Construction Principles and Controllable Fabrication of 3D Graphene Materials (Springer Theses) 9811603553, 9789811603556

Table of contents :
Supervisors’ Foreword
Preface
Acknowledgments
Contents
About the Author
1 Introduction
1.1 Background
1.2 Two-Dimensional Graphene Materials
1.2.1 Structure and Property of Graphene
1.2.2 Synthesis and Application of Graphene
1.2.3 Assembly and Advantages of Three-Dimensional Graphene Materials
1.3 Controllable Synthesis of Three-Dimensional Graphene Materials
1.3.1 Porous Metal Templated Chemical Vapor Deposition
1.3.2 Porous Metal Oxide Templated Chemical Vapor Deposition
1.3.3 Post-Treatment of Graphene Oxides
1.3.4 Templated Polymerization
1.4 Research Approach and Methodology
References
2 Growth Mechanism of 3D Graphene Materials Based on Chemical Vapor Deposition
2.1 Introduction
2.2 Templated CVD Based on Mesoporous MgO
2.2.1 Morphology Replication Based on Mesoporous MgO Templates
2.2.2 Nanostructure Characterization of 3D Mesoporous Graphene Framework
2.2.3 Growth Mechanism of 3D Graphene on the Surface of Mesoporous MgO
2.3 Templated CVD Based on Porous CaO
2.3.1 Synthesis and Characterization of Porous CaO
2.3.2 Templated Growth and Structure Regulation of 3D Porous Graphene
2.3.3 Growth Mechanism of 3D Graphene on the Surface of Porous CaO
2.3.4 Application of 3D Porous Graphene in Lithium–Sulfur Batteries
2.4 Summary
References
3 Construction and Application of 3D Graphene Materials Based on Templated Polymerization
3.1 Introduction
3.2 3D Porous Graphene Mesh
3.2.1 Synthesis of 3D Porous Graphene Mesh
3.2.2 Nanostructure of 3D Porous Graphene Mesh
3.2.3 Heteroatom Doping of 3D Porous Graphene Mesh
3.2.4 Electrocatalytic Performance of 3D Porous Graphene Mesh
3.2.5 Electrocatalytic Activity Mechanism of 3D Porous Graphene Mesh
3.3 Atomic Metal Sites Anchored in 3D Porous Graphene
3.3.1 Material Synthesis and Characterization
3.3.2 Defect Engineering Towards Atomic Co–Nx–C Sites
3.3.3 Bifunctional ORR/OER Activities
3.4 Summary
References
4 Nano-Confined Hybridization and Electrocatalytic Application Based on 3D Mesoporous Graphene Framework
4.1 Introduction
4.2 Construction Principles of Hierarchical Graphene-Based Hybrid Electrocatalysts
4.3 Nano-Confined Electrocatalysts Based on 3D Mesoporous Graphene Framework
4.3.1 Material Synthesis and Characterization
4.3.2 Mechanism of Nano-Confined Hybridization
4.3.3 Electrocatalytic OER Activity
4.3.4 Structure-Property Relation and Activity Mechanism
4.4 Intrinsic Activity Mechanism Study Based on Nano-Confined Electrocatalysts
4.4.1 Nano-Confined Hybridization of Various Active Phases
4.4.2 Structure Transformation and Characterization of Confined Active Phases
4.4.3 Electrocatalytic OER Activity
4.4.4 Guest–Host Chemistry of Multi-Metal Electrocatalysts
4.5 Summary
References
5 Design Principles and Synthesis of 3D Graphene-Analogous Materials and van der Waals Heterostructures
5.1 Introduction
5.2 Synthesis of 3D Molybdenum Disulfide via MgO-Templated CVD
5.3 Controllable Synthesis of 3D Mesoporous Graphene/MoS2 vdW Heterostructure
5.3.1 Rational Design and Synthesis of the 3D vdW Heterostructure
5.3.2 Structure Characterization of the 3D vdW Heterostructure
5.3.3 Growth Mechanism of the 3D vdW Heterostructure
5.4 Electrocatalytic Application of 3D Mesoporous Graphene/MoS2 vdW Heterostructure
5.4.1 Hydrogen Evolution Reaction Activity
5.4.2 Triple Functional Electrocatalytic Activity
5.5 Summary
References
6 Conclusions
Publications and Awards

Citation preview

Springer Theses Recognizing Outstanding Ph.D. Research

Cheng Tang

Construction Principles and Controllable Fabrication of 3D Graphene Materials

Springer Theses Recognizing Outstanding Ph.D. Research

Aims and Scope The series “Springer Theses” brings together a selection of the very best Ph.D. theses from around the world and across the physical sciences. Nominated and endorsed by two recognized specialists, each published volume has been selected for its scientific excellence and the high impact of its contents for the pertinent field of research. For greater accessibility to non-specialists, the published versions include an extended introduction, as well as a foreword by the student’s supervisor explaining the special relevance of the work for the field. As a whole, the series will provide a valuable resource both for newcomers to the research fields described, and for other scientists seeking detailed background information on special questions. Finally, it provides an accredited documentation of the valuable contributions made by today’s younger generation of scientists.

Theses may be nominated for publication in this series by heads of department at internationally leading universities or institutes and should fulfill all of the following criteria • They must be written in good English. • The topic should fall within the confines of Chemistry, Physics, Earth Sciences, • • • • •

Engineering and related interdisciplinary fields such as Materials, Nanoscience, Chemical Engineering, Complex Systems and Biophysics. The work reported in the thesis must represent a significant scientific advance. If the thesis includes previously published material, permission to reproduce this must be gained from the respective copyright holder (a maximum 30% of the thesis should be a verbatim reproduction from the author’s previous publications). They must have been examined and passed during the 12 months prior to nomination. Each thesis should include a foreword by the supervisor outlining the significance of its content. The theses should have a clearly defined structure including an introduction accessible to new PhD students and scientists not expert in the relevant field.

