescence intensity [134].

Fig. 2.15 (a) Schematic illustration for photocatalytic hydrogen generation over MoS2-graphene/ ZnIn2S4 under visible light irradiation. Reprinted with permission from [131], copyright 2016 Elsevier B.V. (b) The schematic diagrams of the reaction process and carriers transport in ZnO/Ag/ CdS composite photoanode. (Reprinted with permission from [132], copyright 2014 Elsevier B.V.)

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Homojunction

Doping can regulate the energy band structure of a semiconductor, which makes it possible to form a homojunction between different phases with the same chemical composition. Abdi et al. [135] constructed heterojunctions in series based on W-doped BiVO4 to suppress the photocarriers recombination. They fabricated the gradient-doped W:BiVO4 by spray pyrolysis. Two phases with different doping concentrations form a homojunction, and the photogenerated carriers can be separated and diffused to the surface under the driving force of the built-in electric fields in series. Similar junctions were applied to the ZnO nanostructure. Xiao et al. produced a ZnO:Ga/ZnO photoanode by a two-step hydrothermal reaction. The as-prepared ZnO:Ga/ZnO composite with nearly the same lattice structure in the two phases exhibits enhanced light absorption and improved electrical conductivity [136].

2.7

Morphology Control

The morphology of photocatalysts plays an important role in photocatalytic performance. A sufficient surface area is necessary to ensure light absorption and reactive sites. The photocatalysts should grow along special crystal planes due to the crystal anisotropy. The microstructure determines the transfer distance of the carriers from the excitation site to the surface, which should match the lifetime of the photogenerated hole-electron pairs to inhibit the recombination [137]. In such a context, nanostructures are desirable to achieve high performance. Accordingly, various nanostructures such as zero-dimensional (0D), one-dimensional (1D), twodimensional (2D), and three-dimensional (3D) have been proposed. 0D nanostructure – 0D nanostructures are typically nanoparticles, like nanospheres, core-shell structures, hollow spheres. Chen et al. produced submicrometer TiO2 hollow spheres (Fig. 2.16a) by a template-free solvothermal route. The hollow nanostructures exhibited good photocatalytic activity on the degradation of phenol, even superior to Degussa P25 [138]. In general, the smaller the size of the nanoparticle, the more specific surface area can be provided. However, when it is smaller enough to exhibit the quantum confinement effect, the absorption edge will move to the blue side with reduced light absorption. The band bending, a built-in field for the photon-generated carrier separation also disappears [139]. Also, if such nanoparticles are immobilized on a conductive substrate, the recombination of photogenerated carriers will be enhanced by the grain boundaries [140]. 1D nanostructure – Compared to the 0D nanostructures, the 1D nanostructures not only have a large surface area but also can provide an orthogonal direction for

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Fig. 2.16 Examples of 0D and 1D nanostructures: (a) TiO2 nano hollow spheres. Reprinted with permission from [137], copyright 2011 American Chemical Society. (b) TiO2 nanotubes. Reprinted with permission from [147], copyright 2001 Materials Research Society. (c) Zr-doped Fe2O3 nanotubes. (Reprinted with permission from [148], copyright 2017 Wiley-VCH)

charge carrier separation that significantly enhances electron/hole collection at the back-contact in a PEC cell. Based on the high length-to-diameter ratio for the 1D nanostructures, light absorption and scattering are significantly enhanced. 1D nanostructures such as nanotubes (NTs) [59, 141, 142], nanorods (NRs) [143– 145], and nanowires (NWs) [67, 146] have been fabricated by hydrothermal reaction [53], electrodeposition [64], chemical vapor deposition [147], etc. Titanium dioxide nanotubes (Fig. 2.16b) are usually prepared by one-step anodization of titanium foil, which was first reported in 2001 by Gong’s group [148]. Recently, Li et al. [149] prepared Zr-doped Fe2O3 nanotubes with an outer diameter of ca. 100 nm (as displayed in Fig. 2.16c) through a solid-state reaction at a high temperature of 800 C. Such a nanostructure of Zr-doped Fe2O3 provided a shortened charge transfer distance and an enlarged depletion layer at both inner and outer surfaces when contacting with the electrolyte. 2D Nanostructure – 2D nanostructures include nanosheets, nanoflakes, a nanobelt, etc. They exhibit relatively large lateral size and an ultrathin thickness confined to the atomic scale regime that collectively leads to surface areas typically larger than those of 1D nanostructures. Moreover, it is easier to control the exposed active facets in 2D nanostructures. By precise control of the growth conditions, some nanosheets provide almost 100% exposed active sites, such as the (110) facet exposed Fe2O3 nanosheets (Fig. 2.17a, b) obtained through a silica hydrogel-mediated dissolution-recrystallization process [150]. Such ultrathin Fe2O3 nanosheets exhibited superior degradation photoreactivity of an organic contaminant in comparison with the α-Fe2O3 nanoparticles. The impressive enhancement was attributed to a special lamellar nanostructure with high surface area and improved conductivity. Another way to synthesize Fe2O3 with 2D nanostructure has been demonstrated by Yang’s group. They oxidized the iron substrate under the thermal flow of O2 and the Fe2O3 nanobelts were formed on the surface [151].

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3D Nanostructure – It is also called a hierarchical nanostructure. It consists of a conspicuous backbone with nanobranches which are usually nanoparticles, nanowires/rods/tubes, and nanosheets. Thus, the 3D nanostructure often resembles nanotrees or nanoflowers with significantly increased surface area and a much larger number of active sites, owing to the high arrangement provided by the branches [137]. Hierarchical nanostructures can either be composed of a single material or a combination of two or more. For only one material, Liu et al. [152] reported that a flower-like CdS was fabricated through a low-temperature mixed solvothermal strategy. The flower structure constructed of several assembled nanosheets with a highly preferred (002) facet is in favor of light absorption and charge transfer. Thus, the flower-like CdS displayed better performance in photocatalytic hydrogen evolution than nanobelt CdS. Wang and coworkers fabricated ZnO nanoflowers (Fig. 2.17c, d) assembled by ZnO nanorods [153]. It was demonstrated that flowers structured ZnO could be controllable by simply adjusting the basicity in the solution. The as-prepared ZnO nanoflowers had a higher content of oxygen vacancies to serve as active centers, leading to enhanced photocatalytic activity. Cauliflower-type structured Fe2O3 with a minimum feature size of 5–10 nm (Fig. 2.17e, f) was produced by atmospheric pressure chemical vapor deposition [68, 154]. It was found that silicon doping played an important part in the formation of such a fine nanostructure. Hierarchical nanostructures composed of two or more materials are more advantageous. The backbone and branches can be made up of different compositions, thus it is possible to construct heterojunctions: WO3-Fe2O3 [155, 156], CdS-TiO2 [157, 158], ZnO-Fe2O3 [159, 160], and Fe2O3-TiO2 [161–163] assemblies have been reported. Han et al. [164] fabricated a 3D hierarchical hetero-nanostructures composed of thin α-Fe2O3 nanoflakes branched on TiO2 nanotubes, which exhibited enhanced activity in photoelectrochemical water splitting and photocatalytic degradation of an organic dye compared with pure TiO2 nanotubes due to the optimized light absorption and the enhanced separation of photogenerated charge carriers at the interface. Yang’s group constructed tree-like heterostructures comprising Si nanowires as backbones and TiO2 nanorods as branches as showed in Fig. 2.18 [165]. The nanotree heterostructure displayed much higher photoactivity than the configuration where a TiO2 thin film is partially covered onto Si nanowire due to light absorbers with large surface area and complementary electrical properties of two materials.

2.8

Loading Cocatalysts

The photocatalytic effect of semiconductors originates from the photoexcited holes with oxidizing ability and electrons with reducing ability. However, photogenerated carriers on the surface tend to recombine rather than react due to the surface states.

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Fig. 2.17 Examples of 2D and 3D nanostructures: (a, b) Fe2O3 nanosheets. Reprinted with permission from [149], copyright 2017 Elsevier B.V. (c, d) ZnO nanoflowers. Reprinted with permission from [152], copyright 2017 Elsevier B.V. (e, f) Fe2O3 cauliflower-type nanostructure. Reprinted with permission from [67], copyright 2006 American Chemical Society

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Fig. 2.18 Examples of hierarchical nanostructures: (a) Schematic of the Si-TiO2 tree-like heterarchical structure. (b, c) SEM images of Si-TiO2 nanotrees. (Reprinted with permission from [164], copyright 2017 American Chemical Society)

Also, slow surface reaction kinetics accelerates the surface recombination, which severely inhibits the efficiency of the photocatalytic reaction. A common and often useful strategy to solve those problems is to integrate a cocatalyst with a photocatalyst. Generally, suitable cocatalysts can accelerate the surface reaction kinetics and inhibit the surface electron-hole recombination, that is, the cocatalysts can capture the holes/electrons from the semiconductors and then transfer them to the liquid, providing a shortcut for the surface reaction. Some semiconductors are electrochemically unstable in the electrolyte environment or easy to be corroded by the photoinduced holes. The cocatalysts can act as physically protective layers or can extract photogenerated carriers from the semiconductors and help them participate in the reaction rather than accumulate on the semiconductor surface to prevent semiconductors from photocorrosion [166]. For PEC water splitting, cocatalysts can be derived from electrocatalysts such as metal phosphates/phosphides [167–169], metal oxides [170–172], (oxy)-hydroxides [173–175]. They can be prepared by atomic layer deposition (ALD), electrochemical deposition, spin coating, dip casting, etc. Cobalt phosphate (Co-Pi) is one of the typical electrocatalyst [176]. Many groups have successfully integrated Co-Pi onto photocatalysts [73, 177, 178], and achieved significantly improved PEC water-splitting performance. However, there are different perspectives in terms of the role of Co-Pi. Klahr et al. [179] deduced that Co-Pi could collect the photogenerated holes from the Fe2O3 photoanode. Co(III) species in Co-Pi film was oxidized to Co(IV) which is more beneficial for water oxidation, leading to reduced resistance and accelerated surface reaction kinetics (Fig. 2.19a, b). However, based on the impedance and PEC measurement, Carroll et al. [180] considered that slow charge transfer between the solid-liquid interfaces limited the photocurrent of Co-Pi/Fe2O3 photoanode while a reduced hole-electron recombination was displayed compared to bare Fe2O3. And the latter was more dominant to obtain a net enhancement of water-oxidation efficiency. In short, those different viewpoints probably are originated from the differences in the loading amount of Co-Pi cocatalysts and the microstructure of the underlying photoanodes [181]. Besides Co-Pi, Fe-Pi is also an efficient cocatalyst for Fe2O3 photoanode.

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Fig. 2.19 (a) Nyquist plots measured under the illumination of a bare hematite electrode (red circles) and with 15 mC cm2 CoPi catalyst (green triangles) measured at 1.25 V versus RHE. (b) Diagram of a CoPi-coated hematite electrode under illumination and applied bias. Reprinted with permission from [178], copyright 2012 American Chemical Society. (c) Applied voltage versus densities of surface states (DOS) of the electrodes. Solid lines correspond to Lorentzian fit results. (d) Proposed kinetic scheme of the hole transfer pathways from the VB of Fe2O3 to the surface states for the Fe2O3 and FeCoW/Fe2O3 photoanodes under illumination at a bias of 1.0 V versus RHE. Shaded and blank areas represent electron filled or empty states, respectively. (Reprinted with permission from [190], copyright 2017 Elsevier B.V.)

When the surface of Fe2O3 was transformed into amorphous iron phosphate, the oxygen atoms are covalently fixed in phosphate. Accordingly, the oxygen vacancies are decreased and the surface states are effectively suppressed, leading to improved photocatalytic activity [167]. Some metal oxides also have similar passivation effect [182, 183]. Formal et al. [183] deposited an ultrathin alumina overlayer on hematite by ALD. The onset potential of the modified Fe2O3 photoanode was reduced by 100 mV due to the suppression of surface states. Many excellent catalysts (e.g., Co- and Ni-based oxides and (oxy)-hydroxides) for electrocatalytic oxygen evolution reaction (OER) can be employed as cocatalysts for photoanode. Zhan et al. [184] deposited CoOx on the WO3 surface and obtained an increased PEC water-splitting efficiency. The cobalt ion valence cycle (Co3+/Co4+) suppressed the surface carrier recombination and enhanced the reaction kinetics. A similar CoOx/WO3 structure was fabricated by Ding’s group [185]. They attributed the improved charge transfer kinetics and Faradaic efficiency to the formation of pn heterojunctions and the increased oxidation selectivity caused by CoOx.

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Also, CoOx can protect the photoanode as well. Higashi et al. [186] claimed that the stable photocurrent of modified TaON was due to the highly dispersed CoOx nanoparticles that efficiently collected photogenerated holes to prevent TaON from self-oxidative deactivation. A similar effect was found in the NiOx layer-coated Si photoanode [187]. Du et al. [188] reported that NiFeOx-decorated Fe2O3 showed an excellent turn-on voltage of 0.61 V versus reversible hydrogen electrode (RHE). The author explained the negative shift of turn-on voltage compared to bare Fe2O3 from the thermodynamics rather than the kinetics, that is, the NiFeOx layer increased the photovoltage of the Fe2O3 photoanode. An amorphous NiOOH layer on the Fe2O3 photoanode led to ca. 150 mV cathodic shift in turn-on voltage [189]. The photogenerated holes could be trapped by Ni2+/Ni3+ pairs in Ni(OH)2/NiOOH before the surface states, followed by detrapping to oxidize water, resulting in improved photocurrent onset potential when the bias was low. When the bias was up to 1.25 V versus RHE, the oxidation process (Ni3+/Ni4+) occurred which was slow, though Ni4+ could oxidize water fast. A similar effect was found in Ni(OH)2 [190]. Li’s group reported a FeCoW oxyhydroxide gel modified α-Fe2O3 photoanode. The onset voltage of the composite photoanode was a shift to the cathode side by 170 mV. The FeCoW oxyhydroxide gel acted as a hole collector to inhibit surface recombination of Fe2O3 as well as shifted the surface state to a higher energy position to obtain a reduced turn-on voltage (Fig. 2.19c, d) [191].

2.9

Conclusions and Future Trends

Photocatalysis is a green technology because it employs inexhaustible solar energy to realize energy conversion without any toxic and harmful by-products, beneficial to both environmental protection and energy development. However, up to now, no semiconductor can perfectly achieve high solar conversion efficiency. Over the past few decades, many different strategies have been successfully proposed, and significant progress has been made to modify photocatalysts. To develop high-performance photocatalysts, attention can be paid to investigating the synergistic effects of the multimethod combination on the activity for photocatalysis in future researches. Further insight should be taken into detailed thermodynamics and kinetics of electron-hole transfer and reactions at the interface, which is crucial to design more efficient solar-to-fuel conversion systems. Although there are still many challenges ahead, it is believed that in the future, photocatalytic technology will penetrate every aspect of our daily life and provide more convenience for human beings with the unremitting efforts of researchers. Acknowledgments This work is partially supported by Prof. Lingbo Luise Liu. The authors also gratefully acknowledge the helpful comments and suggestions of the reviewers, which have improved the presentation.

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Authors’ Contribution Qiuyang Huang wrote the paper. Dan Kong polished the paper and discussed with Yongdan Li the content of this work and finalized the paper. All authors have approved the final version of the manuscript.

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178. Li Y, Chen H (2016) Facile fire treatment of nanostructured hematite with enhanced photoelectrochemical water splitting performance. J Mater Chem A 4(39):14974–14977 179. Klahr B, Gimenez S, Fabregat-Santiago F, Bisquert J, Hamann TW (2012) Photoelectrochemical and impedance spectroscopic investigation of water oxidation with “Co-Pi”-coated hematite electrodes. J Am Chem Soc 134(40):16693–16700 180. Carroll GM, Gamelin DR (2016) Kinetic analysis of photoelectrochemical water oxidation by mesostructured Co-Pi/alpha-Fe2O3 photoanodes. J Mater Chem A 4(8):2986–2994 181. Carroll GM, Zhong DK, Gamelin DR (2015) Mechanistic insights into solar water oxidation by cobalt-phosphate-modified α-Fe2O3 photoanodes. Energy Environ Sci 8(2):577–584 182. Hisatomi T, Le Formal F, Cornuz M, Brillet J, Tetreault N, Sivula K, Graetzel M (2011) Cathodic shift in onset potential of solar oxygen evolution on hematite by 13-group oxide overlayers. Energy Environ Sci 4(7):2512–2515 183. Le Formal F, Tetreault N, Cornuz M, Moehl T, Graetzel M, Sivula K (2011) Passivating surface states on water splitting hematite photoanodes with alumina overlayers. Chem Sci 2 (4):737–743 184. Zhan F, Liu W, Li W, Li J, Yang Y, Li Y, Chen Q (2016) Efficient solar water oxidation by WO 3 plate arrays film decorated with CoO x electrocatalyst. Int J Hydrog Energy 41 (28):11925–11932 185. Huang J, Zhang Y, Ding Y (2017) Rationally designed/constructed CoOx/WO3 anode for efficient photoelectrochemical water oxidation. ACS Catal 7(3):1841–1845 186. Higashi M, Domen K, Abe R (2012) Highly stable water splitting on oxynitride TaON photoanode system under visible light irradiation. J Am Chem Soc 134(16):6968–6971 187. Sun K, McDowell MT, Nielander AC, Hu S, Shaner MR, Yang F, Brunschwig BS, Lewis NS (2015) Stable solar-driven water oxidation to O2(g) by Ni-oxide-coated silicon Photoanodes. J Phys Chem Lett 6(4):592–598 188. Du C, Yang X, Mayer MT, Hoyt H, Xie J, McMahon G, Bischoping G, Wang D (2013) Hematite-based water splitting with low turn-on voltages. Angew Chem Int Ed 52 (48):12692–12695 189. Malara F, Minguzzi A, Marelli M, Morandi S, Pesaro R, Dal Santo V, Naldoni A (2015) Alpha-Fe2O3/NiOOH: An effective Heterostructure for Photoelectrochemical water oxidation. ACS Catal 5(9):5292–5300 190. Wang G, Ling Y, Lu X, Zhai T, Qian F, Tong Y, Li Y (2013) A mechanistic study into the catalytic effect of Ni(OH)2 on hematite for photoelectrochemical water oxidation. Nanoscale 5 (10):4129–4133 191. Xiao J, Huang H, Huang Q, Li X, Hou X, Zhao L, Ma R, Chen H, Li Y (2017) Remarkable improvement of the turn-on characteristics of a Fe2O3 photoanode for photoelectrochemical water splitting with coating a FeCoW oxy-hydroxide gel. Appl Catal B Environ 212:89–96

Ms. Qiuyang Huang received her bachelor’s degree from Chongqing University in 2016. Currently, she is a Ph.D. candidate under the supervision of Prof. Yongdan Li at the School of Chemical Engineering and Technology at Tianjin University, China. Her Ph.D. research is focused on the photoelectrochemical water splitting into hydrogen.

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Professor Dan Kong obtained her Ph.D. degree in Chemical Engineering in 2019 from the University College London, with Professor Junwang Tang. She was a Postdoctoral Researcher at Aalto University in Finland and researched the field of photocatalysis and photoelectrocatalysis. Currently, she is an Associate Professor at the China University of Mining and Technology in China. Dan’s research focuses on the application of low-cost and tunable polymeric semiconductors for solar energy conversion such as overall water splitting and CO2 conversion. Dan’s interests also include efficient microwave-assisted synthetic techniques, photoelectrocatalysis, and mechanism investigation. Professor Yongdan Li received his Ph.D. degree in 1989 from the Industrial Catalysis Program of Tianjin University, China, with Professor Liu Chang. He spent one year at the University of Twente as a visiting researcher and one and half a year in DCPRENSIC in INPL in Nancy as a post-doc. After that, he got an associate professorship at Tianjin University and after another year, he was promoted to a full professor there. He served as the Chair of the Industrial Catalysis Program and the Chairman of the Department of Catalysis Science and Technology in Tianjin until 2017. In June 2017, he was appointed as the Tenured Full Chair Professor of Industrial Chemistry at the School of Chemical Engineering, Department of Chemical & Metallurgical Engineering of Aalto University, Finland.

Chapter 3

Functional Nucleic Acid Hybrid Materials for Photovoltaic Cells: Design, Fabrication, and Performance Dan Bai, Huhu Feng, Xingchen Yu, Chenxin Ran, and Wei Huang

3.1

Introduction to Nucleic Acid Hybrid Materials for Clean Energy

Coping with the increasingly urgent and complex issues of energy supply for industrial manufacture, transportation, communication, finance, security and biosecurity, information storage and processing, and so forth, how to provide energy sustainably has become a serious problem. For those of us who live in the present and being weaned on petroleum fuels with limited resources, it is necessary to develop materials and technologies for renewable and environmentally friendly energy. Solar power is one of the most popular choices among the sources of sustainable energy. If we can efficiently harvest and utilize the sunlight that falls on the Earth, humanity will have almost infinite energy. However, there are still problems to tackle in the design of materials and devices in solar cells. First generation of silicon-based dye-sensitized solar cells receives sunlight by

D. Bai (*) · C. Ran (*) · W. Huang Frontiers Science Center for Flexible Electronics (FSCFE), Institute of Flexible Electronics (IFE), MIIT Key Laboratory of Flexible Electronics (KLoFE), Northwestern Polytechnical University, Xi’an, Shaanxi, China Xi’an Institute of Biomedical Materials and Engineering (IBME), Xi’an Key Laboratory of Biomedical Materials and Engineering (KLBME), Northwestern Polytechnical University (NPU), Xi’an, Shaanxi, China Research and Development Institute of Northwestern Polytechnical University in Shenzhen, Xi’an, Shaanxi, China e-mail: [email protected]; [email protected] H. Feng · X. Yu Frontiers Science Center for Flexible Electronics (FSCFE), Institute of Flexible Electronics (IFE), MIIT Key Laboratory of Flexible Electronics (KLoFE), Northwestern Polytechnical University, Xi’an, Shaanxi, China © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_3

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chromophore pigments to generate electricity. Silicon solar cells are good batteries with high efficiency of power generation and environmental stability. One of the practical problems with the silicon device is its net weight. An average 4 kW power generation solar cell system builds on the rooftop of a house, for instance, a weight of about 250 kg–400 kg [1, 2]. Besides labor costs, the architectural scaffold for solar cell installation must be designed strong enough to withstand the weight of the solar cell. In the age information computing has posted Moore’s law, the current technologies based on silicon and 2D carbon materials tend to mature and saturate, and the performances of current photoelectronic devices are gradually facing their physical limits. Considering the manufacturing cost alone, even the cost of making silicon solar cells themself is becoming cheaper; without lightweight materials the installation costs would still be above the social welfare coverage. Various approaches have been put into practice for making solar cells more compatible and easier to blend into daily life. With the advancement of biotechnology and bioengineering, biomacromolecules have come to attract attention as next-generation materials for photoelectronic devices. Biomolecule-based molecular devices can be miniaturized and integrated with less energy consumption technologies, therefore attracting much attention as a next-generation intelligent material for electronic energy devices, as a choice of carbon and silicon [3, 4]. Biomaterials are abundant in nature and their production is environmentally friendly. Synthetic biomolecular systems with rational design create intelligent systems with advanced functionalities. In recent years, nucleic acids from a natural source and chemical synthesis are taken up as a biomolecular material that contributes to the development of the sustainable energy field. Nucleic acids have been long utilized as potent material in the field of bionanotechnology, for their synthetic feasibility, versatile functionality, high performance, and environmental friendliness. The unique photoelectronic properties of nucleic acids have attracted much research attention in the application areas of electronic energy transfer, information storage, and logic circuit systems beyond the scope of their biological functionalities [5–7].

3.2

Optoelectronic Properties of Nucleic Acid

3.2.1

Structure and Conformation

3.2.1.1

Molecular Basis and Double-Helix Structure

Nucleic acids are one of nature’s most abundant and renewable biomaterials. Natural DNA/RNA and artificial XNA are substances responsible for coding and transportation of information. In recent years, quantification has come to be carried out in various fields with advances in technology and increasing needs in the industry. Photovoltaic cells using nucleic acid hybrid materials and the manufacture techniques were developed [8, 9]. The development of biomaterial-based devices and

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sensors using nucleic acids can achieve high performance without compromising environmental friendliness. In this regard, nucleic acid hybrid materials are considered as a competent scaffold for hosting functional optoelectronic molecules and materials in the designated platform [10, 11]. Nucleic acid materials are easy to obtain from biomass with low cost and easily decomposed without pollution. Typical nucleic acids form in B/D/Z helical structures following the condition of surrounding microenvironments (Fig. 3.1). The exhibited peculiar double-helix structure consisted of four nucleobases in pairs: adenine (A) with thymine (T) and guanine (G) with cytosine (C). The backbones joined internally by the ester bonds were made of sugar and phosphate groups. The diameter of each helix is about 2 nm and the distance between the two base pairs is 3.4 nm [11, 12]. The base pairs are linked together by strong hydrogen bonds. The macromolecule of nucleic acids exhibits a net negative charge because of the outside phosphate groups, which may compensate by sodium ions and nonlocalized counterions that move freely along the macromolecular chain surface.

Fig. 3.1 The B/D/Z type of nucleic acid helical structures with their parameters. Reprint permission acquired from Chem. Soc. Rev. (2016), the royal society of chemistry publishing [11]

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Higher-Order Structures

The functionalization of biopolymers involves the control of their higher-order structures. Furthermore, the transitions of the higher-order structure extend beyond single strand, and double-helix formation plays important role in advanced biomolecular functions such as signaling and material transport. Nucleic acids are considered a cationic biopolymer because of their phosphate backbones. When nucleic acids form a higher-order structure, the electrostatic interactions are strongly affected [13, 14]. On the other hand, Nucleic acids can also be used to assemble a stable complex with the copolymer electrolyte. Nucleic acid strands can favorably define the orientation of photovoltaic donor-acceptor pairs in spatial organization. The preferred inhibition of exciton recombination could be promoted with nucleic acid incorporated with ionic surfactants such as cetyltrimethylammonium chloride (CTMA) and sensitizer molecules (Figs. 3.1 and 3.2), in the various sub-nanostructure such as films, fibers, and meshes [15–18]. On the other hand, nucleic acids could be functionalized by the intercalation interaction between guest molecules with the minor and/or major grooves in weak van der Waals forces. The formation of covalent bonding between guest molecules with nucleic acids is also possible via a randomized doping process. Of course, chemical synthesis and modification of native nucleic acids could also create a copolymer-like complex with guest molecules aligned in a designated position and orientation [19–21]. Modulating the folding and assembling process is the key to achieve designated functionalities when using biopolymer materials such as nucleic acids. Besides the innate biological activities, the opposite nucleic acids with artificial materials can manipulate the higher-order structures of biopolymers such as nucleic acids. Generally, the electrolyte complexes are formed from nucleic acids and polycation compactly in aqueous solvents. Copolymer with moieties of rich hydrophilic side chain has been found to form a completely soluble complex with nucleic acids (Fig. 3.3). In such cationic codecity, electrostatic repulsion between nucleic acids and their surrounding sphere was alleviated and stabilized the double-helix and triple-stranded nucleic acids. The various approaches in organic synthetic chemistry, photochemistry, and molecular biology have been used to stabilize the higher-order structure of nucleic acid hybrids, such as the synthesis of artificial xeno nucleic acids (XNA) or the use of nucleic acid-binding ligands. In the formed electrolyte complex of nucleic acids, structural transition catalysts activate the release and revalidation of nucleic acid base pairs. Higher-order structure assembly and functional operation of nucleic acids can be adapted in ionic peptide materials as well. The macromolecules of nucleic acid strands could act as electronic donoracceptor pairs, which are electrostatically adsorbed around calcium-titanium crystals rather than forming 2D/3D heterolytic structures. Surface-immobilized nucleic acid molecules could eliminate the ion vacancy on the crystal boundary and regulate the electronic structure of the main 3D crystal, thus modulating the carrier composite channel [22, 23]. For instance, molecular docking calculations of the interactions

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Fig. 3.2 The nucleic acid helical structure incorporated with conjugated cationic surfactant by ion-exchange reaction forms copolymers

between nucleic acid and chromophore compounds such as metal complexes have been proven to adopt compact conformation to bind the major groove of the helix structure [24, 25]. The chelating ligands in metal complexes were positioned at the bottom of the major groove, surrounded by the nucleotides, forming stable hydrophobic bindings. Detailed analysis showed that hydrogen bond lengths between the nucleotide with the metal complexes of the helical nucleic acid were 2.5–2.8 Å [26] (Fig. 3.4). This case proved that a wide range of polyaromatic donor-acceptor molecules including metal complexes could interact with nucleic acids. The mechanism proposed was that the helical structure of nucleic acid could affect the exciton

Fig. 3.3 Nucleic acid hybrid materials in the arrangement of spatial organization incorporated with photovoltaic donor-accepter molecules

Fig. 3.4 Detailed and total view of iridium (III) metal complex interacting with nucleic acid double-helix strands. Reprint permission was acquired from Elsevier [26]

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coupling between donor-acceptor pairs, reducing the process of charge recombination, because the helical arrangement of nucleic acids was nearby of the exciton diffusion length [26–29]. To alternate, the structure of polyion backbone and hydrophilic side chains of nucleic acids and ionic copolymers such as Poly(L-lysine)-g-dextran (PLL-g-dex) have been used to form a soluble inter-polyelectrolyte complex with nucleic acid strands to enhance their photoelectronic performance [30].

3.2.2

Optoelectronic Properties of Nucleic Acid

3.2.2.1

Conductivity, Magnetism, and Spintronics

Nucleic acid was known to be a potent material for photoelectronic devices. The high-speed charge transfers caused by charge hopping between nucleic acid base pairs are suitable for making molecular wires and semiconductors. The stacking structure and the mutual action between the base pairs in nucleic acids form the overlapping p-orbital regions. As required in photoelectronic devices, the insulator/ conductor layers were often made from cyclic compounds with double or triple bonds of p-orbital overlapping with π-conjugated areas. Therefore, electrons can be moved via the π cloud and create a tunable forbidden gap in insulators for optimized semiconductor properties [31, 32]. Modulations on the molecular array orientations in nucleic acids are feasible because of their native double-helix structure and the spontaneous higher-order spatial conformations (Fig. 3.5). As many desired functionalities in optoelectronic devices rely on carrier transportation, carrier mobility (carrier movement speed divided by the electric field strength) is a critical factor. Using nucleic acid materials with self-assembled orientation due to their chemical structural characteristics, it is easy to modulate the electronic carrier mobility of current flow and response speed for practical implementation [33, 34]. The dynamics of nucleic acids in vivo could be described by either soliton or localized modulated waves termed breather wave (isolated wave). These terahertz waves are generated due to nucleic acid folding that resonates with biomacromolecules and organelle membrane activities. In terahertz wave 1 THz (=1012 Hz) corresponds to the wavelength of 300μm. The frequency band of terahertz waves has a characteristic absorption frequency in the molecular interactions such as hydrogen bonds and van der Waals force [35–37]. Biomacromolecule systems such as nucleic acids with higher-order structures dissolved in an aqueous solution are expected to be able to obtain such information about their molecular networks. The quantities of the plasticizer incorporated in nucleic acids were also considered as an important factor in the ionic conductivity values in the strands. In an aqueous solution, the polyelectrolyte properties of nucleic acids are more pronounced because of the presence of the diffusion of ions and their counterions [38–41].

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Fig. 3.5 Synthetic nucleic acid as hole transport materials incorporated in electronic circuits. (Reprint permission acquired from Journal of the American Chemistry Society) [34]

Nucleic acid was firstly considered as materials without magnetism until the report on λ bacteriophage genome DNAs has shown paramagnetism at low temperature [39]. In the lyophilized dry sample of nucleic acids, magnetism property was not found. Magnetism in nucleic acids was considered to originate via the base pairs interacting with their surrounding water molecules, where the free π-electron with antiparallel or parallel spins becomes the source of spin magnetism. With both magnetic and semiconductor properties, nucleic acids are expected to be potent materials for making semiconductor devices [41–43]. Utilizing both the electronic charge and spin of nucleic acids, the technologies of electronics, semiconductor, and magnetics could be integrated into the emerging new field of spintronics and create devices with new features and improved performance. Current spintronic devices use mainly inorganic materials based on metals and silicon. By using nucleic acid materials, which are based on light elements of hydrogen and carbon, it is possible to develop a new device with little loss of deflection of spin [18, 44, 45].

3.2.2.2

Optical Properties

Production and optimization of next-generation display materials rely on a wide range of research fields, mainly from the basic optical properties of liquid crystal applications. To put a dielectric liquid crystal into practical use, it is necessary to eliminate defects in molecular orientation structures to stabilize their electric field response [46–48]. Using materials with high-speed responsiveness enables beautiful video display and significantly lowers energy consumption [49]. Apart from the ferroelectric liquid crystals and liquid crystal polymer network systems, research to explore the relationship between the liquid crystal phase and materials from living organisms was also popular in recent years [50]. The study of the mechanism between the fluidity of liquids and the anisotropic formation of crystals also plays an important role in understanding the formation of life forms. New materials with advanced techniques are required, to optimize the properties desired for practical

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display applications (driving voltage, response time, contrast ratio, etc.) [46]. Nucleic acids were recently been carried under investigation as potent display material for their uniform molecular orientation and unique optical properties. There are various substances such as nucleic acids, peptides, proteins, virus particles, and other amphipathic moieties such as surfactants able to take the liquid crystal state. Liquid crystal was usually termed as a substance form having an orientation order found in a group of organic molecular components. Although inorganic component-based liquid crystals were hardly known, layered inorganic crystals such as graphite could also take information in the liquid crystal state. Therefore organic-inorganic hybrid material-based liquid crystals have also been investigated in recent years [51, 52]. As natural material from living organisms, the nucleic acid has high self-assembly, which makes it easier to modulate in nanoscale structures. Bottom-up nanotechnology applied to nucleic acid enables many functionalized materials with low energy costs. By the manner of molecular array alignment, conventional liquid crystal was classified into three phases, as their opacities increase in the order of nematic phase < smectite phase < cholesteric phase. In the nematic phase, the long axes of the molecules are parallel arranged in one direction within a column, and their ends were staggered at random intervals. In the smectite phase, rodlike molecules are arranged in planar layers, without regularity in the sequence in the layer; the molecules were also distributed in one direction (Fig. 3.6) [53]. In the cholesteric phase, the arrangement was complex, in which the molecules were aligned in layers; concerning the layer above and below, each layer rotates to a certain degree to give a spiral structure. The structure of the cholesteric phase is very much close to the natural structure of the nucleic acid helix, which makes the unique properties. Recently, blue phase liquid crystals with optically isotropic birefringent properties attracted much research attention. The structures of blue phase liquid crystals were characterized in form of double-twisted cylinders, which are arranged periodically in space. Thermotropic liquid crystal is represented by the exclusion volume effect due to anisotropic molecular shape; the riotropic liquid crystal appears due to microphase separation due to differences in intermolecular interactions. The conditions of the liquid crystal phase were classified into two types: thermotropic (temperature transition) and lyotropic (concentration transition). Thermotropic liquid crystal is intended to change the phase only by heat and pressure, such as thermoplastics. The lyotropic liquid crystal consists of a multicomponent phase change depending on the temperature and component composition, which is common in biological tissues such as cell membranes [54]. There are many exemplary liquid crystal states in living organisms. For instance, the lamellar structure is formed from components such as sphingolipids between cells of the stratum corneum of the living body. In addition to the function of preventing moisture evaporation from the body, the cholesteric phase of sphingolipids crucially modulates lateral microdomain organization for functionalities such as color expression, making the body optically active. On the other hand, the nerve pulses in the brain were modulated by the insulation effect of smectic phase phospholipids of the parent medium molecule is made, various functions have

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Fig. 3.6 Arrangement of a nucleic acid molecular array for display devices. Reprinted permission was acquired from Elsevier [53]

been realized. When an aqueous solution of nucleic acids was made into a certain concentration, it could effectuate optical property peculiar to the cholesterol phase [55, 56]. In the research of liquid crystal conditions and thermal stability, there are still unsolved questions. Through the utilization of nucleic acid as functional hybrid materials, methods and mechanisms could be established for liquid crystal material creation, phase structure, and property evaluations. Also, the self-assembly structure of the liquid crystal phases is very similar to the native structure of nucleic acids; the structure-activity relationship study is also interesting [57]. From the viewpoint of using natural biological self-assembly structure with the mechanisms of liquid crystal states, materials found in living organisms and functionalized synthetic materials can be used for each other, reciprocally. Nucleic acids are an ideal material to form a liquid crystal state because of their characteristic helix structure. When the base-pair molecules in nucleic acids deviate from the natural spherical symmetry, nucleic acid strands take a rodlike-shaped structure that satisfies the liquid crystal conditions. The feature helix structure was

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formed by mainly the hydrophobic interactions in vivo, during which the spheres become amphiphilic [58]. Current liquid crystalline materials with long rodlike molecular arrays assemble randomly and often suffer from the disadvantageous phase separation beyond a certain concentration. On the contrary, the helical structures of nucleic acids are stable, not becoming anisotropic even in dilute solution. In the case of dsDNA, the double-helix structure is stabilized by hydrogen bonds of complementary base pairs facing each other. Using nucleic acids, it is possible to adapt their native molecular orientation to create advanced functionalities and go beyond the functions of their sugar-phosphate backbone and base-pair molecules [59].

3.3 3.3.1

Design Principles and Fabrication Techniques Design Principles

Conventional photoelectronic energy conversion devices are fabricated by depositing organic photoactive materials on inorganic substrates, and these devices require many energy costs in their synthesis, functionalization, and production. Notably, biomacromolecules, such as nucleic acids with built-in structure and functionalities in themselves, could achieve comparable or even superior optoelectronic performance with the advantages of environmentally benign and low-cost production. Therefore, the design and fabrication of the nucleic acid-based photovoltaic device are highly expected in the field. To investigate the key principles in designing nucleic acid material-based photoelectronic devices, the metrics of sensitivity, stability, reproducibility, and lifetime should be taken into consideration before putting them into practical use. Through defect heterogeneity elimination and adjustment of the energy bandwidth, the optimized nucleic acid hybrid materials can achieve higher radiation efficiency and much akin to n-type properties (Fig. 3.7) [60]. In this case, the trade-off between short-circuit current density (JSC) and fill factor (FF) does not necessarily be compromised, as long as the prepared composition of hybrid material film achieves the appropriate circuit voltage (VOC) and minimum deficit voltage difference with steady power outputs [61]. During the working of a photovoltaic cell, photonic energy is absorbed by the designated donor molecule to promote the excitation of the ground-state electrons from the highest occupied molecular orbital (HOMO) to the lowest unoccupied molecular orbital (LUMO) across the bandgap of the donor molecule. Subsequently, photoexcited electron-hole pairs are generated. In the process of energy redistribution, the excited electrons can either recombine with the hole or diffuse to the donoracceptor interface, where the separation of Coulomb-bounded electron-hole pairs into free electrons and holes is possible when the LUMO of the acceptor is below the LUMO of the donor [62, 63]. The free electrons then may transfer across the energy barrier between donor-acceptor regions toward the cathode region, while the free

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Fig. 3.7 Mechanistic illustration of nucleic acid hybrid material used in perovskite. Reprint permission acquired from American Chemical Society [60]

holes transport toward the anode region. During this process, the minimized energy level mismatch is required for the free charges transferring efficiently across the Schottky barriers at the interfaces. Most importantly, controlling the phase separation degree of donor-acceptor inter-network within the diffusion distance of exciton has been recognized as the most essential strategy to promote efficient charge transfer. Therefore, it is crucial to precisely modulate the donor-acceptor molecules using organized platforms, such as nucleic acids [64, 65].

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In a typical nucleic acid-based photovoltaic cell, the nucleic acid layer often lies between the anode layer and the cathode layer either in direct contact or indirect contact (Fig. 3.7). Electron/hole transport or hole/electron blocking layers may also include in the device structure as buffer layers to ensure efficient charge collection at the interfaces. The main function of the nucleic acid layer is to orchestrate a plurality of donor-acceptor molecular pairs within the confined space in a certain orientation [66]. The high load of oriented donor-acceptor pairs would make high electron mobility in a preferential direction, allowing the photovoltaic cells to convert photon energy into electric energy with high efficiency. Also, the innate features of nucleic acids, such as self-assembly and self-replication in well-defined nanoscale geometry and macroscale morphologies, make them suitable for efficient optoelectronic interactions with donor-acceptor pairs through various mechanisms. The unique abilities of nucleic acid hybrid materials, therefore, have been considered as preferred platforms for their interactions with photovoltaic dye molecules, chromophores, and donor-acceptor pairs, which promote the directed orientation and prevent the aggregation of these molecules simultaneously [67, 68].

3.3.2

Fabrication Techniques

Perovskite solar cells are based on thin-film photovoltaic technology, which requires the fabrication of high-quality perovskite thin film, where the photovoltaic performance of perovskite solar cells is largely correlated to the quality of perovskite thin film. For example, the heterogeneity polycrystalline of the perovskite film may severely affect the transportation of photogenic carriers [69, 70]. Moreover, the lifetime of photo-generated excitons, the mobility of charge carriers, and the carrier diffusion lengths also determine the power conversion efficiency of perovskite solar cells. Notably, film fabrication techniques play a key role in obtaining high-quality perovskite thin film. The main techniques used for thin-film formation include spin coating, droplet coating, inkjet coating, roll-to-roll printing, and vacuum coating (Fig. 3.8). All of these techniques can be used to prepare the semiconductor film for thin film-based optoelectronic devices, and they possess their features and merits [71]. The spin coating technique is suitable for the preparation of polymer film and solution gel film because of its simple operability, simple equipment, and easy preparation in the ambient atmospheric environment [72]. Also, it only requires a relatively simple material with a certain viscosity and solubility. Moreover, the film can be simply prepared by putting the substrate on the spin coater, dropping the viscous solution in the center of the substrate, and starting the spin coater. The thickness of the prepared film can be easily adjusted by optimizing the speed, time, and acceleration of the spin coater, concentration of the solution, volume of the spincoated solution, solvent volatilization speed, and adhesion of the solution to the substrate [73].

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Fig. 3.8 Manufacturing technologies for solar cell devices using hybrid materials

The drop coating technique is processed by dropping a certain amount of liquid on the surface of the substrate, where the liquid diffuses to the entire surface. Then, the liquid on the substrate naturally volatilizes in the solvent atmosphere or atmospheric atmosphere, finally forming a thin film, as shown in Fig. 3.8 The method employs the principle of molecular self-assembly, in which internal molecules rearrange themselves during volatilization to form a thin film. The key point of controlling the characteristics of the film is the selection of solvent [74, 75]. Solvents with different boiling points have different volatilization rates, which will lead to the change of the morphology and the crystallinity of the prepared thin film. The advantage of this preparation method is obvious that it is easy to prepare thin films in an atmospheric environment. However, it also has disadvantages, such as poor homogeneity and high surface roughness of the deposited thin film [76]. In addition to the solution-based techniques, the physical or chemical deposition of thin film in a vacuum environment is called vacuum coating [77]. This technique is suitable to deposit semiconductor materials that show poor solubility in organic solvents. As a result, a vacuum evaporation method is an alternative option [78]. In the vacuum evaporation process, the solid semiconductor is firstly heated above its sublimation temperature to promote its transition into gas in a vacuum environment. The semiconductor molecules will uniformly despite on the substrate placed above the evaporation source. By controlling the heating temperature and deposition velocity, the order degree, thickness, and carrier transport capacity of the deposited

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film can be well-controlled. The advantage of this technique is that it can prepare a uniform thin film with less impurity and high purity, but the disadvantage is that the operational procedure is complicated [79–81]. In recent years, printing and lithography, such as inkjet printing and roll-to-roll (R2R) printing, have become increasingly significant to fabricate photovoltaic and electronic devices. Inkjet coating technique is processed by ejecting a fixed amount of liquid material on the substrate from an ink storage chamber through a nozzle in the form of droplets [82]. The ejected region of the substrate forms a pattern under the combined action of gravity and air resistance, and the solidification of droplets can be realized using solvent evaporation. In this process, the viscosity and surface tension of the solution are the main factors affecting the quality of the thin film printed by inkjet printing. The advantages of inkjet printing are simple operation, high feasibility, low manufacturing cost, high operational flexibility, and mature technology, which can effectively control the deposition of materials and the formation of thin films [83]. However, its disadvantage is that the development of inkjet-based manufacturing technology is time-consuming. Compared to inkjet printing, R2R printing can also be used to efficiently produce low-cost thin film. Patterned ink on the substrate gradually forms a thin film with pressure as the roll begins to turn. This method is suitable for mass production and can reduce labor costs [84, 85]. For solution-based techniques, NAs could play a crucial role because of their tunable solubility in various solvents. As water-soluble and hydrophilic material, NAs could dissolve in water-immiscible solvents, such as acetonitrile and alcohol. Upon modification with polyethylene glycol or hydrogels, NAs have also been reported to dissolve in most organic solvents. Therefore, the introduction of NAs in solution-processed thin-film techniques could extend the solvent system from organic solvent to water, lowering the cost with environmentally friendly green chemistry. Moreover, for the packaging of advanced semiconductors, nucleic acid materials are also compatible with inkjet printing electronic technologies, which allows the device to process at ambient temperature and atmospheric pressure. As a platform in the film-forming reaction, the introduction of nucleic acid has been proven to influence the deposition process by tuning the growth rate of the thin film. As a result, the size and uniformity of the film can be fundamentally controlled. Furthermore, the presence of nucleic acid could also reduce the complex permittivity and dielectric loss of the deposited film, hence improving the electrical performance and power conversion efficiency of the corresponding electronic devices. In the manufacturing process, it is also possible to develop nucleic acid hybrid materials for optoelectronic devices without the need for a microfabrication process or extreme ultraviolet (EUV) lithography machines. For example, laser oscillation requires a resonator molecular system with a precise periodic structure to reflect circularly polarized light of a particular wavelength, and the native structure of nucleic acid themselves could serve as a distributed feedback-type laser oscillator. Also, nucleic acid-based liquid crystals could be used for the wavelength of the light that is confined spontaneously via the helical structures [86, 87]. Although lowering the energy threshold for oscillation of the liquid crystal laser is still a challenge, nucleic

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acids are considered a promising material for semiconductor laser devices. In practical use, nucleic acid-based cholesteric liquid crystal devices could generate electric current driven by the redistribution of the liquid crystal generated by photoelectronic excitation [88]. Also, nucleic acid materials are compatible with flexible electronic devices. Overall, the merits of NAs and their hybrid materials discussed above enable their great potential in the application of next-generation optoelectronic devices.