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Cheng Tang

Construction Principles and Controllable Fabrication of 3D Graphene Materials Doctoral Thesis accepted by Tsinghua University, Beijing, China

Author Dr. Cheng Tang Tsinghua University Beijing, China

Supervisors Prof. Qiang Zhang Tsinghua University Beijing, China Prof. Fei Wei Tsinghua University Beijing, China

ISSN 2190-5053 ISSN 2190-5061 (electronic) Springer Theses ISBN 978-981-16-0355-6 ISBN 978-981-16-0356-3 (eBook) https://doi.org/10.1007/978-981-16-0356-3 Jointly published with Tsinghua University Press The print edition is not for sale in China (Mainland). Customers from China (Mainland) please order the print book from: Tsinghua University Press. © Tsinghua University Press 2021 This work is subject to copyright. All rights are reserved by the Publishers, 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 publishers, 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 publishers nor the authors or the editors give a warranty, express or implied, with respect to the material contained herein or for any errors or omissions that may have been made. The publishers remain neutral with regard to jurisdictional claims in published maps and institutional affiliations. This Springer imprint is published by the registered company Springer Nature Singapore Pte Ltd. The registered company address is: 152 Beach Road, #21-01/04 Gateway East, Singapore 189721, Singapore

Supervisors’ Foreword

It is a great pleasure to introduce Dr. Cheng Tang’s thesis work, accepted for publication within Springer Theses. Dr. Tang joined our group in late 2012 for his undergraduate thesis on the topic of aligned carbon nanotube/graphene hybrids. Impressively, his dedicated and efficient work rendered significant achievements in this tough project and inspired him to continue his research in our group. Then he turned his research interest into graphene, which is more attractive but highly remaining to be researched both for fundamental science and industrial application. He perceptively focused on the controllable construction of 3D graphene framework with intrinsic sp2 hybridization and hierarchical porosity, aiming at bridging the 2D ideal nanostructure and 3D practical materials. Dr. Tang’s thesis includes a series of innovative works on the growth mechanism, synthesis methodology, customized application, and system promotion of 3D graphene materials. The achievements will no doubt benefit the fundamental research and industrial development of graphene with significantly improved performance, and also inspire the further research of various nanomaterials beyond graphene, which is highlighted as following: • The fundamental understanding of graphene growth behavior on metal oxides promotes the mechanism study and mass production of high-quality graphene for high-performance electrochemical energy applications. • The versatile synthetic strategies and utilization concepts of 3D graphene materials offer new opportunities to deepen the understanding of the structure–function relationship of 3D graphene materials and remarkably enhance their performance in energy electrocatalysis. • The verified 3D concept for various 2D materials opens up a new direction for all-round engineering of nanomaterials toward beneficial properties and highperformance application. A significant part of this thesis has been published in flagship Materials & Chemistry journals including Adv. Mater., Acc. Chem. Res., Adv. Funct. Mater. and J. Mater. Chem. A, as well as 4 granted Chinese patents in the synthesis and application of 3D graphene materials. Dr. Tang was also the group leader of an interdisciplinary v

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Supervisors’ Foreword

innovation team in Tsinghua focusing on the practical application of graphene, and made remarkable efforts in science popularization, including a popular science video and a Chinese book chapter about graphene. These fruitful outputs demonstrate the significant amount of efforts and the important contributions to the scientific community. As a result, Dr. Tang’s thesis has been awarded the First Prize of Excellent Doctoral Dissertation at Tsinghua University and the 2019 Chorafas Foundation Award in Chemistry. Dr. Tang never stops his pace of advance and keeps himself at the leading edge of chemistry and nanotechnology research. He said that he enjoyed the research life full of curiosity, challenge, innovation, and significance as well. We wish Dr. Tang to continue his outstanding scientific career, with more achievements in challenging, important and unexplored topics. We also hope that this thesis can provide some inspiration for other researchers in the field of nanomaterials and energy fields. Beijing, China October 2020

Prof. Qiang Zhang Prof. Fei Wei

Preface

Graphene, as a mystery two-dimensional (2D) nanomaterial with unique and excellent properties, has a wide range of application prospects in many fields. It has been gradually extended from academic research to industrial development; however, the efficient bridge from nanostructure to macroscopic materials fundamentally determines the ultimate demonstration of intrinsic properties. By directly constructing a three-dimensional (3D) graphene framework with sp2 hybridization, it will not only maintain the excellent intrinsic properties of 2D graphene in the macroscopic 3D structure, but also is expected to generate some novel properties and significantly improved performance due to the unique 3D structure. Aiming at the underlying growth mechanism and construction principles of 3D graphene materials, the self-limiting growth behavior of graphene on the defect-rich oxygen-terminated polar MgO (111) crystal surface was ascertained. Under its guidance, the porous CaO was introduced as the template for chemical vapor deposition of 3D graphene materials. The roles of the surface chemistry and porous structure of oxide templates, and detailed reaction conditions on the resultant graphene structure was revealed. Ultimately, it was successfully achieved to facilely tune the number of graphene layers (1 to 10 layers) and porous structure (from micropores, mesopores, to macropores). The application of 3D graphene materials in high-rate lithium– sulfur batteries was also investigated, with specific attention on the impact of porous structure hierarchy. To solve the difficulty in versatilely controlling the structure, composition, and properties of 3D graphene materials, a new method based on metal oxide templated polymerization was developed. A 3D porous graphene mesh material was prepared, with great advantages in the regulation of specific surface area (1655 m2 g–1 ), porous structure (3.39 cm3 g–1 ), and nitrogen doping (7.60 at%). Owing to the unique structure of as-obtained 3D graphene material, the bi-functional electrocatalytic activity of the topological defects in graphene materials for oxygen reduction and oxygen evolution reactions was strongly revealed both experimentally and theoretically. Innovatively, the synthesis strategy of “defect engineering” was proposed to smartly utilize the topological defects to efficiently construct monodispersed Co–Nx –C sites in 3D graphene materials for enhanced electrocatalytic activities. vii