3.4

Enhancement of Performance

In general, organic semiconductors have difficulty obtaining materials with high purity due to the difficulty of the synthetic procedure. The performance of photoelectronic devices is easily affected by impurities [89–91]. Oxygen in the atmosphere may also deteriorate the electrical properties. The chemical and biological synthesis of nucleic acids is relatively easy and low cost, with high purity. Furthermore, through nucleic acid sequence programming, it is possible to develop nucleic acids in a three-dimensional higher-order structure accompanying new functionalities. In accord with the conductivity, magnetism, spintronics, and optical properties described in the previous section, using nucleic acid hybrid materials, the current limitation problems of photoelectronic devices and the efficiency of energy conversion could be tackled [92]. The integration of nucleic acid hybrid materials with advances in photoelectronic devices is challenging. Recent developments of perovskite solar cells (PSCs) expanded the fabrication and application possibilities of devices from the siliconbased solar cells. Perovskite solar cells can be made by applying the materials and technologies of flexible electronics. The thickness of the perovskite itself is about 500 nm (nanometer) and about 1/100th of a single hair. Since the weight of a solar cell mainly depends on the device substrate where the materials are applied to, perovskite solar cells have the potential to be produced on top of platforms that are lightweight with flexible texture, such as wrapping film or paper. Future solar cells could be made in the shape of a vinyl sheet, put on the roof like a curtain. Also, it is possible to use transparent conductive electrodes in the making of translucent solar cells [93]. The thickness of the perovskite can be tuned so that the window glass of a house or building can be made into a solar cell. The emerging print full electronic technologies allow simple and fast manufacture of solar cells anytime, anywhere. Therefore, installation costs of solar cells could be considerably reduced, easy to install to a home in populated areas. The term perovskite was originated from the mineral stone vastly distributed in the Ural Mountains of Russia. Generally, their chemical formula is represented by the composition formula of ABX3, where cation is placed in A and B and anions are placed in X (Fig. 3.9). For instance, the natural combination of perovskite is A: Ca; B: Ti; X:O. Most of the current organic-metal halide-based perovskite materials were modified with other components, which are actively used for applied research

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Fig. 3.9 (Upper). Crystal structure of ABX3-type hybrid halide perovskite. The organic and inorganic monovalent cation is A (black), divalent metal cation B (red), and halogen anion X (blue); (lower) flexible perovskite solar cells. Reprint permission acquired from the American Chemistry Society [60]

on solar cells, consisting mainly of crystals from the combination of A: monovalent organic cations of CH3NH31+(MA1+), NH2CH ¼ NH21+(FA1+), Cs+, etc.; B: divalent metal cation of Pb2+, Sn2+, etc.; and X: Cl/Br/I and other halogen anions. By far perovskite is the most popular substrate for solar cell devices due to its tuneable photoelectric properties, including suitable bandgaps, light absorption coefficient, high photoluminescence quantum yield, narrowband emission, relatively small absorption length, long diffusion length, and low-cost solution processability [94, 95]. Organic metal halide perovskite solar cells were born in 2012 with the development of semiconductor materials. The conversion efficiency of the initial light energy as a photosensitized solar cell using this material as a photosensitized object was only 3.2%, which was not comparable to 26% of silicon-based systems. After a series of improvements, the PCE of perovskite solar cells (PSCs) has now exceeded 25%, and it is thought that it has the potential to exceed the records kept by silicon solar cells. Compared with silicon substrates, perovskite solar cells have received extensive attention and in-depth developments from researchers around the world in recent years [96]. When perovskite solar cells first appeared, light to energy conversion efficiency (power conversion efficiency, PCE) was low and durability was also a problem. For instance, in earlier research, the PCE was modulated by adjusting the perovskite crystal structure, blending ratio of methylammonium ion (MA+), inorganic lead ion (Pb2+), and halogen anion iodine ion (I). In the past decade, the power conversion efficiency of Pb-based perovskite solar cells has increased from 3.80% to 25.2%, and the photoelectric detection efficiency of Pb-based photodiode array detectors has increased to 1015 jones. Also, the ultrahigh

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photoluminescent quantum yield and high Seebeck coefficient value of approximately 1000μV/K were obtained from the Pb-based monocrystalline and Pb-based perovskite quantum dots, respectively. Despite such attractive properties in the photoelectric and thermal electrical application, the toxicity of Pb2+ in their preparation, utilization, and disposal causes many further environmental problems, which is one of the bottlenecks in its practical application. Because the Ca-Ti film can be prepared by the liquid phase method at low temperature (100  C), it is compatible with the processing of flexible electronics, especially suitable for preparation into flexible battery devices [97]. In particular, flexible devices can be produced in large areas at low cost through the roll-to-roll process. Therefore, the development of high-efficiency flexible PSCs has gradually become one of the main topics in the field of sustainable energy research. The characteristic foldable, lightweight, and robust flexible devices are very suitable for integrated design with architecture construction and logistics. The substrate of PSCs generally consists of two types, namely, flexible metal foil and flexible polymer substrate. Polyethylene terephthalate (PETE), polyethylene naphthalate (PEN), and other plastic substrates have the properties of low cost, light quality, good light transmission, and suitable mechanical performance. Yet their low melting temperature limits the carrier transport layer preparation process. The polymer substrate of flexible perovskite battery research hence has been focused on the techniques to prepare high-efficiency carrier transport layer film under low-temperature conditions. In contrast, the foil-type material has excellent ductility, conductivity, and high-temperature resistance, so the vast majority of the hightemperature preparation carrier transport layer process can be directly applied to the metal foil substrate. However, the opacity of the foil substrate requires the device’s back electrode to be transparent, which puts forward new requirements for the preparation process of the transparent electrode. Most flexible perovskite batteries now were metal foil substrate-based, which requires the development of highly efficient transparent electrodes. At the same time, the commonly used transparent conductive material indium tin oxide (In2O3-SnO2, ITO) film is very easy to crack under bending conditions, resulting in a sharp increase in resistance, hence causing severe degradation of the entire device performance Therefore, to improve the flexibility of the device, the development of the more mechanical performance of the transparent conductive electrode is an important direction for breakthrough material and technology development [98–100]. One of the unsolved issues is the material used in perovskite solar cells. Currently in piezoelectric ceramic transducer (PbZrO3, PbTiO3), lead is often used. Since Pb is toxic, it becomes a big problem when the PSCs decomposed some lead melts into the environment. However, the amount of lead per unit area used in PSCs is about the same as the amount contained in 1-cm-thick soil in nature. To be specific, the amount of lead contained in a 1 m2 area of PSCs is approximately 0.4 g. Conventional perovskite solar cells have made it easy to achieve more than 20% energy conversion efficiency by containing cesium and rubidium, but as rare earth elements, they were expensive. On the other hand, by finding a method of utilizing naturally abundant raw materials, PSCs could be made at low cost with higher performance. Also, there

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is an advantage that the manufacture of bioinspired raw material is easy. Besides the flexible substrate and electrodes, the enhancement of efficiency relies heavily on the materials of the carrier transport layer. Enhancement of power conversion efficiency (PCE) has been a worldwide race for the performance of solar cells, emerging the highest record of >40% until the year 2020. Various technologies have stimulated the rising PCE in organic cells and dye-sensitized cells record of >14% up to date. Our goal is to produce perovskite solar cells with an energy conversion efficiency of more than 25%, which would contribute to the developing countries suffering from poverty by bringing solar cells to everyday life [1, 2]. The economic pressure would be alleviated once the electric power charge is reduced for industry and home. Apart from the problem of environmental stability, the current challenge in the research of solar cells is the significantly reduced hysteresis behaviors. The reason for the hysteresis occurrence in solar cells is still not fully understood. Hysteresis is commonly referred to as a historical phenomenon or historical effect, in which past events affect the status quo and represent a state that has not been undone. In the case of solar cells, during the power generation process, the current-voltage characteristics under the light irradiation should usually take the same value when the voltage is raised and lowered [101–103]. For perovskite solar cells, the hysteresis has also been reported to originate from the photocurrent density-voltage (J-V) responses. The perovskite material causing a large capacitive effect was also believed that hysteresis depends on. Ion migrations across the device and the effects of electronic defects during measurements have been ascribed to the hysteresis behavior of currentvoltage (I-V) curves in terms of mechanisms. Many cases of hysteresis also reported the different current values detected by the direction of the voltage sweep in solar cells. For electrical hysteresis, when measuring the I-V characteristics by voltage scanning on its terminals, there is a difference between the forward scanning I-V curve and the reverse scanning. The two scanned I-V curves do not coincide with each other even for a relatively low scanning rate of 100 mV/s. Hysteresis behavior was found to be correlated with the scanning rate of solar cell voltage, as the voltage scanning rate increases the hysteresis which becomes pronounced. The discovery of hysteresis motivated research on hybrid materials [103–107]. The advantage of this phenomenon is that the output value of power is stable considering the practical use of solar cells. On the other hand, from the viewpoint of basic research, it has been found that the hysteresis phenomenon could be improved by material design. Modulation of structure-activity relationship provides a great hint to think about the mechanism by which hysteresis occurs, to produce hysteresis-free PSCs with synergetic performance optimization before the commercialization [107–111]. Additionally, using nucleic acid hybrid materials has been proven to increase the photostability of the dye sensitizers and also improve the efficiency of the photovoltaic cell. Unlike conventional polymers, nucleic acids complexed with donoracceptor molecules can impose a defined spatial organization and fixed orientation [111–113]. Nucleic acid material formed with ionic surfactant or lipids with the ionic head group could form regular arrangements such as lamellar structures or parallel aligned layers that alternate through the self-assembly process. The solubility of

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nucleic acid materials in aqueous and various solvents improved their manufacturing processability with environmental friendliness.

3.5

Summary and Outlook

In resource-limited areas where fuel transportation or grid connection is limited by high cost, it is crucial to develop solar-generated electricity. The absence of environmental pollution during use and minimal maintenance are also key performance indicators for the technology and materials of solar power conversion devices in remote areas such as satellites, ocean vessels, mountains, and deserts. As described in this chapter, nucleic acid hybrid materials can assist the improvement of efficient photovoltaic cell performance attribute to their optoelectronic properties and their flexibility to interact with donor-acceptor molecular pairs. The advantage of manufacturing process is being solution-processable, thus reducing the cost while enhancing the durability, using nucleic acid as hybrid biomaterial with green technology. Acknowledgments This work is financially supported by the Science Technology and Innovation Commission of Shenzhen Municipality (深创-20205354JCYJ20190806153018791), Natural Science Foundation of Zhejiang Province (LGF19H200005), Japan-China Medical Association (国卫-2018920), the Frontiers Science Center for Flexible Electronics (FSCFE) and Shaanxi Institute of Biomedical Materials and Engineering (SIBME) is further acknowledged in providing space, equipment, and services. Author Contribution The research work referred to in part in the Results section was undertaken by DB, CXR (who undertook the experimental design, data analysis, and discussion; unpublished data excluded). DB, HHF, and RCX wrote the first manuscript of this book chapter and completed the editing of subsequent drafts.

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Dan Bai acquired her Ph.D. in Molecular Photonics from Newcastle University, UK, in the year 2013, and was appointed as a lecturer and then associated professor in the School of Medicine, Xi’an Jiaotong University, in the year 2014–2019, followed by a current post in Institute of Flexible Electronics, Northwestern Polytechnical University. As PI and co-PI, Dan’s recent publications include the papers in Dyes & Pigments, Molecular TherapyNucleic Acids, JMCC, Biotech.&Bioeng., and ACS Appl.Mat.Int. Appl.OrgMet.Chem., alongside two book chapters; she received several grants with a sum of over 2 million including the ones from the Natural Science Foundation of China and the Japan-China Medical Association. Her group now focuses on the design and creation of molecular systems for photonic and optoelectronic applications. Huhu Feng is a first-year graduate student majoring in materials and chemistry at the Institute of Flexible Electronics, Northwestern Polytechnical University. He is interested in the interdisciplinary research of biohybrid materials for optoelectronic devices. Currently, he has four papers in press.

Xingchen Yu is currently a junior undergraduate student at the Queen Mary Engineering School, Northwestern Polytechnical University. He was awarded the Outstanding Student (2019), Outstanding Youth League Members (2020), and other distinguished achievements.

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Chenxin Ran received his B.S. in Applied Chemistry from the Xidian University in 2009 and obtained his Ph.D. in Electronic Science and Technology from the Xi’an Jiaotong University in 2016 under Professor Minqiang Wang’s supervision. He worked as a visiting scholar from 2014 to 2015 with Professor Liming Dai at Case Western Reserve University. He worked as a postdoc fellow in Xi’an Jiaotong University with Professor Zhaoxin Wu from 2016 to 2020. Now he works as an associate professor at Northwestern Polytechnical University. His research interests include photophysics, structural control, and property optimization of Pb/Pb-free-based perovskite materials and their optoelectronic applications.

Academician Huang Wei received his BSc in Chemistry from Peking University in 1983 and his MSc and Ph.D. in Physical Chemistry from the same university. He did his postdoctoral research in the Department of Chemistry with the National University of Singapore (NUS) where he participated in the foundation of the Institute of Materials Research and Engineering (IMRE) since 1995. In 2001, he became Chair Professor at Fudan University, where he founded and chaired the Institute of Advanced Materials (IAM). In June 2006, he was appointed as the Deputy President of the Nanjing University of Posts and Telecommunications, where he initiated the Institute of Advanced Materials (IAM) and the Key Laboratory for Organic Electronics and Information Displays (KLOEID). In November 2011, he was elected as Academician of the Chinese Academy of Sciences (CAS). In July 2012, he was appointed as the President of the Nanjing Tech University (NanjingTech, 2011 University) and founded the Institute of Advanced Materials (IAM) and the Key Laboratory of Flexible Electronics (KLOFE) with Nanjing Tech. In 2013, he was elected as Director-General of Jiangsu National Synergetic Innovation Center for Advanced Materials (SICAM). In 2015, he founded the International Research Center of Flexible Electronics (CoFE) and the Joint International Laboratory of Flexible Electronics (LoFE). In October 2016, he led the construction of the National Innovation Talent Cultivation Base of Flexible Electronics (National 111 project). In April 2017, he was appointed as the Executive Vice President of Northwestern Polytechnical University. In October 2017, he founded the Institute of Flexible Electronics (IFE) at Northwestern Polytechnical University.

Chapter 4

First-Principles Calculations for the Interfaces of Perovskite Solar Cells Jun-Peng An, Ying Tian, Hong-Tao Xue, Jun-Chen Li, Jun-Qiang Ren, Xue-Feng Lu, and Fu-Ling Tang

4.1

Introduction

Simulation is becoming an important tool for the research of solar cells, which is a useful supplement to experimental research. First-principles calculation based on density functional theory (DFT) is performed on solar cells to understand the relationship between material structure and device performance. DFT [1–3] has become an effective method in fields of physics, quantum chemistry, and computational materials science, which has reasonably predicted the structure and properties of many substances. After several years of development, the all-solid-state perovskite solar cells have achieved a power conversion efficiency of 25.6% [4], which proves that their usage prospect is huge. Perovskite solar cells consist of multilayer structures [5]: substrate material | electron transport layer | light absorption layer | hole transport layer | metal electrode. The key material to a complete device is the perovskite layer, also known as a light-absorbing layer. It has the general chemical formula ABX3, where A is CH3NH3+, HN ¼ CHNH3+, or HC (NH2)2+; B is generally occupied by Pb or Sn; and X is for I, Cl, Br, or I1-xClx [6, 7]. MAPbI3 is a typical organic-inorganic perovskite. Since elemental substitution can be carried out at MA, Pb, or I, other organic-inorganic perovskites are considered derivatives of MAPbI3. The temperature and pressure can change the distance between atoms and change the bonding mode of the materials, thereby causing the phase change of the CH3NH3PbI3 structure. Many x-ray diffraction (XRD) [8–11] results indicate that the pressure will cause the CH3NH3PbI3 phase change from the tetragonal structure

J.-P. An · Y. Tian · H.-T. Xue · J.-C. Li · J.-Q. Ren · X.-F. Lu · F.-L. Tang (*) State Key Laboratory of Advanced Processing and Recycling of Non-ferrous Metals, Department of Materials Science and Engineering, Lanzhou University of Technology, Lanzhou, People’s Republic of China e-mail: [email protected]; tfl@lut.edu.cn © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_4

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to the orthorhombic structure (space group Imm2). Jiang et al. [12] used DFT to study the interaction between the high-pressure cubic structure of CH3NH3PbI3 and the bandgap. Moreover, the charge density and band structure of orthorhombic CH3NH3PbI3 at different pressures were studied by Kong et al. [10]. Our work used DFT to study the structure, electronic, and optical properties of CH3NH3PbI3 at the highest pressure value of 0.4 GPa. Nowadays, the main problem of this MAPbX3 perovskite material is that Pb is highly toxic and it is not friendly to the environment. Hence, from the long-term project to protect the environment, the promotion and application of such lightabsorbing materials will inevitably be restricted. This requires us to develop less or free-lead perovskite materials. In recent research on element doping of organicinorganic perovskites, In, Ag, Al, and Na are used to replace Pb [13–15]. We used first principles to study the changes of energy bands and optical properties caused by Ag+ doping for Pb2+. For the doping research for perovskite batteries, Ag has excellent performance among similar substitution elements (Sr, Sn, In, and so on) and is an environmentally friendly material. It has a similar ion radius (129 pm) compared with Pb2+ (133 pm) [16], so it will not greatly change the crystal structure. Recently, pure inorganic cesium lead halide perovskite CsPbX3 (X ¼ Cl, Br, and I) has been prepared due to its good thermal stability and high charge carrier mobility. Among them, cubic (α)-CsPbI3 with the bandgap 1.73 eV is the most suitable one for the photovoltaic industry. Unfortunately, in a humid environment even at room temperature, the α-CsPbI3 phase is unstable, making it transform into a nonperovskite orthorhombic phase (δ) with lower photovoltaic efficiency” ([17, 18] with permission). The work by Bai et al. [19] confirmed that Mn2+ ions will be inserted into the gaps of the CsPbI2Br lattice during film growth, thereby reducing the recombination loss and increasing the hole extraction efficiency. After adopting the interstitial Mn-doped CsPbI2Br film, the best device efficiency record of this solar cell was broken in 2018, reaching 13.47%. According to the experimental results of CsPbI2Br thin film, it was found that Mn doping seems to have substitution and interstitial doping structures. Understanding the mechanism of Mn doping will provide meaningful help for energy band engineering and the preparation of stable CsPbI2Br perovskites. In addition to the material properties of each layer of the solar cell, the surface and interface structure and their electronic properties are also important to photovoltaic properties. A tremendous number of researchers have pursued higher power conversion efficiency (PCE) through experiments of surface (interface) regulation [20– 23]. In perovskites, defects in the bulk phase, grain boundaries, surface, and interface will much affect surface charge recombination mechanism and the carrier transport, and this hinders the further improvement of the efficiency of polymer solar cells (PSCs). Dangling bonds on the surface result in new electronic states (energy levels) in the bandgap [24]. The interface or surface states are one of the reasons why perovskite batteries have lower conversion efficiency and poor stability under this

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crystal plane orientation. Therefore, it is also the main method to improve device stability and battery efficiency by adopting an effective passivation method. Therefore, within this chapter, the lattice structures and electrical properties of phenethyl ammonium (PEA+) adsorbed-MAPbI3 (110) surface, WZ-ZnO (100)/ MAPbI3 (112) interface, WZ-ZnO/MAPbI3 interface, NiO (110)/MAPbI3 (100) interface, MAPbI3/Au interface, SnO2 (110)/MAPbI3 (100) interface, CsPbI3/SnO2 interface, and CsPbI2Br/SnO2 interface are investigated. For these materials and surfaces or interfaces in them, their lattice structures the surface energy, or interface energies are provided. We calculated the density of states (DOS), charge density difference, Bader charge, and optical properties (including absorption coefficient, refractive indices, and extinction coefficient). Our purposes are to obtain quantitative atomic and electronic information in the battery materials and then to understand the relationship between composition, structure, and performance.

4.2

Methods

In this chapter, most of the calculation work was performed in the Vienna ab initio simulation package (VASP) [25, 26] “based on first-principles density functional theory (DFT) with the Perdew-Burke-Ernzerhof (PBE) [27] version of the generalized gradient approximation (GGA) and the more complex Heyd-ScuseriaErnzerhof (HSE06) [28] hybrid functional. The projector augmented wave (PAW) [29] method was applied to describe the pseudopotential. The electronic configurations are [Xe]6s1, [Xe]5d106s26p2, [Kr]4d105s25p5, [Ar]3d104s24p5,” ([30] with permission) [Kr]4d105s25p2, [Ar]3d54s2, [Kr]4d105s1, O2s22p4, N2s22p3, C2s22p2, H1s1 for cesium, plumbum, iodine, bromine, stannum, manganese, silver, oxygen, nitrogen, carbon, and hydrogen, respectively. The lattice structure optimization adopts the conjugate gradient (CG) [31] method. The interface binding energy, band structure, and density of states (DOS) in the systems are obtained based on the tetrahedral method of Bloch correction [32].

4.3 4.3.1

CH3NH3PbI3 Phase Transformation Under Pressure Lattice Structures of CH3NH3PbI3 Under Pressure

At room temperature and atmospheric pressure, CH3NH3PbI3 is tetragonal, but it may be the orthorhombic structure or the cubic structure under high pressure. Figure 4.1 shows the different models of tetragonal [33–35], cubic [12], and orthorhombic structures [10] of CH3NH3PbI3.

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Fig. 4.1 Lattice structures of CH3NH3PbI3: (a) tetragonal I4/mcm phase, (b) cubic Im3 phase, and (c) orthorhombic Imm2 phase at atmospheric and high pressures. (Ref. [36] with permission. Copyright Elsevier Press)

Fig. 4.2 The projected energy band structures: the tetragonal phase (a), cubic phase (b), and orthorhombic phase (c). The energy orbitals of I-5p, Pb-6s, and Pb-6p are represented by green, blue, and red circles, respectively. (Ref. [36] with permission. Copyright Elsevier Press)

4.3.2

Projected Band Structures for Three Phases of CH3NH3PbI3

“The projected band diagram aims to find out the contribution of the atomic orbit to the energy band near the Fermi level according to the projection of the energy band on each atom and each orbit and its ratio. CH3NH3PbI3’s three-phase projection band structure is shown in Fig. 4.2, and the Fermi levels in Fig. 4.2 are set to 0 eV. The energy bands formed by different orbits are represented by different colors, and each point represents the proportion of orbit projection” ([36] with permission). “Figure 4.2 shows that the conduction band minimum (CBM) of these three structures of CH3NH3PbI3 is contributed by Pb-6p orbitals and valence band maximum (VBM) by I-5p and Pb-6s orbitals. For the two high-pressure phases, the contribution of the Pb-6s orbitals in the cubic phase is greater than the orthorhombic phase. The CBM and VBM of the three structures shown in Fig. 4.2 are at point G, which shows that CH3NH3PbI3 is a direct bandgap semiconductor material regardless of pressure” ([36] with permission).

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Under normal temperature and pressure, the calculated bandgap of the tetragonal structure is “1.76 eV, which is larger than the experimental value of 1.62 eV [8]. The calculation results show that the bandgap values of the three structures at the phase transition point are higher than the corresponding experimental values [10], and the error range of the two is 8%. In the 3 to 0 eV part, the bonding orbitals of Pb-6s and I-5p” ([36] with permission) determine the large fluctuation of the energy band, while in the range of 1.5–5 eV, the antibonding orbitals of Pb-6p and I-5p contribute to the conduction bands.

4.3.3

Variation of the Bandgap of CH3NH3PbI3 with Pressure

We not only investigated “the energy band structure of CH3NH3PbI3 at the phase transition point but also analyzed the effect of different pressure values near the phase transition point on the bandgap. It can be seen from Fig. 4.3 that before the phase transition” ([36] with permission), the bandgap of the tetragonal structure decreases with increasing pressure, while at the phase transition point, the bandgap suddenly increases. After the phase transition occurs, the band gaps of the two possible high-pressure structural phases are larger than those of the normal-pressure tetragonal phase, and after the pressure is gradually increased, the band gaps of the two possible high-voltage structural phases gradually decrease.

Fig. 4.3 Bandgap change of CH3NH3PbI3with pressure. (Ref. [36] with permission. Copyright Elsevier Press)

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Fig. 4.4 CBM and VBM of three CH3NH3PbI3 phases under pressure. (Ref. [36] with permission. Copyright Elsevier Press)

To further study the change of the bandgap of three CH3NH3PbI3 phase structures under high pressure, we give the displacements of CBM and VBM of each phase structure under pressure, as shown in Fig. 4.4. “Before the phase transition, the increasing pressure causes the CBM of the CH3NH3PbI3 tetragonal phase to move downward and its VBM as a whole tends to decrease. Therefore, the pressure before the phase change introduces the bandgap of the tetragonal phase to gradually decrease. At the phase transition point of 0.4 GPa, the CBMs of the two highpressure phases move upward, while the corresponding VBMs move downward, increasing the bandgap of the two structures at the phase transition point” ([36] with permission). After the phase change point, the CBM of the two high-pressure phases tends to decrease as a whole and their VBM tends to increase as the pressure increases, indicating that the bandgap gradually decreases under the effect of pressure. Some studies [8] found that the increase and decrease of the bandgap with the change of pressure are mainly due to the octahedron structural deformation of [PbI6]4. The CH3NH3+ organic cation located in the gap of the octahedron does not significantly change the bandgap. However, the deformation mechanism of [PbI6]4 octahedrons and the principle that the distortion structure changes the bandgap are still unclear. In some literatures [37–41], the bond lengths and their average value are used to describe the change of the degree of polyhedron distortion or the symmetry change of the crystal structure. Based on this, we calculate the “average Pb-I bond length in [PbI6]4 octahedrons and the octahedron deformation degree σ of [PbI6]4 to determine the structure deformation and analyze the reason why pressure affects the bandgap” ([36] with permission). Although the “[PbI6]4 octahedrons in CH3NH3PbI3 shrinks at the same pressure, there is no significant difference in its [PbI6]4 octahedrons between atmospheric pressure and the phase transition point of each phase. Figure 4.5 shows [PbI6]4

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Fig. 4.5 The octahedrons [PbI6]4 in CH3NH3PbI3 at the phase transition point: (a) tetragonal phase, (b) cubic phase, and (c) orthorhombic phase. (Ref. [36] with permission. Copyright Elsevier Press)

Fig. 4.6 The change of average length of Pb-I bonds (a) of [PbI6]4 octahedrons and the standard deviation σ (b) with pressure. (Ref. [36] with permission. Copyright Elsevier Press)

octahedrons with three-phase structures at the phase transition point” ([36] with permission). We use the average length of the Pb-I bonds in [PbI6]4 octahedrons in Fig. 4.6a to study and analyze the relationship between the optimized bond length and pressure in the structures of CH3NH3PbI3. We found that the gradual increase in pressure reduces the average length of all Pb-I bonds in the three [PbI6]4 octahedrons, which shows that the pressure causes the distance between Pb and I atoms to

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shorten and causes [PbI6]4 octahedrons to shrink slightly. This may be the reason for the bandgap reduction of these three structures with pressure. Besides, to quantify “the dispersion of a set of data, the standard deviation σ of Pb-I bonds in [PbI6]4 octahedrons can be expressed, which can be obtained according to the following Eq. (4.1) to characterize the degree of different structures’ deformation of the octahedron. rffiffiffiffi N 1X σ¼ ðx  μÞ N i¼1 i

ð4:1Þ

In Eq. (4.1), xi represents the Pb-I bond length, N is the number of Pb-I bonds, and the average length of Pb-I bonds is expressed in μ. The lower the σ value, the closer the Pb-I bond length in [PbI6]4 octahedron to its average bond length, indicating the smaller degree of distortion of the octahedron. The higher the σ value, the longer the Pb-I bond lengths in a wider range of values, meaning greater distortion of the octahedron” ([36] with permission). Figure 4.6b shows σ, the change of [PbI6]4 octahedron with the pressure in three CH3NH3PbI3 structures. According to Fig. 4.6b, in the possible cubic and orthorhombic phases, the distortion degrees σ of [PbI6]4 octahedron is greater than that of atmosphericpressure tetragonal phase [PbI6]4 octahedron distortion. From Fig. 4.3, we find that the band gaps of the atmospheric-pressure tetragonal phase are smaller than those of the two high-pressure phases. Therefore, the pressure promotes the phase change of CH3NH3PbI3, and the enhancement of its bandgap is due to the increase of the distortion degree σ of [PbI6]4 octahedron. That is to say, for the three structural phases of CH3NH3PbI3, the higher the degree of distortion σ, the larger the bandgap.

4.3.4

Optical Properties of CH3NH3PbI3 Under Pressure

Figure 4.7 shows the optical properties of the three CH3NH3PbI3 at the phase transition point. The real ε1(ω) and imaginary ε2(ω) of the complex dielectric function are shown. The photon energy range is 0–23 eV, and function ε1(ω) can characterize the electronic polarizability. In Fig. 4.7a, when the energy is zero, the static dielectric constants of the cubic phase, tetragonal phase, and orthorhombic phase are 4.32, 4.45, and 4.5, respectively. Subsequently, these values gradually increased, reaching the maximum at 8.28 at 2.4 eV for the tetragonal phase, the maximum at 8.34 at 2.4 eV for the cubic phase, and the maximum at 7.75 at 2.42 eV for the orthorhombic phase. Their value gradually decreases below zero. It can be seen in Fig. 4.7b that these three structures have two main complex dielectric function peaks. The tetragonal phase shows peaks at 3.05 eV and 6.58 eV, respectively, with ε2(ω) “values of 8.49 and 3.02. For the cubic phase, the two peaks are 7.86 and 2.63, while the orthorhombic phase has peaked at 3.03 eV and 6.22 eV. The intensity of the peaks of the cubic phase is the lowest among the three phases, and the threshold energy of these three phases is close to 1.8 eV” ([36] with permission).

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Fig. 4.7 The complex dielectric function real: ε1(ω) (a) and imaginary ε2(ω) parts (b) for three structures of CH3NH3PbI3 at the phase transition point. (Ref. [36] with permission. Copyright Elsevier Press)

Fig. 4.8 The extinction coefficient k(ω) (a) and refractive index n(ω) (b) at the phase transition point (0.4 GPa). (Ref. [36] with permission. Copyright Elsevier Press)

Figure 4.8a shows three structures’ extinction coefficient k(ω) at the phase transition point. Specifically, the “variation trends of k(ω) with light energy are similar to the imaginary parts ε2(ω). Figure 4.8b shows the refractive indexes n(ω). When photons penetrate the material, they slow down due to their interaction with electrons, so the refractive index is greater than 1. The static refractive index n(0) of the cubic phase, tetragonal phase, and orthorhombic phase are 2.08, 2.11, and 2.12, respectively. The refractive index of the tetragonal phase reaches a maximum of

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Fig. 4.9 The optical reflectivity R(ω) (a) and absorption coefficient α(ω) (b) at the phase transition point (0.4 GPa). (Ref. [36] with permission. Copyright Elsevier Press)

2.95 at 2.53 eV. For the other two possible high-pressure phases, the maximum refractive indexes are 2.86 and 2.96, respectively” ([36] with permission). Figure 4.9a shows the optical reflectivity R(ω) at the phase transition point. When the energy is zero, the tetragonal phase has the reflective value “12.7%, and the maximum reflectivity value is about 59.4%” ([36] with permission) occurring at an energy of 7.86 eV. For the orthorhombic phase and the cubic phase, the reflectivity at energy zero is 12.9% and 12.3%, respectively, while the maximum values are 57.4% and 57%, respectively. The maximum reflectivity appears in the “light energy region of 7–8 eV” ([36] with permission). The absorption coefficients α(ω) of the three crystal structures of CH3NH3PbI3 “at the phase transition point are shown in Fig. 4.9b. The material absorbs incident light after the incident light energy exceeds 1.8 eV. As the incident light increases to 2.0 eV, the absorption coefficient of this type of material increases significantly. For these three structures, when the photon energy is greater than 23 eV, the absorption coefficient tends to be almost zero and no longer absorbs light. For the tetragonal phase and cubic phase, the maximum absorption peak appears at 6.80 eV, while for the orthorhombic phase, the maximum absorption peak appears at 6.68 eV” ([36] with permission).

4.4 4.4.1

Perovskite Solar Cell Materials Upon Doping: Ag-Doped MAPbI3: MAPb1-xAgxI3 Structural Properties

For studying the optical and electronic properties of the MAPb1-xAgxI3 system, we optimized the cubic MAPbI3 cell. The optimized lattice parameters are 6.33 Å, which is in good agreement with the experiment value of 6.295 Å [42–45]. A larger

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2  4  5 calculation model is built to obtain a low concentration doped structure with lattice parameters a ¼ 12.660 Å, b ¼ 25.320 Å, and c ¼ 31.650 Å.

4.4.2

Electronic Properties

“Figure 4.10 shows the band structure and total density of states (DOS) of pure MAPbI3. From Fig. 4.10a, we see that the MAPbI3 is a direct bandgap semiconductor material. Although density functional theory (DFT, such as GGA + PBE) often underestimates the bandgap of MAPbI3 perovskites [46–48], our calculated bandgap 1.64 eV is consistent with the simulated results (1.68 eV) [49] and experimental reports (1.60 eV) [50, 51]. Figure 4.10b shows the VBM, mainly contributed by I-5s and Pb-6s orbitals, while the CBM is mainly contributed by I-5p and Pb-6p orbitals” ([52] with permission). The highest experimental Ag doping concentration is 5% [16]. Figure 4.11 shows the band structure and DOS of MAAgI3. Figure 4.11a shows few-electron states nearby the Fermi level, and MAAgI3 has a certain metallicity. We analyzed the total DOS and some DOS of MAAgI3. Figure 4.11b shows the electronic states, mainly attributed by Ag-4d and I-5p orbits contribution nearby the Fermi level. The orbital contribution of I-5p could be introduced by lattice distortion because the ionic radius of Ag+ (129 pm) is less than that of Pb2+ (133 pm). To evaluate MAPb1-xAgxI3’s stability, we calculated the formation of the energy of the MAPb1-xAgxI3 system by Eq. (4.2). E¼

  1 ½Etot ðMAPbI3 Þ þ xEðAgÞ  Etot ðMAPb1x Agx I3 Þ  xE ðPbÞ x

ð4:2Þ

Fig. 4.10 The band structure (a) and density of states (b) of single-cell CH3NH3PbI3. (Ref. [52] with permission. Copyright American Institute of Physics)

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Fig. 4.11 The band structure (a) and density of states (b) of single-cell CH3NH3AgI3. (Ref. [52] with permission. Copyright American Institute of Physics)

Equation (4.2) [53] represents the total energy difference of the system before and after the doping of Pb by Ag. “Dividing by x shows the average formation energy of each Ag atom. Etot(MAPbI3) is the total energy of the undoped lattice, Etot(MAPb1xAgxI3) is the total energy of the supercell cell after doping, and x is the number of Ag substitution atoms. E(Ag) and E(Pb) are the corresponding energies of a single Ag and Pb atom, respectively. The formation of energy changes with the doping concentration in Fig. 4.12. With increasing x, the formation energy increases. The formation energy for Ag substitution is negative, indicating that the process will absorb energy as it progresses. Based on the above relationship, the optical band gap of MAPb1-xAgxI3 varies with doping concentration x in Fig. 4.13. The optical bandgap of MAPbI3 is 1.56 eV, and it agrees well with the experimental values (1.51 eV) [35, 54]. The optical bandgap of MAPb0.975Ag0.025I3 (1.57 eV) increases slightly with doping a small amount of Ag. When the Ag adding concentration increased from 0% to 30%, the MAPb1-xAgxI3 optical band gap increases from 1.56 eV to about 3.11 eV. The literature [16, 55] gave that if x < 0.025, the MAPb1-xAgxI3 power conversion efficiency is better. Then, we will show the MAPb1-xAgxI3 photoelectric properties with the low concentration doping (x  0.1)” (Copyright [52] with permission). We calculate the MAPb1-xAgxI3 system’s total DOS and the partial DOS to investigate the electronic structure changes in Fig. 4.14. Figure 4.14a shows that the total DOS of MAPb1-xAgxI3 system varies with x ¼ 0, 0.025, 0.05, 0.075, and 0.100. When x reaches 2.5%, the MAPb0.975Ag0.025I3’s total DOS is almost unchanged. “Comparing the total DOS of MAPb0.950Ag0.050I3 and those of MAPb0.975Ag0.025I3 in Fig. 4.14a, we find that the valence band moves to a higher energy level, and the energy enters the valence band with x reaches 2.5%, which shows MAPb0.95Ag0.05I3 is a p-type conductor. The CBM is almost unchanged. The bandgap is relatively decreased by 0.013 eV. From the total DOS of MAPb0.975Ag0.025I3 to the MAPb0.9Ag0.1I3 in Fig. 4.14a, the total DOS of MAPb1-xAgxI3 is almost unchanged with x increasing” ([52] with permission).

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Fig. 4.12 The total energy of the supercell CH3NH3Pb1-xAgxI3 (x ¼ 0.025, 0.05, 0.075, 0.10, 0.15, 0.20, 0.25, 0.30). (Ref. [52] with permission. Copyright American Institute of Physics)

Fig. 4.13 The optical band gap of the CH3NH3Pb1xAgxI3 with doping concentration. (Ref. [52] with permission. Copyright American Institute of Physics)

Figure 4.14a shows MAPbI3’s total DOS. Combining with Fig. 4.14b, we can find that I-5p and Pb-6s orbits determine pure MAPbI3’s VBM on about 0.21 eV. The MAPbI3’s CBM is controlled by Pb-6p orbits. To study the MAPbI3’s electronic properties variation with Ag doping concentration, the MAPbI3’s supercell was doped with a different number of Ag atom atoms, and their corresponding “doping concentration x increases from 2.5% to 10%. Partial DOS is analyzed for different doping concentrations x. According to Fig. 4.14b, c, the electron orbital contribution of the Pb and I atom is almost unchanged at 2.5% concentration. The slight difference comes from the contribution of the Ag-4d orbit. Figure 4.14a shows that the MAPb0.950Ag0.050I3’s valence band shifts toward high-energy levels at x ¼ 5% concentration. In Fig. 4.14d, there is a slight local phenomenon on the Pb-6p, I-5p, and Ag-4d orbital from 3.2 to 0 eV” ([52] with permission). By analyzing Fig. 4.14d–f, it can be seen that as the doping concentration x increases, the valence band energy from Ag atoms moves to a lower energy

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Fig. 4.14 The total density of states of (a) CH3NH3Pb1-xAgxI3; (b) the partial density of states of CH3NH3PbI3, (c) CH3NH3Pb0.975Ag0.025I3, (d) CH3NH3 Pb0.950Ag0.050I3, (e) CH3NH3 Pb0.925Ag0.075I3, and (f) CH3NH3Pb0.900Ag0.100I3. (Ref. [52] with permission. Copyright American Institute of Physics)

level. The conduction band introduced by Ag atoms hardly changes, and the bandgap slightly increases as the doping concentration x increases from 5% to 10%.

4.4.3

Optical Properties

The “optical properties of MAPb1-xAgxI3 (x ¼ 0, 0.025, 0.05, 0.075, 0.100) system are investigated through the absorption coefficient α, dielectric function ε, reflectivity R, extinction coefficient k, and refractive indices n. In Fig. 4.15a, it can be seen that there are four main absorption peaks at 1.56, 2.19, 2.81, and 6.26 eV, respectively. The first absorption peak is contributed by the electronic transition from the

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Fig. 4.15 Imaginary part (a) and real part of dielectric function (b) of CH3NH3Pb1-xAgxI3 (x ¼ 0, 0.025, 0.05, 0.075, 0.100) model. (Ref. [52] with permission. Copyright American Institute of Physics)

I-5p or Pb-6s orbital (the valence band) to the Pb-6p orbital (the conduction band) “([52] with permission). The second absorption peak is contributed by the electronic transition from the Pb-6p orbital to the Pb-6p orbital. The third absorption peak is contributed by the I-5p or Pb-6p orbit. The above analysis also explains the absorption peaks in absorption coefficients and refractive index. “The MAPb1-xAgxI3 (x ¼ 0.025, 0.050, 0.075, 0.100) system has three main absorption peaks. Three main absorption peaks have the position of nearby 2.19, 2.81, and 6.26 eV, and with increasing doping concentration (x  0.1), their absorption intensity decreases. This shows that the doping of Ag elements has a certain effect on the absorption strength” ([52] with permission). Figure 4.15b shows the real part of MAPb1-xAgxI3’s dielectric function (x ¼ 0, 0.025, 0.050, 0.075, 0.100). For undoped MAPbI3, it has “good dielectric behavior when the energy range is less than 8.76 eV but shows a certain metal character in the energy range of larger than 8.76 eV. MAPbI3’s high-frequency dielectric constant is 7.19, which agrees very well with the experimental values of 7.1 [56], 6.6 [57], 5.5 [58], 5.6 [59], and 6.5 [60, 61]. The high-frequency dielectric constant is 6.76 at x ¼ 2.5%, lower than that of pure MAPbI3. MAPb1-xAgxI3’s highfrequency dielectric constant (x ¼ 0.050, 0.075, 0.100) is higher than that of pure MAPbI3: 7.6, 10.61, and 7.24, respectively. We note that MAPb0.925Ag0.075I3 has the highest high-frequency dielectric constant. We think that the introduction of Ag elements has a certain impact on it” ([62] with permission). Figure 4.16 shows MAPb1-xAgxI3’s “absorption coefficients (x ¼ 0, 0.025, 0.05, 0.075, 0.100) system. For a solar cell material, the absorption coefficient is a critical parameter of the light absorption layer. The larger the absorption coefficient, the thinner the light absorption layer. From Fig. 4.16, as the Ag doping concentration x increases, the absorption coefficient decreases slightly in the visible light region. Besides, if where the light absorption layer itself is thin, the electron-hole pairs excited by light can be efficiently extracted from the outside, which would achieve a high conversion efficiency. It can be noted that the absorption coefficient rises earlier nearby the basic absorption edge, and a few micron thicknesses of the light absorption layer” ([52] with permission) can fully absorb sunlight.

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Fig. 4.16 Absorption coefficient of CH3NH3Pb1xAgxI3 (x ¼ 0, 0.025, 0.05, 0.075, 0.100) model. (Ref. [52] with permission. Copyright American Institute of Physics)

As shown in Fig. 4.16, in the five models, MAPb1-xAgxI3’s absorption peaks and absorption intensities vary greatly “in the energy range of less than 5.57 eV. The absorption peak is relatively broad and strong in the range of 2-7 eV. The intensity of the absorption peaks begins to decrease when the energy is greater than 7 eV. This shows that the material has a certain range of light absorption. Beyond this range, light absorption will not occur. MAPb1-xAgxI3 (x ¼ 0.025, 0.050, 0.075, 0.100) model at 1.56 eV has no absorption peak. The absorption peak of MAPb0.9Ag0.1I3 (x ¼ 0.1) model is lower than that of the MAPb1-xAgxI3 (x ¼ 0, 0.025, 0.050, 0.075) model in the range of 2-2.83 eV, but the intensity of the absorption peak is higher than that of MAPb1-xAgxI3 (x ¼ 0, 0.025, 0.050, 0.075) model at 3.09 eV. This means that it can generate electron-hole pairs under low-energy irradiation” ([52] with permission). Figure 4.17a shows MAPb1-xAgxI3’s reflectance spectrum (x ¼ 0, 0.025, 0.05, 0.075, 0.10) model. The “absorption peaks decrease as doping concentration increases. The absorption peak is mainly introduced by the imaginary part of the dielectric function in the range of 1-10 eV” ([62] with permission). Figure 4.17b, c shows MAPb1-xAgxI3’s refractive index and extinction coefficient. The relevant data are all from the dielectric function. Now, MAPb1-xAgxI3’s data of the reflectivity, refractive index, and extinction coefficient in the experiment are relatively rare. We hope that our calculation data would be a certain reference for the experiment.

4.5 4.5.1

Mn-Doped CsPbI2Br: Structural, Electronic, and Optical Properties Structural Properties

At room temperature, “CsPbI3 has a cubic crystal structure with a space group of Pm3m (No. 221), as shown in the unit cell structure in Fig. 4.18a. We found that the lattice constants obtained from the experiment are a ¼ b ¼ c ¼ 6.289 Å, α ¼ β ¼ γ ¼

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Fig. 4.17 Reflectivity R (a), refractive indices n (b), and extinction coefficient k (c) of CH3NH3Pb1xAgxI3 (x ¼ 0, 0.025, 0.05, 0.075, 0.100) model. (Ref. [52] with permission. Copyright American Institute of Physics)

Fig. 4.18 Lattice structure: CsPbI3 (a) and CsPbI2Br (b). (Ref. [63] with permission. Copyright Royal Society of Chemistry)

90 [64]. In our work, the lattice constants of CsPbI3 from first-principles calculations are 6.408 Å, similar to other theoretical calculation results (6.394 Å) [65]. Based on the optimized CsPbI3 lattice model, CsPbI2Br’s lattice model is constructed by substitution doping (see Fig. 4.18b). After optimization, the lattice constant of CsPbI2Br is a ¼ b ¼ c ¼ 6.297 Å. Compared with CsPbI3, the lattice of CsPbI2Br shrinks because the radius of the Br atom is smaller than that of the I atom” ([63] with permission). The “2  2  2 CsPbI2Br supercell is composed of 40 atoms. When one of the Pb atoms is replaced by an Mn atom, the doping concentration is 12.5%, and the lattice parameters reduce from 12.595 Å to 12.392 Å” ([63] with permission). Compared with Pb atoms’ radius, the reason for the lattice shrinkage of the lattice is the smaller radius of Mn atoms [66]. When Mn atoms occupy one of the octahedral gaps, the doping concentration is 2.5%, and then the lattice parameters rise to 12.619 Å. Therefore, only the octahedral gap of the CsPbI2Br supercell is taken into consideration for interstitial Mn doping. Figure 4.19 shows the most stable structure of Mn substitutional doping and Mn interstitial doping in CsPbI2Br. “To evaluate the thermodynamic stability of Mn-doping in CsPbI2Br, we obtain its binding energy by Eq. (4.3)” [67]:

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Fig. 4.19 Mn substitutional doped CsPbI2Br (a) and Mn interstitial doped CsPbI2Br (b). (Ref. [63] with permission. Copyright Royal Society of Chemistry)



Pb I Br Mn E tot  nCs E Cs atom  nPb E atom  nI E atom  nBr E atom  nMn E atom Eb ¼ ðnCs þ nPb þ nI þ nBr þ nMn Þ

 ð4:3Þ

where Eb and Etot are the binding energy and “total energy of Mn-doped CsPbI2Br, EA atom is the energy of free atoms i (i ¼ Cs, Pb, I, Br, Mn), and ni is the number of atoms i in the supercell. The calculated binding energies of Mn substitutional and interstitial doped CsPbI2Br are 2.5761 eV/atom and 2.5789 eV/atom, respectively. The negative binding energies and the small energy difference of 2.8 meV indicate that they are two stable doping structures. We can understand the fact that the Mn substitutional and interstitial doped CsPbI2Br had been prepared by experiments [68–70]. Mn interstitial doped CsPbI2Br seems slightly more stable due to its lower binding energy” ([63] with permission).

4.5.2

Electronic Properties

Figure 4.20a shows that the CBM and “the VBM of CsPbI2Br are at the G point. This indicates that CsPbI2Br is a direct bandgap semiconductor material, and it has a bandgap of 1.42 eV. The VBM of CsPbI2Br is mainly contributed by the I-5p and Br-4p orbitals, while the CBM is mainly contributed by the Pb-6p orbital. The band structure of Mn substitutional doped CsPbI2Br is denser than that of CsPbI2Br (Fig. 4.20b). Also, several impurity bands are appearing near the Fermi level, caused by Mn-3d orbitals. As the Mn-3d orbital is half-filled, these impurity bands are also called intermediate bands (IBs). The IBs can act as effective stepping stones to help transport valence electrons into the conduction bands, that is, the existence of IBs helps additional electronic transitions from VB to IB and from IB to CB, apart from the direct electronic transition from VB to CB. IB reduces the photon absorption energy threshold and increases the photocurrent, while the output voltage of the solar cell is not limited by this threshold. CsPbI2Br has a relatively wide bandgap which is

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Fig. 4.20 Band structures and TDOS: (a) undoped, (b) Mn substitutional doped CsPbI2Br, and (c) Mn interstitial doped CsPbI2Br. (Ref. [63] with permission. Copyright Royal Society of Chemistry)

not suitable for direct use as a light-absorbing material for solar cells, and Mn-doped CsPbI2Br can become a promising one for high-efficiency solar cells [71]. Also, there are many novel intermediate-band photovoltaic materials, for example, Ti or Cr-substituted CuGaS2 [72, 73] and Mn-doped AlGaN [74], which have been successfully fabricated experimentally by doping the transition elements. CBM and VBM at the G-point indicate that Mn substitutional doped CsPbI2Br is still a direct bandgap semiconductor. The Mn interstitial doped CsPbI2Br’s Fermi level passes through CBM, and the impurity level introduced by Mn doping is located above the bandgap and near the bottom of the conduction band (Fig. 4.20c). This shows its characteristics of an n-type semiconductor. Several flat bands near the Fermi level are expected to produce a sharp peak in the DOS, mainly from the Mn-3d orbital” ([63] with permission). The bonding characteristics of CsPbI2Br and Mn doping CsPbI2Br (100) planes were shown in Fig. 4.21. In Fig. 4.21a, the charge density around Pb atoms has a strong locality and the bonding characteristics are ionic. Stronger than the covalent bond, the bonding feature between Pb atoms and I or Br atoms is a mixed one. In Fig. 4.21b, the bond length of Mn-I or Br becomes shorter, which is beneficial to the charge transfer between them. The Mn substitutional doped CsPbI2Br has similar binding characteristics as with CsPbI2Br. In Fig. 4.21c, Mn atoms enter the octahedral gap of CsPbI2Br, I, and Br atoms in the PbI4Br2 octahedron are attracted by the Mn atom to cause charge change. The enhanced binding capacity indicates that the doping enhances electron transport and conductivity. PbI4Br2’s octahedron is found deformed.

4.5.3

Optical Properties

The response of ion and electron motion to dipole displacement in the external electric field can be expressed by the dielectric constant [75]. Comparing the DOS and the imaginary part of the dielectric function, we can reveal the interband transition process of excited electrons. The imaginary part of the dielectric function of CsPbI2Br has

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Fig. 4.21 The charge densities on the (100) plane of the undoped (a), Mn substitutional doped (b), and Mn interstitial doped CsPbI2Br (c). (Ref. [63] with permission. Copyright Royal Society of Chemistry)

Fig. 4.22 Real and imaginary parts of the dielectric function: undoped (a), Mn substitutional doped (b), and Mn interstitial doped CsPbI2Br (c). (Ref. [63] with permission. Copyright Royal Society of Chemistry)

three obvious peaks (Fig. 4.22a): the peak at 2.84 eV mainly comes from “the VBM (I-5p and Br-4p) to the CBM (Pb-6p) (see Fig. 4.20a). Figure 4.22b, c shows, in the high-energy range (>2.884 eV), the imaginary part of the dielectric function of Mn substitutional and interstitial doped CsPbI2Br is the same as that of the CsPbI2Br. In the low-energy range (0–2 eV), the imaginary part of the dielectric function of CsPbI2Br is different from that of Mn-doped CsPbI2Br. The imaginary part of the dielectric function of Mn-doped CsPbI2Br produces a new peak at 0.32 eV with a peak value of 0.25, while Mn-doped CsPbI2Br has a new peak at 0.39 eV with a peak value of 2.25, mainly derived from the electron transition between the VBM and the

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Fig. 4.23 Absorption coefficient spectrum: undoped (a), Mn substitutional doped (b), and Mn interstitial doped CsPbI2Br (c). (Ref. [63] with permission. Copyright Royal Society of Chemistry)

impurity band (see Fig. 4.20c). The static dielectric constant (SDC) is the refractive index measured at a frequency higher than the lattice vibration frequency [76] and can be obtained from the real part of the dielectric function. Figure 4.22 shows that the SDC of the undoped, Mn substitutional, and interstitial doped CsPbI2Br are 4.97, 5.99, and 10.27, respectively” ([63] with permission). As shown in Fig. 4.23a, as the “photon energy greater than 1.36 eV, the absorption coefficient of the undoped CsPbI2Br increases to the maximum at 6.65 eV. The absorption spectra of Mn-doped CsPbI2Br in Fig. 4.23b, c show that the light absorption capacity is the same in the high-energy range (>1.28 eV). And in the low-energy range (0–1.28 eV), the light absorption capacity of Mn-doped CsPbI2Br is stronger than the other two. We found that an absorption peak appears at 0.61 eV, which means that the photon energy of 0.61 eV can be absorbed, thereby promoting the transition of electrons from the impurity band near the Fermi level to the CBM” ([63] with permission). The impurity band of Mn-doped III-V compounds had similar reports [77–79].

4.6 4.6.1

Surface in Perovskite Solar Cell Materials: PCBM-Adsorbed MAPbI3 Surface Structure of PCBM-Adsorbed Surface

Figure 4.24a shows the surface adsorption model: the upper half is the adsorbed phenyl-C61-butyric acid methyl ester (PCBM) molecule, and the lower half is the

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MAPbI3 (100) plane. The structure contains 332 atoms, including 122 H, 18 N, 90 C, 66 I, 24 Pb, and 2 O atoms. Since the adjacent unit cells will have periodic effects, we add 30 Å vacuum layers along the c-axis to the adsorption system. During optimization for the structure, the positions of the four layers of atoms away from the surface of MAPbI3 are fixed, while the remaining atoms and PCBM molecules are relaxed. The single-point calculation determines the lowest energy point of the adsorption structure at different adsorption distances. In Fig. 4.25, when PCBM adsorbs to the MA+ site (see Fig. 4.24b) and the adsorption distance is 3.708 Å, the structure has the lowest energy. With structure optimization, the atom position and bond length change as shown by the red arrow in Fig. 4.24a. Comparing it with the clean MAPbI3 surface, the Pb and I atoms in the first layer move down 0.017-0.03 Å and 0.3550.437 Å along the Z-axis, respectively. The C-N atomic bond length in the second layer reduces from 1.510 Å to 1.489 Å. The I and Pb atoms in the third layer move slightly. Also, the C–C bond length in the PCBM molecular cage increased slightly, and the C-C-C angle connecting the two branches to C60 increases by 3 . This is

a

Vacuum

b

Realxed

surface layer1 layer2 layer3 layer4

Fixed

layer5 layer6 layer7 +

MA

C

I

H

Pb

O

adsorption sites: 1 above-I 2 above-Pb

z

3 bridge x

4 above-MA+

Fig. 4.24 Adsorption structure (a) and sites (b) of PCBM on MAPbI3 surface. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

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–1693.85

–1693.86 Total energy/eV

Fig. 4.25 Total energy at different adsorption distances. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

117

–1693.87

–1693.88

–1693.89 (3.708,-1693.895)

–1693.90 3.50 3.55 3.60 3.65 3.70 3.75 3.80 3.85 3.90 Molecule - surface distace/Å

because PCBM molecules adsorbed on the surface cause periodic damage, also cause interaction between atoms.

4.6.2

Adsorption Energy of PCBM on MAPbI3 Surface

Adsorption energy determines the stability of the adsorption system [80]. The adsorption energy of PCBM molecule on the surface of MAPbI3 can be calculated by Eq. (4.4):   E ¼ E atom=MAPbI3ð100Þ  E MAPbI3ð100Þ  NE atom =N

ð4:4Þ

In the above expression, Eatom/MAPbI3 (100) and EMAPbI3 (100) represent the total energy of the system before and after adsorption, respectively. Eatom represents the total energy of PCBM, and N (where N ¼ 1) represents the number of molecules adsorbed atoms. The calculated adsorption energy is 0.87 eV/PCBM molecule. It is a large adsorption binding capacity and the PCBM molecule has good stability on the surface” (Copyright [18] with permission).

4.6.3

Electronic Properties of PCBM-Adsorbed MAPbI3 Surface Model

Figure 4.26 shows the total DOS and (local density of states) LDOS of the adsorption surface. The bandgap is about 0.9 eV when the surface adsorbs PCBM, and the system shows semiconductor characteristics. Moreover, no new electronic state is generated near the Fermi level 0.3 to 0.2 eV, which is conducive to generate photogenerated carriers. The new surface state has a deleterious effect on the solar

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A- Total B- Layer1 C- Layer2 D- Layer3 E- Layer4 F- Layer5

200

100

DOS

150

50 0 A B C D E F –5

–4

–3

–2

–1

0

1

2

3

4

5

Energy/eV

Fig. 4.26 The total DOS (a) and the LDOS (b–f) of the PCBM-adsorbed MAPbI3 surface. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

cell’s filling factor and photovoltaic conversion efficiency. Comparing the LDOS of the first layer MAPbI3 (100) in Fig. 4.26b and the LDOS of the fifth layer in Fig. 4.26f (property is similar to the body), no new electrons appear near the Fermi level state. In summary, PCBM is a material suitable for electronic transmission materials. By calculating the differential charge, the charge transfer of the system is visually displayed. The calculation by Eq. (4.5) is as follows [81]: Δρ ¼ ρMAPbI3=PCBM  ρMAPbI3  ρPCBM

ð4:5Þ

where ρMAPbI3/PCBM is the total charge density of the structure, while ρMAPbI3 and ρPCBM are the charge density of MAPbI3 or PCBM separately in the same adsorption system. In other words, when calculating the charge density of MAPbI3, the PCBM in the model will be deleted as a vacuum, and vice versa. The differential charge of this system adsorbed by PCBM is shown in Fig. 4.27, and green and yellow indicate the decreasing (negative) or increasing (positive) charge density, respectively. The right part of Fig. 4.27 is the charge transfer distribution, and the left part is partially enlarged. The atoms closer to the surface have a “higher electron density, and the farther away from the surface, the smaller the electron transfer. The redistributed electrons promote the bonding. Specifically, for PCBM molecules, the number of electrons of C atoms and H atoms near I atoms decreases, while the number of electrons of C atoms and H atoms near Pb atoms increases. In the first layer of MAPbI3, the number of electrons on both I and Pb atoms increases. This difference in electron transfer also indicates the good bonding ability between the PCBM and MAPbI3 atoms” (Copyright [18] with permission).

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Fig. 4.27 The differential charge of the adsorbed surface. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

4.6.4

Optical Absorption Properties of PCBM-Adsorbed MAPbI3 Surface

The absorption coefficient spectra of the bulk MAPbI3, MAPbI3 (100), and “MAPbI3 surface adsorbed by PCBM are given in Fig. 4.28. There are optical band gaps on all model surfaces. When the photon energy value is at 3.75 eV, the absorption coefficient of the bulk MAPbI3 reaches its maximum peak, and its absorption edge is 1.85 eV. In the photon energy range of 1.5-2.5 eV, the absorption coefficients of the MAPbI3 (100) and the adsorbed surface have a consistent change trend. But the MAPbI3 surface adsorbed by PCBM shows a lower absorption coefficient when the photon energy is less than 3.0 eV. It shows a higher light absorption intensity in the energy range of 3.15–5.0 eV.” (Copyright [18] with permission). Our work provides additional information for experimental research and guides the design of experimental research, and may help design the efficiency of perovskite-type solar cells to seek effective control methods.

4.7 4.7.1

PEA+-Adsorbed MAPbI3 Surface Structure of PEA+-Adsorbed Surface

PEA+-adsorbed MAPbI3 (110) surface contains 84 atoms in total (Fig. 4.29). Passivated H atoms were added to the left of the model to eliminate the effects of dangling bonds of surface atoms. We added a 10 Å vacuum layer to avoid the interaction for the periodic unit cell. The single-point energy calculation method obtains the optimal adsorption distance of the passivating agent (the distance from the N atom in PEA+ to the I atom in the center of the MAPbI3 surface), and then we determine the most stable adsorption structure. The relationship between the total

120 1.2x104

Naked MAPbI3 surface PCBM-adsorbed MAPbI3 surface

1.0x104

MAPbI3 bulk

Absorption/cm–1

Fig. 4.28 The absorption coefficient spectrum of the MAPbI3 bulk, clean MAPbI3 surface, and PCBM-adsorbed MAPbI3 surface. (Ref. [18] with permission. Copyright IOP Publishing Ltd)

J.-P. An et al.

4

8.0x10

6.0x104 4.0x104 2.0x104 0.0 0

1

2

3 Energy/eV

4

5

Fig. 4.29 The structure of PEA+-adsorbed MAPbI3 (110) and the changes in the positions of atoms after optimization. (Ref. [82] with permission. Copyright IOP Publishing Ltd)

energy and the adsorption distance is shown in Fig. 4.30. When PEA+ is 3.393 Å from the center I atom, the total energy of the entire system is 389.377 eV, which is the lowest. In the optimization of the model shown in Fig. 4.30, the first and second layers of atoms near the surface of MAPbI3 (110) and PEA+ are relaxed, and the remaining layers of atoms away from the surface of MAPbI3 (110) are fixed to simulate the bulk. “The red arrow in Fig. 4.30 shows the direction of movement of the atomic position after the relaxation. In the first layer, the I atoms close to the adsorption surface approach the center, while the I atoms far away from the surface move in the opposite direction. The values of movement are 0.204 Å for the former and 0.022 Å for the latter. Also, the dPb-I bond lengths are 3.160 Å and 3.179 Å, which are shortened by 0.025 Å and 0.007 Å compared with before optimization. Compared with the clean surface, I atoms move in the opposite direction, relating to the adsorption of the passivating agent on the surface. In the second layer, the I and

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Fig. 4.30 The total energy change of the model under different adsorption distance. (Ref. [82] with permission. Copyright IOP Publishing Ltd)

Pb atoms moved by 0.057 Å and 0.008 Å along the Z-axis in the negative and positive directions, respectively. The average bond length of dPb-I formed is 3.1755 Å. The changes of bond angle and bond length in the MA+ cation in this layer are greater than the corresponding changes in the first layer. After the relaxation, the bond length (angle) of PEA+ also increased or decreased. The most significant is that the C–N bond length near the surface increases by 0.018 Å, and the N-C-C bond angle increases by 2.078 ” (Copyright [18] with permission). For the MAPbI3 part, the Pb atoms, cations, and I atoms at the central position are adsorbed by the passivating agent PEA+, and a larger movement will occur. Due to the periodic destruction of the lattice structure, the atoms closer to the surface move more significantly during relaxation.

4.7.2

Adsorption Energy of PEA+ on MAPbI3 Surface

The adsorption energy of PEA+ on the surface of MAPbI3 (110) characterizes the stability of the adsorbed substance on the surface (Eq. 4.4) [80], which is 3.898 eV. The adsorption binding shows that the adsorbent has good stability on this surface.

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Electronic Properties of PEA+-Adsorbed MAPbI3 Surface Model

According to the density of states, the passivation of the electronic states on the surface of MAPbI3 (110) by PEA+ can be analyzed. Figure 4.31a shows the total density of states (DOS) of the adsorption system. PEA+ has greatly reduced electronic states present on the surface of MAPbI3, and its DOS value has decreased from 11.5 to about 6.5. Also, PEA+ has the function of adjusting the bandgap while maintaining the semiconductor characteristics of MAPbI3. Compared with the band gap value of the clean MAPbI3 (110) surface, the surface value of adsorbed PEA+ will be significantly reduced. The adsorption of PEA+ does not introduce a new surface state, and the passivation effect is good as shown in Fig. 4.31b. Comparing the local density of clean surfaces, it can be seen from Fig. 4.31c–f that the addition of adsorbed molecules makes the first to fourth layers have no peak value nearby the Fermi. In particular, the passivation effect of PEA+ on the first and third layers is the most obvious. As mentioned earlier, because the bottom passivation H introduces some electronic states, the fifth layer of atoms still has small peaks near the Fermi level 0.5 to 0.3 eV. Figure 4.31i–n is the partial wave state densities of C, N, I, and Pb atoms near the surface and PEA+. Combining with Fig. 4.31i, l and n, at 1.5–2.5 eV from the Fermi level, the C-2p electron orbit in PEA+, and the N-2s electron orbit in MAPbI3, Pb-6p electron orbit exists hybridization. Figure 4.31i, m shows that there are hybridizations between the C-2p electron orbit in PEA+ and the I-5p electron orbit in MAPbI3. The passivation effect of the adsorption of PEA+ on the surface of MAPbI3 (110) is good on its surface state, and the electron orbital hybridization of atoms near the surface also promotes the interaction between the atoms. Figure 4.32 shows the differential charge of PEA+ adsorbed on the surface of MAPbI3 (110). The decrease (negative) or increase (positive) of the charge density is

Fig. 4.31 Adsorption structure of PEA+-MAPbI3: total state density (a), local state density of PEA+ (b), local state density maps of the first to fifth layers (e–g); localization of the passivated H layer domain state density (h); C, N, atomic partial wave state density (i–j) of PEA+ and C, N, I, Pb atom partial wave state density (k–n) at the surface. (Ref. [82] with permission. Copyright IOP Publishing Ltd)

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Fig. 4.32 Differential charge of PEA+-MAPbI3 adsorption structure. (Ref. [82] with permission. Copyright IOP Publishing Ltd)

represented by cyan and yellow, respectively. Figure 4.32a shows the charge transfer as a whole, and Fig. 4.32b is a partially enlarged view. As can be seen from Fig. 4.32, charge transfer occurs in the entire system, but charge transfer between atoms near the surface is more obvious. The farther from the surface, the less the electron transfer. The electron density at the surface is the highest, and the charge is transferred to both sides of the surface. Electron redistribution, which promotes the bonding between the atoms on the surface, indicates that the interaction between the atoms on the surface is strong. According to Fig. 4.32b, in the first layer of MAPbI3, the number of electrons on the I atom increases, and the number of electrons on the Pb atom decreases. For the PEA+ part, the number of electrons of C atoms near the surface decreases, while the number of electrons on N atoms increases.

4.8 4.8.1

Interfaces in Perovskite Solar Cells Materials: MAPbI3/ WZ-ZnO Interface Interface Structure

The interface model of WZ-ZnO (100)/MAPbI3 (112) is shown in “Fig. 4.33. The interface is composed of two parts, on the left are six monoatomic layers in the WZ-ZnO (100) plane, and on the right are six diatomic layers of the MAPbI3 (112) plane” ([84] with permission). To avoid the influence of surface dangling bonds on the interface performance, 12 and 4 passivation H were added on both sides of WZ-ZnO (100) and MAPbI3 (100) surfaces, respectively. Moreover, a ¼ 10.69 Å, b ¼ 9.35 Å, and c ¼ 51.14 Å are the lattice parameters of the WZ-ZnO (100)/ MAPbI3 (112) interface. “There are 36 Zn atoms, 36 O atoms, 6 N atoms, 6 C atoms,

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Fig. 4.33 WZ-ZnO (100)/MAPbI3 (112) interface model. (Ref. [83] with permission. Copyright Springer Nature Group)

Fig. 4.34 The total energy of the interface varies with the interface distance. (Ref. [83] with permission. Copyright Springer Nature Group)

6 Pb atoms, 18 I atoms, 36 H atoms, and 16 passivated H atoms in the interface model. To eliminate the adverse effects caused by the periodic boundary conditions of atoms at the surface, a vacuum layer of 30 Å was added to the model” ([84] with permission). Also, we determined a reasonable interface distance by a single-point calculation. As shown in Fig. 4.34, when the interface distance is 2.582 Å, the model has the lowest total energy of 636.05 eV. The interface binding energy can measure the stability of the WZ-ZnO (100)/ MAPbI3 (112) interface. We use Eq. (4.4) to calculate the interface binding energy. The interface binding energy is 0.164 J/m2. This result is smaller than the values of other systems in the literature, such as WZ-ZnO/CdS (0.61 J/m2) [85], Cu2ZnSnS4/ WZ-ZnO (0.21 J/m2) [86], and WZ-CuInS2/WZ-CdS (0.68 J/m2) [87]. This shows that WZ-ZnO/MAPbI3 is a weakly bonded interface.

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4.8.2

125

Electronic Properties

The electrical properties of the WZ-ZnO (100)/MAPbI3 (112) interface were investigated, specifically, the TDOS, PDOS, the LDOS, and differential charge density of the interface. The TDOS of the WZ-ZnO (100)/MAPbI3 (112) interface is shown in Fig. 4.35a; it is easy to see that there is a new electronic state near the Fermi level. This is because the formation of the heterojunction causes the periodical structure of the crystal to be destroyed so that an additional electron energy level is generated between the conduction band and the valence band, which is also called the interface state. Figure 4.35b, c is the LDOS of the fifth and first layers of the WZ-ZnO (100) plane. Because the fifth layer is far away from the interface, it is less affected by periodicity and dangling bonds. And its properties are closest to the bulk phase, so it is used to analogize the bulk phase WZ-ZnO. From the LDOS in the first layer of the WZ-ZnO (100) plane, a new energy density state is found at 0.5 eV to 0.3 eV near the Fermi level. Figure 4.35d, e is the LDOS of the fifth and first layers of the MAPbI3 (112) plane. Similarly, we use the fifth layer away from the interface to analogize the bulk phase. After a comparative analysis, “we found that new electronic states also appeared at the first layer of MAPbI3 (112), from 1.0 eV to 0 eV near the Fermi level. The above analysis shows that the first layers of WZ-ZnO (100)

Fig. 4.35 The total density of states of WZ-ZnO (100)/MAPbI3 (112) interface (a); the local density of states of WZ-ZnO layer 5 (b) and layer 1 (c); the local density of states of MAPbI3 layer 5 (d) and layer 1 (e); the partial density of states of Zn (f), O (g), CH3NH3 (h), I (i), Pb (j) on the interface layer 1. (Ref. [83] with permission. Copyright Springer Nature Group)

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and MAPbI3 (112) have interface states at 1.0 eV to 0 eV and 0.5 eV to 0.3 eV, respectively” ([88] with permission). The PDOS of the interface system is helpful for further understanding of the WZ-ZnO (100)/MAPbI3 (112) interface. Figure 4.35f–j is the PDOS of Zn, O, I, Pb, CH3NH3 atoms at this interface. Figure 4.35f, g shows those near the Fermi level 0.5 eV to 0.3 eV, in which electronic states appear on the WZ-ZnO (100) surface side mainly from O-2p orbital contributions. Also, the orbital hybridization phenomenon exists in the Zn-3d and O-2p electron orbits at 6 eV to 4 eV near the Fermi level. Figure 4.35h, j shows that near the Fermi level, new electronic states appear on the MAPbI3 (112) surface side mainly from the contributions of I-5p and Pb-6s electron orbitals. According to Fig. 4.35f, j, at the Fermi level, 3 eV to 0 eV, the O-2p electron orbit, I-5p electron orbit, Pb-6s electron orbit, and Pb-6p electron orbit on the MAPbI3 | WZ-ZnO interface also had orbital hybridization. The PDOS of all atoms contributing to the new electronic state at the interface is shown in Fig. 4.36, which is the partial wave state density of Zn1-Zn6, Ox (x ¼ 1–6), and I1-I3 atoms. From Fig. 4.35c, f, it can be seen that from 1.0 eV to 3.0 eV from the Fermi level, the new electronic state on the surface of WZ-ZnO (100) is contributed by Zn atoms on the first layer. Combined with the analysis of Fig. 4.36a–f, it is mainly derived from the 4s orbital of Zn2 atom, Zn4 atom, and Zn6 atom, and this is because these three atoms are affected by the H atom in

Fig. 4.36 The partial density of states of Zn1 (a), Zn2 (b), Zn3 (c), Zn4 (d), Zn5 (e), Zn6 (f), Ox (x ¼ 1–6) (g) atoms on the WZ-ZnO layer 1, I1 (h), I2 (i), I3 (j) atoms on the MAPbI3 layer 1. (Ref. [83] with permission. Copyright Springer Nature Group)

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Fig. 4.37 The density charge difference (a) and partially enlarged view (b) of the WZ-ZnO (100)/ MAPbI3 (112) interface. (Ref. [83] with permission. Copyright Springer Nature Group)

CH3NH3 ion [89]. From Fig. 4.36g, it is found that the contribution of six O atoms at different positions at the interface to the interface state is the same. Figure 4.36h–j clearly shows that the 5p orbital of the I3 atom contributes more to the interface state, while the 5p orbital of the I1 and I2 atoms contributes to the interface state on a comparable level. To study the electronic properties of the WZ-ZnO (100)/MAPbI3 (112) interface more comprehensively, we calculated the differential charge density of the interface system according to Eq. (4.5). Figure 4.37a qualitatively describes the redistribution of charges in the WZ-ZnO (100)/MAPbI3 (112) interface, and Fig. 4.37b is a partially enlarged view. We use yellow to indicate electron increase and bluegreen to indicate electron loss. From Fig. 4.37, it is found that the charge transfers at the interface will promote the bonding near the interface. Moreover, the closer the atoms to the interface, the greater the degree of charge transfer.

4.8.3

Optical Properties

As a material of the light absorption layer, MAPbI3 is a necessary one to understand the optical properties of its interface. If the energy of the photon is greater than the bandgap of the material, the valence band electrons are excited to the conduction band by absorbing the photon, and light absorption may occur. From Fig. 4.38, “the optical band gap of bulk MAPbI3 is about 1.81 eV, which is close to the reference values 1.62 eV [90] and 1.65 eV [91]. The absorption edge of bulk MAPbI3 is calculated to be approximately 1.6 eV, which is higher than the experimental value of 1.5 eV and the calculated value based on SOC-GW ( n > 9) vanadium oxides, exhibit valences of 4+, +5, and +3, +4 respectively. For the

Carbon composite; hierarchical structure Carbon-based composite; hierarchical structure

Graphene-based composite and special morphology

Carbon composite and special morphology Graphene coated and special morphology Carbon composite and special morphology Carbon composite and super hierarchical

Dandelion-like V2O3/C V2O3@NSCNF

V2O3/NG

Porous shuttle-like V2O3/C G-V2O3

Hierarchical porous metallic V2O3@C

V2O3@PNCNFs

V2O3/rGO

V2O3/C

Strategy Carbon composite and special morphology Carbon composite and special morphology Graphene-based composite

Sample V2O3@C HS

Table 13.4 Some applications of V2O3 in this field

Cathode/ AZIBs

Li-S batteries Anode/KIBs

Anode of LIBs and SIBs Anode/SIBs

Anode/LIBs

Anode/LIBs

Anode/LIBs

Anode/LIBs

Application Anode/LIBs

Electrospinning and heat treatment Hydrothermal calcination

Plasma-enhanced CVD

Calcination of MOFs

Hydrothermal and annealing Hydrothermal; electrospinning and calcination Hydrothermal and calcination

Hydrothermal and annealing Hydrothermal

Method Hydrothermal

181mAh/g@2A/g, 68% retention after 1000 cycles 1000mAh/g@2C/g, ~50% retention after 1000 cycles ~230mAh/g@50 mA/g, 95.8% retention after 500 cycles ~200mAh/g@4A/g, 90% retention after 4000 cycles

435mAh/[email protected]/g, after 250 cycles (LIB); 154mAh/[email protected]/g after 500 cycles (SIB)

Performance 583mAh/g@2A/g no obvious decrease after 800 cycles 732mAh/g@100 mA/g, 95% retention after 100 cycles 675mAh/[email protected]/g, no obvious decrease after 300 cycles 315mAh/g@2A/g, 94% retention after 1000 cycles 800mAh/[email protected]/g; no obvious decrease after 100 cycles

[104]

[103]

[102]

[101]

[100]

[99]

[98]

[97]

[83]

References [96]

13 The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . . 369

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Magnéli vanadium oxides, there are few reports on the application of energy storage and electrocatalysis. So in this section, we will introduce Wadsley vanadium oxides in the aspect. Wadsley phase vanadium oxide is a series of mixed-valence vanadium oxides. There are mainly three types, namely V3O7, V4O9, and V6O13. All of them are single- and double-layered compounds. V3O7 and V6O13 have some reports in the field of energy storage such as supercapacitors [105, 106] and metal-ion batteries [107], while V4O9 is rarely studied [108]. Its main application strategy in electrochemistry is to prepare Wadsley phase vanadium oxide with certain morphology or pore structure by hydrothermal method (partial post-heat treatment) [105] or to combine it with carbon carrier for application [109].

13.2.1 V3O7 Li et al. [110] fabricated a well crystalline V3O7H2O by a high-temperature mixing method (HTMM). To compare with V3O7H2O, two types of VO2 were synthesized in the same way. All the vanadium oxides were applied in cathode material of lithium-ion batteries and the test results showed that V3O7H2O has the best performance: 250 mAh/g at a current density of 100 mA/g and 80% retention after 100 cycles test. The authors also discussed the reason for the good performance of V3O7H2O in the view of crystallography. According to the results of XRD refinement, the crystal model was obtained and suggested that well crystalline was benefited in the electrical conductivity of the material. With the presence of the crystal water, the distance of the crystal plane became larger, the lithium-ion could migrate more freely between the interlayer (V3O7H2O), compared with the diffusion in tunnels (two types of VO2), and it is conducive to the lithium-ion deintercalation performance. After that, Liu et al. [111] further carried out with surface carbon coating on the material and obtained the V3O7H2O@C by the same method and found that this material could reverently insert 3.3 moles of lithium ions per mole and had stable cycling performance under different rate conditions. Shen et al. [109] composited V3O7H2O with rGO by hydrothermal method and obtained rod-shaped nanomaterials (as shown in Fig. 13.4). This material was applied in the cathode material of aqueous zinc-ion batteries and the Zn anode also was decorated by rGO. The initial capacity was about 250 mAh/g at a current density of 300 mA/g, the retention was 79% after 1000 cycles. The authors suggested that the electrode combined with rGO could not only inhibit the dendrite growth of the anode but also had better ion deintercalation performance.

13.2.2 V6O13 Like V3O7 and V4O9, V6O13 is a mixed-valence vanadium oxide of +4 and + 5. It belongs to a monoclinic crystal system and also has a single- to double-layered

13

The Application in Energy Storage and Electrocatalyst of Vanadium (Based). . .

371

Fig. 13.4 (a) XRD data of V3O7H2O/rGO. (b) SEM image of V3O7H2O/rGO nanobelts. (c) Typical diffraction pattern and (d) representative TEM image of a single nanoribbon, revealing a width of 100–200 nm. (e) Lattice-resolved HRTEM image [from the region highlighted in (d)]. (f) SEM image and corresponding elemental mapping of (g) carbon, (h) oxygen, and (i) vanadium [109]

alternating layered structure, so it also can have the performance of deintercalation metal ions [6]. The lithium-ion deintercalation reaction is as follows: V6 O13 þ x Liþ þ x e ¼ Lix V6 O13 : Since V6O13 has the highest oxygen potential in the series of vanadium oxide (VO2.16), more lithium ions can be embedded. In the earlier study, Murphy et al. [112] obtained the maximum embedded lithium-ion content of 8 mol per mole of V6O13, with a theoretical specific capacity of about 420 mAh/g, and energy density of 890 Wh/kg was determined [113]. They are higher than the current commercial lithium-ion batteries cathode materials [33], so V6O13 is also one of the potential cathode materials of lithium-ion batteries [114]. Although V6O13 has great application potential, there are still some obstacles for the material which restrict its application. Firstly, V6O13 has a low MIT temperature [115] and it is metallic at room temperature [116]. During the discharge process, the material would be embedded with lithium ions and forms LixV6O13, which greatly reduces the electrical conductivity of the material and thus limits its performance. Secondly, in the charging and discharging process, due to the removal of lithium

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ions, the material is expected to undergo an irreversible phase change but also lattice distortion. These changes would contribute to a volume change of the cell up to 13.6% [117]. This volume effect would also degrade the performance of the material. Therefore, in the present research, the morphology control, doping and surface modification of the material, has become the main strategy for the electrochemical application of the material. Tian et al. [118] prepared a flower-like NH4V4O10 precursor by hydrothermal method, and then prepared ultra-thin pre-lithified V6O13 nanosheet through low-temperature exfoliation and chemical lithium intercalation. This material was applied to the cathode material of LIBs. It could deliver the first lap capacity of 300 mA/h (100 mA/g), and reach 180mAh/g at the current density of 1000 mA/g. The retention was 98% after 150 cycles. Through the electrochemical impedance and in situ conductance test, it was found that the pre-lithiated V6O13 (12 Ω) can enhance the electrical conductivity compared with the non-pre-lithiated material (27 Ω), thus reducing the charge transfer resistance and enhancing the material performance. Ding et al. [119] synthesized a 3D nano-textile-structured V6O13 by the simple liquid redox self-assemble method. With the vanadyl(IV) sulfate [VOSO4 (H2O) where 0  x  6] as oxidant used in varying amounts, different diameter of V6O13 nano-textiles was obtained and used as cathode materials of lithium-ion batteries. The first cycle capacity was 360 mAh/g at the current density of 20 mA/g, and it delivered a capacity of 215 and 134 mAh/g at the current density of 200 mA/g and 500 mA/g, respectively. The retention was about 95% after 100 laps, and the energy density could reach about 780 Wh/kg. The authors suggested the reason for this performance was that the nano-textiles have a conductive structure, which could enhance electron and ion transport. He et al. [120] prepared a Cu-doped V6O13 by a simple hydrothermal method and obtained a belt-like CuxV5-xO13. The test result showed that the Cu0.1V5.9O13 has the best performance: the first cycle capacity was about 380 mAh/g at the rate of 0.1C and 82% retention after 50 cycles. Lai et al. [121] synthesized a flower-like V6O13 which was assembled by ultrathin nanosheet and used as the cathode material of aqueous zinc-ion batteries. It showed an outstanding performance: 395 mAh/g at the current density of 100 mA/g and 300 mAh/g at 5 A/g, after 1000 times charge and discharge; the retention is 87%. The authors thought that crystal water could enhance the intercalating chemical process. Other applications of Wadsley phase vanadium oxides in this field are shown in Table 13.5.

13.3

Vanadium-Based Oxides

Vanadium oxides also are with the adsorption of some species, such as OH[17], etc. Thus they have the potential application of electrocatalysis; however, the catalytic process is complex which includes the adsorption, reduction/oxidation, desorption steps, and so on. So it is difficult for the single vanadium oxides to be

Cathode/LIBs Cathode of AZIBs Cathode of AZIBs

Graphene composite and special morphology Special morphology

Special morphology

Special morphology

Special morphology

Carbon coated and special morphology Mn-doped

– –

V3O7H2O

V6O13NS

β-V6O13

CNPs-V6O13

Mn0.02V5.98O13

V6O13 V6O13

Cathode/LIBs

Cathode/LIBs

Cathode/LIBs

Microwave solvothermal

The cathode of ZIBs and AZIBs Cathode of AZIBs

Hydrothermal Hydrothermal and annealing

Hydrothermal

Hydrothermal

Hydrothermal

In situ electrochemical oxidation Hydrothermal

Hydrothermal

Method High-temperature mixing hydrothermal method Hydrothermal

Cathode of AZIBs

Cathode/LIBs

V3O7H2O/ rGO V3O7H2O

H2V3O8

Application Cathode/LIBs

Strategy Carbon composite and special morphology Special morphology

Sample V3O7H2O@C

Table 13.5 Some applications of Wadsley phase vanadium oxides in this field Performance 262mAh/[email protected]/g, 94% retention after 100 cycles 250mAh/[email protected] A/g, 90% retention after 50 cycles 250mAh/[email protected]/g 79% retention after 1000 cycles 275mAh/g@3A/g, 80% retention after 200 cycles (AZIB) 171mAh/g@5A/g, 85% retention after 1000 cycles 330mAh/g@42 mA/g, 70% retention after 50 cycles 335mAh/g@50 mA/g, 73% retention after 75 cycles 288mAh/g@1C, 82% retention after 100 cycles 350mAh/[email protected], 95% retention after 50 cycles 206mAh/g@10A/g after 3000 cycles 245mAh/g@4A/g, no obvious decrease after 2000 cycles

[129] [130]

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used as an electrocatalyst. They are usually combined with other transition metal (hydro) oxides or noble metals to form the vanadium-based oxides which could be applied in energy electrocatalysis [131–133]. This is what we will introduce in this section.

13.3.1 Electrocatalysis As transition metal oxides, vanadium oxide and vanadium-based oxides also have some reports on energy electrocatalysis application; in this section, we will introduce the application in water splitting [131], oxygen reduction reaction (ORR) [132], and small molecule fuel’s oxidation catalysis [133] of vanadium-based oxide briefly. Shi et al. [131] compared the reports of VS2 and NiV-layered double hydroxide (NiVLDH) and suggested that the oxygen-group compounds of vanadium had certain effects on hydrogen evolution. Then the VOOH catalyst was synthesized by the hydrothermal method (as shown in Fig. 13.5). The catalyst was used as an electrocatalyst for water splitting under the alkaline system. The excellent performance showed that the HER overpotential was 164 mV; OER was 270 mV, and the overall reaction voltage was 1.62 V. The sustainable test was kept for 25 h, which implied excellent durability. Mu et al. [132] prepared spinel CoV2O4 by cation exchange method and used an oxygen reduction reaction electrocatalyst. The performance not only showed that under alkaline conditions, the half-wave potential could reach 0.82 V vs. RHE, which is better than cobalt oxide CoO and Co3O4 alone, almost comparable to

Fig. 13.5 VOOH hollow nanospheres (sample III) for overall water splitting in 1 M KOH. (a) Schematic of the overall water splitting setup. (b) Polarization curve of the VOOH || VOOH catalyst along with Ni foam || Ni foam and Pt/C || RuO2 for comparison. Inset shows the stability test at a constant current density of 50 mAcm2. Catalyst loading: about 0.8 mgcm2 [131]

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commercial carbonized platinum, but also had better durability. Subsequently, it was applied to the cathode catalyst of zinc-air batteries. It was found that the power density could reach 380 mW/cm2, and the voltage was stable after continuous testing for 150 h under the current density of 20 mA/ cm2. The authors believed that this spinel structure could enhance the electrical conductivity of the V-V chain, thus significantly improve the activity of cobalt oxide. Amiddedin et al. [134] synthesized Pt/VOx-MWCNTs by adding the amorphous VOx to Pt-MWCNTs by microwave-assist method and applied in the electrocatalyst of methanol oxidation reaction (MOR) in the acidic system. Compared with the Pt-MWCNTs, this catalyst improved the activity of MOR obviously. The authors suggested that the existence of VOx could help to generate more -OH, and thus improved oxidation reaction of absorbed CO and enhanced the performance of catalyst eventually. Other applications of vanadium-based oxides in electrocatalysis are shown in Table 13.6.

13.4

Summary and Outlook

This chapter mainly introduced the application of vanadium (based) oxides in energy storage and electrocatalysis, mainly focusing on metal-ion batteries and water splitting, oxygen reduction reaction, and catalytic oxidation of small molecular fuels. At present, there are many reports on the preparation, performance, and energy storage or catalytic mechanism of most vanadium oxides. However, there are also some deficiencies, which are summarized as follows. The application of vanadium oxide could be catalogued in two areas. Structural, where the vanadium oxide usually possesses layered structure which intercalates metal ions. Under these conditions, the modified oxide could be used in applications related to ion insertion such as metal-ion batteries or supercapacitors. The other areas are related to the materials adsorption characteristics. Here, the adsorption of vanadium oxide to species, such as -OH, are evaluated in terms of material electrochemistry and include: 1. Prepare vanadium (based) oxides with special morphology or hierarchical structure to increase the activity area. 2. Improve the electrical conductivity of prepared materials: generally, vanadium oxide is a semiconductor or poor electrical conductivity at room temperature, which restrict its performance. Therefore, researchers often composite it with carbon- or graphene-based materials and dope with other elements to improve the electrical conductivity. 3. Composite with other metal or metal oxides and synthesize the vanadium-based oxide electrocatalyst like NiVLDH, etc. using the synergistic effect between elements or compounds. Moreover, in the research progress of vanadium oxide in this field, there are still the following shortcomings that have not been addressed:

V2O5/Ni (OH)2 Pt/V2O5-C

CoV1.5Fe0.5O4

Pt loaded

Ni(OH)2 composite

Special morphology Ni, NiS composite and special morphology Ni composite, hierarchical structure NiS composite, hierarchical structure Fe doped

V10O24nH2O VOx / Ni3S2@NF Ni/V2O3

VOx/NiS/Ni

Strategy Ni composite

Sample Ni/V2O3

Water splitting MOR (acidic)

OER, ORR

OER

Microwave assist, and reduction

Thermal decomposition Hydrothermal

Hydrothermal, calcination Hydrothermal

HER

HER OER

Method Hydrothermal, calcination Hydro-thermal Hydrothermal

Application HER

Table 13.6 Some applications of vanadium-based oxides electrocatalysis

η ¼ 330 mV (50 mA/cm2); Tafel slope ¼ 121 mV/dec @1MKOH η ~ 300 mV, Tafel slope ¼ 38 mV/dec (OER), E1/ 2 ~ 0.7 V vs. RHE η ~ 40 mV, Tafel slope ¼ 44 mV/dec (HER@1 M KOH); 1.59 V for water splitting@6 M KOH 17.4 mA/cm2 at forward sweep peak(Pt@C:12.25)

η ¼ 118 mV; Tafel slope ¼ 101 mV/dec @1MKOH η ¼ 358 mV (100 mA/cm2), Tafel slope ¼ 82 mV/dec @1MKOH η ¼ 22 mV (20 mA/cm2); @1MKOH

Performance η ¼ 44 mV; Tafel slope ¼ 38 mV/dec @1MKOH

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1. The phase transition problem of vanadium oxide which is used in electrode materials of metal-ion batteries has not been solved, which greatly restricts the large-scale application of vanadium oxides. 2. The influence of the change of valence of vanadium in vanadium oxide on catalysis when they are used for electrocatalysts. 3. The development of other mixed-valence vanadium oxides, such as V4O9 in Wadsley phase vanadium oxides, and Magnéli phase vanadium oxides are still rarely reported. In particular, a series of studies on Magnéli phase vanadium oxides are still limited to thermodynamic synthesis. The electrochemical applications for the preparation of special morphology have not yet been reported. Acknowledgments F. L. and Y. W. sincerely thank the National Nature Science Foundation of China (NSFC 52074016) for the financial support. J. L. Liu and S. B. would also like to acknowledge the support from the National Science Foundation (NSF-MRI, CBET 0821370) and R. Welch Foundation (AC-0006) from the Texas A&M University-Kingsville. Author Contribution F. L. and Y. W. collectively conceived this topic. J. L. Liu completed the literature survey and edited the manuscript. S. B. checked the data and the hkl assignment(s). Y. W. generated the data table and wrote the first draft of the paper. All authors reviewed the final edit and approved the final version of the manuscript.

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Yue Wang received his Master’s degree in physical chemistry at the Beijing University of Technology in 2020. He is currently a Ph.D. candidate under Prof. Fan Li’s supervise. His research interests focus on the thermodynamic, kinetic, synthesis, and application in energy electrocatalysis (MOR, ORR and related electrochemistry) of mixed-valence transition metal oxides, such as vanadium oxides, titanium oxides, tungsten oxides, and bandgap materials. He aims to study the thermodynamic, kinetic mechanism of synthesized these variable valence oxides and their applications in energy electrochemical.

Fan Li received his B.S. in Physical Chemistry in Metallurgy from the University of Science and Technology Beijing (USTB) in 1996. He received his M.E. in Chemical Technology from the Institute of Process Engineering, Chinese Academy of Sciences (IPE, CAS). He received his Ph.D. in Materials Science and Engineering from Royal Institute of Technology (KTH), Sweden 2007, and Physical Chemistry in Metallurgy from the University of Science and Technology Beijing (USTB) in 2008. He is a full professor at the Department of Chemistry and Chemical Engineering at Beijing University of Technology (BJUT). Dr. Li is focusing on the electrochemical properties of the transition metal oxides with mixed-valence states loaded Nobel metals catalyst, and applied them in polymer electrolyte membrane fuel cell. The preparation and characterization of nano-structured electro-catalyst with low content noble metals is his main study field. Sajid Bashir received his Ph.D. in Analytical Chemistry from the University of Warwick, England, in 2001. He was a postgraduate research associate at Cornell University to research in the field of plant genetics. Currently, he is a Full Professor at Texas A&M University-Kingsville (TAMUK) and a Faculty Fellow at the US Air Force. He has directed and participated in more than 20 projects supported by the Welch Foundation, TAMUK, Texas Workforce Commission, and US National Institute of Health. He has co-authored >80 book chapters and peer-reviewed journal articles. He is a fellow of the Royal Society of Chemistry, and Chartered Chemist & Chartered Scientists of Science Council. During his service in TAMUK, he trained more than 3000 students on both undergraduate and graduate levels. He created online courses and established safety training protocols in conjunction with Risk Management. Currently, he collaborated with local law enforcement as a consultant.

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Jingbo Louise Liu received her Ph.D. in Materials Science and Engineering from the University of Science and Technology Beijing in 2001. She is a Full Professor at Texas A&M UniversityKingsville (TAMUK) and focused on materials preparation, characterization, and applications. She is a Fellow of the Royal Society of Chemistry and a DEBI faculty fellow at the US Air Force Research Laboratory. She has authored and co-authored, books, book chapters, and peer-reviewed journal articles (>100). During her 12.5-year services in TAMUK, she taught >8700 students and trained about 150 students and scholars to conduct leading-edge research. She directed and/or participated in the projects (>40) supported by the NSF (US and China), NSERC (CANADA), ACS Petroleum Research Funds, and Department of Education as PI, Co-PI, and senior personnel. She was recently elected as the Division Chair-Elect of Energy and Fuels, American Chemical Society.

Chapter 14

Water-Stable Metal-Organic Frameworks for Water Adsorption Xuan Wang and Charles Lee

14.1

Fundamental Basics About Water-Stable MOFs

The study of metal-organic frameworks (MOFs) has boomed in the past two decades, and MOFs are such attractive advanced materials due to their crystallinity, porosity, tunability, and ease of functionality. In the earliest stages of development, many MOFs show significant instability in the presence of a small amount of water content, which is common in real-life applications. While the relative labile coordination between metal ions and organic linkers limits the stability of MOF, a strong electrostatic interaction resulted from high positive charge density (Z/r), like Ti4+, Zr4+, Al3+, Cr3+, Fe3+ ions, and carboxylate linkers, prevents the framework from the attack of any guest species. Among all the water-stable MOF series, the MIL series, the ZIF series, the UiO series, the CAU series, the PCN series, and the MOF-800 series stand out due to their robustness in the water environment. The underlying basics that the formation of a stable framework relies on strong coordination bonds aligns with the theory of Pearson’s Hard Soft Acid-Base (HSAB) (Fig. 14.1) [1, 2]. Synthetic cases of stable MOFs have been reported through direct synthesis and post-synthetic metathesis, in which the latter process involves the transformation from labile coordination to strong coordination by partial or complete replacement of the metal ions [3–5]. Another factor needing significant consideration for good stability is the rigidity of the framework. In general, an increasing length of organic linkers will decrease the rigidity of the framework and thus compromise the stability of the framework. Moreover, the decoration of hydrophobic functional groups on the surfaces or interfaces provides such immunity for water attacks [3, 6, 7]. More detailed discussions on the synthesis of stable MOFs could be found in several recent reviews [2, 3].

X. Wang (*) · C. Lee Texas A&M Higher Education at McAllen, Texas A&M University, McAllen, TX, USA e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_14

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Fig. 14.1 The hard-soft acid-base theory directs the construction of stable MOFs and some representative MOF examples

Stable MOFs have been widely used in all applications, namely gas sorption and separations, catalysis, energy generation, biomedical applications, and environmental science [8–18]. Water-stable MOFs represent a subclass of stable MOFs, in which the framework display resistance to water. Those MOFs advance other functional materials by their potentials in many practical applications especially in the presence of water [19, 20]. This chapter will introduce the structural fundamentals of the selected prototype water-stable MOF series and then present the development of the water adsorption using water-stable MOFs as adsorbents.

14.2

Prototypes of Water-Stable MOF Series

There has been an ever-growing effort to synthesize and utilize water-stable MOFs in their related applications, and many interesting examples of those MOFs have continuously been reported, especially through isoreticular chemistry or generation of composites displaying increased stability. In this section, several series of stable MOFs—the Materials of Institut Lavoisier (MIL) series, the zeolitic imidazolate framework (ZIF) series, the University of Oslo (UiO) series, the Christian-AlbrechtsUniversität zu Kiel, (CAU) series, the porous coordinate network (PCN) series, and the MOF-800 series—are selected, and some representative examples of each series that are of interest are presented in detail.