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In order to fully exploit and make use of the advantages of 3D graphene materials, the concept of “nanoreactor” was creatively proposed to efficiently prepare heterogeneous electrocatalysts. Under the combined effect of defect-anchored nucleation and spatially confined growth, the obtained active phases were in quantum dot size, uniformly dispersed in the 3D graphene framework, and strongly coupled with graphene, resulting in greatly improved electrocatalytic activities for oxygen evolution reaction. This highly efficient synthetic strategy and unique structural design made it possible to probe the fundamental electrocatalytic activity mechanisms, thereby ascertaining the guest–host chemistry of the dual-metal hydroxides in oxygen evolution electrocatalysis. Considering the wide variety of graphene-analogous 2D nanomaterials with different properties, the concept of 3D nanostructures has been generalized. By introducing graphene and ammonia to regulate the substrate surface characteristics and partial pressure during chemical vapor deposition, the 3D assembly of molybdenum disulfide nanosheets was successfully achieved, and a 3D mesoporous van der Waals heterostructure with excellent trifunctional electrocatalytic activities was successfully synthesized. To summary up, the 3D graphene material’s growth mechanism, synthesis methodology, customized application and system promotion were gradually developed, which is expected to deepen the understanding of the structure–function relationship of 3D graphene materials, and open up new ideas for versatile functionalization and high-performance application of graphene and other 2D nanomaterials. Beijing, China December 2020

Dr. Cheng Tang

Acknowledgments

Time flies, and in a blink of an eye, the five-year doctoral career and nine-year study at Tsinghua University (THU) are coming to an end. Along the way, I have grown and gained a lot. First of all, grateful acknowledgment is made to my supervisors, Prof. Fei Wei and Prof. Qiang Zhang. Their erudition and noble character have deeply affected me. I saw the demeanor of a scholar in them, and was impressed by their tireless academic enthusiasm and meticulous scientific spirit. I also sincerely hope that I can inherit a little bit as their student. I would like to express my deepest gratitude to Prof. Zhang for the great trust and expectation that he has always placed on me. He is my mentor and also my friend, not only guiding and urging me in scientific research, but also caring about my future development and my life. All my achievements in scientific research are inseparable from the encouragement of Prof. Zhang. I would like to express heartfelt gratitude to Prof. Maria Magdalena Titirici of Queen Mary University of London for her supervision during my exchange visit, to Prof. Bingsen Zhang of Institute of Metal Research, Chinese Academy of Sciences for his help and guidance in my experiment, and to Prof. Ovidiu Ersen of University of Strasbourg for helping with the 3D electron tomography. I would also like to thank Prof. Wei-Zhong Qian, Prof. Yao Wang, Prof. Guo-Hua Luo, and others at Fluidization Laboratory of Tsinghua University for their guidance and help. In addition, I would like to express my gratitude to Dr. Meng-Qiang Zhao, Dr. Jing-Qi Nie, Prof. Jia-Qi Huang, Hao-Fan Wang, Bo-Quan Li, Hong-Jie Peng, XinBing Cheng, Ling Zhong, Han-Sen Wang, Lin Zhu, Dr. Bin Wang, Dr. Xiaoyang Cui, Xiang Chen, Dr. Xiao Chen, and all students in our research group. It is precisely because of being in such an excellent group that I can grow together with everyone. Last but not the least, my gratitude also extends to my family for their unconditional support and love, especially my dear wife Shao Chen. The future is promising and the road ahead is long. I can only strive diligently and never stop the steps forward to repay the caring and dedication of the abovementioned people. The project was funded by National Natural Science Foundation of China and National Key Research and Development Program of China. I would like to express my sincere thanks here for the support. ix

Contents

1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.1 Background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.2 Two-Dimensional Graphene Materials . . . . . . . . . . . . . . . . . . . . . . . . . 1.2.1 Structure and Property of Graphene . . . . . . . . . . . . . . . . . . . . . 1.2.2 Synthesis and Application of Graphene . . . . . . . . . . . . . . . . . . 1.2.3 Assembly and Advantages of Three-Dimensional Graphene Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.3 Controllable Synthesis of Three-Dimensional Graphene Materials . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.3.1 Porous Metal Templated Chemical Vapor Deposition . . . . . . 1.3.2 Porous Metal Oxide Templated Chemical Vapor Deposition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.3.3 Post-Treatment of Graphene Oxides . . . . . . . . . . . . . . . . . . . . 1.3.4 Templated Polymerization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1.4 Research Approach and Methodology . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2 Growth Mechanism of 3D Graphene Materials Based on Chemical Vapor Deposition . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2 Templated CVD Based on Mesoporous MgO . . . . . . . . . . . . . . . . . . . 2.2.1 Morphology Replication Based on Mesoporous MgO Templates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.2 Nanostructure Characterization of 3D Mesoporous Graphene Framework . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2.3 Growth Mechanism of 3D Graphene on the Surface of Mesoporous MgO . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3 Templated CVD Based on Porous CaO . . . . . . . . . . . . . . . . . . . . . . . . 2.3.1 Synthesis and Characterization of Porous CaO . . . . . . . . . . . 2.3.2 Templated Growth and Structure Regulation of 3D Porous Graphene . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1 1 4 4 5 7 10 11 15 18 21 23 25 35 35 36 36 38 41 42 42 44