14.3

The MIL Series

The MIL series, developed by the Materials of Institut Lavoisier, is one novel category of water-stable MOFs. Credited to be one of the more diverse MOF series in existence, the discussion relating to the MIL series will be held to water-stable MIL MOFs, particularly Al-MIL-53, Fe-MIL-100, and MIL-101-NH2.

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14.3.1 MIL-53 Al-MIL-53 ([Al(OH)(COO)2]n) is a hydrothermally-synthesized water-stable MOF, formed using the BDC linkers and Al3+ metal units [21]. Constructed from the fusion of trans chains of octahedra units of AlO4(OH)2 and BDC linkers [22], this MOF maintains a nanoporous framework with a pore size of 8.5 Å (Fig. 14.2). Additionally, Al-MIL-53 demonstrates remarkable thermal stability, as its Tdecomp is approximately 500
C. Unlike other MOFs, Al-MIL-53 is novel in its ability to “breathe” after being exposed to hydration and dehydration experiments. As the diffusion of water enters and exits the MOF, the interactions between the corner-sharing octahedra units and the water molecules via hydrogen bonding will fluctuate in such a manner that it will reduce the angle size of the BDC-Al-BDC (defined in the study as angle a) by 44%, and follow by the breaking of the hydrogen bonds, which will revert to the original angle size. More detailed structural changes will be discussed in the next section. Al-MIL-53 has a Langmuir surface area, and a BET surface area is 1590 and 1181 m2 g1, respectively. Because of its novel design and stability, Al-MIL-53 has been used in a broad spectrum of water-based applications.

14.3.2 MIL-100 Another novel member of the MIL series is Fe-MIL-100 [23]. First synthesized in 2007, the MOF is the result of a chemical marriage between Fe(III) and benzene1,3,5-tricarboxylate (BTC) as the inorganic units [24]. Fe-MIL-100 crystallizes into a hybridized super tetrahedra cubic system. As an isostructural complement to Cr-MIL-100, it has the structural capabilities to stabilize into two types of cages, Fig. 14.2 Structure demonstrations of MIL-53 from the infinite chains of corner-sharing octahedral metal units (blue polyhedra) connected through BDC (Reprinted with permission from Millange et al. [22]. Copyright @ 2010 American Chemical Society)

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Fig. 14.3 Structure of Fe-MIL-100 consisted of trimers of the Fe SBUs (blue polyhedral) and H3BTC with two types of cages and pentagonal and hexagonal windows (Reprinted with permission from Wang et al. [24]. Copyright @ 2015 American Chemical Society)

each with varying apertures (ca. 25 and 29 Å) and varying window pore sizes (ca. 5.5 and 8.6 Å, respectively) (Fig. 14.3). The N2 adsorption experiments indicate this water-stable MOF has a Langmuir surface area greater than 2800 m2•g1. The thermal stability of Fe-MIL-100 is comparable, as its Tdecomp is approximately 270
C.

14.3.3 MIL-101 Additionally, Al-MIL-101-NH2, another member of the MIL series, has also attracted many interests in a variety of applications. This MOF is synthesized in solvothermal conditions by the addition of Al3+ ions and H2BDC-NH2 [23]. As a reticular structure to Cr-MIL-191 and Fe-MIL-101, the formation of Al-MIL-101NH2 requires aminoterephthalate ligands, while the previous experience using terephthalic acid as a linker yielded Al-MIL-53. It is important to note that the formation of Al-MIL-101-NH2 requires aminoterephthalate ligands, as previous experiments using terephthalic acid as a linker yielded Al-MIL-53. Furthermore, the utilization of these rigid linkers in conjunction with Al3+ octahedral clusters creates super-tetrahedral building units. Thus, this allows for the formation of two cage types, each with quasi-spherical mesoporous qualities [24]. One cage type, defined as a medium cavity, exhibits a 1.2 nm pentagonal window type and has 12 pentagonal faces, and the other cage type, named large cavity, exhibits a 1.6 nm hexagonal window type with 16 pentagonal face types. The MOF has thermal stability like its Cr-MIL-101reticular counterpart; its Tdecomp is approximately 650 K. The BET surface area for this MOF is approximately 2100 m2•g1.

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The ZIF Series

The ZIF series demonstrates an alternate method of utilizing the HSAB theory to develop water-stable MOFs. As previously mentioned, stable MOFs are usually a combination of either hard acids and hard bases or soft acids and soft bases. The ZIF series are generally constructed from a soft acidic metal ion, notably Zn2+, and soft basic azolate ligands, such as imidazolates and tetrazoles. ZIF-8, ZIF-68, ZIF-69, and ZIF-70 are chosen for our discussion of water-stable MOFs.

14.4.1 ZIF-8 By undergoing solvothermal reactions, ZIF-8 utilizes imidazolate linkers and Zn2+ ions to form a 3D structure maintaining a sod topology [25]. Furthermore, the MOF exhibits permanent porosity with a BET surface area of 1630 m2•g1 and a Langmuir surface area of 1810 m2•g1. ZIF-8 is also credited with having a comparable pore size of ~11.6 Å, along with a small aperture, with the calculated size of 3.4 Å [26] (Fig. 14.4). The water-stable MOF is also characterized to have excellent thermal stability with Tdecomp above 500
C. Additionally, when exposed to various environments, especially aqueous environments, ZIF-8 demonstrated remarkable water-stability by maintaining its structure in a heated water environment.

14.4.2 ZIF-68 Another subclass of the ZIF series is ZIF-68-70 [27]. ZIF-68 also utilizes solvothermal conditions as a means of fusing bIM ligands with Zn2+ ions. The resulting chemical marriage forms a gme topological 3D structure with its largest Fig. 14.4 The single crystal structure of ZIF-8 with large cages (in blue) (Reprinted with permission from Gallaba et al. [26]. Copyright @ 2016 American Chemical Society)

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Fig. 14.5 The network of ZIF-68 (Reprinted with permission from Van der et al. [28]. Copyright @ 2015 American Chemical Society)

cage exhibiting approximately 24 vertices [28] (Fig. 14.5). Additionally, ZIF-68 has a 10.2 Å pore size and 7.5 Å apertures. Concerning stability, the MOF has remarkable thermal and chemical stability. Specifically, the TGA analysis shows ZIF-68 has a Tdecomp of 390
C, and exposure to various environments, notably heated benzene, methanol, and water revealed that the MOF maintained its structure. The porosity of the MOF demonstrated a Langmuir surface area of 1220 m2•g1.

14.4.3 ZIF-69 In addition to ZIF-68, an isoreticular structure in the ZIF series is ZIF-69 [27]. This MOF shows similar characteristics to ZIF-68. Specifically, both share a similar topology (gme) and the number of vertices in the largest cage. They also share a Tdecomp of 390
C. However, in comparison to the pore size, ZIF-69 has a smaller pore size, measuring at 7.2 Å. Additionally, the MOF has a smaller Langmuir surface area of 1070 m2•g1. As previously addressed with ZIF-68, this MOF shares similar thermal and chemical stabilization qualities. Particularly, when exposed to the same environments and stimuli, the structural integrity of ZIF-69 was maintained.

14.4.4 ZIF-70 ZIF-70 is an isoreticular structure of ZIF-68 and ZIF-69 [27]. As addressed prior, many of the same qualities exhibited by the other two ZIF structures, including topology and the number of vertices present in the largest cage size, are exhibited in this structure. Specifically, some of the differences between them include pore size and Langmuir surface area, which are 15.9 Å and 1970 m2•g1, respectively.

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The UiO Series

Arguably one of the most recognizable series in water-stable MOF development is the UiO series, pioneered by the University of Oslo. The UiO series demonstrates the effectiveness of the HSAB theory; specifically, the UiO series enhances the coordination bond strength between the carboxylate-based ligands and the hard acid nature of the Zr4+ ions.

14.5.1 UiO-66 to UiO-68 This combination results in the formation of two incredibly stable frameworks, particularly in water environments. Since all three MOFs are reported with a similar structure as well as other properties, we will discuss them together. The UiO series, namely UiO-66, UiO-67, and UiO-68, can be assembled through Zr ions and ditopic acids with different lengths to form two types of cages (Fig. 14.6), the smaller deemed as a super-tetrahedron and the larger deemed as a super-octahedron [29]. UiO-66 is the classic representation of the entire series. UiO-66 (Zr6(p.3-O)4(μ3OH)4(COO)12) is a cubic close-packed MOF built from Zr4+ ions and 1,4-benzenedicarboxylate (BDC) units through solvothermal reactions. Notably, UiO-66 is characterized by its incredible stability, partly due to its 12-coordinated system. In the 3D structure, UiO-66 has a window opening of 6 Å, as well as a Langmuir surface area of 1187 m2•g1. Furthermore, as the most restrictive MOF in the UiO series, it was experimentally derived to have a temperature of decomposition (Tdecomp) of 540
C. UiO-67 utilizes 4,40 -biphenyldicarboxylate (BPDC) linkers and Zr4+ ions, whose use increases the Langmuir surface area to 3000 m2•g1. Because of the significantly larger surface area, UiO-67 boasts an 8 Å cage size. UiO-68 is constructed out of Zr4+ ions and terphenyldicarboxylate (TPDC) units. It

Fig. 14.6 Two types of cages in the UiO series: the super-tetrahedron (a) and super-octahedron (b) (Reprinted with permission from Cavka et al. [29]. Copyright @ 2008 American Chemical Society)

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has a Langmuir surface area of 4170 m2•g1, consequently allowing for an expanded cage size of 10 Å.

14.6

The CAU Series

Another class of representative MOFs is named after Christian Albrechts Universitat (CAU), the German university that housed the MOFs discoverers, the Stock group [30]. Among all the reported porous MOFs, our interests focus on CAU-3 and CAU-10.

14.6.1 CAU-3 The high-throughput synthesis of Al ions with nonfunctionalized H2BDC produces another framework, CAU-3, in a highly diluted basic environment [31]. Under the optimized synthetic condition, the researchers obtained another two isoreticular MOFs with H2BDC-NH2 and 2,6-naphthalene dicarboxylic acid (H2NDC). Overall, CAU-3 is an fcu net that is built from 12-connected [Al12(OCH3)24(COO)12] SBUs and BDC linkers. This network includes strongly distorted tetrahedral and octahedral cavities with a pore size of 10 and 11 Å, respectively. The BET surface area of the frameworks has been measured at 1250 m2•g1 (CAU-3-BDC-NH2), 1550 m2•g1 (CAU-3-BDC), and 2320 m2•g1 (CAU-3-NDC).

14.6.2 CAU-10 The Stock group later successfully obtained a series of isoreticular Al-based MOFs, classified as CAU-10-X (X ¼ H, CH3, OCH3, NO2, NH2, OH at the 5-position of the benzene ring of 1,3-dibenzo dicarboxylates) through high-throughput synthesis [32]. The CAU-10-X series are microporous frameworks with square-shaped, sinusoidal channels with a diameter of 7 Å.

14.7

The PCN Series

Porous Coordination Network (PCN) represents MOFs synthesized by the Zhou Group. One of their primary research focuses is to synthesize stable MOFs with large pores. Aligning the purposes of discussion for this chapter, PCN-222 (may also be referred to as MOF-545), PCN-224, PCN-228-230, PCN-250, PCN-333, PCN-601, and PCN-777 are reviewed below.

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14.7.1 PCN-222/MOF-545 PCN-222 ([Zr6(μ3-OH)8(OH)8(TCPP4)2]) [33], also reported as MOF-545 [34], is a solvothermally synthesized MOF utilizing metal-chelated meso-tetrakis (4-carboxypheny) porphyrin (M-TCPP) units and Zr4+ ions. Structurally, the MOF resembles 3D needle-like crystals with a BET surface area of 2220 m2•g1 and hexagonal channels with a diameter of 3.7 nm [35] (Fig. 14.7). Additionally, the MOF crystallizes in the P6/mmm space group. PCN-222 is incredibly stable in a variety of environments, especially in acidic environments, as demonstrated by an isostructural compound, PCN-222(Fe).

14.7.2 PCN-224 Following suit, PCN-224, [Zr6(μ3-OH)8(H2O/OH)12]2[TCPP-M]3n, is another water-stable MOF harnessing the abilities of square planar M-TCPP units and the six-connected Zr6 clusters through the linker-elimination strategy [1] (Fig. 14.8). Like PCN-222, it formulates a 3D crystal in a space group of m3m, thus creating a structural design that facilitates the formation of channels with a diameter of 19 Å. In addition to its 3D design, the MOF has a BET surface area of 2600 m2 g1, arguably the highest surface area reported by porphyrin-based MOFs, and a pore volume of 1.59 cm3•g1.

14.7.3 PCN-228-230 Using MOF topology as the guiding criterion, the Zhou group further elongated the pore aperture through linker extensions of the porphyrinic units and created a series of porphyritic MOFs, classified as PCN-228-230 [36]. The BET surface areas range from 4510 m2•g1 (PCN-228), 4619 m2•g1 (PCN-229), and 4455 m2•g1

Fig. 14.7 Crystal structure of PCN-222. The TCPP (yellow) is connected to four 8-connected Zr6 clusters (green cuboid) with a twisted angle to generate a 3D network with 1D large channels (Reprinted with permission from Zhang et al. [35]. Copyright @ 2016 American Chemical Society)

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Fig. 14.8 Microscope images of (a) PCN-224(Ni) and (b) PCN-224(no metal). Crystal structure, structural components, and underlying network topology of PCN-224(Ni): (c) the 6-connected D3d symmetric Zr6 in PCN-224. (d) Tetratopic TCPP ligands (violet square) with twisted dihedral angles generate a framework with 3-D nanochannels (e, f). (Reprinted with permission from Feng et al. [1]. Copyright @ 2013 American Chemical Society)

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(PCN-230) along with an increasing pore size ranging from 2.5 to 3.8 nm. Notably, PCN-230 displays excellent stability under the treatment of aqueous acidic and basic solutions with pH levels ranging from 0 to 12.

14.7.4 PCN-250 To further the development of chemically- and thermally-stable MOFs, Feng and coworkers developed a kinetically tuned dimensional augmentation synthetic route with the use of a preformed [Fe2X (μ3-O) (CH3COO)6] (M could be Fe2+/3+, Co2+, Ni2+, Mn2+, Zn2+) building block [37]. In this approach, the use of premade secondary building units (SBUs) simplified the crystal growth process by easing the structure reorganization between metal clusters and organic linkers. As a result, the researchers reported 34 different MOFs from 30 different multitopic ligands (ditopic, tritopic, and tetratopic carboxylic acids) and mixed ligands in the presence of a modulator acetic acid. Among all the MOFs, PCN-250, constructed from the 6-connected [Fe2X (μ3-O)] cluster and the azobenzene tetracarboxylic acid, stands out for the ultrastability. Interestingly, it was also reported that variation in the arrangement of linkers produced two framework isomers [38] (Fig. 14.9). Despite the introduction of a softer Lewis acid M(II) in the cluster, PCN-250 remains robust after undergoing treatment of glacier acid and aqueous solution with a pH ranging from 1 to 11 lasting over 24 h. Moreover, the chemical stability test approves that PCN-250(Fe2Co) remains intact after being exposed to the water treatment more than 6 months.

Fig. 14.9 Structure representation of two framework isomers for PCN-250 (Reprinted with permission from Yuan et al. [38]. Copyright @2017 Elsevier Inc.)

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14.7.5 PCN-333 After carefully considering the pore shape and symmetry of MOFs with large pore size, the Zhou group applied the ligand extension to construct PCN-332 and PCN-333 [39], which are isoreticular to MIL-100 [21]. While both MOFs display mesoporous cages, PCN-333 boasts the largest pore size of 55 Å and the highest void volume of 3.84 cm3 g1 (Fig. 14.10). Slight modifications to the framework result in exceptional water stability, as both Al-PCN-333 and Fe-PCN-333 maintain both the crystallinity and porosity upon the contact of either pure water or aqueous solution with pH from 3 to 9.

14.7.6 PCN-601 Contrary to the acid resistance of high-valent metal-based MOFs, many carboxylatebased MOFs collapse in diluted basic solutions. The Zhou group adopted the pyraolzate motif in the organic linker to increase the basic resistance of the framework [40]. PCN-601 was constructed from the Ni ions and 5,10,15,20-tetra (1H-pyrazole-4-yl)-porphyrin (H4TPP) ligands (Fig. 14.11). Nitrogen isotherms indicate that PCN-601 has a BET surface area of 1309 m2•g1 with the simulated pore size of 1.1 nm. The stability tests reveal that PCN-601 could survive not only in water but also in a saturated NaOH solution at 100
C.

Fig. 14.10 Structure drawing of PCN-332 and PCN-333 (Reprinted with permission from Feng et al. [39]. Copyright @ 2015 Macmillan Publishers Limited)

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Fig. 14.11 PCN-601 is an ftw-a network (a, d), which was 12-connected [Ni8(OH)4(H2O)2 pyrozolate12] nodes shown as the blue polyhedron (b, e) and 4-connected planar TPP4 ligand as the red plane (c f) (Reprinted with permission from Wang et al. [40]. Copyright @2016 American Chemical Society)

14.7.7 PCN-777 Under careful structural rationalization, the Zhou group reported another Zr-based MOF, PCN-777 [41]. This framework results from 6-connected Zr6 clusters and 4,40 ,400 -s-triazine-2,4,6-triyl- tribenzoate with the aids of a competing reagent of trifluoroacetic acid (Fig. 14.12). Moreover, PCN-777 poses a mesoporous cage of 38 Å with a BET surface area of 2008 m2•g1 and a void volume of 2.28 cm3 g1 [42]. Also, PCN-777 is immune to the treatment of pure water and aqueous solutions of pH ranging from 3 to 11.

14.8

The MOF-800 Series

Dr. Omar Yaghi from the University of California-Berkeley is one of the earliest pioneers who initiated the study of MOFs [43] as well as establish the isoreticular chemistry [44–46]. His group has reported many milestone frameworks in MOF research. Their water-stable MOF-800 series is another novel class of Zr (IV)-based MOFs [47]. Synthesized in 2014, each MOF in the series are similar in nature with some synthetic and architectural differences. However, every MOF in the series naturally exhibits water absorption properties as well as efficient usability. For the discussion related to this family, emphasis will be made on MOF-801 and MOF-841 (Fig. 14.13).

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Fig. 14.12 The coordination of 6-connected Zr6 units and organic linker TATB gives rise to the PCN-777 with large mesoporous cages (Reprinted with permission from Balestri et al. [42]. Copyright @2017 American Chemical Society)

14.8.1 MOF-801 Using solvothermal conditions, MOF-801 ([Zr6O4(OH)4(fumarate)6]∙6(H2O)) is synthetically derived using fumaric acid [47]. The MOF is characterized as the smallest member of the MOF-800 series, but it is novel in its construction. MOF-801 is reported with the synthetic capability to crystallize into a large crystal type (MOF-801-SC) or a microcrystalline powder (MOF-801-P). As the consequence of increasing amounts of defects in the powder format, MOF-801-P demonstrates a BET and Langmuir surface area of 990 and 1070 m2 g1, respectively. Conversely, MOF-801-SC has a measured BET and Langmuir surface area of 690 and 770 m2•g1, respectively. MOF-801 can exhibit two tetrahedral cavity diameters of 5.6 and 4.8 Å along with an octahedral diameter of 7.4 Å. Regardless of its size, MOF-801 uses Zr6O4(OH)4(CO2) n SBUs, which allows it to maintain a permanently porous structure.

14.8.2 MOF-841 The last MOF of discussion for this series, MOF-841 is the major product of the synthesis involving the production of the by-product MOF-812 [47]. Like the previously discussed MOF, MOF-841 (Zr6O4(OH)4(MTB)2(HCOO)4(H2O)2)

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Fig. 14.13 Zr6O4(OH)4(-CO2)n SBUs with different coordination numbers listed under each of the SBUs as n-c (left column) are connected with organic linkers (middle) to form MOFs (right column) of various topologies (three-letter codes in the right column) (Reprinted with permission from Furukawa et al. [47]. Copyright @ 2014 American Chemical Society)

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undergoes its reaction in solvothermal conditions, utilizing H4MTB organic linkers in developing the crystal structure. There is only one cage type that has a measured diameter of 11.6 Å in this framework. MOF-841 has a measured BET surface area of 1390 m2∙g1and a Langmuir surface area of 1540 m2∙g1. Besides the aforementioned MOF series, there are many other water-stable MOFs (JUC-62 [48], Bio MOF [49], ZJU-101 [50], SLUG-21 [51], SLUG-35 [52], FIR-53 [53], NU-1000 [54], and other series based on the same framework) that are interesting but not discussed here.

14.9

Water-Stable MOFs for Water Adsorption

In the second section, we described that a high surface area is one of the most advantageous properties for MOFs. This becomes particularly important when MOFs possess a high water capacity. Comparing to many commercial water adsorbents, like silica and zeolites, some MOFs display high water uptake in the region of low pressures. Besides, studies have approved that the guest molecules are attracted through physical adsorption [55–57], which ease the regeneration process and ultimately promise MOF reusability. Furthermore, finely tuning the hydrophilicity of the framework through linker design has a significant difference in many other guest molecules [57–59]. Since the evaluation of the use of MOFs in water-related applications is vast and complex, there have been many comprehensive reviews on other water-related applications, such as water separation and purification [60– 62]. Since we have a broad interest in the water crisis, we herein aim at the relationship between pure water and MOFs, where water either plays an essential role in the structure of the framework or is attracted to the framework through hostguest interactions. We will describe these studies based on developmental history, from the early stages in the structural characterization, to comprehensive investigations on adsorption and mechanisms, finally the most recent fabrication of MOF-based water harvester in real-world applications.

14.10

Early Stage for Structure Characterization

One of the very first examples of using water-stable MOFs could date back to 2002 [63]. The researchers noticed that the flexible MIL-53 was adopted to the guest molecular by a reversible swelling through atomic displacement, known as the “breathing effect” [21, 63] (Fig. 14.14). The heating at 275
C of the as-synthesized Al-MIL-53 yielded the anhydrous form Al-MIL-53ht at high temperature, in which the dimensions 7.3  7.7 Å2 of the original tunnels changed to channels with dimensions of 8.5  8.5 Å2 after removing the protonated and disordered BDC molecules. Notably, this transformation between the two forms is reversible. Upon the adsorption of one water molecule per Al at room temperature,

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Fig. 14.14 Views of the 3D structure of Al-MIL-53 showing the channel system: (a) as-synthesized Al-MIL-53as, in which the channels are occupied by free disordered H2BDC molecules; (b) calcined form Al-MIL-53ht with empty channels; (c) room temperature form (Reprinted with permission from Serre et al. [63]. Copyright @2002 American Chemical Society)

Al-MIL-53lt was observed to change channel dimension, creating a channel with a dimension of 2.6  13.6 A2. The IR, NMR, and modeling data revealed that the water hydrogen bonding between hydration and dehydration was responsible for the structural switch between Al-MIL-53lt and Al-MIL-53ht. Four years later, Millange and coworkers demonstrated that the Cr-MIL-53 featured a similar breathing effect to the Al-MIL-53 [21] during the hydration and dehydration process [64]. However, unlike its counterparts, Fe-MIL-53 remained in a narrow pore form at high water loading through a metastable anhydrous intermediate Fe-MIL-53int. The Devautour-Vinot group further investigated the energetic features of the breathing effect of Cr-MIL-53, Fe-MIL-53, and V-MIL-47 [65]. The TGA data revealed that all the MIL materials could be dehydrated near 370 K, which was much lower than that of commercial adsorbent zeolites. Both TGA and DSC results suggested a two-step process comprised the water adsorption.

14.11

Investigations of Adsorption and Mechanism

With the development of characterizing technology and more collaborations between scientific fields, more insights about the adsorption have been discovered [66]. Kaskel and researchers compared the water adsorption isotherms of several benchmark MOFs—HKUST-1, MIL-100, MIL-101, ZIF-8, and DUT-4. Both HKUST-1 and DUT-4 collapsed upon the long-term water contact, while ZIF-8 and MIL-101 were retained; this indicates the importance of water stability when using MOFs as water sorbents. The hydrophobic characteristics of ZIF-8 were substantiated by the increasing water adsorption at increasing pressures. Particularly, the MOF initially showed insignificant water adsorption up to the relative pressure of 0.6, then a slight increase at p/p0 of 0.8, and a steep increase for the water condensation in the high-pressure region. MIL-100 and MIL-101 both displayed hysteresis loops in their water sorption isotherms. Additionally, quantitative analysis revealed that the calculated heat of adsorption for MIL-100 and MIL-101 was 48.83

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Fig. 14.15 Bridging effect: a sketch of the formation of hydrogen bridging in the pores of the adsorbent and the resulting voids. (a) Adsorption of single water molecules to the hydrophilic centers. (b) Formation of water clusters due to hydrogen bonding. (c) Filled pore with free voids between the water agglomerates. M and L denote the metal cluster and the organic linker, respectively. Black dots depict the water molecules (Reprinted with permission from Küsgens et al. [66] Copyright @ 2009 Elsevier Inc.)

and 45.13 kJ/mol, respectively. These values were close to the energy required to generate hydrogen bridges between single water molecules. The researchers explained the different sorption behavior for the five MOFs in terms of bridging effect; the adsorbed water molecules on the first layer would generate more hydrogen bridges with additional water molecules, which ultimately formed water clusters in the pore (Fig. 14.15). In the same year, Henninger, Janiak, and coworkers reported a rare water-stable Ni-based MOF, SIE-1, which showed a remarkable loading spread of 280 g/kg at 150
C for the duration of over 10 cycles [67]. Besides previous structural studies of flexible MIL-53, the Bein research group investigated the water sorption behavior for another flexible Fe-MIL-88B with the aid of in situ X-ray diffraction (XRD) [68]. The data revealed that the relative vapor pressure increased during the adsorption; they found the lattice parameter c slightly decreased, while the other parameter increased (Fig. 14.16). These changes in crystal parameters contributed to a more than 40% increase in the cell volume, which ultimately resulted in a two-step adsorption process with a large hysteresis in the isotherms (Fig. 14.17). Since the MIL-53 network has been so attractive to many scientists, Stock and coworkers grafted five functional groups (–Cl, –Br, –CH3, –NO2, (OH)2) to the terephthalate linkers on Al-MIL-53 [69]. Among the five MOFs, the hydroxy MOF analog displayed a wide hysteresis loop. This phenomenon was translated into a two-step adsorption process. Adsorption was initiated by a fast uptake at low pressure, with increasing vapor pressure leading to pore opening within the framework and, consequently, taking up additional water molecules into the pores. To study an additional aspect of the MIL series, four isoreticular MIL-101 MOFs were evaluated by the Kitagawa group [70]. They studied the influence of functional groups (–H, –NO2, –NH2, –SO3H) on their water sorption behaviors. The adsorbed water amounts ranged from 0.8 to 1.2 g g1 for the MOFs. Strong interactions

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Fig. 14.16 Complete diffraction pattern under p/p0 (left) and schematic of change of lattice parameters a and c at the relative pressure of water of 0 and 0.85 (right) (Reprinted with permission from Scherb et al. [68]. Copyright @ 2010 The Royal Society of Chemistry) 300 Adsorbed Volume / cm3g–1

Fig. 14.17 Water sorption isotherms of activated MIL-88B with hysteresis (Reprinted with permission from Scherb et al. [68]. Copyright @ 2010 The Royal Society of Chemistry)

adsorption desorption

250 200 150 100 50 0 0.0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 p/p0 (H2O)

between framework and water molecules were observed in the cases of the sulfonic and amino group with evidence of the increase of isosteric heats of adsorption of water. Tuning the pendant groups on the linkers based on hydrophobicity was approved to be an effective way to change the water sorption uptake. A similar evaluation of the influence of functionalized MOFs on water adsorption was reported by Walton and coworkers [71]. They studied the water adsorption for a series of isoreticular UiO-66 MOFs, with amino, nitro, methoxy, and naphthyl substitutes. UiO-66-1,4-Naphyl displayed the lowest water uptake, as evidenced by a very small hysteresis when compared to those of the other four materials. This could be explained by the inhibition of the water filling caused by the hydrophobic nature of the naphthyl ring. Interestingly, the UiO-66-OMe displayed the highest

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water uptake, which was certainly not expected due to the methoxy groups being a high electron-donating functional group. The authors explained that the methoxy group acted as a directing agent for the pore filling of water, which resulted in more efficient water packing in the pore. In another Walton study, the researchers evaluated the water adsorption for seven classic MOFs—HKUST-1, Mg-MOF-74, UiO-66, UiO-66-NH2, DMOF-1, DMOF1-NH2, and UMCM-1—and the subsequent structural analysis using PXRD and nitrogen isotherms [72]. The Zn-based MOFs (DMOF-1, DMOF-1-NH2, and UMCM-1) lost crystallinity after the water vapor absorption, while others retained good crystallinity according to the PXRD patterns. Moreover, only the UiO-66 and the amino analog displayed negligible loss of surface area. In the same year, Chang and coworkers demonstrated a series of elaborate comparisons of the water sorption properties of mesoporous Cr-MIL-101 and Fe-MIL-100 to a selection of commercial water adsorbents (Zellites NaX, silicoaluminophosphate SAPO-34, and silica gel) in the context of water uptake, dehydration conditions, and adsorption sites [73]. The water sorption isotherms at 30
C at a relatively high partial pressure indicated the water uptakes of the MIL MOFs were 2.5–5.0 times higher than those of the commercial adsorbents. Additionally, most of the pore volume could be dehydrated at 30
C under vacuum, which was significantly lower than the high-temperature requirement of dehydration for the commercial adsorbents. The desorption profiles indicated that the desorption rates of these two MOFs were distinctively higher than those of commercial adsorbents. Moreover, the enthalpy profiles measured by microcalorimeter revealed two adsorption regions, in which the region below 0.3 g g1 corresponded to the formation of adsorbed water monolayer from the unsaturated metal sites and pore surface, which was similar to those of for NaX and SAPO-34, and the region above 0.3 g g1 was ascribed to the capillary condensation in the mesopores. Despite cases supporting carboxylate-based MOFs, Dinca and coworkers reported the water stable pyrazole-based MOFs, Zn(NDI-X), where the substituent X could be –H, –NHEt, or –SEt. This series of MOFs were mesoporous MOFs with ~16 A-wide channels and constructed from the 4-coordinated zinc chain with the functionalized dipyrazolate linkers [74]. Due to the largely hydrophobic surface of the framework, all three MOFs displayed Type V isotherms, which are described as little water adsorption at a low vapor pressure ( p/p0 ¼ 0–0.4) but a sharp increase in the region of high pressure. To increase hydrophilicity on the surface of the MOF, the researchers post-synthetically modified the surface of Zn(NDI-SEt) with more polar and hydrophilic sulfoxide(-SOEt) and sulfone(-SO2Et) groups in the presence of excess amounts of dimethyldioxirane. As a result, the sharp increase in water uptake began at p/p0 at 0.2–0.3 and 0.3–0.4 for Zn(NDI-SOEt) and Zn(NDI-SO2Et), respectively. This work demonstrated the use of a modular synthetic approach in the heating and cooling process in real devices. In another study about the CAU-10-X series, pronounced differences were observed in the water adsorption isotherms at 25
C as shown in Fig. 14.18 [32]. The amine and hydroxyl functional groups provided additional adsorption sites by forming strong hydrogen bonding with the adsorbed water molecules. The

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Fig. 14.18 Adsorption isotherms of CAU-10-X at 25
C (Reprinted with permission from Reinsch et al. [32]. Copyright @ 2013 American Chemical Society)

dehydration could be completed by vacuum at room temperature without any structural change. In Stock’s study, CAU-10-H was later investigated using in situ powder XRD (PXRD) during the hydration/dehydration process [75]. The PXRD pattern showed a reversible structural change upon the water sorption process at ~20% RH. After 100 cycles of water adsorption, the water capacity of the CAU-10-H remained as 0.34 g g1 without any visible loss. The framework retained its crystallinity and its integrity after 700 adsorption/desorption cycles, even when exposed to water. With the expectation of improved stability between water sorption cycles, Henninger, Janiak, and coworkers deposited the microporous aluminum fumarate (μp-AF) onto aluminum metal sheets via the thermal gradient approach [76]. The bulk of μp-AF displayed a maximum water uptake of ~0.45 g g1 with a small hysteresis between the relative pressures of 0.2 and 0.35 in the water adsorption isotherms. The water capacity of the μp-AF coating was 0.37 g g1, even after 4500 cycles of water adsorption and desorption. After the modulator-regulated mature synthesis of Zr-based MOFs, the Yaghi group extensively studied the water adsorption of several water-stable MOFs [47]. They compared 20 MOFs, 10 of which were Zr-based MOFs, along with an additional three benchmark porous materials, hypothesizing that the hydrophilicity of organic linkers would dictate the efficacy of MOF water adsorption at low pressure. However, the pore size played a primary role in the overall water uptake. Among all the porous materials, MOF-801 and MOF-841 were the best performers in terms of water adsorption, ease of regeneration, and recyclability. The adsorption study indicated that the powder form of MOF-801 showed a higher affinity for water in comparison to the single crystal form, which was ascribed to the existence of more crystal defects in the powder MOF. The heat of adsorption of water for MOF-801-P

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and MOF-841 was calculated to be approximately 60 and 50 kJ mol1, respectively. In this study, temperature tended to have little effect on the maximum water uptake of those materials. The single-crystal XRD and neutron diffraction studies also demonstrated that the primary water adsorption sites in low-pressure environments in MOF-801 were the tetrahedral cavities, formed by hydrogen bonding between the μ3-OH in the Zr SBUs and adsorbed water molecules. As the pressure increased, more guest water molecules began to populate into the octahedral cavities by forming new hydrogen bonds with the adsorbed water molecules in the primary binding sites. Farrusseng and coworkers also rationalized the mechanism of water adsorption for 15 MOFs in terms of the Henry constant, the pressure when pore filling occurs, and the maximum water uptake [77]. The mesoporous MIL-101 showed a total capacity of ~0.9 g g1 with hysteresis in the water adsorption isotherm due to the irreversible capillary condensation. The Cr-MIL-101-NH2 displayed a wider hysteresis than the Cr-MIL-101 and Cr-MIL-101-NO2 and a larger amount of water adsorbed due to the more hydrophilic nature of amine groups. Such hysteresis phenomena were absent in the water adsorption isotherms of the microporous MOFs. The breathable behavior of the Al/Ga-MIL-53 series enclosed a hysteresis loop, which resulted from the greater flexibility. Using the same topology as the frameworks, the researchers learned that the linker functionalization of amino groups increased the Henry constants as it decreased the pore filling pressure. Also, they found that the water adsorption of a mesoporous MOF was initially predominated by the physical adsorption of water that occurred at the metal cluster sites at lower pressure but changes to capillary condensation at higher pressure, as evidenced by hysteresis loops. The ideal adsorbents for heat pumps and chillers should have high water uptakes at low pressures with robust stabilities [78]. Bearing this in mind, Chang, Serre, and coworkers designed a water-stable Al-based MOF, MIL-160, with a 2,5-furan dicarboxylic acid, in which two oxygen atoms were utilized to increase the hydrophilicity of the framework. The AlO6 octahedral chain coordinated to the carboxylates from the linkers, forming square-shaped sinusoidal 5 Å-wide channels. The furan-rich MIL-160 endorsed the framework with a relatively higher adsorption enthalpy of 54 kJ/mol. Both the computational simulation and proton MAS NMR results indicated the primary adsorption sites at the hydroxyl groups on the chains. Furthermore, a second NMR peak rising from the weak interactions between the furan oxygen and the adsorbed water molecules was also observed. MIL-160 possesses a high water capacity of 235 g kg1, in which most of the water molecules were adsorbed at p/p0 < 0.18 without any hysteresis and could be removed below 373 K. Thanks to these features, they have helped MIL-160 to be a promising alternative as a water adsorbent in solar cooling applications. No weight loss was noticed after 10 cycles of water adsorption and desorption using MIL-160.

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409

Development of MOF-Based Device in Practical Applications

Due to the high-water capacity, excellent reusability, and ease synthesis of MOF-801, Yaghi and Wang group further developed a water harvester system (Fig. 14.19), which consisted of copper foam, a layer of MOF-801 powder, and a condenser interfaced with a thermoelectric cooler [79]. Vapor from the environmental humidity was adsorbed into the MOF layer through physical diffusion, followed by a low-grade heat from the sunshine on the absorber driven by the desorption process, and concluding with those vapors being further condensed into the water collector. Their MOF-based device demonstrated the capability to capture 2.8 L of water per kg of MOF daily at 20% humidity level from the atmosphere in ambient conditions with sunlight at a flux of less than 1 sun (1 kW/m2). Overall, this design not only maximized the yield by applying the sunlight-driven vapor-desorption but also required no additional energy input. After discovering the potential of MOFs in the application of water production, the same research group implemented the MOF-based device in the practical water capture from desert air [80]. An industry scale of MOF (~1.2 kg MOF-801) was used to build the device, but no liquid water had been collected at the Arizona desert testing spot until the device was slightly modified to use the exterior insulation at a tilted position. This device produced 100 g of water per kg of MOF per one full cycle with no energy input. After carefully evaluating the parameters that affected the harvesting efficiency, they developed a next-generation device with MOF-303, which was constructed from an infinite aluminum-oxo chain and 1H-pyrazole-3,5dicarboxylate. MOF-303 was selected for its unique structure that not only facilitated

Fig. 14.19 The MOF-based water-harvesting prototype

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a water capacity of 0.48 g.g1 but also displayed excellent water stability and good temperature response. As expected, the new-generation device gave a remarkable increase of 114% water production. After demonstrating an effective proof of concept for the water harvesting device, the Yaghi group and Wang group further optimized the engineering aspect of the MOF-based water harvester by coating the MOF layer black with an OTTI-aerogel cover as the optically transparent thermal insulator [81]. With this advancement, they made full advantage of the day-time heat to adsorb the water vapor and the night-time radiative cooling to increase the effective relative humidity (RH) during the swing from 10% (daytime at desert areas) to 20–40% (night time). The thermal efficiency for solar input to water conversion was reported to be ~14%, and the highfidelity computational simulations predicted that this device could deliver ~0.25 L water/kg MOF-801 for a single day cycle. Later, the same groups compared the kinetics on water sorption of MOF-303 to the other four commercial sorbents [82]. The kinetic study of MOF-303 displayed not only showed the fastest water uptakes at 30% and 40% but also the fastest dehydration within minutes upon mild heating. Based on the findings, the Yaghi and Wang collaboration improved their adsorbent-assisted water harvesting cycle from once a day to multiple cycles with the aid of solar power. This enhancement endorsed the MOF-303-based device with the capability to generate 1.3 L per kg of MOF every day in an arid environment of 27
C and 32% RH. In the three-day collection tested in the Mojave Desert, this device produced 0.7 L per kg of MOF per day. To avoid the irreversible capillary condensation of water in a large pore, the Dinca group took into consideration the critical diameter of water capillary action (20.76 Å) in the selection of MOFs for water capture [83]. MOFs with pore size below this critical diameter were predicted to be more likely to display continuous pore filling of water [71]. A series of mesoporous MOFs, M2Cl2(BTDD) (M ¼ Mn, Co, Ni), with an inherited 22 Å wide windows, were studied for the water isotherms. The Co-based MOF turned out to be a superior water absorber with the total water uptake of 0.968 g g1 at 94% RH with no hysteresis of adsorption as well as a calculated heat of adsorption of water at ~55 kJ/mol in zero coverage. Similar to the MOF-801 filled water harvester, this Co-based MOF was projected to capture 0.87 L water per kg of MOF, which agreed with the simulation results (0.82 g g1 MOF). Furthermore, the simulation also predicted that the water capture for this device only dropped by 5.1% over 6 cycles. They also found that this material had a cooling capacity of 400 kWh∙m3 per cycle, which corresponded to a 20
C temperature lift. Maurin, Chang, Serre, and collaborators presented a stable MOF called MIP-200 (MIP ¼ Materials of Institute of Porous Materials from Paris), which consisted of Zr6 oxo clusters and tetracarboxylic linkers [84]. MIP-200 was synthesized under a topology-guided strategy and a Kagome-type framework with both hexagonal (~13 Å in diameter) and triangular channels (6.8 Å in diameter). Both the water adsorption isotherms and TGA profile revealed MIP-200 adsorbed water in a two-step process, with the majority of adsorbed water being removed below 65
C. Negligible differences in weight loss were noticed in the 50 cycle-water

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sorption tests. Also, the GCMC simulation showed that the hydroxyl group on the adjacent Zr6 SBUs coordinated to the adsorbed water molecules in both channels, resulting in the MOF exhibiting high hydrophilicity. Initially, the simulations predicted a simulated heat of adsorption at ~70 kJ/mol, which was evidenced by a steep increase in low relative pressure ( p/po < 0.1). However, as a network of hydrogen bonds formed during the pore filling stage, before the saturation stage at p/po ¼ 0.15–2.0, the adsorption of enthalpy was ascribed to be 55 kJ/mol. No decrease in the working capacity for gravimetric water adsorption was noted in the cyclability tests of MIP-200. Conclusively, MIP-200 demonstrated excellent chemical and mechanical stability, which shows its tremendous promise in its application for adsorption-driven refrigeration. Another MOF, CAU-23, was shown to exhibit properties that would make it suitable for application in adsorption-driven chillers, substantiated by the collaboration of the Henning, Maurin, Zou, and Stock groups [85]. CAU-23 is constructed from the aluminum-oxo chain with the TDC linkers, resulting in square channels with a side length of 7.6 Å. CAU-23 possesses a water adsorption capacity of 0.375 g g1 at p/po ¼ 0.3 and demonstrates a cooling temperature of 10
C. Like other Al-based MOFs, the adsorption sites are located at the μ-OH groups of metal SBUs, in which the formation of hydrogen bonds initiates the adsorption process. Following this, the water molecules occupied the channels until saturation at p/ po ¼ 0.3 by forming networks of hydrogen bonds between the adsorbed water molecules. After 5000 water sorption cycles, CAU-23 retained favorable crystallinity and water capacity. Additionally, CAU-23 displayed a coefficient of performance for cooling of 0.8, even at a very low-desorption temperature of ~50
C.

14.13

Outlook

The chapter has provided a brief overview of the fundamental basics of water-stable MOFs in terms of their synthesis and structural features as well as their applications in water storage. After the extensive development of MOFs in the last two decades, both the stability and the liability of MOFs have been utilized in various fields, with new synthetic strategies for synthesizing water-stable MOFs being discovered in recent years. Therefore, the study of these materials in all aspects is quite limited. Further investigations are needed promptly to fully utilize their potentials in energyrelated applications, especially in the area of water. Acknowledgments The authors thank Texas A&M Higher Education Center and College of Science of Texas A&M University for their additional support. Author contributions X. W. planned and prepared the manuscript and C. L. assisted in writing part of section 2.

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X. Wang and C. Lee Xuan Wang is an instructional assistant professor at Texas A&M Higher Education at McAllen, where she holds appointments in the College of Science at Texas A&M University. Her research interests include the synthesis and development of metalorganic frameworks and porous coordination polymers for chemical sensing and the application in the environmental science field. Her academic training includes undergraduate study at Oklahoma Christian University (B.S., 2011) and graduate study under the direction of Professor Hong-Cai Zhou at Texas A&M University (Ph.D., 2016). She began an academic career as a visiting assistant professor at Colorado State University-Pueblo after graduation.

Mr. Charles Lee is an undergraduate student at Texas A&M Higher Education Center at McAllen. He is pursuing a degree in Biomedical Science. An emerging physician-scientist, Charles’s developing research interests combine biochemical and chemical principles in drug discovery and translational medicine applications. Outside of school, Charles is an accomplished equestrian competing with his beloved American Quarter Horses.

Chapter 15

Supercapacitors: History, Theory, Emerging Technologies, and Applications Yiyang Liu, Paul R. Shearing, Guanjie He, and Dan J. L. Brett

15.1

Introduction

With the development of a global economy, rapid population increase, and the implications of global warming, traditional energy sources will not be able to meet the demand and increasing deployment of renewable energy and transition of electrochemical power systems for vehicle propulsion calls for alternative methods of energy storage [108]. It is particularly important to seek environmentally friendly and sustainable energy resources, such as solar, wind, and geothermal energy; however, an urgent issue is to achieve efficient storage of these renewables [21, 55]. Currently, the development of novel electrochemical energy storage devices, including batteries, supercapacitors (SCs), and fuel cells, is being highly valued by researchers and enterprises. During the past three decades, the applications of rechargeable batteries have surged in many fields, from mobile electronic devices to grid-scale energy storage systems [58]. For example, lithium-ion batteries (LIBs) can achieve a high energy density of ~180 Wh
kg1 (Fig. 15.1), and this continues to

Y. Liu Electrochemical Innovation Lab (EIL), Department of Chemical Engineering, University College London (UCL), London, UK P. R. Shearing · D. J. L. Brett Electrochemical Innovation Lab (EIL), Department of Chemical Engineering, University College London (UCL), London, UK The Faraday Institution, Didcot, UK e-mail: [email protected] G. He (*) Electrochemical Innovation Lab (EIL), Department of Chemical Engineering, University College London (UCL), London, UK School of Chemistry, University of Lincoln, Lincoln, UK e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_15

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Fig. 15.1 Ragone plot comparing electrochemical energy conversion and storage devices

increase through rapid advancement of the field. However, for many applications, their power density is unsatisfactory, which leads to a slow charging rate. Meanwhile, the shortcomings of LIBs, including high cost, insufficient lithium reserve, and potential safety issues, may retard LIB deployment and their long-term prospects. In addition to ultra-high power density (10 ~ 100 kW kg1) compared to other energy conversion and storage devices, SCs have merits including operation over a wide range of temperatures (40 ~ 80  C), high efficiency, and fast charge/discharge rates (in seconds) [3, 4, 34]. Meanwhile, compared with some commercial technologies, such as fuel cells, SCs possess a much lower cost and longer cycle lives (> 100,000 cycles). As energy storage devices, the properties of SCs sit between traditional capacitors and rechargeable batteries. As shown in Fig. 15.2, supercapacitors can be used as both quick-start power supplies for electrical vehicles and balanced power supplies for lifting devices; they can also be used as traction energy for hybrid electric vehicles, internal combustion engines, and trackless vehicles, as well as power supplies for other equipment. It is worth mentioning that SCs are of great significance in plug-in hybrid vehicles. Generally, it is difficult to simultaneously deliver high energy density, power density, and long cycle life in a single power system, since the rechargeable batteries must make a trade-off among these properties. The adoption of auxiliary energy systems using SCs can address this challenge: the main energy system (rechargeable batteries) with high energy density provides a longer driving range, while the auxiliary energy system (SCs) with the high power density and cycle life provides the short-term auxiliary power in scenarios such as acceleration and slope climbing. To provide the necessary and comprehensive understanding for new researchers in the fields related to SCs, this chapter provides a concise introduction and discussion covering: (a) the development history, and fundamental theories and principles

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Fig. 15.2 Application scenarios of supercapacitors

(Sect. 15.1); (b) electrolytes (Sect. 15.2); (c) electrode candidates, including carbon materials (Sect. 15.3); transition metal compounds (Sect. 15.4) and novel electrode materials (Sect. 15.5); and (d) conclusion and perspectives (Sect. 15.6).

15.2

The History of Supercapacitors

The “Leyden jar” is the earliest report of a capacitor. Invented in 1746 by Prof. Pieter Van Musschenbroek at the University of Leiden, it comprises a glass jar filled with water into which a brass rod is placed [84]. An early “Leyden jar” can be charged to a high voltage of 20,000 ~ 60,000 V and has a typical capacitance of 1 nF per pint of size. However, SCs represent an energy storage device positioned between traditional capacitors and rechargeable batteries, which possess high power densities and relatively high energy. In 1853, German physicist Hermann Von Helmholtz put forward a theoretical model of an interfacial double layer (DL): under the action of a specific potential, two layers of charge with the same amount but the opposite charges will be generated at the interface between electrode materials and electrolyte, thus forming an electrical double layer [35]. Based on insight from this theory and research on porous carbon electrode materials for fuel cells and rechargeable batteries, H. Becker from General Electric developed a “low voltage electrolytic capacitor with porous carbon electrodes” [11]. Although the underlying mechanisms for energy storage in porous carbon electrodes were not clear, the apparatus demonstrated a high energy

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density similar to rechargeable batteries at that time. Inspired by Becker’s research, Donald L. Boos from Standard Oil of Ohio (SOHIO) further developed this technology and made the first attempt to market SCs [14]. Then, SOHIO licensed this technology to NEC Corporation (Nippon Denki Kabushiki-gaisha), who first successfully commercialized this electric double-layer capacitor (EDLC) as backup power for computer memory [7, 8]. However, the market of first-generation EDLCs expanded slowly, due to their unsatisfactory energy density, low operating current, and high internal resistance. To address these challenges, Trasatti and Buzzanaca found that ruthenium oxides (RuO2) demonstrated outstanding capacitive performance, which set off a wave of research on pseudocapacitors based on metal oxides as active materials [83]. Between 1975 ~ 1980, Brian E. Conway carried out extensive fundamental studies and commercial efforts on electrochemical capacitors using RuO2 electrodes, and significantly enriched the knowledge of the electrochemical capacitor field. He redefined the term “supercapacitor” and proposed the capacitance comes from two parts: (a) the electrical charges stored in the Helmholtz electric double layer and (b) the results of Faradic reactions with “pseudocapacitance” charge transfer, the mechanisms of which include redox reactions, intercalation, and electrosorption [19]. Since then, pseudocapacitive SCs using RuO2 as electrodes have been used in various applications ranging from automobiles to the military. Since the 1990s, a series of inexpensive transition metal (e.g. Mn, Ni, Co, and V) oxides, conductive polymers, and other electrode materials have also been extensively studied.

15.3

Working Principles and Classification

As shown in Fig. 15.3, the structure of SCs consists of four components: cathode, anode, electrolyte, and separator. Defined by the different types of electrode materials used, supercapacitors can be divided into three categories: electric double-layer

Fig. 15.3 Typical structure of supercapacitors

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Fig. 15.4 Types of supercapacitors. (a) EDLCs; (b) pseudocapacitive SCs (pseudocapacitors); (c) hybrid SCs

capacitors (EDLCs, Fig. 15.4a), pseudocapacitive SCs (pseudocapacitors, Fig. 15.4b), and hybrid SCs (Fig. 15.4c).

15.3.1 Electric Double-Layer Capacitors (EDLCs) As shown in Fig. 15.4a, charge storage in EDLCs is achieved by accumulating charges at the interface between electrodes and the electrolyte. Under the action of an external electric field, equal amounts of charges carried by anions and cations move to the cathode and anode, respectively, thus forming a potential difference. When a charge accumulation occurs on the surface of an electrode, ions with opposite charges in the nearby electrolyte migrate to the electrode and accumulate on its surface, thus forming an electric double layer. When the external electric field is withdrawn, the electric double layer still exists and stabilizes the voltage due to the attraction between opposite electric charges. After the supercapacitor is connected to an external electrical appliance and a closed circuit is formed, the charged ions absorbed on the electrodes will migrate directionally and form a current in the external circuit until the electrolyte returns to electrically neutral. The charge and discharge of an EDLC are highly, or even completely, reversible. This is due to the fast ion electric absorption and desorption process and almost no chemical reactions involved [97]. Since 1853, a series of theories and models were developed, among which the most representative ones are the Helmholtz model, Gouy-Chapman (G-C) model, and the Gouy-Chapman-Stern (G-C-S) model. The central concept of the Helmholtz model is that opposite charges are equally distributed on both sides of the interface; furthermore, this structure can be equivalent to a flat capacitor (Fig. 15.5a).

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Fig. 15.5 Models of the EDL. (a) Helmholtz model; (b) G-C model; (c) G-C-S model; (d) modified G-C-S model

Therefore, the relationship between the unilateral charge density (σ) and the potential between two layers (V ) can be determined by Eq. (15.1) [35]: σ¼

εr ε0 V d

ð15:1Þ

where εr and ε0 are the electrolyte and vacuum dielectric constant, respectively, and d represents the effective thickness of the electrical double layer. Subsequently, the capacitance (Cd) of the capacitor can be expressed as Cd ¼

∂σ εr ε0 ¼ d ∂V

ð15:2Þ

So far, the Helmholtz model has successfully abstracted a general electrochemical scene into two basic formulas. However, there is an apparent defect in this model: it can be inferred from Eq. (15.2) that Cd is a constant value, while in the experimental observation, Cd will be affected by factors including electrolyte concentration and relative potential [65]. Two apparent trends can be observed in a typical capacitance curve in the double-layer region (Fig. 15.6): (a) Cd presents a V-shaped symmetric distribution relative to the potential V and (b) the value of Cd enlarges with the increase of electrolyte concentration. To optimize the Helmholtz model, the G-C model (Fig. 15.5b) was proposed by introducing a diffuse layer model of the DL [100]. Near the electrode, the charges are strictly distributed on the surface of the electrode; while in the electrolyte, the charges will diffuse into the bulk solution far from electrode–electrolyte interfaces, due to the interaction between different ions. Therefore, when the potential difference between two sides of the interfaces is large, more ions will be compressed near the electrode; while the ions can also reach charge equilibrium with the electrodes in a small space when the electrolyte concentration is high. The improvement of the G-C model changes the effective thickness (d ) in the Helmholtz model from a constant to a variable, thus interpreting the trend of Cd more appropriately.

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Fig. 15.6 Typical capacitance vs. a potential difference (Helmholtz model) in the double-layer region [10]

Nevertheless, the G-C model still deviates from the actual situation. A typical G-C model for Cd prediction is confronted with three major issues: (a) at both ends of the curve, Cd approaches infinity, which is contrary to the actual tests; (b) the predicted value of Cd is much higher than the measured value; (c) the G-S model fails for highly charged DLs. In the G-C model, the deviation of Cd prediction occurs because the charge is abstracted into a mass point, which is a standard method in physics. Under the large potential difference, these “point charges” will be infinitely close to the electrode surface; therefore, the distance between positive and negative charges will approach zero, resulting in Cd approaching infinity. Later, Otto Stern proposed the Gouy-Chapman-Stern (G-C-S) model by combining the Helmholtz model and G-C model and considering the size of ions (Fig. 15.5c) [78]. In the G-C-S model, there are two ion distribution regions at the electrode–electrolyte interface: some ions adhere to the electrode as suggested by Helmholtz, forming a Helmholtz layer, while others form the Gouy-Chapman diffusion layer. In the diffusion layer, the capacitance (Cdiff) is produced by the action of thermal motion of electrolyte ions; in the Stern layer, ions are absorbed on the electrode surface and result in a constant capacitance (CH), which is independent of the potential difference. Therefore, the capacitance of the whole double layer (Cdl) can be expressed by Eq. (15.3): 1 1 1 ¼ þ C dl CH Cdiff

ð15:3Þ

According to Eq. (15.3), the value of Cdl is always smaller than that for CH and Cdiff. When the potential difference is low, the value of Cdiff is very small, and Cdl is mainly affected by Cdiff and has a V-shaped curve versus potential difference; when under high potential difference, the value of Cdiff is large, its contribution to the Cdl is negligible. In this case, the value of Cdl approaches CH. However, there remain issues with the G-C-S model, such as a) the ions in the electrolyte will be surrounded by solvent ions to form solvated ions; b) if the adsorption force at the interface

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surpasses the electrostatic force, then even the similar charges can be stable at the interface. In 1947, David C. Grahame modified the G-C-S model and suggested that some ions or uncharged substances could penetrate the Stern layer [29]. In the modified GC-S model, the term “specifically absorbed ions” refer to the ions losing their solvation shell and directly contacting with the electrode when approaching the electrode. As shown in Fig. 15.5d, the modified model suggested the existence of three regions, including the inner Helmholtz plane (IHP), outer Helmholtz plane (OHP), and diffuse layer. The IHP refers to the closest distance of specifically adsorbed ions, and the OHP refers to the distance of non-specifically absorbed ions. Although the theories and models are increasingly accurate and sophisticated, debate and further exploration continues. Some recent studies revealed that the capacitance of EDLCs is also related to the active specific surface and porosity of electrode materials, wettability between electrode and electrolyte and the acidity and alkalinity of the electrolyte solution, etc [2, 42, 87]. Currently, most of the electrode materials suitable for EDLCs are carbon materials with excellent electrical conductivity, such as graphene, porous carbon, and carbon nanotubes [37, 67, 80].

15.3.2 Pseudocapacitive Supercapacitors (SCs) The energy storage of pseudocapacitive SCs is realized through the highly reversible chemisorption or redox reactions involving electron transfer, which will trigger underpotential deposition (UPD) on the surface or the (quasi-) two-dimensional space in the bulk phase of the electrochemically active materials. Like rechargeable batteries, the Faradaic process occurs on the surface of electrodes for SCs. The electrode reactions in the rechargeable battery involve the transfer of localized valence electrons, which follows the Nernst equation. In contrast, charge storage in pseudocapacitive SCs is achieved by the transfer of delocalized valence electrons, and the corresponding electrode potential is proportional to the electric charge passing through the electrode. A typical charging and discharging process for a transition metal oxide in the aqueous electrolyte are shown as follows: In acidic electrolyte: MOx + H+ + e , MOx  1(OH) In alkaline electrolyte: MOx + OH  e , MOx(OH) Besides, adding redox reactive ions into the electrolyte can also increase pseudocapacitance [50]. For instance, Ren et al. applied a redox-active electrolyte (1 M H2SO4 electrolyte adding 0.8 M Fe2+/Fe3+ additives) in a supercapacitor with a porous nanoflower polyaniline (PANI) electrode and reported a high capacitance (1062 F g1 at 2 A g1), energy density, and cycling stability (93% at 5 A g1 over 10,000 cycles) [70]. In the as-assembled SC, two redox systems are originating from: (a) the conjugated double bonds in the polymer networks (Eqs. (15.4) and (15.5)) and (b) a Fe2+/Fe3+ active redox couple (Eq. (15.6)).

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   Pxþ m Ax þ xe , Pm þ xA

ð15:4Þ

þ  þ Pyþ m C y þ ye , Pm þ yC

ð15:5Þ

Fe3þ þ e , Fe2þ

ð15:6Þ

where Pm refers to PANI, m is the polymerization degree; A and C+ refer to the anions and cations, respectively. Generally, pseudocapacitive SCs possess superior specific capacitance and energy density compared to EDLCs. However, the bulk phase of electrode materials or electrolyte composition will change, with the occurrence of redox reactions, especially the pseudocapacitive effect on the surface of the electrode. Therefore, material fatigue damage is inevitable during long-term charge and discharge cycles, resulting in poor cycle life. Therefore, the lifespan of pseudocapacitors is shorter than that of EDLCs, while still longer than the batteries. Early studies on pseudocapacitive electrodes focused on transition metal oxides (TMOs); however, the low electrical conductivity and unsatisfactory Faradaic charge storage capability of TMOs limited their application. One common approach to gain a high-performance pseudocapacitive electrode is to grow a layer of active material with the capability of pseudocapacitive or battery-type charge storage on the surface of EDLC materials. This method has been reported in many studies. At present, the materials widely studied as the electrode candidates for pseudocapacitive SCs include hetero-atom doped (N, P, O, S, B, etc.) carbon materials, transition metal oxides/hydroxides (MnO2, NiOOH, etc.), two-dimensional transition metal carbides and nitrides (MXene), and conductive polymer materials (polyaniline, polypyrrole, etc.), among others.

15.3.3 Hybrid Supercapacitors (SCs) According to Chen et al., there are three types of charge storage mechanisms: non-Faradaic capacitive storage (NFCS) for EDLC storage, capacitive Faradic storage (CFS) for pseudocapacitive storage, and non-capacitive Faradaic storage (NCFS) for battery-type storage [97]. The effective combination of an EDLC electrode and a pseudocapacitive or battery electrode is a promising methodology to improve the energy density, power density, and cycle life of SCs, as well as significantly broadening the operation potential window [17]. In early studies, hybrid SCs were given various names, such as asymmetric SCs and hybrid ion capacitors, which may lead to confusion. Therefore, based on Chen’s definition, this chapter classifies different types of hybrid SCs based on the charge storage mechanisms of the electrodes (Fig. 15.1 and Table 15.1) [97].

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Table 15.1 Classification of supercapacitors based on the charge storage mechanisms EDLCs

Pseudocapacitve

Hybrid SCs

Rechargeable

SCs

Batteries

Electrode 1

NFCS

CFS

NFCS

NFCS

CFS

NCFS

Electrode 2

NFCS

CFS

CFS

NCFS

NCFS

NCFS

15.4

Electrolyte

The electrolyte is one of the critical components of SCs and provides ionic conductivity to facilitate charge compensation on each electrode. One of the substantial challenges for SCs is the unsatisfactory energy density: for a typical EDLC, the energy density is less than 5 W h kg1; for pseudocapacitive SCs or hybrid SCs, their energy density is commonly less than 15 W h kg1 [105]. This energy density cannot fully satisfy current commercial and industrial demands. The relationships among energy density of SCs (E), capacitance (C), and the operation voltage (V ) can be expressed by Eq. (15.7): 1 E ¼ CV 2 2

ð15:7Þ

According to Eq. 15.7, increasing either capacitance or both capacitances and operation voltage is an effective method to enhance the energy density of SCs. In addition to the energy density, the power density (P) is also closely related to the operating voltage, as shown in Eq. (15.8): P¼

1 V2 4W TS Rcell

ð15:8Þ

where Rcell relates to the equivalent series resistance (ESR) of SCs (Ω). It is worth noting that increasing the operating voltage will be more effective than increasing the capacitance to enhance both the energy and power density because the energy/ power density is proportional to the square of the operation voltage. Within the stable electrochemical voltage range of the electrode materials, the operation voltage is mostly dependent on the electrochemical stable potential window (ESPW) of the electrolyte. Figure 15.7 demonstrates the typical electrochemical stable potential window of commonly used electrolytes, including aqueous electrolyte, organic electrolyte, ionic liquid, and (quasi-)solid-state electrolyte. The application of electrolytes with higher ionic conductivity is essential to reduce the ESR of SCs (Rcell). Generally, the ionic conductivity (σ, S cm1) can be expressed by Eq. (15.9):

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Fig. 15.7 Electrochemical stable potential window of different types of electrolytes

σ¼

  nðZeÞ2 D E exp  a kT kT

ð15:9Þ

where n refers to the concentration of the carrier, Ze is the charge of the carrier (for protons, Z ¼ 1), D relates to the self-diffusion coefficients, and Ea refers to the activation energy of ion diffusion. Adopting electrolytes with high self-diffusion coefficients and low activation energy, as well as increasing the number of charge carriers, will enhance the ionic conductivity. In addition to a wide potential window and high ionic conductivity, an ideal electrolyte possesses the following characteristics: (a) high (electro)chemical stability and wide operation range; (b) well-matched with other SC components (electrodes, current collectors, separators, etc.); (c) low flammability, volatility, and toxicity; (d) low cost; and (e) environmentally friendly nature. It is very challenging to satisfy these requirements simultaneously because different types of electrolytes possess their intrinsic strength and weakness. The following content in this section provides an overview of the different types of electrolytes to construct a basic understanding of the development of the electrolytes for SCs.

15.4.1 Aqueous Electrolyte Generally, aqueous electrolytes possess significant merits such as low cost and high ionic conductivity (e.g. ~ 0.8 S cm1 for 1 M H2SO4 at 298.15 K) [28]. Meanwhile, the small hydrated ion (2.80 ~ 4.28 Å) of the aqueous electrolyte can quickly immerse in the micropores of the electrode, so that the large size distribution and specific surface area of the electrode materials can be effectively used [105]. The aqueous electrolyte can be divided into three categories: acidic, alkaline, and neutral electrolytes. Currently, the most commonly used acidic, alkaline, and neutral

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aqueous electrolytes are H2SO4 solution, KOH solution, and alkali metal salt solution (e.g. Na2SO4, KCl), respectively [1, 46, 86]. Compared with neutral aqueous electrolytes, strong acid and alkaline aqueous electrolytes provide higher ionic conductivity and lower internal resistance; however, their high corrosiveness may lead to the structural instability of SCs, which increase the risk of electrolyte leakage and cause severe environmental pollution. However, aqueous electrolytes are a relatively unpopular choice for current commercial SCs, due to a critical issue: the narrow ESPW limited its application under high operation voltage. This is because the ESPW for pure water is only 1.23 V vs. standard hydrogen electrode (SHE); under high operating voltage, the hydrogen evolution reaction (HER) and oxygen evolution reaction (OER) can occur, leading to water splitting. Although the ESPW of aqueous electrolytes can be elevated via adding salt/acid/alkali or applying “water-in-salt” aqueous electrolytes, the effects were quite limited (< 2.2 V vs. SHE) [25].

15.4.2 Organic Electrolyte Organic electrolytes currently dominate the commercial SC market due to several merits, including a) high ESPW (2 ~ 4 V vs. SHE), b) wide operating condition range (temperature and humidity), c) low corrosiveness to SC components, and d) high electrochemical stability. Additionally, the use of organic electrolytes allows low-cost materials (e.g. Al) to be used in current collectors and packages. However, organic electrolytes also face some substantial challenges. Firstly, compared to the aqueous electrolytes, organic electrolytes generally possess larger ionic radii, lower ionic conductivity, and higher internal resistance, which demands electrode materials with larger pore size and leads to the low surface area utilization of the electrode materials. Meanwhile, higher voltage and current density may lead to a lower ion concentration in the organic electrolyte, thus resulting in low efficiency and potential safety issues. Therefore, although organic electrolytes possess a high ESPW (2 ~ 4 V vs. SHE), their operation voltage is generally limited to 2.5 ~ 2.8 V. Besides, SCs using organic electrolytes require a complicated fabrication process in a tightly controlled environment due to the volatility, flammability, and toxicity of the organic electrolyte. Take lithium-ion-based SCs for example, the commonly used organic electrolyte salts include lithium chloride/perchlorate, lithium bis(trifluoromethanesulfonyl)imide (LiTFSI), lithium hexafluorophosphate (LiPF6), and tetraethylammonium tetrafluoroborate (TEABF4). The widely studied solvents are ethylene carbonate (EC), propylene carbonate (PC), dimethyl carbonate (DMC), ethylene dicarbonate (EDC), acetonitrile (ACN), N, N-dimethylacetamide (DMF), γ-butyrolactone (GBL), etc [69, 72, 98]. A single solvent system may have several issues. For instance, Japan has banned the use of commercial electrolytes based on ACN solvent in SCs, due to their low flash point and high toxicity [76]. To noticeably change both the viscosity and ionic conductivity of organic electrolytes, research has focused on

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solvents based on a combination of different organic solvents or additives, such as PC-TMC, EC-DMC, and EC-EDC [82, 102]. The further development of electrolyte salts and solvents mainly focuses on improving ionic conductivity and reducing viscosity, to enable the electrolytes to maintain excellent electrochemical performance under extreme conditions (e.g. high voltage and low temperature) [56].

15.4.3 Ionic Liquid Ionic liquid electrolytes are salts that are liquid within a specific temperature range (< 373.15 K). Ionic liquids are considered as suitable alternative electrolytes owing to the following advantages: (a) high tunability due to the variety in anions and cations, (b) high thermal and (electro)chemical stability, and (c) negligible volatility and non-flammability. Many SCs based on ionic liquid electrolytes can achieve operation voltages above 3 V; under this voltage, some commercial electrolytes (e.g. ACN and PC) may face severe electrochemical decomposition issues [52, 105]. Researchers can effectively regulate and optimize ionic liquid electrolytes to satisfy the demands of applications in SCs, in terms of operation condition (temperature, voltage) and internal resistance [26]. Unfortunately, ionic liquid electrolytes are currently unavailable for commercial use, resulting from their unsatisfactory characteristics such as low ionic conductivity, undesirable viscosity, and high cost. Even for ionic liquids with relatively high ionic conductivity, such as 1-butyl-3-methylimidazolium tetrafluoroborate ([EEIM] [BF4], 0.014 S cm1 at 298.15 K), their ionic conductivity is still much lower than that for organic electrolytes (e.g. ~ 0.059 S cm1 for TEABF4/ACN at 298.15 K) and aqueous electrolytes (e.g. ~ 0.8 S cm1 for 1 M H2SO4 at 298.15 K) [28, 41, 72]. Besides, compared to aqueous and organic electrolytes, the viscosity (40 ~ 220 cp) of ionic electrolytes is much higher, which leads to poor electrochemical performance under high current densities [18]. Recent studies suggested that SCs using ionic liquid electrolyte can be improved based on two key approaches: (a) modification of anions or cations to possess similar structure to dispersants and prevent cations from aggregation, thus further improving the ionic conductivity of the electrolytes [47]; and (b) further optimization of the pore structure of electrode materials and wettability between electrodes and electrolytes [24].

15.4.4 (Quasi-)Solid-State Electrolyte With the development of wearable and printable electronic devices, flexible electrochemical energy storage devices have received increasing attention [21]. Besides, to serve as ionically conductive media, (quasi-)solid-state electrolytes also work as

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separators, thus simplifying the fabrication process of SCs without electrolyte leakage issues. At present, major types of (quasi-)solid-state electrolytes are based on polymer electrolytes, which can be divided into three categories: inorganic solid electrolytes (ISEs), solid polymer electrolytes (SPEs), and gel polymer electrolytes (GPLs). ISEs are a particular type of inorganic materials in the crystalline or glass state, which conduct ions by diffusion through the lattice [9]; SPEs are composed of polymers and solvent-free salts, which conducts ions through the polymer chain [61]; in contrast, GPEs consist of polymers, aqueous electrolytes, or conductive salts dissolved in solvents [57]. Generally, only GPEs can be applied in flexible devices, because ISEs and SPEs are not bendable and possess almost no flexibility. Among these three electrolytes, GPEs are the most commonly used electrolytes for (quasi-)solid-state SCs because of their relatively high ionic conductivity (103 ~ 102 S cm1) compared to ISEs (104 ~ 102 S cm1) and SPEs (105 ~ 104 S cm1) [12, 22, 57]. However, GPEs are confronted with certain limitations: (a) their weak mechanical strength may lead to the internal short circuit, and (b) the presence of water narrows the operation voltage window. Besides, a common challenge for (quasi-)solid-state electrolytes is their limited contact surface area with the electrode materials, especially for porous nanomaterials, which will lower the utilization of the electrode active materials. This issue may increase the equivalent series resistance (ESR) of SCs and lead to a reduction in rate performance and poor specific capacitances. At present, the development of (quasi-)solid-state electrolytes is not mature, and no such electrolyte can satisfy all of the following key requirements: (a) high ionic conductivity, (b) high thermal and (electro)chemical stability, and (c) sufficient mechanical strength.

15.5

Carbon Material Electrodes

Various types of carbon materials are currently used for SC electrodes, due to their low cost, high surface area, excellent electrical conductivity, etc. Since it was first used in SCs in 1957, researchers have made numerous efforts to enhance the electrochemical performance of carbon materials, which mainly focus on surface modification, pore structure and surface area optimization, hetero-atom doping, etc. [32, 53]. This section will provide a brief introduction covering the primary carbon materials, including activated carbon (AC), carbon nanotubes (CNTs), and graphene.

15.5.1 Activated Carbon (AC) AC is the earliest widely used active electrode material for SCs, owning to the large specific surface area (> 1000 m2 g1), high electrical conductivity, and low cost

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[100]. AC can be obtained through the carbonization of carbon-rich organic precursors under the protection of inert atmosphere (N2, Ar), and physical or chemical activation. These precursors originate from a wide range of sources, including (a) natural renewable resources (e.g. cellulose, coconut shells, discarded rubber, and even human hair) [69], (b) fossil fuels and their derivatives (e.g. coal, asphalt, and coke) [5], and (c) synthetic precursors (e.g. polymers, metal-organic frameworks, and covalent organic frameworks) [88]. The mainstream activation methods include physical activation, chemical activation, and microwave activation. Physical activation usually takes place at 600 ~ 1200  C in CO2 or a water vapor atmosphere [44]. Usually, it can be divided into two steps: a) precursor pyrolyzed at 400 ~ 900  C in an inert gas atmosphere, and b) oxidation at 350 ~ 1000  C to increase the porosity and specific surface area [73]. The chemical activation only contains one step: the precursors are mixed with chemical substances such as strong base (e.g. KOH, NaOH), strong acid (e.g. H2SO4, H3PO4), carbonate (e.g. K2CO3, Na2CO3), chloride (e.g. ZnCl2, FeCl2), and activated at 400 ~ 900  C [48]. The microwave activation causes internal dipole rotation and ion transfer by heating the precursor with microwave radiation, thus generating the activated carbon. All three activation methodologies have their advantages and disadvantages. For instance, the significant merits of chemical activation are lower operating temperature, less time consumption, and higher specific surface area (more suitable for EDLCs applications). Although some chemical agents (e.g. KOH) can provide hierarchically porous structure and high specific surface area (> 3000 m2 g1), their potential hazards and toxicity limited further applications [85]. Additionally, heat is transferred from the surface of the precursors to the interior during physical and chemical activation, thus inevitably creating a temperature gradient. This temperature gradient will result in the amorphous and heterogeneous microstructure of the prepared AC. In contrast, microwave activation transfers heat from the interior to the surface, thus allowing for volumetric heating and delivering energy evenly to the bulk materials. In practical applications, it is necessary to combine multiple activation methods to integrate their advantages and avoid shortcomings. Besides activation methods, the chemical and physical properties of the AC are also related to many other factors, such as the composition of the precursor, activation temperature, and activation time. For instance, a longer activation time may lead to higher porosity and broader pore size distribution [33]. The heteroatom doping and pore structure optimization are the most common strategies for the performance enhancement of porous carbon-based SCs. However, the current understanding of how these two strategies influence the ion dynamics, charge storage mechanisms, and electrochemical performance is still minimal. Many experimental and computational studies have demonstrated that hetero-atom doping in carbon materials can induce higher charge delocalization and donor density near Fermi levels, expand interlayer spacing, and enhance the wettability of active materials [68, 93]. However, the effect of different elements on the properties of carbon materials may vary due to the differences in atomic size and electronegativity. In general, the mainstream views on the mechanisms and functionalities of

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Fig. 15.8 Carbon materials with different hetero-atom doping. (a) pristine, (b) O-doping, (c) N-doping, (d) P-doping, (e) B-doping

widely studied doping elements (O, N, P, and B) are (a) O-doping (Fig. 15.8b) can effectively change the surface wettability and provide more active sites; (b) N-doping (Fig. 15.8c) can distort the structure of carbon materials, thus providing additional defects and active sites; and (c) similar to nitrogen, P-doping (Fig. 15.8d) can distort the carbon structure to a greater extent due to the larger radius. As a result, it will achieve an electron-donating ability, better electron delocalization, more defects, and higher wettability; d) B-doping (Fig. 15.8e) will result in a shift of Fermi level to the conduction band, thus modulating the electronic structure of carbon materials and providing a higher electrical conductivity [13, 43]. Also, compared to single heteroatom doping, some recent studies have shown that double or multiple hetero-atom doping further optimizes the nano-/microstructure of the electrodes, thus significantly enhancing the electrochemical performance [49].

15.5.2 Carbon Nanotubes (CNTs) Both single-walled CNTs (SWCNTs) and multi-walled CNTs (MWCNTs) are promising candidates as electrode materials, due to their unique structure, excellent electrical conductivity, and good stability (chemical, electrochemical, and thermal). Although CNTs have a large theoretical specific surface area (1315 m2 g1) and high electrical conductivity (350 ~ 55,000 S cm1), their specific capacitance is only 20 ~ 80 F g1, because there are relatively few micropores that facilitate charge transfer [51, 66]. For MWCNTs, although their volume of micropores can be increased by activation, the specific capacitance is still inferior to that of AC.

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There are many reports of CNT-based electrode candidates with remarkable specific capacitance; however, a critical obstacle to commercialization is their low packing density (100 ~ 250 kg m3). Densification of CNT-based electrodes to an industrial level without severely compromising ion transport is a necessary consideration to realize commercial SCs. Wei et al. proposed a strategy to shorten the transport distance of electrolyte ions via hierarchical structures, which is constructed by using 90% granulated double-walled CNTs (DWCNTs) or triple-walled CNTs (TWCNTs) and 10% MWCNTs (10μm long) as the linker (Fig. 15.9) [104]. After CO2 etching, the as-fabricated membrane with hierarchical structure demonstrates remarkable packing density (420 kg m3) and high specific surface area (871 m2 g1). Owning to the excellent mechanical properties and high electrical conductivity, CNTs are often used as supporting substrate materials to grow pseudocapacitive or EDLC electrode materials. Huang and co-workers designed and synthesized the amorphous MnO2@MWCNTs fiber, in which MnO2 was incorporated into MWCNTs fiber uniformly [75]. The as-fabricated fiber-based supercapacitor exhibits a high power density (1.5 mWh cm3) and a remarkable specific capacitance of 8 F g1 at 1 A cm3 (~12.9 times higher than that of MWCNT fiber-based devices, 0.62 F g1). The excellent electrochemical properties are as follows: (a) compared with crystalline MnO2, the amorphous MnO2 possesses a highly disordered structure, which is more conducive to rapid ion-diffusion and fast electrode kinetics; (b) the uniformly distributed MnO2 nanoparticles inhibited the stacking of MWCNTs, further facilitating the ion transportation; and (c) the wellaligned MWCNT network offers a higher electron transport performance.

MWCNTs

Granulated DWCNTs/TWCNTs DWCNTs CO2

850°C TWCNTs

Acid washing

Dispersion & filtration

Granulated DWCNTs/TWCNTs

MWCNT network

Fig. 15.9 Schematic diagram of the fabrication of CNT membrane consisted of dominant granulated DWCNTs/TWCNTs and long MWCNTs as the backbone. Reproduced with permission from Elsevier [104]

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15.5.3 Graphene Since Geim and Novoselov separated graphene from graphite by the mechanical stripping method in 2004, graphene has attracted extensive attention from researchers in various fields due to its excellent electrochemical properties, strong thermal and (electro)chemical stability, high electron mobility (15,000 cm2 V1 s1), large specific surface area (2630 m2 g1), “intriguing” thermal conductivity (~5000 W m1 K1), and “muscular” mechanical strength (~ 1 TPa) [64, 71]. Graphene is a two-dimensional planar structure material composed of sp2 hybridized monolayer carbon atoms arranged in a honeycomb structure and can be reciprocally converted with CNTs, fullerenes, and other carbon materials. There are many methods to obtain graphene, including chemical vapor deposition (CVD), carbon dioxide reduction, supersonic spray, ion implantation, and laser treatment. [80]. Chemically exfoliating graphite to graphene oxides (GOs), followed by a controlled reduction of GOs to graphene via reductants (e.g. hydrazine hydrate), is generally considered to be the most efficient and economical method [36]. Theoretically, the specific capacitance of single-layer graphene is ~550 F g1, when the G entire surface area is fully utilized [91]. However, the presence of π–π bonds can cause severe aggregation, thus reducing specific surface area and suppressing the actual capacitance. Besides, there are many reports of the incorporation of graphene-based and pseudocapacitive materials as a promising strategy to elevate the electrochemical performance of electrodes [16]. However, increasing active materials loading may inhibit electron transfer and ion diffusion, resulting in a severe decline in the capacitance of SCs [38]. To address both of these challenges, Yao and colleagues synthesized a graphene aerogel with ultra-high MnO2 loading (182.2 mg cm2) via 3D printing (Fig. 15.10) and achieved a record-high areal capacitance of 44.13 F cm2 [96]. It is worth noting that the capacitance of this unique 3D structure is not limited by ion diffusion at high active materials loading, which is impossible for conventional bulk electrodes.

15.6

Transition Metal Compounds Electrode Materials

According to the definition, the theoretical specific capacitance (CT) of transition metal compounds (TMCs) can be determined by (Eq. 15.10): CT ¼

Q It ¼ mV mV

ð15:10Þ

where I refers to the applied current during the charge/discharge (A), m is the total weight of the electrode,t is the discharge time, and V is the operating voltage window. Also, CT can be expressed based on the molecular weight of TMCs (M, g mol1) as Eq. (15.11):

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Fig. 15.10 Schematic diagram of the fabrication of 3D-printed graphene aerogel/MnO2 electrode. Reproduced with permission from Elsevier [96]

CT ¼

nF MV

ð15:11Þ

where n refers to the number of electrons transferred during the redox reaction, and F is Faraday’s constant. Therefore, the TMCs with lower molecular weight and more electron transfer usually have a higher specific capacitance. This section will provide a brief introduction to several extensively studied transition metal compound electrode candidates: transition metal oxides (TMOs), MXene, and other transition metal compounds (hydroxides, sulfides, phosphides, and selenides).

15.6.1 Transition Metal Oxides (TMOs) RuO2, as a precious metal oxide pseudocapacitive electrode with high specific capacitance and power density, is the first material studied and applied on SCs for applications in national defense and aerospace [83]. Although RuO2 is one of the electrode materials that possess the highest electrochemical performance at present, its negative characteristics such as high cost (~ 7140 $ per ton for ruthenium), toxicity, and need to use with strong acid electrolyte (H2SO4) has limited its civil use [7, 8, 101]. Therefore, some inexpensive and environmentally friendly TMO pseudocapacitive electrode materials (e.g. MnO2, Fe2O3) have received considerable attention in recent years [54]. TMO electrodes are confronted with two major challenges: (a) the unsatisfactory electrical conductivity results inactive materials being only partially utilized; (b) during the charging and discharging cycles, the recurring redox reactions on the surface of the active materials lead to poor cycling stability [95].

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MnOOH

MnOOH

Beta-MnO2 MnOOH Birnessite

MnOOH/Beta-MnO2/Birnessite

Beta-MnO2 Birnessite

Beta-MnO2/Birnessite

MnOOH Beta-MnO2 Birnessite

Fig. 15.11 Schematic diagram of the synthesis of β-MnO2/parallel birnessite core/shell nanorod. Reproduced with permission from the American Chemical Society [107]

MnO2 has been widely studied for SCs, due to factors that include high theoretical capacitance (1370 F g1), low cost, and environmentally friendly properties [59]. Zhu et al. synthesized a β-MnO2/birnessite core-shell hybrid TMO, which using β-MnO2 as the core and highly ordered birnessite sheets (Fig. 15.11) [107]. The utilization of the MnO6 unit in the electrode was improved because the parallel highly ordered shell structure provides an effective transport channel for the ions in the electrolyte. In the 1 M Na2SO4 electrolyte, the as-assembled asymmetric SC exhibits an impressive specific capacitance of 657 F g1 at 0.1 A g1 (based on parallel birnessite) and a high energy density of 17.6 Wh kg1. Also, it is an effective methodology to composite TMOs with carbon materials, because (a) the charge transfer resistance during the (dis)charging processes can be effectively reduced and (b) the addition of carbon materials can enhance the electrical double-layer capacitance. Besides, some other strategies are promising to enhance electrochemical performance, such as nanostructure regulation, defects engineering, and heteroatom doping [45, 108].

15.7

Transition Metal Carbides and Nitrides (MXene)

First reported in 2011, Gogotsi et al. prepared two-dimensional layered Ti3C2 by hydrofluoric acid and ultrasonic method; a series of novel two-dimensional transition metal carbides and nitrides (MXene) have been developed and applied as SC electrode materials [62]. MXenes are defined as Mn + 1AXn, where M represents transition metals, A refers to elements in IIIA or IVA group, X represents carbon or nitrogen, and n ¼ 1, 2, or 3 [63]. In recent years, the MXene has drawn increasing attention in energy storage applications, owing to its excellent electrochemical performance originated from characteristics including high electrical conductivity, unique structure, and the feasibility of electron transfer due to the variable oxidation state of transition metal M [40].

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Further increasing the interlayer spacing of MXene will promote ion (de)insertion kinetics and offer more effective surface exposure to electrolytes, thus enhancing the specific capacitance. The synthetic methods can affect the structure of MXene, for instance, etching with suitable concentration (< 6 M) hydrofluoric acid (HF) and the morphologies of MXene transferred into multilayer lamellas [39]. Compared with high concentration HF, the diluted HF allows the presence of more high-mobility water molecules between the voids of MXene, thus resulting in a larger permanent interlayer spacing and higher capacitance. Then, to obtain a graphene-like monolayer MXene, several methods are developed for the exfoliation of multilayered HF etched-MXene, such as scotch tape delamination, high power ultrasonication, and chemical intercalation delamination [40].

15.8

Other Transition Metal Compounds (Hydroxides, Sulfides, Phosphides, and Selenides)

The low electrical conductivity of TMOs limits their rate performance and energy density under high current densities. By replacing the anions (O) in TMOs with elements such as S, P, and Se, the electrical conductivity of TMCs can be significantly improved. A series of TMCs beyond TMOs have been developed and applied to SCs, such as CuS, MoS2, and TiS2. In recent years, transition metal sulfides with the structure of MCo2S4 (M ¼ Ni, Zn, and Cu) have attracted particular attention, due to their faster electron transport, higher electrochemical activity, and specific capacitance [30]. Chen and colleagues prepared graphene-coated NiCo2O4 core-shell materials (NiCo2O4@G) via a hydrothermal method, which demonstrates a high specific capacitance (1432 F g1 at 1 A g1) and a stable cycling performance (83.4% capacitance retention after 5000 cycles) [99]. The authors suggested that the functionalities of the graphene shell can be summarized in two aspects: (a) mitigate the potential volume change during the operation and (b) protect NiCo2S4 from dissolution.

15.9

Emerging Electrode Materials

In addition to traditional electrode materials introduced above, a variety of novel materials with excellent electrochemical properties have emerged in recent years. This section provides a brief introduction of several representative materials with great potential for use as electrodes for both EDLCs and pseudocapacitive SCs, including black phosphorus, metal-organic frameworks (MOFs), covalent organic frameworks (COFs), and conductive polymers (CPs).

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15.9.1 Black Phosphorus (BP) As early as 1914, black phosphorus (BP) was synthesized by heating white phosphorus under high pressure (12,000 atmospheres); however, it was not until 2014 that it was rediscovered as a member of two-dimensional materials [92]. Similar to graphite, bulk BP has a black appearance, layered structure, high mechanical strength (~ 94 GPa), good density (2.69 g cm3), and electrical conductivity (~100 Sm1) (Fig. 15.12a) [79]. Also, the lamellar BP (phosphorene) has a graphene-like lattice of interlinked six-membered rings where each atom is bonded to three other atoms (Fig. 15.12b), thus forming an orthorhombic pleated honeycomb structure [15]. Phosphorene is prepared in a similar way to graphenes, such as scotch-tape delamination, liquid exfoliation in organic solvents (e.g. acetone, N-methyl pyrrolidone), and high-energy ball milling. [90]. At present, there are few studies based on BP. Through a modified electrochemical approach under ambient atmosphere, Yu et al. synthesized a 3D BP sponge based all-solid-state SC (Fig. 15.12c), which comprise ultrathin (< 4 nm), large (> 10μm) and high-quality (unoxidized) BP nanosheets as the basic unit. During the cyclic voltammetry test, the capacitance is 80 F g1 and 28 F g1 at the scanning rate of 10 mV s1 and 100 mV s1, respectively [89]. With in-depth study, researchers

Fig. 15.12 (a), (b) Structure of black phosphorus (BP); (c) schematic diagram of the synthesis of 3D BP sponge all-solid-state SCs. Reproduced with permission from the Royal Society of Chemistry [89]

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attempt to exfoliate arsenic from the same main group with phosphorus as the electrodes for SCs, which demonstrates an impressive specific capacitance of 1578 F g1 at 14 A g1 [60].

15.9.2 Metal-Organic Frameworks (MOFs) Metal-organic frameworks (MOFs), also known as porous coordination polymers (PCPs), are organic–inorganic hybrid materials with intramolecular pores by selfassembly of organic bridging ligands and metal ions or clusters through coordination bonds. Compared with conventional porous materials, MOF materials have a series of characteristics that are conducive to the application of SCs: (a) higher porosity (e.g. 94%) and specific surface area (e.g. 7140 m2 g1 for NU-110E), due to the longer organic bridging ligands; (b) controllable chemical and physical properties and porous structure, owning to the diversity of skeletons and the high modifiability of organic ligands and metal centers; (c) the highly crystallized compounds facilitate the determination of the precise spatial structure by single crystal or polycrystalline diffraction [23, 27]. However, MOFs and their derivatives also have some common weaknesses, such as unsatisfactory (electro)chemical stability and relatively low electrical conductivity (107 ~ 102 S cm1) [94]. Currently, the following types of materials applied in SCs can be derived from MOFs: bimetallic oxides, metal oxide composites, carbon composites, and highly conductive materials. Some MOFs with good electrical conductivity can be directly used as electrode materials. By using Ni3(2,3,6,7,10.11-hexaiminotriphenylene)2 (Ni3(HITP)2) as the sole electrode for EDLCs (Fig. 15.13), Dennis et al. reported the first example of SCs entirely from MOFs without conductive additives or binders [74]. Some characteristics of Ni3(HITP)2 are conducive to its applications in energy storage, such as high electrical conductivity (50 S cm1), high specific surface area (630 m2 g1), and regular open channels (1.5 nm one-dimensional cylindrical channels). However, there is a dilemma that higher porosity tends to result in relatively lower electrical conductivity. To enhance electrical conductivity without sacrificing porosity, a practical methodology is to composite MOFs with carbon materials [103]. Besides, using MOFs as precursors, the bimetallic oxides can be efficiently designed and prepared through carbonization. Compared with single transition metal oxides, transition bimetallic oxides are equipped with superior electrochemical performances, such as (a) the co-existence of two different cations in a single crystal structure produces more electrons than a single crystal, thus enhancing the electrical conductivity; (b) more redox-active sites can be provided by the synergistic effects of transition bimetallic oxides [6].

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Fig. 15.13 (a) Molecular structure of Ni3(HITP)2; (b) relative size of pore and electrolyte in the ideal Ni3(HITP)2 [74]

15.9.3 Covalent Organic Frameworks (COFs) As a novel electrode candidate for SCs, the earliest covalent organic framework (COF) was designed and synthesized by Yaghi and co-workers in 2005 [20]. Like MOFs, the COFs are equipped with high specific surface area, controllable porous structure, and excellent molecular designability. COFs can easily introduce the heteroatoms into the frameworks and change the physical and chemical properties of the materials, such as boron condensation (Fig. 15.14a) and triazine-based trimerization (Fig. 15.14b). Although COFs can be used as electrodes for both EDLCs and pseudocapacitive SCs, their poor electrical conductivity retards their deployment in SCs. Banerjee and co-workers successfully synthesized a redox-active TpOMe-DAQ thin sheet via an imine condensation, based on 2,4,6-trimethoxy-1,3,5-benzene-tri carbaldehyde (TpOMe) and 2,6-diaminoanthraquinone (DAQ) [31]. The pristine TpOMe-DAQ electrode achieved an excellent areal capacitance of 1600 mF cm2 at 10 mA cm2 and cyclic stability (> 100,000 cycles) in the 3 M H2SO4 aqueous electrolyte, originated from the existence of interlayer hydrogen bonds and the redox activity due to the quinone/hydroquinone transformation (DAQ amine).

15.9.4 Conductive Polymers (CPs) Conductive polymers (CPs) or intrinsically conductive polymers (ICPs) are electrically conductive organic polymers. Generally, the single and double bonds appear alternately in the backbone of CPs, thus forming conductive conjugated π-π bonds. When conductive polymers are used as electrode materials for SCs, they generally possess the following advantages: (a) high electrical conductivity

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Fig. 15.14 COFs synthesis. (a) Boron condensation; (b) triazine-based trimerization

(102 ~ 105 S cm1), (b) wide voltage windows, (c) controllable electrochemical performance, (d) high theoretical specific capacitance, (e) excellent stability and reversibility, and (f) environmentally friendly nature [81, 106]. As shown in Fig. 15.14, the extensively studied CPs include poly(thiophene)s (PTH), polyaniline (PANI), polypyrrole (PPy), poly (p-phenylene vinylene) (PPV), and poly(3,4-ethylene dioxythiophene) (PEDOT) [77]. Among these CPs, PANI, PTH, PPy, and their derivatives are the most widely applied pseudocapacitive electrode materials for SCs. CPs store charges mainly through the redox reaction between electrons in the material and ions in the electrolyte. Generally, CPs possess an excellent microporous structure, good wettability, and a much higher electrical conductivity compared to TMCs. However, it is still essential to further enhance their electrical conductivity, since the electrochemical reactions occur throughout the CPs-based electrode during charging and discharging. Meanwhile, CPs may swell and shrink during the operation, which leads to unsatisfactory electrochemical cycle lives. To solve these issues, there are many promising strategies: (a) incorporating carbon nanomaterials to reduce the probability of deformation and further improve the electrical conductivity; (b) optimizing the structure and morphology of electrode materials to increase the specific surface area and shorten the electron transport pathways; and (c) assembling the hybrid SCs using EDLC electrode as the anode and CP as the cathode, to extend the device cycle life.

15.10

Conclusion

This chapter provides a succinct introduction and discussion of the development history, fundamental mechanisms and theories, and extensively studied electrolytes and electrodes of SCs. Although numerous encouraging achievements have been

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attained in this field, it is still at an early stage to realize the employment of highperformance SCs. The development and perspectives of SCs can be briefly summarized as follows: (a) The main factors affecting the performance of SCs by electrolytes are operation range (temperature and voltage), ion mobility, ionic conductivity, thermal and (electro)chemical stability, and compatibility with electrodes. Although electrolytes for SCs have achieved much encouraging progress in recent years, the inherent challenges for each type of electrolyte have not yet been solved. Currently, there is no such electrolyte that can meet all requirements. For example, aqueous electrolytes have high ionic conductivity, while suffering from their narrow ESPV. More efforts are required in energy storage mechanism studies, advanced characterization techniques, and optimization of the electrode–electrolyte interface. (b) Further optimization of conventional materials or development of novel materials is the key to improve the (electro)chemical performance of SCs. An ideal electrode candidate is expected to be highly conductive, low cost, favorable structure, high specific surface area, high (electro)chemical and thermal stability, and environmentally friendly properties; materials with all these conducive characteristics are very challenging to develop. Currently, many effective strategies have been developed, such as structural regulation, hetero-atom doping, and composite materials engineering. (c) Compared with EDLCs and pseudocapacitive SCs, hybrid SCs combine the characteristics of the rechargeable batteries and capacitors, which is promising to achieve a high energy density and power density simultaneously. Current studies on hybrid SCs are limited to the laboratory stage, which is mainly focusing on the selection and optimization of the electrolyte and electrode materials. This chapter suggested that the design and optimization of new cell configuration, as well as the further tests and verification based on industrial scenarios and application requirements, are the research direction for hybrid SCs development. Acknowledgments The authors would like to thank the Engineering and Physical Sciences Research Council (EPSRC, EP/V027433/1, EP/533581/1), the Royal Society (RGS\R1\211080; IEC\NSFC\201261) and Faraday Institution (EP/S003053/1) Degradation project (FIRG001) for financial support. Author Contribution Mr. Yiyang Liu visualized and wrote the original draft. Prof. Paul R. Shearing wrote, reviewed, and edited the manuscript. Dr. Guanjie He conceived, supervised, wrote, reviewed, and edited the manuscript. Prof. Dan J. Brett helped in funding acquisition, supervision, writing, review, and editing.

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Mr. Yiyang Liu received the B.Eng degree in chemical engineering from the University of Nottingham and the MS degree from University College London in 2019 and 2020, respectively. He is currently a Ph.D. candidate in the Electrochemical Innovation Lab, Department of Chemical Engineering, University College London (UCL), London, UK, under the supervision of Prof. Dan Brett and Dr. Guanjie He. His current research focuses on the material design and synthesis for the application in hybrid battery systems.