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2.3.3 Growth Mechanism of 3D Graphene on the Surface of Porous CaO . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3.4 Application of 3D Porous Graphene in Lithium– Sulfur Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3 Construction and Application of 3D Graphene Materials Based on Templated Polymerization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2 3D Porous Graphene Mesh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.1 Synthesis of 3D Porous Graphene Mesh . . . . . . . . . . . . . . . . . 3.2.2 Nanostructure of 3D Porous Graphene Mesh . . . . . . . . . . . . . 3.2.3 Heteroatom Doping of 3D Porous Graphene Mesh . . . . . . . . 3.2.4 Electrocatalytic Performance of 3D Porous Graphene Mesh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.2.5 Electrocatalytic Activity Mechanism of 3D Porous Graphene Mesh . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 3.3 Atomic Metal Sites Anchored in 3D Porous Graphene . . . . . . . . . . . 3.3.1 Material Synthesis and Characterization . . . . . . . . . . . . . . . . . 3.3.2 Defect Engineering Towards Atomic Co–Nx –C Sites . . . . . 3.3.3 Bifunctional ORR/OER Activities . . . . . . . . . . . . . . . . . . . . . . 3.4 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4 Nano-Confined Hybridization and Electrocatalytic Application Based on 3D Mesoporous Graphene Framework . . . . . . . . . . . . . . . . . . 4.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.2 Construction Principles of Hierarchical Graphene-Based Hybrid Electrocatalysts . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3 Nano-Confined Electrocatalysts Based on 3D Mesoporous Graphene Framework . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.3.1 Material Synthesis and Characterization . . . . . . . . . . . . . . . . . 4.3.2 Mechanism of Nano-Confined Hybridization . . . . . . . . . . . . . 4.3.3 Electrocatalytic OER Activity . . . . . . . . . . . . . . . . . . . . . . . . . 4.3.4 Structure-Property Relation and Activity Mechanism . . . . . . 4.4 Intrinsic Activity Mechanism Study Based on Nano-Confined Electrocatalysts . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.4.1 Nano-Confined Hybridization of Various Active Phases . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.4.2 Structure Transformation and Characterization of Confined Active Phases . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 4.4.3 Electrocatalytic OER Activity . . . . . . . . . . . . . . . . . . . . . . . . . 4.4.4 Guest–Host Chemistry of Multi-Metal Electrocatalysts . . . . 4.5 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

49 50 54 55 57 57 58 58 60 64 68 73 76 77 80 83 85 86 89 89 90 93 93 95 98 100 105 105 108 111 113 115 116

Contents

5 Design Principles and Synthesis of 3D Graphene-Analogous Materials and van der Waals Heterostructures . . . . . . . . . . . . . . . . . . . . 5.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.2 Synthesis of 3D Molybdenum Disulfide via MgO-Templated CVD . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3 Controllable Synthesis of 3D Mesoporous Graphene/MoS2 vdW Heterostructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.1 Rational Design and Synthesis of the 3D vdW Heterostructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.2 Structure Characterization of the 3D vdW Heterostructure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 5.3.3 Growth Mechanism of the 3D vdW Heterostructure . . . . . . . 5.4 Electrocatalytic Application of 3D Mesoporous Graphene/MoS2 vdW Heterostructure . . . . . . . . . . . . . . . . . . . . . . . . . 5.4.1 Hydrogen Evolution Reaction Activity . . . . . . . . . . . . . . . . . . 5.4.2 Triple Functional Electrocatalytic Activity . . . . . . . . . . . . . . . 5.5 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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119 119 120 123 123 124 127 129 129 133 135 136

6 Conclusions . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 139 Publications and Awards . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 143

About the Author

Dr. Cheng Tang received his B.Eng. and Ph.D. from the Department of Chemical Engineering, Tsinghua University in 2013 and 2018, respectively, under the supervision of Prof. Qiang Zhang and Prof. Fei Wei. He worked as a Postgraduate Research Associate at Queen Mary University of London in Prof. Maria-Magdalena Titirici’s Research Group from May to August 2015. During his Ph.D. research, Dr. Tang focused on the design and synthesis of hierarchical porous carbon-based materials and their applications to electrocatalysis and batteries. Since August 2018 he has worked as an ARC Research Associate at The University of Adelaide, where he turned research interest to atomic-level design and engineering of nanomaterials for high-performance electrochemical production of fuels and chemicals. His research activities focus on the development of functional nanomaterials for key reactions in various electrocatalysis and electrosynthesis technologies, including oxygen reduction/evolution reactions (ORR/OER), hydrogen evolution reaction (HER), nitrogen reduction reaction (NRR), and CO2 reduction reaction (CRR). His major contributions involve material synthesis, mechanism study, and catalyst design for targeted electrochemical reactions. Dr. Tang has co-published 1 book chapter and 67 refereed journal papers (to 30 October, 2020), including 3 ESI Hot paper (top 0.1%), 17 ESI Highly Cited papers (top 1%, 13 as the first author), and 8 cover-featured papers (7 as the first author). These include 32 papers as the first/co-first authors in flagship Materials & Chemistry journals including Angew. Chem. Int. Ed. (1), Adv. Mater. (9), Sci. Adv. (1), Chem. Soc. Rev. (1), Acc. Chem. Res. (1), Adv. Funct. Mater. (1), ACS Catal. (1), J. Mater. Chem. A (5), and others. His h-index = 40 and total citations > 6,050 on Google Scholar, and h-index = 39 and total citations > 5,240 on Web of Science. He was awarded the 2020 Clarivate Analytics Highly Cited Researchers in the ‘Cross-Field’ category.