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Paul Shearing is a professor in chemical engineering at University College London where he holds The Royal Academy of Engineering Chair in Emerging Battery Technologies. His research interests cover a broad range of electrochemical engineering themes with a particular interest in the relationship between performance and microstructure for energy materials.He co-directs UCL’s Electrochemical Innovation Lab and leads the UK STFC Global Challenge Network in Batteries and Electrochemical Devices. He was a founding investigator of the UK’s Faraday Institution, where he chairs the Training and Diversity Panel and is PI of the LiSTAR program investigating Li-sulfur battery technologies. Dr. Guanjie He is a senior lecturer in chemistry at the University of Lincoln, and an Honorary Lecturer in the Department of Chemistry and Department of Chemical Engineering at the University College London (UCL). His research focused on materials for electrochemical energy storage and conversion applications, especially electrode materials in aqueous electrolyte systems. He has published over 90 papers in peer-reviewed journals (citation > 2500, h-index of 27). He is serving as a guest editor for BioMed Research International, Frontiers in Materials, Frontiers in Chemistry, and Frontiers in Bioengineering and Biotechnology, a topic editor for Coatings, and an editorial board member for Batteries. He received his Ph.D. degree from the Chemistry Department, UCL, under the supervision of Prof. Ivan Parkin. Before this, he received his BSc in College of Materials Science and Engineering, Donghua University. During 2018–2019, he worked in Electrochemical Innovation Lab in the Department of Chemical Engineering, UCL, as a research fellow with Prof. Dan Brett and Prof. Paul Shearing. Prof. Dan Brett is a professor in chemical engineering at University College London (UCL) and an expert in electrochemical materials science and technology development. He received a Ph.D. from Imperial College London and a Bachelor of Science from King’s College London. His research is in the area of electrochemical engineering and technology. This includes electrochemical energy conversion and storage (fuel cells, batteries, supercapacitors, electrolyzers); and electrochemical sensors; electroanalysis, hybrid vehicles, and micro-generation technologies. He specializes in developing novel diagnostic techniques for the study of high- and low-temperature fuel cells and their materials. He is also active in modeling, instrumentation development, engineering design, device fabrication, materials development, and techno-economic analysis of electrochemical energy conversion technologies. He has published >400 peer-reviewed journal papers (citation >10000, h-index of 52) and is in the top 40 published academics in the world over the last 10 years in electrochemical power sources (combined field of fuel cells, batteries, and supercapacitors – Scopus). He was awarded the 2009 De Nora Foundation Prize and the 2011 Baker Medal (Institute of Civil Engineers) for his work on fuel cells.

Chapter 16

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy Storage Jianmin Luo

16.1

Introduction

With the ever-increasing depletion of fossil-fuel resources and energy consumption requirements, sustainable energy technologies are urgently needed. Therefore, considerable attention has been focused on the development of novel materials for highperformance electrochemical energy storage applications. Two-dimensional (2D) nanomaterials including graphene [1], hexagonal boron nitrides [2], transition metal dichalcogenides (TMDs) [3], black phosphorus (BP) [4], and metal-organic frameworks (MOFs) [5] have been widely researched and regarded as the promising candidates for electrochemical energy storage over the past decade owing to their attractive properties. 2D transition metal carbides and nitrides (MXenes), fast-growing nanomaterials in the 2D materials family, were first discovered by exfoliating from threedimensional (3D) MAX phases (M is the early transition metal, A is III or IV A-group element, X is a carbon (C) or nitrogen (N)) with hydrofluoric acid (HF) [6]. During the etching process, A layers in the MAX phase are etched and hydrophilic groups (OH, O, and/or F groups) are bonded to the outer M layers. In their general formula of Mn + 1XnTx, Tx stands for the surface hydrophilic groups [7]. Up to now, more than 30 types of MXenes, including Ti3C2, Ti2C, Ti4N3, and Mo2C, with a variety of compositions and structures have been successfully synthesized, and dozens more have been explored by computational methods [8– 11]. MXenes have a unique combination of physical and chemical properties, including high electrical conductivity, mechanical properties, enabling stable colloidal solutions in water, and efficient absorption of electromagnetic waves, which have led to a large number of applications ranging from energy storage [12–21],

J. Luo (*) Thayer School of Engineering, Dartmouth College, Hanover, NH, USA e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_16

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electromagnetic interference shielding [22], environment and catalysis [23–26], medicine [27], and electronics [28–30]. Besides, the unique characteristics of MXenes, including large redox-active surface area, high electrical conductivity, rich surface chemistry, and tunable layered structure, have endowed them as promising candidates for electrochemical energy storage applications [31]. However, similar to other 2D materials, the electrochemical energy storage performances of MXenes are still restricted by their trend to be restacked and agglomerated, which limits the ionic and electronic transport in electrode materials [32]. For the sake of the full utilization of their electrochemical performance, strategies including introducing interlayer spacers between MXene layers and creating porous structures of MXene were proposed [33–35]. Since MXenes have energy storage space between layers, if the interlayer space can be rationally exploited and utilized, the overall electrochemical performance of MXene for energy storage will be enhanced [36]. Therefore, interlayer structural engineering of MXene with enlarged interlayer spacing is highly necessary. Also, there have some published works that summarized the synthesis techniques, properties, and applications of MXenes, less attention has been paid to the interlayer structural engineering and regulation mechanisms of MXene-based architectures [11, 37–41]. Consequently, it is highly desirable to have a comprehensive understanding of interlayer structural engineering of MXene-based nanomaterials for electrochemical energy storage applications. In this chapter, we focus on the interlayer structural engineering of MXenes, and provide an in-depth discussion regarding the recent layered structure design with expanded and engineered interlayer structures and their electrochemical energy storage applications. The effect of different intercalation agent types and interlayer structural engineering methods related to electrochemical energy storage are discussed (Fig. 16.1). In the first section, we briefly summarize the synthesis techniques of MXenes and go over the different types of layered structure design of MXenes by selecting different intercalation agents. Besides, synthesis-dependent structural properties of MXenes with enlarged and engineered interlayer space are also investigated. In the second section, a series of interlayer structure designs along with the electrochemical energy storage performance is systematically overviewed. In the last section, the challenges and prospects of the interlayer structural engineering of MXene-based nanostructure are addressed. Accordingly, the investigation of interlayer structural engineering of MXene would be greatly conducive to the further optimization of MXene for electrochemical energy storage applications.

16.2

Interlayer Structural Engineering of 2D MXene

Before the layered structure design of MXene, different synthesis techniques are needed to prepare MXene with high quality. The intrinsic properties of MXenes are closely related to their synthesis techniques [39]. Therefore, synthesis conditions can

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Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

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Fig. 16.1 The overview representative interlayer structural design of MXenes, synthetic method, and their electrochemical energy storage applications

directly influence the layered structure design of MXenes and their properties and energy storage performances. At the beginning of this section, we briefly summarize the synthesis techniques of MXenes and illuminate the relationship between synthesis techniques and the structural properties of MXenes. Electrodes made from pristine MXene with small interlayer spacing exhibit unsatisfied electrochemical performance [31]. Significantly, MXene with expanded and engineered interlayer structures for excellent storage capability is confirmed [42–44]. Therefore, in the following, we review the different types of layered structure design of MXenes by selecting different intercalation agents, including heteroatoms, alkali metal ions, other metal ions, polymers, larger cations, and 1D/2D/3D nanomaterials. Importantly, the synthesis-dependent structural properties of MXenes with expanded and engineered interlayer structures are also studied.

16.2.1 Synthesis of MXenes In general, two primary strategies, wet etching synthesis and non-etching synthesis, have been explored to synthesize MXenes [11]. Wet etching synthesis refers to a method of extraction of MnXn
1 (n ¼ 2, 3, 4) from their layered precursors MAX by removing the A atoms’ layers [8]. MAX phase is a large family of ternary carbides and nitrides. Until now, more than 70 different kinds of MAX phases have been reported, and the possible constituent atoms are shown in Fig. 16.2a [37].

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Fig. 16.2 (a) Elements in periodic tables used to synthesize MAX phases. (b) Three typical structures of MAX phases (M4AX3, M3AX2, M2AX) and the corresponding MXenes. (c) The molecular structure model of single layer M3X2Tx. (d) Publication trends in MXene and MXene for electrochemical energy storage (Source: Web of Science. Search index: [topic ¼ MXene or MXene energy storage])

Three typical structure MAX phases (M4AX3, M3AX2, and M2AX) and their corresponding products MXenes are shown in Fig. 16.2b. After removing the A atom layers in MAX by an aqueous fluoride-containing acidic solution, different

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terminal groups would be bonded with interfacial M atoms (Fig. 16.2c). For example, the HF etching method was first used to delaminate Ti3AlC2 MAX in 2011, and many hydrophilic groups (–OH, –F) were bonded with interfacial Ti atoms [6]. It should be noted that different etching conditions, including different etchants concentration and etching temperature, can strongly influence the properties and quality of as-synthesized MXene [39]. So far, various types of etchants (HF, HCl + LiF, NaOH, NH4HF2, and TMAOH) have been used and some novel routes have been developed [6, 45–47]. Besides, a non-etching method such as chemical vapor deposition (CVD) should also be possible for MXene synthesis [48]. In the previous research results, ultrathin 2D α-Mo2C orthorhombic crystals with up to 100-μm lateral size can be produced by using methane as the precursor and copper foil above a molybdenum foil as the substrate. MXenes with a large lateral size and few defects would be synthesized by this method, which facilitates the study of their intrinsic properties [48]. In this part, we focus on the two etching synthesis techniques of MXenes by fluoride-containing acidic solutions, which are the most widely used methods for MXenes synthesis. HF etching method is one of the most common synthesis methods for synthesizing MXenes due to its flexibility to selectively etch away the A layers from most of the precursor MAX phases [6]. Taking Ti3AlC2 as an example, the reaction process can be described by the following equations [6]. Ti3 AlC2 þ 3HF ¼ Ti3 C2 þ AlF3 þ 3=2H2

ð16:1Þ

Ti3 C2 þ 2H2 O ¼ Ti3 C2 ðOHÞ2 þ H2

ð16:2Þ

Ti3 C2 þ 2HF ¼ Ti3 C2 F2 þ H2

ð16:3Þ

The etching conditions for MAX phases vary from one to another, which depends on the compositions, structure, atomic bonding, and particle size of the material [39]. Based on the previous research results, larger n in Mn + 1CnTx requires stronger etching and longer etching time [39]. After HF etching, multilayered MXene is often obtained. Therefore, to obtain separated MXene sheets, typical intercalation and delamination processes are required [49]. In the intercalation process, organic molecules (e.g., dimethyl sulfoxide (DMSO) and tetrabutylammonium hydroxide (TBAOH)) are usually used to expand the interlayer spacing of multilayered MXene [47, 49–51]. Also, sonication treatment is often required in the followed delamination process. Besides, the disadvantage of the HF etching method is the highly corrosive nature of HF, which would lead to the formation of rich structural defects in the obtained MXene [39]. The hazardous environmental and safety concern also exists during the utilization process of HF. To avoid the use of highly toxic and dangerous HF, researchers have made a lot of effort to explore safer ways to synthesize MXenes. In addition to HF, modified fluoride-based acid based on a mixture solution of a strong acid and a fluoride salt can also be used to selectively remove A atoms layers from MAX precursor [45]. It can be found that strong acids can react with fluoride salts to form in situ HF, which

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not only can selectively etch A atoms, but lead to the intercalation of cations (e.g., Li+, Na+, and K+) and water between MXene layers. Therefore, the interlayer spacing of MXene increases, and the interaction between MXene layers weakens after using this in situ HF method. It should be noted that both the concentration of fluoride salts and strong acid can influence the quality and size of obtained MXene flakes [39]. When using the minimally intensive layer delamination (MILD) method (12 M LiF/9 M HCl), the simple washing process alone is enough to delaminate multilayered Ti3C2 MXene, synthesizing single- and few-layer MXene with few defect and larger lateral size [52]. Besides, difluorides such as KHF2 or NH4HF2 can also be used to selectively etch the Ti3AlC2 MAX phase [53]. In general, compared to HF etching, modified fluoride-based acid etching exhibits different properties: (1) Synthesized MXenes have larger lateral sizes and fewer defects; (2) Larger interlayer spacing of the obtained MXenes facilitates the delamination; (3) Higher content of –O terminal groups and less content of –F terminal groups.

16.2.2 Layered Structure with an Enlarged Interlayer Spacing It is well known that the performances of energy storage applications are highly dependent on the electrode materials. As a result, the design and synthesis of highquality materials become an essential step before constructing high-performance electrodes. Until now, MXenes have been widely studied for energy storage applications with the enhanced electrochemical performance [11]. The layered structure of MXene with enlarged interlayer spacing is demonstrated with increased energy storage space for charge-carrying ions intercalation and storage as well as excellent ion transport property due to shorter diffusion pathways [31]. These favorable characteristics endow MXenes with much improved electrochemical performance, which facilitates their practical applications for energy storage applications. In this following section, we will go over the recent progress in layered structure design of MXenes with enlarged interlayer spacing by selecting different intercalation agents, including heteroatoms, alkali metal ions, other metal ions, polymers, larger cations, and 1D/2D/3D nanomaterials. After HF or in situ HF etching, MXenes have a highly negatively charged interface and form stable colloidal solutions in water as confirmed by zeta potential measurements [43, 51]. Therefore, cations can be spontaneously intercalated between MXene layers by electrostatic interaction, increasing the interlayer spacing of MXene [43, 50, 54–56]. It is noteworthy that the new chemical bonding between intercalant and MXene can be formed during the intercalation process of intercalants, which can induce the modification of MXene interfacial properties [44]. Gogotsi et al. studied the intercalation behavior of cations (e.g., Li+, Na+, K+, NH4+, Mg2+, and Al3+) from aqueous salt solutions between Ti3C2 MXene layers [54]. The results show that the intercalation of the above cations is spontaneous, and

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the intercalation behavior depends on pH and the nature of the cations. When Ti3C2 MXene is treated by potassium hydroxide (KOH) and ammonium hydroxide (NH4OH) solutions, the interlayer spacing of Ti3C2 can be increased from 1 to 1.27 nm. Besides, the alkaline solution treatment can not only increase the interlayer spacing of MXene but also change the terminal groups in the interfaces between MXene layers. Peng and co-workers used an NaOH solution to treat Ti3C2 [23]. After alkalization intercalation, Ti3C2 MXene exhibits unique heavy metal ion (e.g., Pb2+) adsorption behavior. The NaOH treatment process is conducive to the transformation of –F to –OH/–ONa groups. Besides, the surface [Ti-O]
H+ groups exhibit strong metal–ligand interaction with Pb2+, and the [Ti-O]
H+ and the [Ti-O]
Na+ sites are propitious to the enhancement of Pb2+ uptake. Except for the abovementioned alkali metal ions, other metal ions (e.g., Sn4+ and 2+ Co ) can also be intercalated into the interlayer of MXene [44, 57]. Luo et al. first synthesized Sn4+ pillared Ti3C2 MXene by the use of alkali metal ions (e.g., Li+, Na+, and K+) to pre-intercalate between Ti3C2 layers, followed by the Sn4+ pillaring process based on the ion-exchange interaction between Sn4+ and pre-intercalated alkali metal ions [44]. After Sn4+ intercalation, the interlayer spacing of Ti3C2 increases from 0.98 to 1.28 nm. Note that the strong Ti-O-Sn bonding can be confirmed after the intercalation of Sn4+ due to the formation of an inner sphere complex on the interface of Ti3C2. By the same method, Song and co-workers fabricated Co2+ intercalated V2C MXene (V2C@Co) [57]. After Co2+ intercalation, the interlayer spacing of V2C increases from 0.74 to 0.95 nm by the formation of V–O–Co bonds. Although metal ions can be successfully intercalated into the interlayer of MXene by electrostatic interaction or ion exchange interaction, the added values of interlayer spacing are limited owing to the small size of metal ions [43]. Luo and co-workers used large volume cationic surfactants as the intercalation agent to intercalate between MXene layers to increase the interlayer spacing of MXene (Fig. 16.3a) [43]. Significantly, the interlayer spacing of MXene can be fine-tuned by creating pillared structures based on the spontaneous intercalation of different size of cationic surfactant (dodecyltrimethylammonium bromide (DTAB), tetradecyltrimethylammonium bromide (TTAB), cetyltrimethylammonium bromide (CTAB), stearyl trimethylammonium bromide (STAB), and dioctadecyldimethylammonium chloride (DDAC)) under different temperatures (30–70  C) (Fig. 16.3b, c) [43]. When using cetyltrimethylammonium bromide (CTAB) as the intercalation agent under the temperature of 40  C, the interlayer spacing of Ti3C2 increases from 0.98 to 2.23 nm. Since CTAB has a long-chain hydrophobic tail (Fig. 16.3d), the possible position of CTA+ in the interlayer space of Ti3C2 is shown in Fig. 16.3e [43]. When using the larger size of stearyl STAB as the intercalation agent at 50  C, the interlayer spacing of Ti3C2 can increase to 2.71 nm, a nearly 177% value-added compared with that of pristine Ti3C2 (0.98 nm). More importantly, ion storage sites for other cations can be provided after the intercalation of cationic surfactant between MXene layers. According to the ion-exchange mechanism, different pillared structure MXene can be fabricated by a facile cationic surfactant pre-pillaring followed by the cations pillaring method. Based on this method, Luo et al. successfully synthesized Sn4+ pillared Ti3C2 MXene (CTAB-Sn(IV)@Ti3C2) by CTAB pre-pillaring and subsequent Sn4+ pillaring process (Fig. 16.3a) [43]. Scanning

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Fig. 16.3 (a) Schematic illustration of the preparation of Sn4+ pillared Ti3C2 (CTAB-Sn(IV) @Ti3C2) by CTAB pre-pillaring process followed by a method of Sn4+ pillaring. (b) XRD patterns of cationic surfactants pre-pillaring Ti3C2 at 40  C. (c) Interlayer spacing of cationic surfactants pre-pillared Ti3C2 at different treatment temperatures (30–70  C). (d) Schematic drawing about the structure of CTAB. (e) Possible position of CTA+ in the interlayer space of Ti3C2 MXene. (f) STEM image of CTAB-Sn(IV)@Ti3C2 and corresponding elemental mapping of Ti, C, and Sn (insets). Reproduced with permission-Copyright 2016, American Chemical Society [43]. (g) Schematic illustration of pyrrole polymerization using MXene. (h) Cross-sectional TEM image of aligned polypyrrole chains (bright layers) between MXene sheets (darker layers). Reproduced with permission-Copyright 2016, John Wiley and Sons [60]. (i) Schematic illustration of the preparation of S atoms intercalated Ti3C2 (CT-S@Ti3C2). (j) The change of interlayer spacings about S atoms intercalated Ti3C2 under different annealing temperatures. STEM image of CT-S@Ti3C2-450 and corresponding elemental mappings of Ti, C, and S. Reproduced with permission-Copyright 2019, John Wiley and Sons [36]. (k) Schematic illustration of the layer-by-layer self-assembly MXenebased multilayer films onto planar substrates. Reproduced with permission-Copyright 2019, Springer Nature [62]

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transmission electron microscopy (STEM) image of CTAB-Sn(IV)@Ti3C2 combined with elemental mapping of Ti, C, and Sn confirms the successful intercalation of Sn4+ between Ti3C2 MXene layers (Fig. 16.3f). In addition to the metal ions and cationic surfactants, other large ions such as amine ions can also be used as the intercalation agents. Gogotsi et al. and co-workers proposed an isopropylamine (i-PrA) intercalation method to delaminate Nb2C MXene [50]. When placing Nb2C MXene in an i-PrA solution, the R-NH3+ cations in the solution can spontaneously intercalate between MXene layers due to the electrostatic interaction, and the corresponding interlayer spacing of Nb2C increases from 1.04 to 2.27 nm. The successful intercalation of R-NH3+ cations can effectively weaken the interlayer interactions, which helps to the delamination of Nb2C. Additionally, some polymers have been confirmed to be the promising intercalation agents for MXenes [58–60]. Gogotsi’s group developed a facile strategy for the in situ polymerization of pyrrole confined between Ti3C2 MXene layers (Fig. 16.3g) [60]. The hydrogen bonding between the anion and the surface of the electrode contributes to the self-align of polypyrrole (PPy) between the MXene layers; Ti3C2 MXene has a strongly pronounced acidic character after HF etching, which can interact with pyrrole and result in a protonated pyrrole molecule. The protonated pyrrole molecule continuously reacts with unprotonated pyrrole, forming a long chain of PPy (chain propagation). Therefore, no oxidant is needed in the polymerization process of the pyrrole in this work. Cross-sectional transmission electron microscopy (TEM) images of the PPy/Ti3C2 film in Fig. 16.3h show that PPy chains are not only intercalated but also well aligned between the MXene layers, producing a periodic pattern. Besides, the formed PPy becomes conductive after doping with the fluorine from the termination of the MXene layers. After PPy intercalation, the interlayer spacing of Ti3C2 increases from 1.4 to 2.4 nm. This increased value is comparable to that of large size cations. Except for oxidant-free polymerization, electrochemical polymerization can also be used to synthesize PPy intercalated MXene. Zhi’s group proposed an electrochemical polymerization method to intercalate PPy between Ti3C2 layers [58]. During the 60 s electrochemical polymerization process, PPy first intercalates in the interlayers of multilayered Ti3C2 and then wraps the Ti3C2 particles. After PPy intercalation, the interlayer distance of Ti3C2 increases from 0.94 to 0.96 nm. Besides, Ling et al. used a simple mixture method to mix polydiallyldimethylammoniumchloride (PDDA) and polyvinyl alcohol (PVA) with Ti3C2 MXene to produce Ti3C2/polymer composites [59]. The successful intercalation and confinement of the above polymers in the interlayers of Ti3C2 can increase the interlayer spacing of Ti3C2. The as-fabricated composites are flexible and high electrical conductivities. Moreover, facile precursor intercalation and the annealing process are effective ways to fabricate heteroatom intercalated MXene, which were confirmed to be one of the MXene-based nanostructures with enlarged interlayer spacing [36, 61]. Tao’s group successfully synthesized sulfur (S) atoms intercalated Ti3C2 MXene by a facile CTAB pretreatment, thermal diffusion with elemental S, and finally annealing and carbon disulfide (CS2) washing process (Fig. 16.3i) [36]. After CTAB pretreatment, the interlayer spacing of Ti3C2 increases from 0.98 to 2.23 nm,

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which helps to the intercalation of elemental S between Ti3C2 layers. The intercalated elemental S can react with the interfacial titanium (Ti) atoms to form Ti-S bonds with expanded interlayer spacing during the annealing process. CS2 is used to wash the residual elemental S. After annealing at 500  C, the interlayer spacing of S atoms intercalated Ti3C2 (CT-S@Ti3C2) with the help of CTAB pretreatment, 1.37 nm is larger than that the cases without CTAB pretreatment (0.98 nm) and CTAB pretreatment (1.15 nm), which indicates the successful intercalation of S atoms between Ti3C2 layers (Fig. 16.3j). The STEM image of CT-S@Ti3C2 combined with elemental mapping of S confirms the successful intercalation of S atoms between Ti3C2 MXene layers (Fig. 16.3j). Wang and co-workers fabricated nitrogen (N)-doped Ti3C2 MXene (N-Ti3C2) by annealing the Ti3C2 under an ammonia (NH3) atmosphere [61]. After annealing treatment by NH3, N is introduced to substitute C atoms in the MXene, resulting in the increased interlayer spacing of Ti3C2 from 0.96 to 1.23 nm at 200  C. In addition to the above methods, the layer-by-layer assembly can also be used for constructing MXene-based nanostructures with enlarged interlayer spacing. Compared with liquid phase pre-pillaring and pillaring method, precursor intercalation and annealing/polymerization method, single layer MXenes are often required in layer-by-layer assembly method, which may need an extra exfoliation process. However, the layer-by-layer assembly can be used to fabricate free-standing electrodes, and it can assemble MXene with different dimensional nanomaterials, which has excellent flexibility. Hamedi et al. used a layer-by-layer self-assembly method to fabricate pillared Ti3C2 MXene/tris(2-aminoethyl) amine (TAEA) multilayers (Fig. 16.3k) [62]. The obtained pillared structure is highly ordered, and this structure reinforces the interconnection between MXene and TAEA. Notably, the TAEApillared MXene multilayers exhibit a high electronic conductivity of 7.3*104 Sm
1. The high electronic conductivity is ascribed to the face-to-face quasi-intimate interface contact between MXene and TAEA. Moreover, Gogotsi et al. prepared a flexible and free-standing sandwich-like Ti3C2 MXene/carbon nanotube (CNT) paper by layer-by-layer assembly method [33]. After CNT intercalation, the interlayer spacing of Ti3C2 increases from 1.21 to 1.42 nm. Furthermore, onionlike carbon, reduced graphene oxide (rGO), activated carbon, and hard carbon can also be fabricated by assembling with MXene [63–65]. The resulting MXene-based papers have good electrical conductivities, high surface areas accessible to ions, and mechanical robustness. Given the abovementioned samples, we summed up the state-of-the-art research works about designing MXene-based nanostructures with enlarged interlayer spacing via the introduction of intercalation agents. In general, the intercalation agents can be classified into the following categories: (1) Heteroatoms; (2) Alkali metal ions; (3) Non-alkali metal ions; (4) Polymers; (5) Large cations; (6) 1D/2D/3D nanomaterials. The main fabrication methods of MXene-based nanostructures with enlarged interlayer spacing, including (1) liquid phase pre-pillaring and pillaring; (2) precursor intercalation and annealing/polymerization; (3) layer-by-layer assembly. After the intercalation of intercalation agents, MXene-based nanostructure with enlarged interlayer spacing shows obvious characteristics for energy storage

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applications: (1) large interlayer space for charge-carrying ions intercalation and storage; (2) excellent ion transport due to short diffusion pathways; (3) limitation of volume expansion of active materials by the confinement effect of a 2D structure during cycling. Besides, by carefully selecting the types of surfactant molecules, pillared MXene structures with controllable interlayer spacings from 1 to 2.7 nm can be prepared to match the sizes of the ions in the electrolyte, maximizing the electrochemical performance [43]. Furthermore, this fabrication method of pillared MXene is conducive to the controllable modification of the surface chemistry about MXene, which is one of the main challenges in designing MXene-based nanostructure [31].

16.3

Interlayer Structural Engineering of 2D MXene for Electrochemical Energy Storage Applications

MXene has a wide chemical and structural variety, which makes them competitive with other 2D materials for energy storage applications. After HF or in situ HF etching, MXenes with a unique layered structure and interfacial characteristics can be obtained. Taking the case of Ti3AlC2, after HF etching, the corresponding interlayer spacing of Ti3C2 is 0.98 nm, and the –OH and/or –F groups present at the surface of the Ti3C2 layers. MXene has interlayer space for cations intercalation and storage. If the interlayer space of MXenes can be further increased and utilized, the electrochemical performance of MXene-based electrodes will be further enhanced [43]. Theoretical studies confirm that lithium (Li) atoms can directly be adsorbed on the O terminated MXene monolayers (Fig. 16.4a), and the theoretical Li capacity of Ti3C2O2 is 268 mAh g
1 based on the configuration of Ti3C2O2Li2 [66]. The valence electron localization function (ELF) of Ti3C2O2 clearly shows the electron transfer between C and Ti and between Li and O (Fig. 16.4a). There are few electrons localized above the adsorbed Li layer. Importantly, the extra Li layers adsorption is also studied (Fig. 16.4b) [66]. With the extra Li layers, the interaction between the different Li layers is visible in the ELF, which indicates that the extra Li layer can bind to the already lithiated MXene. The structure of MXene with additional Li layers is stable (Fig. 16.4b) [66]. In the MXenes, the first extra Li layer forms about 2.8 Å above the lithiated surface. Besides, the second and third extra Li layers require 2.3 Å for each layer. Therefore, a larger space is needed to accommodate extra Li layers, which indicates MXene with enlarged interlayer spacing can store extra Li layers. Also, theoretical Na storage capacities of Ti3C2, Ti3C2F2, and Ti3C2O2, and change of Ti3C2 MXene volume after Na adsorption are explored (Fig. 16.4c–e) [42]. The results show that enlarged interlayer spacing of bare and O-functionalized Ti3C2 MXene enables stable multilayer adsorption and therefore significantly enhances their theoretical capacities. Furthermore, two layers

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Fig. 16.4 (a) Side view of lithiated monolayers of Ti3C2O2 and valence ELF of (110) sections of lithiated Ti3C2O2Li2 monolayer. (b) Side view of an extra metallic Li layer adsorbed on top of lithiated Ti3C2O2Li2 monolayers and valence ELF of (110) sections of lithiated Ti3C2O2Li2 monolayer with an extra Li layer. Reproduced with permission-Copyright 2014, American Chemical Society [66]. (c) Geometric structures for bilayer adsorption on both sides of bare and O-functionalized Ti3C2 MXenes. (d) Capacity as a function of x in Ti3C2T2Nax with T ¼ bare, F, and O. (e) Percentage changes of volume as a function of x in Ti3C2T2Nax with T ¼ bare, F, and O. Reproduced with permission-Copyright 2016, American Chemical Society [42]. (f) HAADF image of Ti3C2X electrode upon Na intercalation with the cutoff potential of 0 V. Reproduced with permission-Copyright 2015, American Chemical Society [67]

of sodium (Na) intercalation between Ti3C2 MXene layers have been confirmed in the HAADF image of the Ti3C2X electrode (Fig. 16.4f) [67]. In general, the design of MXene-based nanostructure with enlarged interlayer spacing provides a new route to improve their energy storage performance. In this section, a series of interlayer structure design along with the electrochemical energy storage performance are systematically overviewed, including Li-ion batteries (LIBs), Na-ion batteries (SIBs), supercapacitors, Li/Na-ion capacitors (LICs/SICs), Na metal batteries (SMBs), potassium-ion batteries (PIBs), and magnesium ion batteries (MIBs).

16.4

Supercapacitors

Supercapacitors are regarded as one type of very promising energy storage technology since they possess higher power densities than batteries while their energy densities are superior to conventional capacitors [68]. Supercapacitors could play a role to bridge the energy and power gap between conventional capacitors and batteries [68]. MXenes’ unique structure renders them attractive for supercapacitors

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applications [69], including (1) A conductive inner transition metal carbide layer enables fast electron supply to electrochemically active sites; (2) A transition metal oxide-like surface is redox-active; (3) A 2D morphology and pre-intercalated water enable fast ion transport. Therefore, MXenes are promising candidates as electrodes for energy storage applications. Gogotsi et al. demonstrate both the spontaneous and electrochemical intercalation of cations, including Na+, K+, NH4+, Mg2+, and Al3+, from aqueous salt solutions between 2D Ti3C2 MXene layers (Fig. 16.5a) [54]. After cations intercalation, the interlayer spacing of Ti3C2 MXene increases. In basic solutions, such as KOH and NaOH, the binder-free Ti3C2 paper exhibits highly flexible and yields volumetric capacitance of up to 350 F cm
3, showing notable intercalation capacitances (Fig. 16.5b, c). Besides, the fact that a variety of ions, including Na+ and Al3+, can be intercalated between the MXene layers may also enable MXene use in batteries as well as in metal-ion capacitors (battery–supercapacitor hybrids) [54]. However, the sizes of metal ions are not large enough. If larger size molecules are used as intercalation agents, the interlayer spacing of MXene can be further increased, and the corresponding electrochemical performance can be further improved [43]. Yang and Gogotsi et al. used the mechanical shearing of a discotic lamellar liquid crystal phase of Ti3C2 to assemble vertically aligned 2D Ti3C2 (Fig. 16.5d) [70]. Notably, a non-ionic surfactant, hexamethylene glycol monododecyl ether (C12E6), was introduced to enhance molecular interactions between the vertically aligned MXene nanosheets, therefore increasing the interlayer spacing to 5.8 nm (Fig. 16.5e). Compared with the traditional electrode that has limited ion transport in thick films, the vertical alignment of MXene flakes enables directional ion transport, which can lead to thickness-independent electrochemical performances in thick films. The rate performance of the vertically aligned MXene films decreases only slightly when the thickness of film increases from 40 to 200μm, especially for the scan rates below 2000 mV s
1 (Fig. 16.5f) [70]. Therefore, the vertical alignment of functional nanomaterials with enlarged interlayer spacing through the manipulation of their liquid crystal mesophase confirmed here provides a new and powerful technique to construct advanced architectures with exceptional performance. Besides, polymers can also be used as intercalation agents between MXene layers. Ling et al. mixed a charged poly-diallyl dimethylammonium chloride (PDDA) or polyvinyl alcohol (PVA) and Ti3C2 MXene to produce Ti3C2/polymer composites (Fig. 16.5g) [59]. The obtained polymers intercalated MXene composite films that have excellent flexibility, good tensile and compressive strengths, and electrical conductivity. The successful intercalation of polymers between Ti3C2 layers, which increases the interlayer spacing of MXene and can be confirmed by the high-resolution transmission electron microscopy (HRTEM) (Fig. 16.5h). Used as electrodes for supercapacitors, typical cyclic voltammograms (CVs) of Ti3C2, Ti3C2/ PDDA, and Ti3C2/PVA-KOH films are shown in Fig. 16.5i. Compared with Ti3C2 film, the enhancement in the capacitance of Ti3C2/PVA-KOH film can be ascribed to the intercalation of PVA between MXene flakes with enlarged interlayer space, improving access to deep trap sites [59]. In general, the composite films exhibit impressive volumetric capacitance values as high as 528 F cm
3 at 2 mV s
1 and

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Fig. 16.5 (a) Schematic illustration of the intercalation of cations between Ti3C2Tx MXene layers. (b) SEM image of the Ti3C2Tx MXene paper electrode. (c) CV curve of Ti3C2Tx paper in the KOH electrolyte. (d) Ion transport in vertically aligned Ti3C2Tx MXene films. (e) SEM image of vertical nanosheets on the horizontally aligned MXene current collector. (f) Rate performance of MXLLC films at scan rates ranging from 10 to 100,000 mV s
1. Reproduced with permission-Copyright 2019, Springer Nature [70]. (g) Schematic illustration of MXene/PVA film with outstanding properties. (h) TEM images of Ti3C2Tx/PVA. (i) CV curves of electrodes were obtained at a scan rate of 2 mV s
1. (j) Schematic illustration of the structure of polypyrrole-Ti3C2Tx. (k) Crosssectional TEM image of polypyrrole-Ti3C2Tx. (i) Rate performance of tested compositions and comparisons of their capacitances with previously reported Ti3C2Tx electrode. Reproduced with permission-Copyright 2016, John Wiley and Sons [60]

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306 F cm
3 at 100 mV s
1. Moreover, Boota et al. synthesized polypyrrole (PPy) intercalated Ti3C2 composite by simultaneous intercalation, alignment, and metalfree polymerization of pyrrole between Ti3C2 layers (Fig. 16.5j, k) [60]. When tested as supercapacitor electrodes, a PPy/Ti3C2 composite yields a high volumetric capacitance of ~1000 F cm
3 and capacitance retention of 92% after 25,000 cycles (Fig. 16.5l). The high electrochemical performance and excellent cycling performance of PPy/Ti3C2 composite are ascribed to the highly conductive nanostructure of the PPy/Ti3C2 film in which the aligned PPy chains between the metallic Ti3C2 layers provide low electronic resistance and create aligned channels for facile ionic transport. 2D conductive Ti3C2 layers may protect the intercalated PPy chains and minimize their degradation. Besides, other intercalation agents, such as N heteroatom, activated carbon, CNT, and rGO, can also be used to design MXene-based nanostructure with enlarged interlayer spacing [33, 61, 63, 64]. The above MXene-based electrodes all exhibit high capacitances and perform well at high rates due to additional and fast diffusion paths for electrolyte ions as well as large interlayer space for cations intercalation.

16.5

Li-ion Batteries/Capacitors

There are continuous efforts in the pursuit of high-performance Li-ion batteries/ capacitors with high energy/power density to meet the increasing demand for modern electronics and transportation [43]. In the family of MXenes, Ti3C2 is the most studied material as anode for LIBs [12, 44, 71]. Naguib et al. first used Ti3C2 MXene as the anodes for LIBs [71]. At a current density of 1 C, a stable capacity of 110 mAh g
1 can be obtained after 80 cycles. When the current density increases to 10 C, a capacity of 70 mAh g
1 can still be obtained for the Ti3C2 electrode. Although the capacity of Ti3C2 MXene is relatively low at the initial stage of research, there is a huge room for improvement. Luo et al. first synthesized the Sn4+ pillared MXene (PVP-Sn(IV)@Ti3C2) by the use of alkali metal ions to pre-pillar Ti3C2 MXene, followed by the Sn4+ pillaring process (Fig. 16.6a) [44]. By the ion-exchange interaction between Sn4+ and pre-intercalated alkali metal ions under the help of polyvinylpyrrolidone (PVP), Sn4+ can be successfully intercalated between Ti3C2 layers. The intercalation of Sn4+ in the interlayer of Ti3C2 MXene increases the interlayer spacing of Ti3C2 from 0.98 to 1.28 nm by forming strong Ti-O-Sn bonding. The SEM image shows that Sn4+, in the form of Sn (IV) nanocomplex, is uniformly distributed on the Ti3C2 matrix (Fig. 16.6b). Used as the anodes for LIBs, the PVP-Sn(IV)@Ti3C2 nanocomposites exhibit a superior reversible volumetric capacity of 1375 mAh cm
3 at 100 mA g
1 after 100 cycles (Fig. 16.6c). Even at a high current density of 3 A g
1, a stable specific capacity of 504.5 mAh cm
3 can be retained. The excellent electrochemical performance can be ascribed to the “pillar effect.” During the lithiation process, Li+ can alloy with the intercalated Sn4+ to form Li-Sn alloy, causing the volume expansion, which can effectively prop the MXene layers open and endow the MXene with increased

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Fig. 16.6 (a) Schematic illustration of the structure of Sn4+ pillared Ti3C2 MXene (PVP-Sn(IV) @Ti3C2). (b) SEM image of PVP-Sn(IV)@Ti3C2. (c) Rate performance of PVP-Sn(IV)@Ti3C2. Reproduced with permission-Copyright 2016, American Chemical Society [44]. (d) Schematic illustration of the intercalation of Co2+ between V2C MXene layers. (e) SEM image of Co2+ intercalated V2C (V2C@Co). (f) Rate performance of Co2+ intercalated V2C. Reproduced with permission-Copyright 2018, John Wiley and Sons [57]. (g) Schematic illustration of the structure of Sn4+ pillared Ti3C2 MXene (CTAB-Sn(IV)@Ti3C2). (h) STEM image of CTAB-Sn(IV)@Ti3C2. Inset is the elemental mapping of Sn. (i) Cycling performance of CTAB-Sn(IV)@Ti3C2. Reproduced with permission-Copyright 2016, American Chemical Society [43]. (j) Schematic illustration of the sandwich-like MXene/CNT papers. (k) Cross-sectional SEM image of Nb2CTx/ CNT composite paper. (l) Cycling stability of the Nb2CTx/CNT paper electrode at different cycling rates. Reproduced with permission-Copyright 2015, John Wiley and Sons [50]

interlayer space for Li+ intercalation, namely the “pillar effect” [44]. Besides, Co2+ can also be intercalated between the MXene layers [57]. Similarly, Song et al. fabricated Co2+ intercalated V2C MXene (V2C@Co) (Fig. 16.6d) [57]. After Co2+

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intercalation, strong V–O–Co bonds can be formed between Ti3C2 MXene layers. The SEM image of V2C@Co displays no obvious layer stripes, mainly attributes to the intercalation of Co2+ (Fig. 16.6e). Owing to the unique layered structure and pseudocapacitive contribution of Co species, the obtained V2C@Co exhibits improved Li-ion capacity and delivers a high capacity of 1117.3 mAh g
1 at 0.1 A g
1 and a significantly ultralong cycling stability over 15,000 cycles (Fig. 16.6f) [57]. Although alkali metal ion can pre-pillar and increase the interlayer spacing of MXene, the increments of interlayer spacing are limited due to their small size. Luo et al. selected different sizes of cationic surfactants to intercalate between MXene layers and controlled the interlayer spacing of MXene from 1 to 2.71 nm [43]. The intercalated cationic surfactants in the interlayer of MXene can provide the ion-exchange sites for other metal ions. Tao et al. successfully fabricated Sn4+ pillared Ti3C2 MXene (CTAB-Sn(IV)@Ti3C2) by using liquid-phase CTAB pre-pillaring and subsequent Sn4+ pillaring process (Fig. 16.6g) [43]. STEM image of CTAB-Sn(IV)@Ti3C2 confirms the successful intercalation of Sn4+ between Ti3C2 layers (Fig. 16.6h). As anode for lithium-ion storage, the obtained CTABSn(IV)@Ti3C2 electrode delivers a high reversible capacity of 765 mAh g
1 at 0.1 A g
1 after 100 cycles, which exhibits a higher capacity of CTAB intercalated Ti3C2 (248 mAh g
1) (Fig. 16.6i). The excellent Li-ion storage performance of CTAB-Sn(IV)@Ti3C2 is ascribed to the “pillar effect.” Besides, the CTAB-Sn(IV) @Ti3C2 anode exhibits good Li+ storage kinetics. Coupling the CTAB-Sn(IV) @Ti3C2 anode with commercial AC cathode, the assembled LICs exhibits a high energy density of 239.50 Wh kg
1 [43]. Besides, Gogotsi’s group used a layer-by-layer method to assemble a CNT/Nb2C MXene composite “paper” electrode (Fig. 16.6j) [50]. THE cross-sectional SEM image of Nb2C/CNT composites exhibits an open structure (Fig. 16.6k). After CNT intercalation between Nb2C layers, the obtained Nb2C/CNT papers exhibit good electrical conductivity, high surface areas accessible to ions, and mechanical robustness, suggesting their promising application for Li-ion storage. The obtained CNT/MXene composite “paper” electrode exhibits a high capacitance of 325 F cm
3 as an electrode for LICs (Fig. 16.6l). Also, other metal ions and nanomaterials can be used as the intercalation agents to build MXene-based nanostructures with enlarged interlayer spacing and used as electrodes for LIBs and LICs [72, 73]. Compared with aqueous electrolytes, the size of the charge carrying ions in the organic electrolyte is larger [31]. Therefore, constructing MXene-based nanostructures with enlarged interlayer spacing is crucially important to maximize their electrochemical performance.

16.6

Na-ion Batteries/Capacitors

Na-ion storage is a promising alternative to Li-ion storage due to the similar working mechanism and significant advantages as for elemental abundance and low cost of Na resources [36, 74–76]. However, compared with the smaller radius of Li ions

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Fig. 16.7 (a) The structure of MXene/CNT composite. (b) SEM image of porous Ti3C2Tx/CNTSA. (c) Volumetric and gravimetric rate performances. Reproduced with permission-Copyright 2019, John Wiley and Sons [36]. (d) The structure of S atoms pillared Ti3C2 MXene (CT-S@Ti3C2). (e) STEM image of CT-S@Ti3C2-450. Inset is the elemental mapping of S. (f) Rate performance of CT-S@Ti3C2-450. Reproduced with permission-Copyright 2015, Springer Nature [17]. (g) Schematic illustration of the reaction mechanism of Ti2CTx by electrochemical activation. (h) TEM image of activated Ti2CTx after the first CV. (i) Rate capability for Ti2CTx with other reported electrodes. Reproduced with permission-Copyright 2016, Elsevier [79]

(0.76 Å), Na ions have a much larger radius (1.02 Å), which makes it difficult to find appropriate host materials for Na ions storage due to the large volume expansion and sluggish reaction kinetics. Therefore, seeking novel electrode materials for Na-ion storage is still desirable. Fortunately, MXene is reported to be a promising anode material for SIBs and SICs [36]. Heteroatoms intercalated MXene can be fabricated with enhanced Na-ion storage performance. Luo et al. successfully intercalated S atoms into the interlayer of Ti3C2 MXene (CT-S@Ti3C2) by a facile CTAB pretreatment, thermal diffusion with elemental S, and subsequent annealing process (Fig. 16.7a) [36]. After annealing under the Ar atmosphere, S atoms pillared Ti3C2 with interlayer-expanded structure via Ti-S bonding can be developed. The interlayer spacing of CT-S@Ti3C2 is about 1.37 nm, which is larger than that of CTAB pretreated Ti3C2 after annealing under

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Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

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the same temperature (1.15 nm). The successful intercalation of S atoms between Ti3C2 MXene layers can be confirmed by the STEM image and the corresponding elemental mapping results of CT-S@Ti3C2 (Fig. 16.7b). Used as the anode for Na-ion storage, the CT-S@Ti3C2 delivers a high capacity of 550 mAh g
1 at 0.1 A g
1v (120 mAh g
1 at 15 A g
1) and excellent cycling stability over 5000 cycles at 10 A g
1 (Fig. 16.7c). The excellent electrochemical performance is ascribed to the strong Ti-S bonds, which not only can react with the intercalated Na that contribute to the “pillar effect” but also can provide stable double-layer Na ions adsorption that is confirmed by DFT calculation [36]. Furthermore, the obtained CT-S@Ti3C2 exhibits fast Na ion storage kinetics, which has been confirmed by the kinetic analysis. A high energy density of 263.2 Wh kg
1 can be obtained for the assembled SICs by coupling CT-S@Ti3C2 anode with commercial AC cathode. In addition to the introduction of heteroatoms, alkali metal ions can also be intercalated between the MXene layers and alter their electrochemical behaviors [17, 77]. Yamada et al. realized the usage of MXene as the electrode in a nonaqueous Na+ electrolyte. The results show that after the initial Na+ intercalation and activation processes, the Ti2C MXene electrode exhibits pseudocapacitor behavior (Fig. 16.7d) [17]. After the initial Na+ intercalation, the interlayer spacing of Ti2C increases from 0.8 to 1.0 nm, which can be confirmed by the TEM image (Fig. 16.7e). Compared with the rate capability of Ti2C with other electrode materials, such as hard carbon, expanded graphite, and P2-Na0.66[Li0.22Ti0.78]O2, Ti2C delivers higher capacity at higher current density. Therefore, Ti2C is a high-performance electrode material with a high capacity, stability, and high power for Na-ion storage. Luo et al. studied the Na+ storage behaviors of different alkali metal ions (e.g., Li+, Na+, K+) pillared Ti3C2 MXene [78]. Meanwhile, they annealed the Na+ pillared Ti3C2 MXene at different temperatures (e.g., 450 and 700  C) to regulate the terminal groups in Ti3C2 (Fig. 16.7f). The results show that Na+ intercalation during the pillaring process can increase the number of active sites for Na+ storage, and it also can decrease the Na+ diffusion barrier. In general, the results confirm that the existence of –OH groups can provide active sites for Na+ storage, while decrease the Na+ storage kinetics [78]. Compared with the most used liquid phase intercalation method to build MXenebased nanostructure with enlarged interlayer spacing, layer-by-layer assembly is more flexible because many nanomaterials can be used as intercalation agents by this method. Xie et al. prepared Ti3C2 MXene/CNT films by self-assembly of negatively charged Ti3C2 MXene and positively charged CNTs as intercalation spacers (Fig. 16.7g) [79]. In the SEM image of Ti3C2/CNT, an open sheet arrangement and CNTs can be seen sandwiched between Ti3C2 sheets, indicating the successful placement of CNTs as intercalation spacers in the Ti3C2 paper (Fig. 16.7h). The 2D/1D hybridization between MXene nanosheets and CNTs reduces the stacking of nanosheets produces hierarchical films with a porous structure, thereby improving the accessibility of MXene nanosheets to the electrolyte. Used as free-standing electrodes for Na-ion storage, the porous MXene/CNT papers exhibit a high volumetric capacity of 421 mAh cm
3 at 20 mA g
1 (Fig. 16.7i) [79].

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Other Energy Storage Systems

Except for supercapacitors, LIBs, LICs, SIBs, and SICs, MXene-based nanostructures with enlarged interlayer spacing can also be used as the electrodes for other energy storage applications, including SMBs, PIBs, and MIBs [65, 80, 81]. Luo et al. used the liquid phase CTAB pre-pillaring and Sn2+ pillaring method to fabricate Sn2+ pillared Ti3C2 MXene (CT-Sn(II)@Ti3C2) and serve as a stable matrix for Na metal anode [80]. After Sn2+ intercalation, the interlayer spacing of Ti3C2 MXene increases from 0.98 to 1.9 nm. The intercalated Sn2+ in this work not only induces Na to nucleate and grow between Ti3C2 MXene layers but also endows Ti3C2 with larger interlayer space to accommodate the deposited Na under the effect of “pillar effect,” thus resulting in the uniform Na deposition [80]. Xu et al. used Ti3C2 MXene as a multifunctional binder to fabricate a free-standing and flexible hard carbon film electrode by a simple vacuum-assisted filtration strategy [65]. In the MXene-bonded hard carbon film, hard carbon film particles, as the active materials, are embedded in the 3D conductive network of MXene sheets. MXene nanosheets with enlarged interlayer spacing can effectively stabilize the electrode structure and accommodate the volume expansion of hard carbon particles during the K+ intercalation. Used as the anode for PIBs, the MXene-bonded hard carbon anode delivers a high reversible capacity of 210 mAh g
1 at the current density of 50 mA g
1 after 100 cycles with a capacity retention of 84%. Besides, Yan et al. used CTAB to intercalate between Ti3C2 MXene layers and increased the interlayer spacing of Ti3C2 [81]. The obtained CTAB intercalated Ti3C2 MXene exhibits a high Mg2+ storage capability. DFT simulations verify that the intercalated CTA+ cations reduce the diffusion barrier of Mg2+ on the MXene surface. Consequently, the MXene electrode exhibits a desirable volumetric specific capacity of 300 mAh cm
3 at 50 mA g
1.

16.8

Conclusions and Perspectives

MXenes have been widely applied in the field of electrochemical energy storage due to their high electrical conductivity, large redox-active surface area, rich surface chemistry, and tunable structures. Demonstrably, in recent years, there has been a rapid increase in the number of publications on MXenes and MXene for electrochemical energy storage applications (Fig. 16.2d). In this chapter, we first briefly summarize the synthesis techniques of MXenes, mainly engross in the HF/in situ HF etching synthesis. Different synthesis techniques used to synthesize MXene directly influence the structure and properties of

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Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . .

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the obtained MXenes. In the research process of layered structure MXene for electrochemical energy storage, a better comprehension of the layered structure and interfacial chemistry of MXene is requisite. Up to now, the interfacial structure design of MXene-nanostructures with enlarged interlayer spacing can be classified according to different intercalation agents, including (1) heteroatoms; (2) alkali metal ions; (3) non-alkali metal ions; (4) polymers; (5) large cations; (6) 1D/2D/ 3D nanomaterials. The main fabrication methods of MXene-based nanostructures with enlarged interlayer spacing, including (1) liquid phase pre-pillaring and pillaring; (2) precursor intercalation and annealing/polymerization; (3) layer-bylayer assembly. After the intercalation of intercalation agents, MXene-based nanostructure with enlarged interlayer spacing shows obvious characteristics for energy storage applications: (1) large interlayer space for charge-carrying ions intercalation and storage; (2) excellent ion transport due to short diffusion pathways; (3) limitation of volume expansion of active materials by the confinement effect of a 2D structure during cycling. Besides, by carefully selecting the types of surfactant molecules, pillared MXene structures with controllable interlayer spacings from 1 to 2.7 nm can be prepared to match the sizes of the ions in the electrolyte, maximizing the electrochemical performance. Furthermore, this fabrication method of pillared MXene is conducive to the controllable modification of the surface chemistry about MXene, which is one of the main challenges in designing MXene-based nanostructure. Interlayer structural engineering of 2D MXene, synthesis methods, and their electrochemical energy storage applications are summarized in Table 16.1. Although the research of designing MXene-based nanostructure for electrochemical energy storage applications has been widely studied, some remaining challenges still need to be overcome for the future development of MXene. For future research, the following aspects are particularly important. (1) Achieving controllable and uniform terminations on the interface of MXene is necessary. (2) A thorough understanding and systematical study on the properties of intercalation agents (e.g., ionic, molecular species), including transport, stability, and bonding environment, and their effect on the electrochemical and physicochemical properties of layered structure MXenes is necessary. (3) A basic understanding of the property of active centers (interfacial transition metal layers and terminations) for ion transport and intercalation, and the energy storage mechanism by the cooperation between theoretical calculation and experiment lays solid foundations for the development of MXenes in electrochemical energy storage.

Polymer (PVA, PDDA, PPy)

Metal ions (Sn4+, Co2+, Sn2+)

Alkali metal ions

Intercalation agent types Heteroatoms (N, S)

PPy/l-Ti3C2

PPy-Ti3C2Tx

Ti3C2Tx/ PDDA Ti3C2Tx/PVA

CT-Sn(II) @Ti3C2

Multilayer Ti3C2 PVP-Sn(IV) @Ti3C2 V2C@Co

Materials N-Ti3C2Tx CT-S@Ti3C2

Liquid phase mixture (RT) Liquid phase mixture (RT) Precursor intercalation and in situ polymerization Electrochemical polymerization

Methods NH3 annealing (200  C) Precursor intercalation and annealing (450  C) Liquid phase intercalation (RT) Intercalation by ion exchange (RT) Intercalation by ion exchange (RT) Intercalation by ion exchange (RT)

0.96

Supercapacitors

Supercapacitors

Supercapacitors

2.25 2.40

Supercapacitors

1.84

Na metal matrix

Li-ion batteries

0.95 1.90

Li-ion batteries

Applications Supercapacitors Na-ion capacitors Supercapacitors

1.28

1.25

Interlayer spacing 1.23 1.37

59

530 F cm
3 at 2 mV s
1

35 mF cm
2

~1000 F cm
3 with capacitance retention of 92% after 25,000 cycles

59

58

60

80

57

44

54

References 61 36

1375 mAh cm
3 (635 mAh g
1) at 0.1 A g
1 1117.3 mAh g
1 at 0.1 A g
1, 199.9 mAh g
1 at 20 A g
1 High coulombic efficiency over 500 cycles (up to 10 mA cm
2 and 5 mAh cm
2) ~296 F cm
3 at 2 mV s
1

Performance 192 F g
1 in 1 M H2SO4 550 mAh g
1 at 0.1 A g
1 and ~120 mAh g
1 at 15 A g
1 ~350 F cm
3 after 10 K cycles

Table 16.1 Interlayer structural engineering of 2D MXene, synthesis methods, and their electrochemical energy storage applications

472 J. Luo

1D/2D nanomaterials

Large cations (CTA+, STA+, DDA+, amine ions)

MoS2-in-Ti3C2

Nb2CTx + iPrA + CNTs Ti3C2Tx/ MWCNT MXene-rGO

DDAC@Ti3C2

STAB-Sn(IV) @Ti3C2

CTAB-Sn(IV) @Ti3C2

Electrostatic selfassembly Annealing (500  C)

Liquid phase pre-pillaring and pillaring (40  C) Liquid phase pre-pillaring and pillaring (50  C) Liquid phase intercalation (70  C) Liquid phase mixture (RT) Lay-by-layer assembly –

2.10

1.22

1.67

1.42

Li/Na-ion batteries

Supercapacitors

Supercapacitors

Li-ion batteries

Li-ion capacitors

2.71

2.27

Li-ion capacitors

2.23

43

~450 mAh g
1 at 1 A g
1

63

1040 F cm
3 at 2 mV s
1

72

33

435 F cm
3 at 2 mV s
1

340 mAh g
1 at 20 A g
1 for Li+ storage; 310 mAh g
1 at 1 A g
1 for Na+ storage

50

~420 mAh g
1 at 0.5 C

–

43

765 mAh g
1 at 0.1 A g
1

16 Interlayer Structural Engineering of 2D MXene for Electrochemical Energy. . . 473

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J. Luo

Acknowledgments I wish to thank the support of the Thayer School of Engineering, Dartmouth College, 14 Engineering Drive, Hanover, New Hampshire. Author Contribution The article was conceived and written by JL.

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59. Ling Z, Ren CE, Zhao MQ, Yang J, Giammarco JM, Qiu JS, Barsoum MW, Gogotsi Y (2014) Flexible and conductive MXene films and nanocomposites with high capacitance. Proc Natl Acad Sci U S A 111:16676–16681 60. Boota M, Anasori B, Voigt C, Zhao MQ, Barsoum MW, Gogotsi Y (2016) Pseudocapacitive electrodes produced by oxidant-free polymerization of pyrrole between the layers of 2D titanium carbide (MXene). Adv Mater 28:1517–1522 61. Wen YY, Rufford TE, Chen XZ, Li N, Lyu MQ, Dai LM, Wang LZ (2017) Nitrogen-doped Ti3C2Tx MXene electrodes for high-performance supercapacitors. Nano Energy 38:368–376 62. Tian WQ, VahidMohammadi A, Wang Z, Ouyang LQ, Beidaghi M, Hamedi MM (2019) Layer-by-layer self-assembly of pillared two-dimensional multilayers. Nat Commun 10:2558 63. Yan J, Ren CE, Maleski K, Hatter CB, Anasori B, Urbankowski P, Sarycheva A, Gogotsi Y (2017) Flexible MXene/graphene films for ultrafast supercapacitors with outstanding volumetric capacitance. Adv Funct Mater 27:1701264 64. Yu LY, Hu LF, Anasori B, Liu YT, Zhu QZ, Zhang P, Gogotsi Y, Xu B (2018) MXene-bonded activated carbon as a flexible electrode for high-performance supercapacitors. ACS Energy Lett 3:1597–1603 65. Sun N, Zhu Q, Anasori B, Zhang P, Liu H, Gogotsi Y, Xu B (2019) MXene-bonded flexible hard carbon film as anode for stable Na/K-ion storage. Adv Funct Mater 29:1906282 66. Xie Y, Naguib M, Mochalin VN et al (2014) Role of surface structure on Li-ion energy storage capacity of two-dimensional transition-metal carbides. J Am Chem Soc 136:6385–6394 67. Wang XF, Shen X, Gao YR, Wang ZX, Yu RC, Chen LQ (2015) Atomic-scale recognition of surface structure and intercalation mechanism of Ti3C2X. J Am Chem Soc 137:2715–2721 68. Simon P, Gogotsi Y, Dunn B (2014) Where do batteries end and supercapacitors begin? Science 343:1210–1211 69. Lukatskaya MR, Kota S, Lin ZF et al (2017) Ultra-high-rate pseudocapacitive energy storage in two-dimensional transition metal carbides. Nat Energy 2:17105 70. Xia Y, Mathis TS, Zhao MQ, Anasori B, Dang A, Zhou ZH, Cho H, Gogotsi Y, Yang S (2018) Thickness-independent capacitance of vertically aligned liquid-crystalline MXenes. Nature 557:409–412 71. Naguib M, Come J, Dyatkin B, Presser V, Taberna P-L, Simon P, Barsoum MW, Gogotsi Y (2012) MXene: a promising transition metal carbide anode for lithium-ion batteries. Electrochem Commun 16:61–64 72. Ma K, Jiang H, Hu YJ, Li CZ (2018) 2D nanospace confined synthesis of pseudocapacitancedominated MoS2-in-Ti3C2 superstructure for ultrafast and stable Li/Na-ion batteries. Adv Funct Mater 28:1804306 73. Zhang CF, Park SN, Seral-Ascaso A, Barwich S, McEvoy N, Boland CS, Coleman JN, Gogotsi Y, Nicolosi V (2019) High capacity silicon anodes enabled by MXene viscous aqueous ink. Nat Commun 10:849 74. Meng RJ, Huang JM, Feng YT et al (2018) Black phosphorus quantum dot/Ti3C2 MXene nanosheet composites for efficient electrochemical lithium/sodium-ion storage. Adv Energy Mater 8:1801514 75. Li WH, Yang ZZ, Li MS, Jiang Y, Wei X, Zhong XW, Gu L, Yu Y (2016) Amorphous red phosphorus embedded in highly ordered mesoporous carbon with superior lithium and sodium storage capacity. Nano Lett 16:1546–1553 76. Chen Z, Augustyn V, Jia XL, Xiao QF, Dunn B, Lu YF (2012) High-performance sodium-ion pseudocapacitors based on hierarchically porous nanowire composites. ACS Nano 6:4319–4327 77. Okubo M, Sugahara A, Kajiyama S, Yamada A (2018) MXene as a charge storage host. Acc Chem Res 51:591–599

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78. Luo JM, Fang C, Jin CB et al (2018) Tunable pseudocapacitance storage of MXene by cation pillaring for high-performance sodium-ion capacitors. J Mater Chem A 6:7794–7806 79. Xie XQ, Zhao MQ, Anasori B et al (2016) Porous heterostructured MXene/carbon nanotube composite paper with high volumetric capacity for sodium-based energy storage devices. Nano Energy 26:513–523 80. Luo JM, Wang CL, Wang H, Hu XF, Matios E, Lu X, Zhang WK, Tao XY, Li WY (2019) Pillared MXene with ultralarge interlayer spacing as a stable matrix for high-performance sodium metal anodes. Adv Funct Mater 29:1805946 81. Xu M, Lei SL, Qi J, Dou QY, Liu LY, Lu YL, Huang Q, Shi SQ, Yan XB (2018) Opening magnesium storage capability of two-dimensional MXene by intercalation of cationic surfactant. ACS Nano 12:3733–3740

Jianmin Luo received his PhD in 2018 from the Zhejiang University of Technology in China. He is currently serving as a postdoctoral associate in the Thayer School of Engineering at Dartmouth College. His research is mainly focused on the design of 2D nanomaterials for energy storage and conversion. Dr. Luo has published more than 30 papers in journals (such as Nano Letters, ACS Nano, Advanced Functional Materials, Chemical Society Review) with a total citation of over 2000 and an H-index of 20. He also peer reviews Journals including ACS Nano, Advanced Functional Materials and Small, etc. Besides, Dr. Luo holds two issued Chinese patents as well.

Chapter 17

The Role of Ex Situ Solid Electrolyte Interphase in Lithium Metal Batteries Rajesh Pathak, Yue Zhou, and Qiquan Qiao

17.1

Introduction

17.1.1 Overview Energy is a key to drive today’s world, from consumer applications such as portable smartphones and laptops to industrial applications such as transportation, large-scale energy storage micro-grids, and medical equipment [1–3]. To fulfill the everincreasing energy demand, fossil fuels are being widely consumed. However, the use of fossil fuels is responsible for global climate change, fluctuation in the price, and concern about the availability of limited fossil fuels shortly. To address these concerns, widespread renewable energy sources such as solar, wind, and hydropower need to be deployed. For the continuous supply of energy, the energy storage device is aiding the integration of renewable energy into the electricity system [4– 8]. As a result, in the upcoming decades, fossil fuels can become history, and electric energy can drive the whole world.

R. Pathak Department of Electrical Engineering and Computer Science, South Dakota State University, Brookings, SD, USA Applied Materials Division, Argonne National Laboratory, Lemont, IL, USA e-mail: [email protected] Y. Zhou Department of Electrical Engineering and Computer Science, South Dakota State University, Brookings, SD, USA e-mail: [email protected] Q. Qiao (*) Mechanical and Aerospace Engineering, Syracuse University, Syracuse, NY, USA e-mail: [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_17

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The lithium (Li)-ion batteries (LIBs) are considered a reliable energy storage device due to their high energy density, power density, and long cycle life. LIBs are widely used in a variety of applications such as cellphones, laptops, electric vehicles, and hybrid electric vehicles. The LIBs are gradually reaching the upper limit of specific energy density (~ 300 Wh kg
1). For the practical applications of high energy battery applications such as long drive range electric vehicles with a single charge, a high specific energy density of more than 500 Wh kg
1 is needed. This requires aggressive electrochemical chemistries between the conversion reaction or high-voltage/high-capacity intercalation cathode and Li as an anode. Li-metal is an ideal anode material for rechargeable batteries owing to its high theoretical specific capacity of 3860 mAh g
1, low density of 0.53 g cm-3, and extremely low redox potential of
3.04 V vs. standard hydrogen electrode [9–12]. Thus, it is imperative to replace the conventional graphite anode with a lithium metal anode (LMA). Figure 17.1a shows that the use of lithium metal as an anode can provide high specific capacity and high operational voltage, leading to the high energy density of the batteries.

17.1.2 Challenges in Lithium Metal Batteries The direct use of Li metal as a node has shown a great challenge owing to the highly reactive nature, the infinite volume change of Li, and the formation of undesired lithium dendrite growth and unstable fragile/unstable SEI. Firstly, the hyperactive lithium spontaneously reacts with electrolyte, active material from the cathode, and other side reactions. The irreversible reaction consumes electrolyte, lithium, and active materials leading to electrolyte dry-out and quick capacity fading [8, 13– 18]. Secondly, the infinite volume expansion of Li leads to the pulverization of the anode, severe lithium corrosion, and formation of dead lithium. The loss of active materials, undesired side products, and the formation of inactive dead lithium increase the impedance of the cell. Thirdly, the nonuniform electric field due to the rough lithium surface induces the hot spots for Li nucleation which continues to grow Li dendrite with increased cycling. The needle or whisker-shaped dendrites can pierce the separator causing the short circuit or can catch fire. Also, the high surface area of Li dendrite growth and continuous formation/deformation of SEI consumes much electrolyte and lithium causing the electrolyte dry-out and low CE. Numerous efforts have been done for dense, reversible, and dendrite-free Li deposition. Figure 17.1b shows the volume expansion issues associated with the lithium metal anode during plating/stripping cycles. The volume expansion issues could lead to further corrosion, cracks on the electrode responsible for further Li dendrite growth. Figure 17.1c demonstrates the formation of Li dendrites and dead Li during plating/ stripping cycles using Li metal as the anode. Thus, it is essential to protect the LMA to buffer the volume expansion and suppress Li dendrite growth in LMBs.

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Fig. 17.1 Importance and challenges in using lithium metal as the anode. (a) Capacity and voltage comparison of typical electrodes. The inset shows the specific energy density comparison between LiBs and LMBs. (Reproduced from Ref. [19] with permission from John Wiley and Sons). (b) The pulverization of LMA due to infinite volume change in a typical Li stripping/plating cycle. (Reproduced from Ref. [20] with permission from National Academy of Sciences). (c) The formation of dendritic, mossy, and dead Li during the plating/stripping cycle. (Reproduced from Ref. [13] with permission from Elsevier)

17.1.3 Strategies to Revive LMA Substantial efforts have been done by many researchers as demonstrated in Fig. 17.2. to address the inherent issues related to the LMA to achieve their potential applications. For guiding the uniform Li deposition and improved cycling performance, the common approaches are (1) the development of lithiophilic and conductive threedimensional (3D) micro-/nanostructured framework to guide the uniform Li deposition and accommodate the volume expansion, (2) the engineering of an artificial protective layer on top of Li by ex situ/in situ which can greatly inhibit the side

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Fig. 17.2 Strategies to address the inherent problem of LMA

reactions, (3) the replacement of inflammable liquid electrolyte with higher Young’s modulus solid-state electrolyte, and (4) other strategies such as modification of separator and defining the battery testing conditions.

17.1.3.1

The Development of 3D Micro-/Nanostructured Li host

The engineering of 3D porous metal-based or carbon-based Li hosts has proven to be a successful technique to accommodate the Li for dendrite-free Li deposition. Figure 17.3a, b shows the schematic representation of the Li-ion flux distribution and the Li plating on the planar and freestanding copper nanowire (CuNW). The uniform distribution of the electric field due to the nanowires leads to the uniform Li-ion flux and homogeneous Li deposition [21]. Besides, the high surface area of the interconnected CuNW framework significantly lowers the effective current density. Moreover, the CuNW accommodates the Li inside the porous nanostructure to impede the Li dendrite growth. In contrast, the planar Li has an uneven hot spot which creates a nonuniform electric field and results in the nonuniform Li-ion flux distribution and inhomogeneous Li deposition. The commonly used metallic or carbon-based porous frameworks show a poor affinity for Li metal. Thus, the poor Li wettability/lithiophobic porous framework suggests high nucleation

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Fig. 17.3 3D nano-/microstructured Li host. (a, b) Schematic representation of the Li-ion flux distribution and Li-plating behavior on planar and CuNW framework. (Reproduced from Ref. [21] with permission from American Chemical Society). (c) Schematic illustration of Li plating on VA-CuO NSs-Cu. (Reproduced from Ref. [22] with permission from John Wiley and Sons). (d) Schematic illustration of Li plating on CNF-Cu. (Reproduced from Ref. [11] with permission from John Wiley and Sons). (e) Schematic illustration of Li plating on ultrafine AgNP decorated CNF. (Reproduced from Ref. [26] with permission from John Wiley and Sons). (f) Schematic illustration of fabricating Li/c-MOF composite. (g, h) Voltage-capacity profile of Li deposition on C-MOF at 10 μA cm
2 with and without lithiophilic Zn clusters. (Reproduced from Ref. [27] with permission from John Wiley and Sons). (i) Schematic representation of fabricating Li/porous carbon composite. (Reproduced from Ref. [20] with permission from National Academy of Science)

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overpotential, voltage hysteresis, and unfavorable Li deposition. As a result, the Li plating/stripping cycles show to lower the energy barrier, provide the nucleation seed for Li growth, and guide the uniform Li deposition in a porous framework; the use of lithiophilic coating or decoration with heterogeneous seed growth is essential. This significantly reduces the nucleation overpotential and voltage hysteresis. To further enhance the Li deposition, the thin layer of vertically aligned (VA) lithiophilic CuO nanosheets (NSs) was grown on the Cu [22]. The lithiophilic CuO-NSs grown on Cu as shown in the schematic Figure 17.3c assure the fast Li-ion diffusion, and the Cu current collector guarantees the fast electron transport. The high Li-ion affinity of lithiophilic CuO promotes steady Li nucleation and deposition. During Li plating/stripping, the reversible conversion reaction facilitates fast Li storage. The reversible conversion (Cu2O, Cu3N, Ni3N) and alloy (ZnO, MgO, Si3N4, Sn, Au, Si, Ag) forming lithiophilic materials have successfully reported enhanced dendrite-free Li deposition, improving the battery cycling performance. Despite the attractive features of the metal-based current collector, the high cost, high mass density (Cu, Ni, etc.), and the complex synthesis process of 3D structure hinder their further practical applications. In contrast, the carbon is naturally abundant, has a lower cost and lower density, is easier for developing a porous framework, and thus is considered as a promising candidate for Li host. The emerging porous carbons such as hollow carbon spheres, carbon nanofibers (CNFs), graphitized CNFs, carbon nanotubes (CNTs), and carbonized MOFs have been demonstrated feasible as Li host owing to their good conductivity, large surface area, flexibility, and sturdiness [23, 24]. The high surface area or voids provide sufficient space for Li deposition, and the flexibility can buffer the volume expansion of Li. Chen et al. designed the vertically aligned CNFs on Cu foil for uniform and dendrite-free Li plating as shown in Figure 17.3d. The conductivity difference between the electron-conducting current collector and less conducting carbonaceous interlayer facilitates the deposition of Li underneath the interlayer [11, 12]. The interlayer should have higher electrolyte wettability, lithiophilicity, and Li-ion diffusivity. The semiconducting interlayer does not favor the Li deposition on top of it as a separate film during initial Li plating. Instead, Li-ion diffuses and starts depositing from the bottom. Recently, mixed ionic conductors/interlayers which allow fast ionic and electronic transport have also attracted great attention in regulating the dendrite-free Li deposition. To avoid the cost and lower the density of the batteries, free-standing porous carbonaceous materials with enhanced conductivity are supposed to be efficient Li hosts. Moreover, the lithiophobic carbon cannot uniformly guide the initial Li nucleation and leads to higher nucleation and voltage overpotential. Thus, the lithiophilic decoration on such porous carbon materials is one of the promising strategies [25]. Yang et al. reported the silver nanoparticles (AgNPs) anchored CNF by treating the silver acetate-CNF by electric joule heating [26]. As lithiophilic heterogeneous silver nanoparticles (AgNPs) have considerable solubility in Li, the Li deposition on Ag-anchored CNF shows zero nucleation overpotential and significantly lowers the overpotential. Consequently, the uniform Li deposition was

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achieved with ultrafine AgNP decorated CNF as shown in Figure 17.3e. However, the in situ electrochemical platings of Li in a porous framework may not have uniform spatial Li deposition control and are complex to disassemble and clean before using the Li/porous structured composite while assembling the full cell battery. The direct infusion of molten Li in such a porous framework for making composites is considered a considered suitable approach. Zhu et al. reported the infusion of molten Li into the porous, and lithiophilic Zn cluster confined carbonized metal organic framework c-MOF [27]. Figure 17.3f shows the illustration of fabricating Li/c-MOF composite. The pre-planted uniformly dispersed Zn as nucleation seeds guides the uniform Li deposition due to the thermodynamic matching with Li. Zn has appreciable solubility with Li and also forms Li-Zn alloy. As a result, the nucleation barrier in c-MOF with the Zn cluster shows an eliminated nucleation barrier compared to the c-MOF without the Zn cluster as shown in Figure 17.3g, h. Liang et al. also reported the infusion of Li into the 3D conducting carbon scaffold coated with lithiophilic Si as shown in Figure 17.2i. The Li/C composite anode demonstrated smooth and uniform confinement of Li inside the porous matrix. Besides, the Li/C composite anode showed a flat and stable voltage plateau because the Li is plated or stripped from the Li itself, which minimizes the nucleation barrier initiating from the lattice mismatch. The lithiophilic decoration on a porous nano-/ microstructured framework not only ensures uniform lithium nucleation but also reduces the effective current density and the degree of interface fluctuations and confines the Li deposition/dendrites within the porous scaffold. Further development of the SEI layer on such a porous framework can be the next leap to push the CE closer to the real battery operation. However, the quality of SEI plays a critical role in the effective and stable operation of LMBs, as several requirements for an ideal SEI need to be satisfied [28]. To further avoid the risk of high temperature during molten infusion of Li into the porous free-standing electrode, the mechanical process of the press has been applied to make the Li/C composite anode. The development of an artificial SEI on such porous Li host/current collector greatly enhances the cycling stability.

17.1.3.2

The Development of Solid-State Electrolyte

The use of flammable liquid electrolyte in LMBs not only compromises the energy density but also stimulates continuous side reactions with LMA, resulting in low Coulombic efficiency (CE) [29]. Thus, it is highly recommended to utilize the solidstate electrolyte (SSE) for addressing the aforementioned issues of LMA and liquid electrolyte. The design of an ideal SSE should consider ionic conductivity, electro/ chemical and mechanical stability, being ecofriendly, and low cost. The SSEs with high Young’s modulus is found to repress the Li dendritic growth and obtain better Li utilization. Moreover, SSE can be integrated with low cost and high capacity cathodes such as sulfur and metal sulfides. The SSE can intrinsically mitigate the shuttle effect and address the metal dissolution of cathode materials. To date,

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75Li2S 25P2S5 glass

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Fig. 17.4 The Arrhenius plot of various SSEs compared to the organic liquid electrolyte (Reproduced from Ref. [30] with permission from John Wiley and Sons)

different types of SSEs such as oxides, sulfides, Li-rich antiperovskites, thin films, and polymers have been developed and investigated. Polymer SSE is a promising candidate because of their facile fabrication process, ultrathin nature, flexibility, high level of safety, and electrochemical stability towards Li metal anode/cathode but suffers from oxidation and thermal stability. The sulfide-based SSE has high conductivity (~ 10
2 S cm
1) which is comparable to the organic liquid electrolyte as shown in Fig. 17.4 but suffers from chemical instability. Similarly, the oxide-based SSE shows higher stability toward anode or cathode but suffers from high interfacial incompatibilities and requires high processing costs and complex device integration. In SSEs, ions migrate utilizing ionic motion across the crystal lattice. Thus, they are expected to higher transference number (~1) compared to the organic liquid electrolyte (~0.4–0.5). To address the issues with SSE, scientists are continuously investigating the facile, low-cost, and scalable method.

17.1.3.3

The Development of Solid Electrolyte Interphase

The surface layer between the LMA and the organic liquid electrolyte plays an important role in determining the morphology of Li deposition. The morphology of Li deposition influences battery safety and performance. The SEI inhibits the direct physical contact between the LMA and the solvent of electrolyte. Beyond that, SEI can regulate the Li-ion distribution from the bulk electrolyte to the LMA. The SEI can be formed in two ways: (1) in situ and (2) ex situ. The in situ formation of SEI is during the battery operation, and the ex situ SEI can be generated before the battery operation.

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The common solvent used in the nonaqueous liquid electrolyte of lithium batteries includes ester and ether. The ester or carbonate electrolyte has comparatively higher oxidation stability but suffers from limited compatibility with LMA. In contrast, the ether or glyme-based electrolyte has comparatively lower reactivity with LMA, providing higher CE and lower voltage hysteresis, but suffers limited oxidation stability. The common Li salts for the nonaqueous liquid electrolyte includes lithium hexafluorophosphate (LiPF6), bis(trifluoromethanesulfonyl)imide (LiTFSI), lithium bis(fluorosulfonyl)imide (LiFSI), etc. The deposition of insulating products on the LMA during the initial cycles due to the side reaction between the liquid electrolyte and hyperactive LMA which is thermodynamically unstable (Gibbs energy
8.670 kJ/mol) in organic solvents/electrolytes or from the non-faradic adsorption and the faradic electrochemical reaction is called SEI as shown in Figure 17.5a. The reduction of Li salt and solvent in the electrolyte (usually

Fig. 17.5 The schematic illustration of the mechanism in SEI. (a) The mechanism of SEI formation due to the direct contact of bulk Li electrode and liquid electrolyte. (Reproduced from Ref. [42] with permission from John Wiley and Sons). (b) The schematic demonstration of ideal SEI. The schematic demonstration on SEI formation in (c) the Peled model, (d) Mosaic model, and (e) Coulombic interaction model, respectively. (Reproduced from Ref. [43] with permission from John Wiley and Sons)

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below 1.0 V vs. Li+/Li) and the hyperactive lithium decompose, leading to the formation of SEI. This SEI layer (30–50 nm) passivates the lithium anode, preventing the further corrosion or consumption of lithium with the electrolyte. Ideally, the SEI should possess (1) high chemical stability, (2) high mechanical stability, (3) sufficient Li-ion diffusion and electron flow prevention, (4) flexibility to buffer Li volume expansion, and (5) charge/discharge reversibility as shown in Figure 17.5b. There are different models to explain the SEI formation mechanism such as the Peled model, Mosaic model, and Coulombic interaction mechanism model. Peled model (Figure 17.5c) tells that the SEI is formed from the surface reaction by the reduction of electrolyte components [31]. Mosaic model (Figure 17.5d) tells that the insoluble and insulating multiphase discharge products generated during the side reaction between the electrolyte and LMA get deposited on negatively charge LMA and the deposited layer is called mosaic layer [32]. The mosaic layers with grains and grain boundaries offer Li-ion transport across it. Coulombic interaction model (Figure 17.5e) tells that the positively charged and partially positively charged Li-ions bring into the line as head and foot, respectively [33]. Thevenin and Muller brought up several updated models which include the polymer electrolyte layer model and compact stratified layer model [34, 35]. The naturally formed SEIs are very fragile, have low Li-ion conductivity and insufficient chemical and mechanical stability, and cannot completely passivate the LMA. Such SEI cannot buffer the volume expansion issues and has weak anchoring strength with the underlying bulk lithium electrode. The continuous formation and deformation of such unstable SEI consumes both liquid electrolyte and LMA and increases the impedance of the battery, leading to premature battery death. Besides, the SEI cannot suppress lithium dendrite growth which can penetrate the separator challenging the safety of the battery. To acquire SEI with nearly ideal properties, as mentioned above, the use of various additives, salts, solvents, and fillers and optimization of the electrolyte concentration have been reported [36–40]. For the reduction, these components are generally chosen in such a way that the lower unoccupied molecular orbital (LUMO) is lower than the electrolyte and its components. Although dendrite-free Li deposition and substantial improvements in the battery cycling have been achieved, the formation of in situ SEI consumes a superfluous amount of electrolyte and bulk Li electrode [41]. Moreover, there is no calculated control on the dimension, distribution, SEI components, and thickness of in situ formed SEI. The ex situ-based method is another way to develop the SEI with more powerful control on the dimension, distribution, SEI components, and thickness. This chapter provides an overview of the efficient ex situ-derived SEI for the protection of LMA. In this chapter, the results of the studies on ex situ SEI that have been carried out to date are pulled out and summarized. Specifically, this chapter discusses the engineering process and efficient working mechanism of the ex situ artificial SEI, which can overcome the challenges associated with in situ-derived SEI. The development of ex situ SEI with sufficient Li-ion conductivity, large flexibility, high adhesion, or anchorage affinity with Li metal surface and high mechanical and chemical stability can be a promising technique. These approaches can potentially lead to much higher

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efficient SEI, based on the simultaneous resolution of the above challenges hindering the battery performance. This study will allow estimating the hands-on metrics and ideal materials, in an attempt to revive the LMBs for dendrite-free Li deposition, improved cycling performance, and high energy density batteries. Finally, the perspectives and critical challenges needed to address the issues in ex situ-formed SEI for further improvement of stable and robust SEI are concluded.

17.1.4 Ex Situ-Based SEI The adding of an artificial layer on top of LMA to protect it from liquid electrolyte before the battery assembling or battery operation is considered as the process of ex situ-based SEI formation. As this is an external process, the electrolyte and plated lithium will not be consumed as in the in situ SEI formation process. Besides, there is an effective control on the thickness, dimension, and distribution, which play a significant role in the battery performance. The strong adhesion on the bulk Li electrode prevents the peel off or corrosion of SEI during Li-plating/Li-stripping cycles. The nonuniform electric field distribution in the bulk Li electrode, due to the rough, crack surface, leads to the nonuniform Li-ion flux. The nonuniform Li-ion flux deposits the Li nonuniformly, which continues to accumulate with longer charge/discharge cycles, acting as a hot spot for Li nucleation. In general, there are two methods to generate ex situ-based SEI, which include (1) physical deposition method and (2) chemical deposition method. The solution, gas, or any other separate solid layer can be utilized to develop the ex situ-based SEI. The smooth surface of SEI leads to the uniform Li-ion flux, improving the Li deposition behavior for longer Li-plating/Li-stripping cycles. Despite the attractive performance of the battery and controlled Li dendrite growth, many of the reports are concentrated on the single aspects of SEI. For example, the reports have shown that improved performance can be attributed to only high Young’s modulus or high Li-ion conductivity. It is imperative to think and consider all the ideal properties of SEI that can significantly contribute to the improved cycling performance of the battery. Besides, the material manufacturing process, cost, safety, and effectiveness are to be deemed.

17.1.4.1

Physical Deposition

The construction of SEI on top of LMA physically adding a layer is considered as a physical deposition. The physical interlayer can be from a few nanometer thicknesses to several micrometers. Depending upon the thickness, the adhesiveness of the interlayer is very important. For example, for the ultrathin film, high adhesivity is required. Otherwise, there are chances of SEI being peeled off and degraded. The separate thick (μm-range) interlayer has also been used to create a physical barrier between the LMA and liquid electrolyte. Such kind of thick interlayer does not need to bear high adhesivity. For example, glass fiber cloth [44], 3D conductive stainless

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steel fibrous metal felt [45], the monolayer of nanostructured/interconnected hollow carbon nanospheres [46], and carbon paper interlayer [47] have been employed as an interlayer to protect the LMA. However, the use of an additional thick extra layer on top of LMA compromises the energy density of the battery. Various methods have been employed to develop the ex situ-based SEI by a physical deposition method such as spin coating, radio-frequency (R-F) sputtering, direct-current (DC) sputtering thermal evaporation, chemical vapor deposition (CVD), pulsed laser deposition (PLD), atomic layer deposition (ALD), drop-casting, or doctor blading techniques. While bridging the academic research to industry, the cost is an important factor. The ultrahigh vacuum needing instruments are expensive and inefficient for large-scale production. The facile and low cost, a solution-based method such as dip coating, drop a cast, and doctor blading can be more practical for large scale productions. In our previous work, the ex situ SEI was developed by depositing the ultrathin bilayer of graphite and SiO2 on the lithium metal [12]. Figure 17.6a shows the Torr combination system for R-F sputtering. Figure 17.6b, c shows the topography SEM images of Li before and after the physical deposition of graphite/SiO2 bilayer on the surface of Li. The ultrathin bilayer of graphite/SiO2 demonstrated outstanding Li-plating/Li-stripping cycles and full cell performance. The graphite provides flexibility and electrically connects the plated Li with the bulk Li underneath the SEI, which reduces the chances of dead Li formation. The high Li-ion conductivity, high mechanical strength, and reversible lithiation/delithiation properties of SiO2 facilitate fast Li-ion transport, suppress the Li dendrite growth, and store Li or slow down the Li dendrite growth, respectively. The R-F deposition of an artificial SEI layer produces a smooth and high-quality film, enabling uniform Li-ion flux on the surface of the Li anode. The rate of deposition and the carrier gas or reactive gas can be easily tuned as per the requirement. Lee et al. used the doctor blading techniques to develop the composite protective layer (CPL) on the LMA as shown in Figure 17.6d [48]. The coated CPL with a thickness of 25 μm using the slurry of Al2O3, PVDF-HFP, DMF, and carbonate electrolyte showed improved long-term battery performance compared to the bare Li. The CPL provides a physical barrier between the Li and liquid electrolyte, enables uniform Li-ion flux and sufficient Li-ion transport, and mechanically suppresses the Li dendrite growth. Figure 17.6e, f shows the surface topography and cross-sectional SEM images of CPL protected Li metal. Table 17.1 summarizes the materials and methods used for the SEI development and the corresponding electrochemical cell test performance achieved with that SEI.

17.1.4.2

Chemical Deposition

The chemical reaction between the hyperactive LMA and the material for SEI formation leads to the strong anchorage between the formed SEI and the LMA. The high anchorage affinity or adhesivity of SEI components leads to the formation

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Fig. 17.6 Schematic of the physical deposition of SEI. (a) Schematic of the R-F sputtering to deposit the target materials on the Li metal surface in the Ar-environment. (b–c) Surface topography of Li before and after the physical deposition of bilayer graphite-SiO2 on the Li metal surface. (Reproduced from Ref. [12] with permission from John Wiley and Sons). (d) Schematic illustration of the CPL preparation by doctor blading method. (e, f) The surface topography and cross-sectional SEM image of the CPL deposited Li metal anode. (Reproduced from Ref. [48] with permission Elsevier)

of dense SEI. The thickness of such chemically formed SEI also varies from a few nanometers to several micrometers. The chemical reaction can be allowed to happen between the Li metal and precursors such as a solution-based, gas-based, or solid-based. The solution-based

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Table 17.1 Summary of the artificial SEI engineered by a surface coating method SEI components The ultrathin bilayer of graphite and SiO2 Li3PO4 Carbon Cu3N Al2O3 MoS2 Al2O3 Al2O3 Double-layer nanodiamond Al2O3 + PVDF-HFP Cu3N + SBR ETPTA+DADMA-TFSI PEDOT-co-PEG Perfluorinated ionomer layer Langmuir-Blodgett Hollow carbon nanosphere Double-layer nanodiamond Trifluoroethanol (TFEA)-modified eggshell membrane (TESM) Metal chloride perovskite film MASnCl3 and MAPbCl3

Method R-F sputtering (40 nm)

Ref. [12]

Magnetron sputtering (30 nm) Magnetron sputtering (80 nm) R-F magnetron sputtering (115 nm) R-F magnetron sputtering (20 nm) R-F magnetron sputtering (20 nm) ALD (14 nm) ALD (~2–3 nm) MPCVD (300 nm) Doctor blade (25 μm) Doctor blade/drop-casting (1 μm) Doctor blade/UV irradiation (1 μm) Immersing LMA into the polymer and nitromethane solution (10 μm) Solvent-casting (20 nm – 9 μm) Roll press (0–3 μm) Pressing (wall thickness – 20 nm) MPCVD (300 nm) Interfacial layer (90 μm)

[49] [50] [51] [52] [53] [54] [55] [56] [48, 57] [58] [59] [60]

Spin coat-solid state transfer-pressing (1 μm and 2 μm)

[61] [62] [46] [56] [63] [64]

reaction includes doctor blading, spin coating, dripping, drop-casting, dip coating, etc. The various parameters such as exposure time, temperature, the concentration of the solution, and other environmental conditions determine the properties (thickness, SEI components, morphology, etc.) of SEI. In our previous work related to the solution-based chemical deposition, we treated the Li metal with the electrolyte containing SnF2 [11]. The bare Li shows the dendritic Li growth, but 3 wt% SnF2pretreated Li shows the dendrite-free Li deposition as illustrated in Figure 17.7a, b. The smooth surface with an artificial SEI thickness of 25 μm (Figure 17.7c) showed outstanding battery performance. Also, in our previous work related to the gas-based chemical deposition, we treated the Li metal with the N2-assisted plasma [13]. Figure 17.7d shows the experimental setup for N2 plasma, and Figure 17.7e shows the cross-sectional SEM image of Li anode with 2 min N2 plasma. The optimized thickness of 8 μm achieved with 2-min N2 plasma showed excellent electrochemical battery performance (Table 17.2).

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Fig. 17.7 Chemical surface treatment of LMA. (a, b) Schematic of the dendritic Li growth and smooth Li deposition in bare Li and SnF2-pretreated Li, respectively. (c) The surface topography SEM image of Li and inset shows its cross-sectional images. (Reproduced from Ref. [11] with permission from Nature Springer). (d) Experimental setup for the developing Li3N on Li by the plasma process. (e) Corresponding cross-sectional images after 2 min of N2 plasma. (Reproduced from Ref. [13] with permission from Elsevier)

17.1.5 SEI Properties and Functionality The conventional approaches were focused only on one of the properties of SEI which cannot fully stabilize the SEI. Consequently, long-term battery performance and fast charge/discharge rates couldn’t be possible. Thus, the advanced strategies of developing SEI need to focus on all the basic requirements to achieve an ideal SEI. Figure 17.3b shows the properties of an ideal SEI which includes high Li-ion conductivity, high mechanical and chemical stability, flexibility, and reversibility. Here the representative properties of the SEI are discussed.

17.1.5.1

Thickness

The various parameters such as deposition time, the concentration of the precursor, and the fabrication method play an important role in determining the morphology and thickness of SEI. The thickness of SEI can be optimized by taking different parameters such as Li-ion diffusivity and mechanical and chemical stability into considerations. For thickness optimization, various structural and corresponding chemical characterizations can be carried out. Structural characterizations such as XRD, Raman, SEM, BET, and XPS and the electrochemical cell performance such as Li plating/stripping, Li deposition morphology, voltage profile analysis, and EIS measurement help to optimize the thickness of SEI. The higher thickness of SEI can suppress the growth of Li dendrites; however, it might sluggish the Li-ion transport. In contrast, the lower thickness of SEI may not fully protect the LMA. The excess

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Table 17.2 Summary of the artificial SEI engineered by surface chemistry method SEI components/thickness of SEI Li3PO4 (50 nm) Cu + LiF (2 μm) Organic/inorganic film (75 nm) Li-rich metal alloy/Li-halides (< 10 μm) Ge, GeOx, Li2CO3, LiOH, LiCl, and Li2O (1.5 μm) LiF nanoparticles (
) MIEC LLTO (8.2 μm) Al-Li alloy/LiCl (10 μm) Cuboid LiF (77.4 μm) LiF, Sn, Sn-Li alloy (25 μm) [LiNBH]n chains (140–160 nm) Li2S (25 nm) Li2S/Li2Se (2 μm) LiF layer (40 nm) LiF coating (380 nm) LiF thin film (40 nm) Li3N (8 μm) Li-O-Si linkage (
)

Materials/method to pretreat the LMA A solution of polyphosphoric acid (PPA) and dimethylsulfoxide (DMSO) Dripping (CuF2 + DME) solution on LMA Soaking the Li metal in the pure solvent of fluoroethylene carbonate (FEC) MClx solution in tetrahydrofuran (THF) for 20 sec

Ref. [65]

Immersing LMA organic GeCl4 – THF steam (1.5 μm)

[69]

Immersing Cu into aqueous solution of LiPF6 Toluene in between LLTO and LMA Soaking LMA in AlCl3-ionic liquid Immersing LMA into NH4HF2 + DMSO Drop-cast SnF2 containing carbonate electrolyte Dehydrogenation reaction between LMA and ammonia borane-NH3BH3 (AB) + THF (Sulfur/240  C) gas – LMA (SeS2/150  C) gas – LMA Nitrogen trifluoride/175  C gas – LMA Fluoropolymer/350 Cgas – LMA/175  C Freon R134a/150  C gas – LMA Nitrogen gas-enhanced plasma – LMA Immersing LMA into tetraethoxysilane for 5 min

[70] [71] [72] [73] [11] [74]

[66] [67] [68]

[75] [76] [77] [78] [79] [13] [80]

use of chemicals or materials to construct SEI compromises the energy density of LMBs. Thus, the materials with low mass density with a high Li-ion diffusion coefficient are to be chosen. The SEI thickness ranging from ultrathin few nanometers to thick micrometer size has shown improvement in the electrochemical battery performance.

17.1.5.2

Transference Number

The high Li-ion transference number (TLi+) increases the energy density and charge rates in LMBs. Recently, scientists are motivated to develop the SEI through the surface coating of polymer which uniquely is cation conducting. The low Li-ion conductivity and low TLi+ of SEI limit the practical application of LMBs at high current density. The enhancement of Li-ion conductivity and Li-ion TLi+ is challenging work. The use of cation conducting membrane as SEI can significantly reduce the loss caused by ion polarization. Tu et al. reported SEI deposited on the LMA surface by the solvent-casting method of perfluorinated ionomer [61]. The ionomer SEI with a thickness of 20 nm demonstrated ionic conductivity of

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Fig. 17.8 (a, b) Comparison of Li-ion conductivity and transference number of different electrolyte systems from the previous reports. (Reproduced from Ref. [61] with permission from John Wiley and Sons). (c) Schematic of armored MIEC SEI developed by the displacement reaction between CuF2 and Li, showing the high modulus, ionic conductivity, and surface energy. (Reproduced from Ref. [66] with permission from John Wiley and Sons). (d) Comparison between bare Li foil and rAGA-Li anodes toward water stability. (Reproduced from Ref. [102] with permission from John Wiley and Sons). (e, f) Antioxidative characteristic of pSEI-Li and bare Li using contact angle measurement, surface tarnish test, and Raman spectroscopy. (Reproduced from Ref. [59] with permission from John Wily and Sons)

~1*10
3 S cm
1 at room temperature and high TLi+ ~ 0.9. Figure 17.8a, b shows the summary and comparison of ionic conductivity and lithium transference number of various electrolytes reported previously. The gray color symbolizes the ionic conductivity measured above room temperature. Solid and open symbols indicate the solid-state and liquid electrolyte. For composite or hybrid electrolytes, half-filled symbols are used. Cho et al. demonstrated ultraviolet (UV)-polymerized ethoxylated trimethylolpropane triacrylate (ETPTA) and diallyldimethylammoniumbis (trifluoromethanesulfonyl)-imide (DADMA-TFSI) as an organic/inorganic mimic SEI [59]. Such SEI demonstrated a high ionic conductivity of 1.2*10
3 S cm
1 and TLi+ of 0.69, which can be attributed to the fast Li-ion transport provided by the inorganic layer. The lone pair electrons (such as C-N bond) and amine polar function

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(such as amine group) in MOF successfully inhibit the anion migration which results in the improved TLi+ [81, 82]. Besides, the anion blocking or negatively charged group such as –SO3
, -COO
, and PO32
containing cation-selective ionomer membrane SEI lessens the chances of direct contact of solvent or anions of electrolyte with LMA. As a result, the TLi+ increases which prolonged the growth of dendritic Li. Moreover, the polar functional bonds on the surface of MOF effectively screen the ions and adsorb only Li-ion, hindering the migration of anion. This results in the improvement of TLi+ [81].