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Chapter 1

Introduction

1.1 Background Apart from air and water, what else is everywhere in our lives? Looking around, we can find that the cup for drinking is made of glass, the clothes we wear are made of fiber products, the books we read are made of paper, the wardrobe at home is made of wood, and the building where we live is made of reinforced concrete. All items are made of different materials. It can be said that materials have already penetrated into our food, clothing, shelter and transportation, serving as the basis for human survival and development. Obviously, today’s human life can no longer leave the materials. The process of human use of materials is almost parallel to the process of human civilization. As early as tens of thousands of years ago, humans began to use one of the earliest materials—stone—to cast artifacts. The birth of stone tools is a manifestation of the ancient wisdom of mankind, and it also marks the beginning of a distinction between mankind and other animals, resulting in an important era—the “Stone Age”. With the continuous advancement of cognition and the development of processing technology, the use of materials by mankind has entered a new stage, from simply polishing the shape to changing its properties, and then discovering new materials. As a result, the “Bronze Age”, “Iron Age”, and “Silicon Age” named after the materials come one after another. Taking a material as a footnote to a section of human civilization is sufficient to illustrate the great significance of materials for human society. It is worth noting that for each new material, from the first discovery to controllable preparation, use by mankind, popularization, and even the promotion of industrial technology revolution, it requires not only initial insight and creativity, that is, the guidance of theoretical scientific knowledge such as chemistry and physics, but also sophisticated and effective up-scaling processes, that is, the application of engineering science and technology such as chemical engineering.

© Tsinghua University Press 2021 C. Tang, Construction Principles and Controllable Fabrication of 3D Graphene Materials, Springer Theses, https://doi.org/10.1007/978-981-16-0356-3_1

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After the international financial crisis, the global industrial competition pattern has undergone major adjustments. Developed countries have re-examined their development strategies and formulated “re-industrialization” strategies centered on revitalizing their manufacturing industry. A new round of technological revolution and industrial transformation characterized by the integration and development of a new generation of information and communication technology and manufacturing is being bred up on a global scale, which has promoted the profound adjustment of the world’s industrial technology and division of labor. Manufacturing has become the commanding heights of global economic competition. As one of the three pillars (new materials, new information, and new energy) of the world’s new technology revolution, the new material industry is undoubtedly the basis and guide for the development of advanced manufacturing, and an important support for industrial upgrading and building new competitive advantages. What is the new material that people are looking forward to? How will it transform our world and lead us to the next stage of development? What important role will chemical engineering and technology play in it? In 2010, the Nobel Prize in Physics was awarded to scientists Andre Geim and Konstantin Novoselov from the University of Manchester, in recognition of their work “for groundbreaking experiments regarding the two-dimensional material graphene”. Since then, graphene, as the “King of new materials”, has set off an upsurge in academia and industry. Andre Heim has publicly expressed the hope that graphene will change our lives like plastic. China’s national strategy “Made in China 2025” clearly pointed out that in the field of new materials, it is necessary to do a good job in the advance layout and development of strategic cutting-edge materials such as graphene. The development of graphene materials has been incorporated into China’s national strategic layout. Graphene is a type of two-dimensional (2D) nanocarbon materials with Dirac carbon structure. Due to its high electron mobility, excellent chemical stability, outstanding electrical, thermal and mechanical properties, it not only creates a new space in theoretical science, but also has shown great application potential in the fields of semiconductors, flexible electronics, sensors, composite materials, energy storage and conversion, biomedicine, environmental protection, and thermal management (Fig. 1.1) [1]. Graphene is currently a hot spot in the field of new materials research and application [1–3]. However, the structure and quality of graphene are greatly affected by different synthesis methods, so the actual physical and chemical properties and resulting performance fluctuate greatly. With the continuous deepening of basic research and the gradual advancement of industrial technology, the development of graphene materials has entered a “climbing period”, and there is an urgent need for conceptual and technological innovation in efficient synthesis and main-material application. The controllable construction of three-dimensional (3D) hierarchical structure is the basis and core for the efficient application of graphene materials. It can effectively avoid the stacking of 2D nanosheets and the agglomeration of nanomaterials, and ensure the full demonstration of theoretical high surface area, electrical conductivity and mechanical strength in macroscale devices. But so far, the construction principles

1.1 Background

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Fig. 1.1 The application fields of graphene materials. Reprinted from ref. [1], copyright 2015, with permission from The Royal Society of Chemistry

of 3D graphene materials for high-performance applications are lacking, and the controllable synthesis methods are limited, which make it difficult to promote the continuous development in academic research and the rapid development of the new material industry. Therefore, this Thesis aims to realize the fine design and controllable synthesis of ideal 3D graphene materials. Using porous metal oxides as the growth template, the research gradually expands from the four aspects of 3D graphene materials including: growth mechanism, synthesis method, customized application, and concept generalization. Firstly, combining the knowledge of surface catalysis and reaction process, a comprehensive understanding of the growth mechanism and construction principles of 3D graphene on the surface of metal oxides was achieved. On this basis, a porous CaO template was proposed for the first time for the chemical vapor deposition (CVD) of graphene, which expanded and improved the metal oxide-based strategies for 3D graphene material synthesis and property regulation. Secondly, a new strategy of templated polymerization was developed to prepare 3D graphene materials, leading to efficient control of the material structure and performance, and a new family of 3D graphene materials. Thirdly, according to the structural characteristics of 3D graphene materials, customized and efficient application strategies were developed. New concepts of “defect engineering” and “nanoreactor” were proposed

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for the controllable synthesis of new functional materials. Meanwhile, using the 3D graphene and its hybrid catalysts as the material platform, the fundamental activity mechanism of several key electrocatalytic reactions was deeply explored, revealing the “defect chemistry” of carbon materials and the “guest–host chemistry” of dualmetal hydroxides. Finally, the concept, principles and strategies of 3D nanostructures were successfully extended to other graphene-analogous 2D material systems. A new kind of 3D mesoporous interlayer van der Waals heterostructure was synthesized for the first time. This Thesis provides important theoretical guidance and innovative concepts for the development of graphene and other 2D nanomaterials, rendering an important impact in many fields such as nanomaterials and energy catalysis.

1.2 Two-Dimensional Graphene Materials 1.2.1 Structure and Property of Graphene Graphene, a 2D crystal of monoatomic layer, is formed by densely arranged carbon atoms in a hexagonal lattice. In 2004, Andre Geim and Konstantin Novoselov and co-workers reported a method of repeatedly sticking and tearing highly oriented pyrolytic graphite with tape, which can obtain high-quality monolayer graphene simply and efficiently [4]. They successfully transferred it to a silicon substrate, and can characterize and locate graphene through the color difference under the optical microscope. They then systematically studied its electrical properties, leading to the discovery that graphene has a bipolar electric field effect, a high carrier concentration and mobility, and sub-micron-scale ballistic transport characteristics [4]. They shared the Nobel Prize in Physics 6 years later for this work. Since then, as the first singleatom-layer 2D material obtained experimentally, this breakthrough has set off the upsurge for investigating graphene and other 2D materials. Consequently, graphene has attracted much scientific interest in the field of materials both academically and commercially in the past decade. Thanks to the unique sp2 hybridized monoatomic layer crystal structure, graphene has many excellent physical properties, as shown in Table 1.1. Theoretically, monolayer graphene is a zero-bandgap semiconductor material. The movement of electrons does not follow the Schrodinger equation, but the Dirac equation. It can be approximated as massless and run at speeds up to 1/300 of the speed of light. The electrons of graphene are strictly confined in a 2D plane, and the quantum Hall effect can be observed even at room temperature. Theoretical calculations show that the carrier mobility of defect-free monolayer graphene is as high as 2 × 105 cm2 V−1 s−1 [5], and the measured value for mechanically exfoliated graphene in a suspended state was similar [6], which is 100 times that of commonly used silicon materials. The thermal conductivity of graphene is theoretically as high as 10,000 W m−1 K−1 [7], and experimentally measured to be 5300 W m−1 K−1 [8], which is close to three times that of diamond. The electrical conductivity of graphene is theoretically as high as 106

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Table 1.1 Intrinsic physical properties of graphene, including theoretical values for monolayer graphene and experimental maximum values Properties

Theoretical values Experimental values Notes

Bandgap (eV)

0

–

Silicon (300 K) 1.1

Electrical conductivity (S cm−1 )

~106

–

Silver 0.6 × 106

Carrier mobility (cm2 V−1 s−1 )

2 × 105

1.5 × 104

Silicon 1.5 × 103

Thermal conductivity (W m−1 K−1 )

~10,000

5300

Diamond 2000

Current carrying capacity (A cm−2 )

~109

–

Copper ~106

Specific surface area (m2 g−1 ) 2630

–

–

Light transmittance

~97.7%

97.4%

–

Young’s modulus (TPa)

1.050

~1.0

Steel 206 GPa

Tensile strength (GPa)

130

130

Steel 1.3

S cm−1 , which is comparable to copper, but the current carrying capacity of graphene can reach 109 A cm−2 , which is 1000 times that of copper. Free-standing graphene is highly transparent and has a transmittance of up to 97.7% of visible light, regardless of wavelength [9]. The areal density of monolayer graphene is 0.77 mg m−2 , and the specific surface area reaches 2600 m2 g−1 . In addition, graphene has a very high Young’s modulus (1050 GPa) [10] and tensile strength (130 GPa) [11], much higher than steel with the same thickness. Compared with other materials, the uniqueness and advantages of graphene are not only the outstanding intrinsic physical properties, but also the integration of many aspects, which has laid an outstanding material foundation for the fundamental research and application development of graphene. However, it should be noted that the above properties are mostly theoretical calculation values or measured results of high-quality monolayer suspended graphene obtained by mechanical exfoliation. Most prepared graphene samples, especially graphene materials prepared in batches, exhibit much poorer properties due to the existence of defects and interfaces, which limits the performance in macroscopic devices. This has become a core challenge that needs to be overcome in the academic research and application development of graphene materials.