17.1.5.3

Young’s Modulus and Flexibility

The mechanical stability and flexibility are very important properties of SEI for stable long-term cycling performance in LMBs. The hostless LMA experiences repeated volumetric change which can cause serious mechanical cracks and repeated breakdown/repair on the SEI during Li plating/stripping. Such cracks behave as hot spots for Li deposition, which eventually cannot effectively guide the uniform and dendrite-free Li deposition. Therefore, the high mechanical strength and flexibility of SEI can suppress the Li dendrite growth and buffers the volume expansion issues for maintaining the interface stabilities and SEI structural integrities. Figure 17.6c shows the schematic illustration of an armored SEI (Cu and LiF) formed by the displacement reaction between CuF2 and Li. The Cu atoms in the SEI improve the Li-ion conductivity by offering more diffusion space. Besides, the high surface energy of such SEI realizes uniform Li-ion distribution, high ionic conductivity, and sufficient Young’s modulus to suppress the growing Li dendrites. Previous studies have suggested that a moderate elastic modulus above 3.0 GPa is sufficient enough for increasing the SEI failure time and suppressing the Li dendrite growth [65, 83, 84]. If Young’s modulus is high enough to improve or maintain the structural uniformity in SEI, the stability of SEI will be enhanced. Extensive research has highlighted the benefit of developing SEI with high mechanical strength, some of which are summarized in Table 17.3. The infinite volume expansion issues of Li can create cracks and repeated formation/deformation of SEI. However, the flexible SEI can accommodate the interfacial fluctuations and further prevent the SEI from corrosion and peeling off. Lie et al. designed 20 nm Li polyacrylic acid (LiPAA) as a self-adapting interface fluctuation SEI by drop-casting the solution containing LiPAA in DMSO [89]. The main organic functional group COOLi on the LiPAA anode not only provides flexibility to the surface but also protects the LMA from the air environment. The use of flexible SEI components such as styrene-butadiene rubber (SBR) [58], thinfilm carbon, or graphite coating on Li is also effective in addressing the infinite relative electrode dimension change during cycling [12, 90]. The mechanically robust metal-organic framework (MOF) SEI cemented by polymer to modify its flexibility is of great interest to develop stable SEI [81]. The use of 3D porous Li hosts such as layered reduced graphene oxide (rGO) with nanoscale interlayer gaps [91], hollow carbon nanospheres [46], and nanostructured unstacked graphene drum

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Table 17.3 Ex situ SEI components with Young’s modulus value SEI components PPA-Li Double-ND Single crystal ND/ND-Li composite LiF + PVDF-HFP LiF + Sn + Li-Sn Graphite + SiO2 Li3N Cu + LiF A single atomic layer of h-BN Al2O3 Al2O3 + PVDF-HFP Nafion + LiCl Li2CO3 + LiF LiDFBOP + FEC derived SEI

Modulus (GPa) ~10.0–11.0 >200 ~170.0/30.0 6.7 55.6 10.7 48.0 12.9 ~1 Tpa 6.4–14.3 3.4 6.1 7.0 64.9

Average thickness ~ 50 nm ~ 300 nm 216 μm 12 μm 25 μm 40 nm 8 μm 2 μm 1 nm 20 nm 25 μm 2.5 μm 50 nm –

Ref. [65] [56] [85] [66] [11] [12] [13] [66] [86] [52] [48] [87] [67] [88]

[92] is another strategy to address the volume change problem of LMA. Besides, the use of ceramics materials is very promising interfacial layers to buffer the volumetric expansion of the anode [93–96]. The SEI components which offer the synergy of high mechanical strength and sufficient flexibility are proven to be effective in the stabilization of SEI [58, 97]. The dual-layered organic/inorganic film was also constructed by soaking the Li metal in the pure solvent of fluoroethylene carbonate (FEC) for different time durations [67]. The electrostatic attraction between the positively charged Li atom and negatively charged F atom can break the C-F bond of FEC to form LiF, and, subsequently, the organic components such as CH2CHOCO2Li (ROCO2Li) and CH2CHOLi (ROLi) also formed. The organic layer provides the flexibility, and the inorganic layer provides high mechanical strength to suppress the Li dendrite growth. The use of graphene oxide (GO) enables high mechanical strength and flexibility, which can tolerate the interfacial fluctuations [98].

17.1.5.4

Mixed Ionic/Electronic Conductor

The SEI with high Li-ion conductivity and poor electronic conductivity allows sufficient Li-ion diffusion but prevents the flow of electrons. The low electronic conductivity enables the Li deposition underneath the SEI (bottom to top) and suppresses the growth of Li dendrites. The lower electronic conductivity of the SEI compared to the current collector prevents the deposition of Li on top of SEI. Further, weak binding and the poor Li wettability of such SEI provide the space between the SEI and the current collector and promotes the Li deposition underneath the SEI protective layer, respectively. Nevertheless, the protective layer with high electronic conductivity which can facilitate fast diffusion of Li-ion on the surface

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and within interlayer gaps has also been considered effective in stabilizing the SEI [91, 92]. The conductive protective framework is supposed to lower the large electrical impedance and strong voltage hysteresis at high current rates. Besides, the mixed Li-ion conductors can also facilitate the high Li-ion migration across the SEI. Li2S has been established as an efficient protective layer owing to its high conductivity (~10
5 S cm
1). Liu et al. designed mixed Li-ion conductors Li2S/ Li2Se where Li2Se shows higher Li-ionic conductivity than Li2S [76]. Recently, the mixed ionic/electronic conductors (MIEC) which enable rapid transport of both Li-ions and electrons have attracted great attention [71, 99]. The MIEC can store the Li at the grain boundary regions of ionic and electronic conductors, ionic conductor regulates the Li-ion, and electronic conductor transfers the electron rapidly. The electronic conductor may also store the Li by forming a Li-rich alloy. For example, in the hybrid interphase of C60 and Mg, Li and Mg form a Li-Mg alloy [100].

17.1.5.5

Antioxidative

The water penetration inside the battery during a rainy day or while washing a car can cause potential safety threats. When Li is in contact with moisture, it gets oxidized and evolves unwanted gas, deteriorating the battery performance and challenging the safety issues. The water repellent hydrophobic SEI layer can shield the LMA from such moisture exposure. Dong et al. engineered the LMA passivation SEI layers using a vertically aligned graphene framework with a hydrophobic roof of graphene tiles [101]. Figure 17.6d shows the Li metal without any protection (left side) and with reduced accordion-like graphene oxide array (rAGA) protection (right side). The reaction between Li and H2O can produce H2 gas and cause corrosion in the Li metal. An air- or waterproof graphene tile roof house like hierarchical architecture, combined with a vertically aligned graphene microstructure channel, substantially protects the LMA and stores excessive Li in its porous structure. Liao et al. demonstrated the treatment of LMA by using GeCl4 which can generate Ge, GeOx, Li2CO3, LiOH, LiCl, and Li2O on Li surface, allowing the stable symmetric and full cell Li-O2 cycling performance at a relative humidity of 45% [69]. There are numerous reports on the use of hydrophobic or water-resistive SEI to protect the LMA. The other report to protect the LMA from the moisture is to use organic components that are hydrophobic or water-resistive [59]. Figure 17.8e–g shows the antioxidative behavior of the printable solid electrolyte interphase (pSEI) deposited on Li under different humid conditions. The bare Li quickly oxidizes when in contact with the water inset of Fig. 17.8f. In contrast, the pSEI-Li has a higher contact angle of 65 indicating higher water-repelling capability. The surface of bare Li quickly tarnished into black upon exposure to a humid environment. In contrast, pSEI-Li maintained higher stability against the humid environment of ~50% relative humidity. The inset shows the photographic image of pSEI-Li and bare Li after exposure to humid air for 10 and 40 s. Figure 17.8g shows the Raman spectra of bare Li and pSEI-Li before and after 120 s exposure to humid air. The distinct and intense peak

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of LiOH and Li2CO3 at 657 and 1060 cm
1 were observed for bare Li. In contrast, the peaks were very weak from the antioxidant pSEI-coated Li.

17.1.5.6

Hybrid SEI

The hybrid SEI which can simultaneously store Li by the reversible plating/stripping and redox reaction mechanism is considered effective in guiding the uniform Li deposition. The components with high Li-ion conductivity allow sufficient Li-ion diffusion, and the chemically active components react with the hyperactive dendritic Li and slow down the rate of dendrite growth during plating/stripping cycles. The SEI components of the hybrid SEI can reversibly store Li by lithiation/delithiation reaction mechanism [12, 53]. In our previous studies, the bilayer of 20 nm graphite and 20 nm SiO2 was deposited on Li foil by the R-F sputtering method. The film deposited at room temperature via magnetron sputtering does not show any diffraction peak in the XRD spectrum, indicating the amorphous phase [12, 49]. The amorphous thin film of SiO2 reacts with plated Li forming Li2Si2O5, Li4SiO4, and Li15Si4. Figure 17.9a, b shows the core level Si 2p and O 1 s peak of as-deposited SiO2, after 50th plating and after 50th stripping. The binding energy (BE) of corelevel Si 2p and O 1 spectra at 103.3 eV and 532.71 eV indicates the deposition of SiO2. The shifting of BE for both Si 2p and O 1 s toward a lower value indicates the lithiation of SiO2. Moreover, the shifting of BE for both Si 2p and O 1 s toward the original (as deposited) value indicates the delithiation/recovery of SiO2. Thus, the phenomenon of redox chemical reaction between the plated Li and thin-film SiO2 suggests the storage of Li which slows down the growth of Li dendrite. The deposition of the protective layer increases the impedance for Li-ion transport. The lithiation in such surface coatings during Li-plating/Li-stripping cycles improves the electronic conductivity, resulting in higher conductivity and lower impedance. Cha et al. sputtered 10-nm-thick semiconducting 2H-MoS2, which transforms to 1 T-MoS2 when the number of intercalated Li per MoS2 exceeds 0.4. The MoS2 film uniformity and high ionic conductivity were achieved with the galvanostatic electroplating. The SEI developed by the surface chemistry route with a chemically active SEI component could also store Li by reversible alloy formation [11, 68, 103]. The reduction of metal chlorides, fluorides, or some similar salts by the Li metal can produce the Li-rich alloy and Li-halides. The formation of Li-rich ion conductive alloy as an SEI component enables fast Li-ion migration as they have a higher Li-ion diffusion coefficient than Li metal. Besides, the resistive Li-halides prevent the reduction of Li-ion on the surface, allow Li deposition underneath the protective layer, and inhibit the electron flow, implying the suppression of Li dendrite growth. In our work, an electrolyte containing SnF2 was drop-cast on the surface of Li to form LiF, Sn, and Li-Sn alloy. The chemically active Sn reversibly stores Li by the alloy formation [11]. Figure 17.9c, d shows the cyclic voltammetry (CV) measurement of the bare Li and artificially fluorinate hybrid SEI with an optimized 25- μmthick (AFH-25) symmetrical cell. The straight line in bare Li indicates the Li-plating/

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Fig. 17.9 (a, b) The XPS analysis of SiO2-hybrid SEI at different conditions. (Reproduced from Ref. [12] with permission from). (c, d) The first cycle CV measurement of symmetrical cells with bare Li and AFH-25. (e) Plating/stripping electrochemical performance of bare Li and AFH-25 symmetrical cell at 0.5 mA cm
2. (Reproduced from Ref. [11] with permission from)

Li-stripping mechanism. In contrast, AFH-25 symmetrical cell shows the reversible redox peaks at ~
0.12/0.12 v, indicating the lithiation/delithiation of active Sn. The AFH-25 symmetrical cell has the longer stable plating/stripping cycles with reduced nucleation overpotential and voltage hysteresis as shown in Fig. 17.9e. The inset of Fig. 17.9e shows the plating/stripping voltage profile for the first five cycles at a current density of 0.5 mA cm
2 to achieve a capacity of 1 mAh cm
2. The lower

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nucleation overpotential and lower voltage hysteresis were observed for AFH-25 Li compared to the bare Li symmetrical cell. Besides the Li-plating and reaction mechanism to store excess Li, the SEI can also be hybrid in terms of SEI components. For example, the use of organic and inorganic SEI components has successfully stabilized the in situ- and ex situ-based SEI [58, 66, 67]. The flexibility of organic layer or components such as PVDF-HFP, PEO, ROCO2Li, and ROLi (where R represents an alkyl group and styrenebutadiene rubber of SEI buffers the volume expansion of Li during Li plating / stripping. To add up the mechanical strength to maintain the structural integrity within the SEI, the inorganic components such as (LiF, Li2CO3, Al2O3, SiO2).

17.1.5.7

Lithiophilic

The strategy of using lithiophilic decorations or coatings which undergoes conversion and/or alloy reaction has greatly improved the Li deposition behavior even at higher plating/stripping rates. Various lithiophilic coatings (such as metal oxides, metal nitrides, and metal sulfides) or seed decoration such as heterogeneous metallic nanoparticles within the SEI enhances the Li-ion adsorption. The lithiophilic coating/seed decoration guides the initial Li deposition that results in the lower nucleation overpotential and lower voltage overpotential [25]. However, the conversion reaction with lithiophilic metal oxides may lower the CE in the initial cycles due to the formation of lithiophilic Li2O. The presence of electrochemically electronic metal conductors in the SEI facilitates uniform Li-ion transport and a fast electron route. For example, the lithiophilic P2O7 and P3O9 serve as favorable sites for the adsorption of Li atoms due to the presence of a free electron pair on oxygen [62]. The other lithiophilic polar functional groups such as (C¼N, C-N, C¼O, O-H) on the 3D oxidized polyacrylonitrile (PAN) nanofiber interlayer serve as a Li nucleation site which prevents the Li deposition at the hot spots [104]. The higher affinity of the polar functional group confirms better electrode/electrolyte contact. The formation of Li-rich alloy and lithiophilic Li2O during alloying reaction and conversion reaction, respectively, can further enhance the interface stability. Besides, the development lithiophilic-lithiophobic gradient on the SEI also enhances the stability and guides the uniform Li deposition [22, 105, 106]. The lithiophobic upper layer of SEI and lithiophilic lower layer of SEI in dual-layered SEI prevent the Li deposition on the top and allow fast Li-ion transport underneath the SEI. Lithiophilic materials help to strongly capture the Li-ion, reducing the chances of dead Li formation. Lithiophilic materials suppress the growth of Li dendrites.

17.1.6 Operation Under Practical Conditions Many of the researchers are focused on the battery material level, regardless of energy density. To achieve a high energy density needed for practical applications of

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Fig. 17.10 (a, b) Schematic of Li protection by PAA-capped matrix and the battery performance comparison between bare Li and PAA@Li under lean electrolyte of 15 μL mAh-1 and cathode mass loading of ~12 mg cm
2. (Reproduced with permission from [Ref. 107]) (c, d) Schematic of Li protection by PVA-modified SEI and the battery performance using bare Li and PVA-protected Li under lean electrolyte of 7.5 μL mAh
1 and cathode mass loading of ~10 mg cm-2. (Reproduced with permission from [Ref. 109]). (e, f) Schematic of Li protection by nanoporous Cu and the battery performance of (Reproduced with permission from [Ref. 108])

LMBs, the crucial cell design principles of considering high mass loading of the cathode, lean electrolyte, and lean lithium play an important role. For example, the long cycling stability with the use of excessive or flooded electrolyte has already been reported. Nevertheless, the crucial practical condition of using lean electrolyte in the cell design principle, which has been overlooked, is a key factor that impacts the energy density of the battery. Feng et al. demonstrated a capped 3D Cu foam by coating the polyacrylic acid (PAA) layer on lithiophilic CuO-Cu foam (PAA@Li matrix) as shown in Fig. 17.10a to offer low interfacial impedance, homogeneous Li-ion flux for smooth and uniform Li deposition, and low nucleation barrier [107]. The chemically and mechanically stable PAA coating, just like a cap, protects the Li and stabilizes the interface for longer cycling stability under high current density (350 h at 5 mA cm
2 for 5 mAh cm
2). The full cell (NMC811/PAA@Li matrix) tested under the electrolyte-to-capacity ratio of 15 μL mAh
1 with high

AU1

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cathode mass loading of 12 mg cm
2 exhibited capacity retention of 75.3% after 200 cycles at 2C as shown in Fig. 17.10b. However, the full cell using bare Li anode showed early capacity decay after 20 cycles. Zhao et al. paired the polyvinyl alcohol (PVA) polymer-protected Li as shown in Fig. 17.10c and NCM as a cathode which demonstrated a stable cycling performance with ~74.5% capacity retention after 120 cycles and an average CE of 99.3% under a lean electrolyte condition (7.5 μL mAh
1) at 0.3 C as shown in Fig. 17.10d [108]. Designing the nanoporous Cu layer on the LMA surface as shown in Fig. 17.10e demonstrated the cycling performance of more than 100 cycles with 7.5 μL mAh
1 in Cu-Li/NCM523 full cell as shown in Fig. 17.10f [109]. Gao et al. used reactive polymer composite (RPC)-poly(vinyl sulfonyl fluoride-ran-2-vinyl-1,3-dioxolane) (P(SF-DOL)), Li fluoride nanoparticles, and graphene oxide sheets to design a stable SEI which suppresses the electrolyte consumption [98]. The RPC-derived SEI-containing polymers, nanoscale LiF particles, and GO nanosheets passivate the Li surface. This leads to the control of the side reaction between the liquid electrolyte and LMA, which has been confirmed from the lower concentration of Li (19.6%) and F (11.2%) in RPC-derived SEI compared to conventional electrolyte-derived SEI (37.3% of Li and 33.6% of F). As a result, the RPC-derived SEI effectively retains 77% of the electrolyte after 180 cycles, but without ex situ SEI it only retains 41% after 50 cycles. Besides, RPC-derived SEI demonstrates high mechanical strength and flexibility due to the presence of GO nanosheets. The polymer-inorganic SEI enabled a high CE of 99.1% at a deposition capacity of 4 mAh cm
2 and stable cycling performance over 200 cycles in 4 V Li/LiNi0.5Co0.2Mn0.3O2 cell under lean electrolyte (7 μl mAh
1), limited Li, and high areal capacity of 3.4 mAh cm
2 conditions. Cha et al. also studied the LMB performance at the lean electrolyte condition (3 g Ah
1, ~ 12 μL) by protecting the LMA through R-F sputtered MoS2 deposition [53]. Yin et al. demonstrated the use of metal chloride perovskite thin films such as MASnCl3 and MAPBCl3 on top of Li metal to shield the LMA from liquid electrolyte [64]. The high symmetry of perovskite provides vertical channels for Li-ion transport, and both the intercalated Li-ions which form Li-M alloy and insulating LiCl layer are beneficial for guiding Li deposition. The perovskite framework only allows Li-ion transport but not the solvent molecules. Such metal chloride perovskite-based interfacial-protected Li, at a lean electrolyte condition of 20 μL mAh
1, enabled a LiCoO2-Li cell with an area capacity of 2.8 mAh cm
2 for more than 100 cycles with a capacity retention of 85% at 0.5C.

17.1.7 Conclusion and Outlook LMA has been considered as a promising anode material in next-generation lithium metal batteries. LMA can deliver a high energy density when paired with conversion and intercalation cathodes such as sulfur, oxygen, metal fluoride, lithium cobalt oxide, lithium nickel manganese cobalt oxide, etc. Despite attractive features, LMA suffers from inherent issues of infinite volume expansion and its hyperactive nature,

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leading the challenges in battery safety and performance. The solution to the LMA/liquid electrolyte interface incompatibilities and instabilities can bring the LMA into practical existence. The design of stable and robust SEI is the key to achieve a practical application of LMA. The development of SEI by physical or chemical deposition method has led to the dendrite-free Li metal deposition, resulting in improved electrochemical battery performance. However, the fundamental understanding of the SEI formation, the structural and electrochemical properties, and the regulation or stabilization to revive the LMA are still inadequate. Several advanced experimental and theoretical techniques have been utilized to understand the fundamental properties of SEI formation and their regulation. The continuous research and development of nanoscience and nanotechnology have led to improving the interfacial instabilities of electrode/electrolyte interphase. The theoretical calculations [110–113] and advanced characterization [114–117] have successfully disclosed the mystery of correlation between electrode/ liquid electrolyte interface chemistry and the electrochemical battery performance. Despite noteworthy advancement, there are several unresolved arguments about the formation, mechanism, and impact of SEI and challenges that remain to be discovered in the field of SEI engineering. Based on the above advanced and impactful reports, characterizations, simulations, and our thinking, we believe that this book chapter will catch additional attention in the upcoming research community focusing on the development of high-quality SEI to revive the LMA. The ultrastable and robust SEIs pave the pathway toward the practical applications of rechargeable LMBs including Li-S batteries, Li-O2 batteries, and other high-capacity/high-voltage conversion cathodes. Acknowledgments We acknowledge the financial support from the SDBOR competitive grant, NSF MRI (1428992), EDA University Center Program (ED18DEN3030025), NSF IUCRC (#1841502). Author Contributions This paper was entirely written by RP. The final draft was reviewed QQ. Conflicts of Interest There are no conflicts to declare.

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Rajesh Pathak received his MSc. in Physics from the GoldenGate International College, Tribhuvan University, Nepal (2009–2011), and received his PhD from South Dakota State University (2016–2020). Currently, he is a postdoctora; appointee in the Applied Materials Division at Argonne National Laboratory, Lemont, USA. He has coauthored more than 45 peerreviewed journal articles including the leading journals such as Nature Communications, 11(1), 1–10, and Advanced Energy Materials, 9(36), 1901486. Currently, his research focuses on the development of a solid electrolyte-electrode interface (SEI) between lithium metal anode and the electrolyte (liquid/solid state) and electrode property analysis to develop next-generation high energy density lithium metal batteries.

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Yue Zhou received his Ph.D. in electrical engineering at Pennsylvania State University in 2015. He was a postdoctorate associate at the Massachusetts Institute of Technology who researched the field of energy storage. Currently, he is an assistant professor at South Dakota State University directing the research on lithium metal batteries, sodium-ion batteries, and supercapacitors.

Qiquan Qiao received his Ph.D. from the Virginia Commonwealth University in 2006. He was a postdoctorate associate at the University of Florida. Currently, he is a full professor at South Dakota State University who directed the research in the field of photovoltaics, lithium metal/ion batteries, sensors, micro- or nanomanufacturing/fabrication, food-energy-water (FEW) sustainability, and precision agriculture technologies. He has directed and participated in many projects supported by NSF, NASA, USAID, EDA, 3M, Agilent, Raven Industries, etc. He has published more than 200 peer-reviewed papers in leading journals including Science, Nature Communications, Energy and Environmental Science, Journal of the American Chemical Society, Advanced Materials, Advanced Energy Materials, Advanced Functional Materials, Nanoscale, Joule, ACS Energy Letters, Nano Energy, etc.

Chapter 18

3D X-Ray Characterization of Energy Storage and Conversion Devices Chun Tan, Andrew S. Leach, Thomas M. M. Heenan, Rhodri Jervis, Dan J. L. Brett, and Paul R. Shearing

18.1

Introduction to X-rays

X-rays form part of the electromagnetic spectrum with wavelengths in the range of ca. 0.1 Å to 100 Å, equivalent to energies between ca. 0.1 keV and 100 keV. At the lower energy end of the spectrum are ‘soft’ X-rays with energies up to around 5 keV, whilst ‘hard’ X-rays typically have energies beyond this. Due to their higher energy compared to visible light, they possess a valuable property – the ability to penetrate matter and hence reveal the inner structure of materials that are not optically transparent. The principal interaction mechanisms of X-rays with matter will be discussed in detail in the following section, and these can be harnessed to probe many material properties. Advances in synchrotron and laboratory-based radiation sources have led to numerous X-ray characterization methods being developed, including X-ray absorption spectroscopy (XAS) [1], X-ray diffraction (XRD) [2], and transmission X-ray microscopy (TXM) [3]. These complementary techniques, spanning multiple time and length scales, provide a wealth of information about the electronic states (XAS), crystalline ordering (XRD), and microstructures (TXM) of materials and are being widely adopted to investigate the mechanisms behind the operation and degradation of electrochemical devices, including lithium-based batteries, fuel cells, and redox flow batteries. Many of these X-ray-based spectroscopy and imaging methods can be extended through a process known as tomography to capture a spatially resolved map of the measured property in 3D, although in practice the beam, sample, and detector configurations of some techniques are easier to implement for tomography than

C. Tan (*) · A. S. Leach · T. M. M. Heenan · R. Jervis · D. J. L. Brett · P. R. Shearing Electrochemical Innovation Lab, Department of Chemical Engineering, University College London, London, UK The Faraday Institution, Quad One, Harwell Science and Innovation Campus, Didcot, UK e-mail: [email protected]; [email protected] © The Author(s), under exclusive license to Springer Nature Switzerland AG 2021 Y.-j. Gao et al. (eds.), Advances in Sustainable Energy, https://doi.org/10.1007/978-3-030-74406-9_18

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others. Regardless, there is a strong motivation for 3D studies of the physical and chemical states of the electrodes within electrochemical devices due to the inherently heterogeneous structures and processes occurring within. Examples of applications include micro-and nano-computed tomography (CT) studies, which have already been widely conducted on a broad range of electrochemical devices to visualize factors such as mechanisms of electrode degradation. Adapted from work on other porous media, parameters such as tortuosity, porosity, and pore and particle size distributions [4–7] may also be extracted from real electrodes, which are valuable metrics, given that they are typically indicative of electrochemical performance. This chapter will begin with an introduction to X-ray interactions with matter (for a more detailed treatment of the topic, the reader is directed to references [8–11], followed by a discussion of the current state of lab-based and synchrotron light sources, including details of their operation and potential future developments. With this in mind, the reader will be guided through most X-ray-based 3D characterization methods that have been applied to energy storage and conversion devices, including more novel techniques that are still under development or are in the early stages of application to this field. The limitations of using X-rays to study energy storage and conversion materials and devices will also be discussed, and how alternative particle sources (e.g. neutrons, electrons, and muons) may provide additional insight, or be more applicable for certain materials. Finally, considering the current knowledge gaps within the literature, the potential direction of future research avenues will be discussed.

18.2

X-ray Interactions with Matter

Like any other form of electromagnetic radiation, X-ray photons can be decomposed into an oscillating set of perpendicular electrical and magnetic fields and exhibit waveparticle duality (i.e. having characteristics of both particles and waves). Both energy and momentum are contained in an electromagnetic wave (or correspondingly, a photon traveling at the speed of light, c [m s
1]) and its energy, E [J s], is inversely proportional to its wavelength, λ [m], with the proportionality constant being the Planck constant h [11], and is given by the following equation: E¼

hc λ

ð18:1Þ

When X-rays propagate through a medium, various interactions with the atoms of the medium result in the absorption, scattering, refraction, and reflection of the incident radiation, as with any other form of electromagnetic radiation. The overall attenuation through absorption or scattering of X-ray photons incident on an object is governed by several principal photon interaction mechanisms, as summarized in Fig. 18.1. An excellent account of these photon interaction mechanisms and their relative significance throughout the electromagnetic spectrum is

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Fig. 18.1 X-ray interaction pathways resulting from an incident X-ray beam on an object, adapted with permission from Banhart [11]

provided at length by Evans [10] or more recently by Banhart [11]. To summarize, the principal photon interaction mechanisms are • • • •

The photoelectric effect. Coherent scattering (Rayleigh/Thomson). Incoherent scattering (Compton). Electron–positron pair production (in the fields of the nucleus and atomic electrons). • Photonuclear absorption (Fig. 18.2).

18.2.1 Absorption The photoelectric absorption of X-rays, as shown in Fig. 18.3, is the dominant effect contributing to the attenuation of incident radiation within the X-ray energy range used in most imaging applications and occurs when an incident X-ray photon interacts with a bound electron in an atom. The probability of an electron occupying a space is generally higher for lower shell numbers (e.g. shell 1 or K), compared to higher shells that are further from the nucleus (e.g. shell 7 or Q) as they have smaller orbital volumes; consequently, the photoelectric effect is more likely to occur for the most tightly bound electrons (i.e. the inner K- and L-shells). This, however, is providing that the incident photon energy is equal to, or exceeds, the binding energy of that electron [10]. If it does, the incident photon is completely absorbed, and the electron is released as a photoelectron with an energy equal to the incident photon less the electron’s binding energy.

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Fig. 18.2 Relative importance of principal photon interaction mechanisms that dominate at different photon energies, where at σ ¼ τ, the relative contributions between coherent and incoherent scattering and photoelectric absorption are equal, and at σ ¼ κ, the relative contributions between coherent and incoherent scattering and pair production are equal. Adapted with permission from Evans [10]

Fig. 18.3 Various scattering effects of an atom on an incident X-ray photon [11]

18.2.2 Scattering Different types of X-ray scattering will be introduced here, which can be broadly categorized into elastic and inelastic scattering. During an elastic scattering event, there is no net transfer of energy, but the direction of propagation may be modified; during an inelastic scattering event, there is a transfer of energy, meaning the scattered X-ray is of lower energy than the incident X-ray. X-ray diffraction and fluorescence will be discussed here as examples of elastic and inelastic scattering, respectively. Other types of scattering are currently not used for imaging, such as Compton scattering, resonant inelastic X-ray scattering (RIXS), and X-ray Raman scattering.

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18.2.3 Elastic Scattering – X-ray Diffraction X-ray diffraction has been one of the most widely used techniques for measuring the structure of condensed matter. Its development has been the cornerstone of other solid-state measurements and a lot of our understanding of chemical bonding. Diffraction occurs when X-rays are elastically (meaning no energy is transferred) scattered off of an atom. When this atom is part of an ordered structure as found in many minerals, salts, inorganic compounds, or alloys, the diffracted X-rays interact with each other. When the spacing of the atoms, energy of the X-rays, and angle of incidence of the X-ray abide by Bragg’s law, given by Eq. 18.2, there is constructive interference of the diffracted X-rays and a peak will show on the diffraction pattern. At other atom spacings or incident angles, the interference is destructive, and this makes up the background noise to the diffraction pattern. nλ ¼ 2d sinθ

ð18:2Þ

1:2398 λ

ð18:3Þ

E¼

where E is energy (eV) and λ is photon wavelength (μm). A schematic of X-ray interaction with an atomic lattice leading to constructive interference of scattered radiation (diffraction), according to Bragg’s law is illustrated in Fig. 18.4 and defined in Eq. 18.2, where λ is wavelength of the X-ray, d is the atomic spacing, and θ is the incident angle of the X-ray and Eq. 18.2 for photon energy (in electronvolts). The patterns collected during these measurements are ‘fingerprints’ for unique crystal structures, linked to the size and shape of the unit cell. Every possible arrangement of atoms within a unit cell is given a unique identifier based on its symmetry, known as a space group (there are 230 unique space group types), and results in a distinct diffraction pattern. The shape of the unit cell determines the arrangement of the peaks in the diffraction pattern, whilst the size of the unit cell determines the relative separation of those peaks. This means that materials with the same space group give the same general diffraction pattern, although they may differ in peak position if the size of the unit cells (or inter-atom spacing) is different. There are databases of known structures to compare unknown samples, and also mixtures of phases and/or compounds can be distinguished by comparing the intensity of peaks. With high-resolution patterns, subtleties such as

Fig. 18.4 Schematic of Bragg’s law

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strain on the unit cell can be determined with a fitting of the shape of the diffraction peaks, known as refinement.

18.2.4 Total Attenuation Each of the principal photon interaction mechanisms can be quantified in terms of cross-sections [cm2] that define the probability of an interaction between a photon and an atom within an area traversed by the incident beam, namely, the atomic photoelectric cross-section, τ, for the photoelectric effect; the coherent scattering cross-section, σ coherent, the incoherent scattering cross-section, σ incoherent; the pair production cross-section, κ, for electron positron pair production; and the photonuclear absorption cross-section, σ photonuclear, for the photonuclear absorption effect. These interactions all result in the attenuation of the incident radiation, and therefore to determine the total attenuation of material, each of the contributions can be combined by summation to yield a total cross-section, σ tot [cm2], as defined by Hubbell [12]: σ tot ¼ σ coherent þ σ incoherent þ τ þ κ þ σ photonuclear

ð18:4Þ

At the range of photon energies used in X-ray imaging (between 1 and 100 keV), the photoelectric effect dominates, and both electron–positron pair production and photonuclear absorption are negligible. Therefore, the κ and σ photonuclear terms can be neglected as they only become significant at energies greater than ca. 5 MeV [10, 12], and the total cross-section can be simplified to σ tot ¼ σ coherent þ σ incoherent þ τ

ð18:5Þ

From the density of the material ρ, total cross-section per atom, σ tot, and the atomic mass of the element of interest, Ar, the linear attenuation coefficient, μ (cm
1) or mass attenuation coefficient, μ/ρ (cm2 g
1), can therefore be calculated, as defined by Hubbell [13]: μ N ¼ σ tot A ρ uAr

ð18:6Þ

The total cross-section per atom, σ tot, accounts for each contribution from the principal photon interactions with the material, uAr is the molecular weight of the material, and NA is Avogadro’s constant [13]. This relation is useful because theoretical estimates of the total cross-section are closely aligned to experimentally measured values [12], enabling material identification given a known density. Assuming a monochromatic and collimated X-ray beam of known energy, E, transmission, T, is related to the linear attenuation coefficient, μ, and thickness of a material, t, by the exponential relationship known as the Beer–Lambert law, as

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displayed in Eq. 18.7. In turn, transmission, T, can be determined by measuring the intensities of both the incident beam, I0, and the transmitted beam, I, as it is simply the ratio between the two, as given by Eq. 18.8. T ¼ exp ð
μt Þ T¼

I I0

ð18:7Þ ð18:8Þ

Transmission is a useful concept in practical terms as the intensities of both incident, and transmitted X-ray beams can be measured using scintillators and cameras in X-ray imaging systems, where the incident beam can be measured without the sample in the field-of-view (known variously as the reference, flatfield, light-field, or open-beam image, etc.).

18.2.5 Refraction The wave-like properties of X-ray photons also result in the refraction of incident radiation when passing through a medium, as visible light would when passing through an optically transparent material. The ratio of the phase velocity of electromagnetic radiation within a material to the speed of light in a vacuum is known as the complex refractive index of the material, n. This contains both real (1
δ) and imaginary (β) components, which, respectively, describe the phase shift and attenuation of the incident radiation [14]: n ¼ 1
δ
iβ

ð18:9Þ

Within the range of X-ray energies where the photoelectric effect dominates, but away from the absorption edges of the elements within the medium of interest, both terms can be shown to reduce to [14] r c λ2 2π λ β ðλÞ ¼ μ ðλÞ 4π δ ð λ Þ ¼ ρe

ð18:10Þ ð18:11Þ

Where ρe is the electron density, rc is the classical electron radius, λ is the wavelength of the incident radiation, and μ is the wavelength (or energy)-dependent attenuation coefficient. It is important to emphasize that these equations are only valid at X-ray energies away from the absorption edge because photon interactions

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around the absorption edge result in a jump in the linear attenuation coefficient that is not accounted for in the term ‘μ’.

18.2.6 Fluorescence Incident beams (e.g. X-ray or electron) can eject atomic electrons and produce fluorescence X-rays. Fluorescence occurs when the hole left by the ejected electron is filled by an electron from a higher orbital, e.g. a K-shell hole populated by an L-shell electron. The energy at the orbital closer to the nucleus is lower; consequently, the surplus energy must be ejected as radiation in the form of an X-ray photon that is characteristic of the material. These types of photons are described as characteristic because their energy is equal to the difference between the two shell energies (e.g. between L and K), which in turn are determined by the atomic number of the element. Fluorescence detectors typically output spectra of photon intensities for a range of emitted X-ray energies. If fluorescence is occurring, several sharp intensity peaks are observed around energies that are characteristic of the material’s well-defined electron orbital energies. These peaks can be distinguished according to the electron transition that has occurred, i.e. which outer shell the electron traveled from to fill the lower shell hole. For instance, the L!K shell transition is generally denoted as Kα, and the M!K is noted as Kβ; and if the hole resides in the L shell, a transition from the M shell would be denoted Lα, and so on. Once the peaks have been allocated to corresponding transitions, the intensity of the characteristic emission can be correlated to the amount of each element, thus the material composition. However, although the range of Kα values for the periodic table as a whole is broad, between 0.1–100 keV, for many useful elements values can be relatively low, e.g. Si Kα is ca. 1.7 keV, and therefore highly susceptible to attenuation, making the detection of fluorescence signals from low Z elements increasingly difficult for increasing sample thickness [15].

18.3

X-ray Sources for Characterization

Two main phenomena are exploited for the generation of X-ray photons used in scientific experiments: the first involves bombarding an anode with electrons accelerated from a cathode in X-ray tubes; and the second is the release of radiation from the acceleration of electrons using magnets at particle accelerators [11]. The former is more prevalent in laboratories, where the anode target (often a metal such as tungsten or copper) produces X-ray photons via two main interaction pathways [16]: Bremsstrahlung radiation that has originated from the deceleration of high-speed electrons by the electric fields of the target nuclei, producing a continuous spectrum of photon energies; or characteristic radiation that produces peaks of

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high photon intensities at discrete energies specific to the target element, via fluorescence [15]. The accelerating voltage of an X-ray tube can be considered equivalent to the kinetic energy of the electrons and dictates the highest possible energy of an X-ray photon; i.e., a 120 kV tube voltage results in the emission of X-ray photons of no more than 120 keV in energy [16]. The second origin of X-ray photons often comes from a particular type of particle accelerator: a synchrotron. These consist of many straight sections connected into a near-circular, polyhedral storage ring. Charged sub-atomic particles are injected into this ring and then circulated under a high vacuum. The electrons are decelerated in a controlled manner using magnetic fields to produce X-rays of a particular energy. Several large-scale synchrotron facilities exist today with the primary purpose of serving as light sources for experimentalists in fields ranging from biology to materials science and engineering. ‘Brilliance’ is often used to quantify beam quality, and unlike the continuous spectrum produced in lab sources, synchrotrons produce high brilliance beams; i.e., they produce a high number of photons per unit time, per unit area, with low divergence, and a high degree of monochromaticity. Moreover, beams can be further differentiated according to geometry; depending on the nature of the light source and source optics, the resulting X-ray illumination may have a pencil-, parallel-, fan-, or cone-beam geometry, as illustrated in Fig. 18.5.

18.3.1 Lab-Based X-ray Sources A lab-based X-ray source typically produces a divergent beam with a cone geometry, as presented in Fig. 18.5. This results in the formation of a penumbra and geometric magnification, M, of the object depending on the distance between the object, source, and detector [17]. The geometric magnification of the object is given by M ¼ (R1 + R2)/R1, where R1 is the source-to-sample distance, and R2 is the sampleto-detector distance, also shown in Fig. 18.5. By reducing the spot size of the beam with a microfocus source, the geometric penumbra can be limited to some extent. A comparison between different spot sizes can be seen in Fig. 18.6. The polychromatic nature of the X-ray beam generated with lab-based sources leads to undesirable effects such as beam-hardening because materials tend to attenuate X-rays at lower photon energies more than X-rays at higher photon energies [19], resulting in the transmitted beam having a different spectral composition from the incident beam [11]. To minimize this effect, a combination of source filters and software correction can be used. Despite these inherent shortcomings, the main advantages of lab-based X-ray sources are the significantly lower cost and greater accessibility when compared to large-scale facilities.

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Fig. 18.5 Pencil-, parallel-, and cone-beam configurations [17], and the penumbra effect arising from the cone beam

18.3.2 Synchrotron Light Sources As previously mentioned, the advantages of synchrotron radiation sources are generally due to their favorable beam brilliance; the beams produced are collimated (parallel), high-flux, and tunable to specific photon energies using complex magnets and monochromator crystals [20]. The parallel beam geometry (Fig. 18.5) of a synchrotron source simplifies 3D reconstruction and removes the issue of penumbra effects. Furthermore, the high flux significantly reduces acquisition times, to the extent that experiments may instead be time-limited by stage motor movements and detector readout frequency [17]. For comparison, a typical acquisition that may require hours with lab-based micro-computed tomography (micro-CT) may be completed within minutes for synchrotron micro-CT. Thus, the high temporal resolution achievable at synchrotron sources opens vast possibilities for operando tomography on energy storage materials, where microstructural changes in the material may occur within relatively short timescales. Finally, the highly tunable X-ray energies available with synchrotron radiation enable numerous possibilities

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Fig. 18.6 Reduction in penumbra effects with a microfocus X-ray source (left) (right) compared to X-ray source with larger spot size (right), reproduced with permission from Withers [18]

for performing measurements around the absorption edges of different elements, enabling the study of electrochemical changes in energy materials [21]. The production of X-ray photons at synchrotron light sources was described in simple terms in the previous section, although the differences in configuration between synchrotron light sources and between beamlines at the same light source vary immensely. For end-users, the experimental hutch is where measurements are performed, although many components exist between this hutch and the storage ring to produce and refine the spatial and temporal coherence of the X-ray beam that is required to be compatible with the characterization method of interest. Close to the main storage ring is the ‘front end’ where X-rays are produced using insertion devices (ID), such as electron wigglers and undulators, or sometimes bending magnets, and may then be focused with a focusing mirror. For spectroscopy and scanning microscopy applications, a ‘pencil’ beam is commonly used, with a well-defined spot size. At X-ray imaging and microscopy end stations, a full-field beam is typically used, the size of which can be controlled by adjustable slits along the beam path. The energy of the beam can be selected using monochromators, although in some cases demanding high throughput, a ‘pink’ or ‘white’ beam is used due to the significantly higher flux attainable when using a wider spectral range. Whilst the spatial divergence of the beam is generally small, especially for beamlines located further away from the storage ring (i.e. a parallel geometry may be approximated), further focusing is possible with X-ray optics such as Fresnel zone plates or Kirkpatrick–Baez (KB) mirrors. The sample is typically mounted on a stage, which may be capable of up to six degrees of spatial freedom to align the sample with the beam and detector: X, Y, and Z translation steps, angular rotation, pitch, and yaw. Downstream from the sample,

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detectors are used to record changes to the beam resulting from its propagation through the sample, by measuring either the transmitted beam or its scattered components. The use of evacuated or helium-filled flight tubes within the beam path may be necessary for measurements conducted with ‘soft’ X-rays due to the significance of air attenuation. Many types of detectors exist; the most common type contains a scintillator made of X-ray fluorescent phosphors, such as CdWO4, that converts the incident X-ray beam into visible light. The image formed on the scintillator may then be magnified through coupled optics or captured directly by a charge-coupled device (CCD) within a camera. Whilst scintillator-based detectors measure incident X-ray radiation indirectly by conversion to visible light and offer the highest spatial resolution, they often have a limited dynamic range and introduce noise into an image from various sources (including electronic and readout noise). Photon counting detectors, such as photomultiplier tubes and semiconductor-based hybrid detectors, as offered by Medipix and PILATUS, have been developed to directly measure ionizing radiation, counting single photons that hit the detector pixel. With these detectors, dynamic range is practically unlimited and the images collected are generally noise-free, although it is still possible to saturate detector pixels and potentially damage them if the incident photon flux is too high. Furthermore, there are limits on the spatial resolution achievable due to the pixel size and pitch on these detector arrays. Most of the control, acquisition, and reconstruction software at synchrotron light sources are customized open-source code, running through either a command line or a graphical user interface. This is necessary due to the highly novel nature of each experiment and the numerous possible permutations in the configuration of the beamline. Despite the obvious advantages of using synchrotron light sources for 3D X-ray studies, several practical limitations exist that confine their broader adoption; these include high capital and running costs and limited accessibility, although extensive user programs exist that ensure affordable or free access for academic research. Additionally, the high flux of the synchrotron X-ray sources may induce dosedependent radiation damage to electrodes within in-situ cells due to material dissolution or degradation of polymers commonly used within energy storage materials [22].

18.4

Applications of 3D Imaging

The various beam geometries achievable within a laboratory and at synchrotron light sources have been discussed in the previous section. Cone- and parallel-beam geometries are most frequently used in full-field imaging techniques, where the sample is captured in a ‘single shot’ within a certain field-of-view. The pencilbeam geometry is mainly used for scanning probe imaging, where the sample is rastered across the beam whilst measurements are made, with each point corresponding to a different spatial position. Finally, novel ‘lens-less’ techniques,

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such as coherent diffraction imaging and ptychography, and their 3D reconstruction, will be discussed.

18.4.1 Full-Field Techniques: X-ray Absorption CT As described previously, X-ray absorption imaging (and CT) detects changes in the imaginary part of the complex refractive index of a material, and within the range of energies typically used in X-ray imaging, photoelectric absorption accounts for a large fraction of the total attenuation. Measurements obtained by absorption imaging provide information about the mass density of material through the relationship between transmission and the mass attenuation coefficient, which permits the visualization of the physical structure of a material. With a priori knowledge of the sample composition, some inferences can be made about the presence or absence of phases throughout a sample, provided there is sufficient contrast. X-ray absorption CT is a relatively mature field stemming from medical and industrial diagnostics. Its application to energy materials, whilst more recent, has proved compelling and particularly suited to the heterogeneous nature of electrode materials – enabling a more informed insight not only into the diagnosis of failure mechanisms in commercial devices but also in developing novel battery chemistries such as Li-S and Li-air through visualization of the phenomena occurring within the materials. The main advantage of tomographic techniques lies in the ability to spatially distinguish between features that may be super-positioned in a throughplane 2D image or a subsurface layer in a 2D surface image. Tomographic techniques encompass a wide range of length scales, enabling the visualization of energy storage devices and materials from the system and pack level, to individual particles, and even to individual grains within a single particle. One of the first few applications of X-ray CT on battery systems was the use of an industrial CT scanner on NiMH spacecraft batteries [23]. The use of X-ray tomography for the 3D reconstruction of battery electrodes was pioneered by Shearing et al. in 2010, who performed high-resolution X-ray micro-tomography (micro-CT) on the graphite negative electrode of a Li-ion battery using a Gatan X-ray ultramicroscope [4]. The voxel dimensions obtained were 480 nm using 41.4 magnification, producing a total reconstructed sample volume of 43  348  478μm3, from which microstructural information such as porosity, tortuosity, surface area, connectivity of pores, and pore and particle size distribution were extracted [4]. Haibel et al. performed the first in-situ micro-CT of a battery on alkaline zinc cells at a synchrotron facility and, in addition to quantifying the particle size and spatial distribution of zinc particles, attempted to correlate changes in the volume fraction of the cathode material during discharge to the mean density of the material and hence the chemical processes that occur which result in changes in mean density [24]. Since then, various research groups have applied micro- and nano-CT to diagnose cell failure mechanisms (in terms of microstructural degradation such as cracks and dislocations formed) [25]; determined the morphological evolution

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(in terms of microstructural parameters and the contact area between the different phases, etc.) of a wide range of Li-ion battery cathodes and anodes [26–30] with cycling; studied the current collectors used in batteries [31]; and characterized the presence and structures of lithium metal at the anode (such as the formation of lithium dendrites and other electrodeposited microstructures using synchrotron radiation sources) [32, 33]. Not limited to battery studies, absorption CT has been used to study the evolving microstructure of solid oxide fuel cell materials and multiple length scales, from the overall geometry of the cell to the macropore structure, and to the microstructure of the composite electrodes themselves (where the contrast between the electrolyte, catalyst, and pore phases is required to characterize the triple-phase boundary density and therefore performance of the material [34]). Although the feature size of polymer electrolyte membrane fuel cell catalysts is too small to resolve with X-ray imaging (