1.2.2 Synthesis and Application of Graphene After more than ten years of research, the synthesis method of graphene has been continuously developed and improved, and the cost has dropped significantly while the quality has gradually improved. Specifically, the synthesis methods of graphene

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Fig. 1.2 a The relationship between cost and quality for different synthesis methods of graphene. Reprinted from ref. [2], copyright 2012, with permission from Springer Nature. b 3 mg mL−1 graphene oxide dispersion. c Transmission electron microscope (TEM) image of graphene oxide. b, c are reprinted from ref. [19], copyright 2015, with permission from Springer Nature. d Schematic diagram of single crystal graphene prepared by CVD method. Reprinted from ref. [27], copyright 2014, with permission from American Chemical Society. e Optical image of a graphene film transferred on SiO2 /Si wafers. Reprinted from ref. [17], copyright 2009, with permission from American Chemical Society

mainly include mechanical exfoliation [4, 6, 8, 11], chemical vapor deposition (CVD) [12–18], oxidation–reduction and liquid phase exfoliation [19–22], SiC epitaxial growth [23], molecular assembly [24, 25] and other methods. Among them, CVD, oxidation-reduction, and liquid phase exfoliation are the most mature methods for large-scale production of graphene. The structure features (e.g., number of layers, size, pore structure) and properties (e.g., defect density, surface functional groups, electrical conductivity, thermal conductivity, mechanical strength) of graphene materials obtained by different methods are quite different, thus suitable for different application areas (Fig. 1.2a) [1, 3]. As shown in Figs. 1.2b and c [19], the graphene micro flakes obtained by oxidationreduction and liquid phase exfoliation have the lowest cost, the largest productivity, and feasibility for processing. However, the obtained graphene sheets are relatively small in size and have many defects. So they are widely used in composite reinforcement materials, anti-corrosion coatings, conductive inks, electrochemical energy storage, biosensing, adsorption and separation, transparent conductive films, and other fields [1, 26]. In contrast, the graphene films obtained by CVD (Fig. 1.2d [27], e [17]) are of high quality, controllable number of layers, and high conductivity. Therefore, they have excellent application prospects in the fields of electronic devices, transparent conductive films, field effect transistors, etc [2, 12]. As a typical 2D nanomaterial, although graphene has excellent basic physical properties and the synthesis process is constantly improving, the obtained graphene materials inevitably encounter the challenges for controllable 3D assembly and demonstration of intrinsic properties in practical devices. The problem of sheet

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stacking and agglomeration, messy and inefficient assembly, and abundant defects derived from interface connections will greatly reduce the light transmittance, electrical conductivity, thermal conductivity, and mechanical strength of graphene materials, thereby limiting their application.

1.2.3 Assembly and Advantages of Three-Dimensional Graphene Materials Constructing 3D graphene materials through a “bottom-up” approach can not only maintain the excellent properties and performance of 2D graphene structural units, such as high specific surface area and high conductivity, but also derive novel characteristics due to the unique macroscopic structures, such as hierarchical porosity, 3D mechanical strength, and nano-scale confinement. Therefore, obtained 3D graphene materials have shown great application prospects in the fields of electrochemical energy storage, electrocatalysis, composite materials, and environmental protection [28–33]. The hierarchical nanostructure of 3D graphene materials can guarantee the rapid transfer of electrons and ions, provide sufficient mechanical flexibility and chemical stability during the reaction process, and provide a large specific surface area for abundant active sites and interfaces. Therefore, 3D graphene materials are widely used in electrochemical energy storage devices, such as lithium-ion batteries [34– 37], lithium–sulfur batteries [38–40], and supercapacitors [29, 41, 42]. Zhao and co-workers prepared a 3D graphene/tin nanoparticle composite material by a onestep CVD method using NaCl as the template and SnCl2 ·2H2 O and citric acid as precursors, which was used as the anode for lithium-ion batteries (Fig. 1.3a) [43]. Tin nanoparticles were 5~30 nm in size and wrapped by a porous graphene framework

Fig. 1.3 a TEM image of 3D graphene/tin nanoparticle composites and b corresponding lithiumion battery performance. Reprinted from ref. [43], copyright 2014, with permission from American Chemical Society

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(layer thickness was about 1 nm). The porous structure of obtained 3D graphene effectively prevented the aggregation of tin nanoparticles and enable their direct contact with the electrolyte. Besides, its excellent mechanical flexibility buffered the huge volume change of the electrode material during the charge and discharge process. Combined with excellent electrical conductivity and structural stability, the rate performance and cycle stability of assembled lithium-ion batteries were greatly improved (Fig. 1.3b). The specific capacity can reach 682 mAh g−1 at a current density of 2 A g−1 , with a retention of 96.3% after 1000 cycles. Zhang and co-workers used 3D mesoporous graphene materials as the cathode scaffold for lithium–sulfur batteries [44–48], in which the mesopores enabled a high loading of sulfur and buffered the volume changes of active materials, the graphene framework provided efficient electron conductive network, and the interface effect enhanced adsorption of intermediate products, thus resulting in greatly improved sulfur utilization, rate performance and cycle stability. In addition, 3D graphene materials can also be used as the “amphiphilic” synergistic interface separator for lithium–sulfur batteries [49] and for the protection of lithium metal anodes [50]. In the field of supercapacitors, 3D graphene materials can be directly used as electrodes for high-performance electric double layer capacitors due to their huge specific surface area, excellent electrical conductivity and porous mass transfer channels. Shi and co-workers assembled graphene oxide (GO) into 3D graphene hydrogel by hydrothermal treatment. When used directly for supercapacitors, the specific capacity can reach 160 F g−1 [51], while after reduction by hydrazine to improve the conductivity, the specific capacity can increase to 222 F g−1 [52]. The performance of electric double layer supercapacitors can be further improved by optimizing the synthesis methods or post-treatment, by improving the conductivity [52] and the specific surface area [53–55], by regulating the pore size distribution [56–58], and by introducing heteroatom dopants [59–62]. For example, Huang and co-workers fabricated a few-layer 3D nitrogen-doped ordered mesoporous graphene-like materials using nickel-assisted CVD on mesoporous SiO2 templates [63]. The obtained material exhibited abundant mesopores (2~4 nm), a large pore volume of 2.20 cm3 g−1 , a high specific surface area of 1580 m2 g−1 , and excellent conductivity as well. When it was used in electrolyte electric double layer supercapacitors with 2 M Li2 SO4 electrolyte, the energy density and power density based on the device mass can be as high as 41.0 Wh kg−1 and 26.0 kW kg−1 . Additionally, the 3D graphene materials can provide an ideal substrate to efficiently hybridize with other electrochemically active components, such as transition metal oxides [64–66] and conductive polymers [67, 68], thereby greatly improving the energy density by increasing pseudo-capacitance. Furthermore, due to its excellent mechanical flexibility and 3D interconnected framework, 3D graphene materials also offer ideal candidates for the development of high-performance flexible and bendable batteries [69–71] and supercapacitors [72–74]. In recent years, gas-involving heterogeneous electrocatalytic reactions, such as oxygen reduction reaction (ORR), oxygen evolution reaction (OER) and hydrogen evolution reaction (HER), have received more and more attention. ORR is the most important electrode reaction for fuel cells and metal-air batteries, while the coupling

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of OER and HER can realize green hydrogen production by electrocatalytic water splitting. They are all of great significance to the development of next-generation energy technologies [1, 75–77]. Nanocarbon materials and nanocarbon-transition metal compound hybrids have been widely demonstrated as excellent alternatives to replace precious metal catalysts due to their high activity and stability, thus promoting the large-scale practical application of related energy devices [78–88]. As shown in Fig. 1.4a, the gas-involving heterogeneous electrocatalytic reactions only occur at the gas-solid-liquid triple-phase boundary regions. Three critical steps are coupled with each other including mass diffusion, electron transfer, and surface catalytic reactions. Correspondingly, the optimal catalyst requires high electrical conductivity, fast mass transfer, and high intrinsic catalytic activity. Compared with conventional nanocarbon materials, 3D graphene materials exhibit more prominent advantages. On one hand, the doping of nitrogen, phosphorus, boron and other heteroatoms can alter the electronic structure of graphene and thus generate highly active catalytic sites. The resulting ORR and OER performance under alkaline conditions can be comparable to that of precious metal catalysts [77, 89–99, 100–103]. On the other hand, the further optimization by 3D nanostructures can expose the active sites to the greatest extent, strengthen the surface contact and reaction at the gas-solid-liquid

Fig. 1.4 a Schematic diagram of gas-involving heterogeneous electrocatalytic reaction. Reprinted from ref. [107], copyright 2018, with permission from American Chemical Society. b TEM image of mesoporous/microporous nitrogen-doped graphene with excellent ORR activity. Reprinted from ref. [104], copyright 2014, with permission from Springer Nature. c TEM image of 3D nitrogendoped graphene/NiCo hydroxide composite catalyst. Reprinted from ref. [108], copyright 2013, with permission from John Wiley and Sons

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triple-phase boundary regions, and further improve the apparent catalytic performance of the 3D graphene materials (Fig. 1.4b) [104]. Through theoretical calculations, Qiao and co-workers revealed that HER activity of dual-doped graphene (such as nitrogen and sulfur co-doped) can be much better than that of singledoped graphene, and the activity continued to improve with the increase in doping content and the specific surface area [105]. For graphene materials with a constant doping content of 5%, the HER activity can exceed that of MoS2 catalysts when the specific surface area increases to 1000 m2 g−1 [105]. Additionally, Hu and co-workers revealed that the topological defects of the 3D graphene material itself, such as zigzag edges and five-membered ring structures, are also a very important type of active sites [106]. The defect-rich dopant-free graphene nanocages were demonstrated to be promising ORR electrocatalysts with comparable activity to nitrogen-doped carbon nanomaterials [106]. Besides the highly active catalytic sites provided by itself, 3D graphene material can also be used as a multifunctional substrate to efficiently hybridize with transition metal compounds, which can regulate the nanostructure and interface coupling of transition metal compounds, and ensure the rapid transfer of electrons and ions [110]. It will help fully demonstrate the catalytic activity of the active phases [111– 115]. Qiao and co-workers reported a 3D nitrogen-doped graphene/NiCo hydroxide composite catalyst using 3D graphene hydrogel as the substrate followed by nitrogen doping and NiCo hydroxide deposition (Fig. 1.4c) [108]. This novel 3D composite exhibited excellent electrical conductivity, hierarchical macroporous/mesoporous structure, abundant nitrogen doping, and strong interfacial coupling, thus leading to significantly improved ORR activity which was even three times the activity of each component. In addition to the hybridization in the pores of 3D graphene materials, the abundant topological defects are also favorable for efficient construction of new kinds of active sites. For example, Yao and co-workers synthesized single-atom nickel sites on graphene by annealing defective graphene and nickel nitrate at 750 °C and acid treatment [109]. Combining scanning transmission electron microscopy (STEM), X-ray absorption spectrum (XAS) analysis and theoretical calculations, the authors revealed that the nickel single atoms were coordinated with the topological defects of graphene and exhibited outstanding catalytic activity. The nickel single atoms were coordinated with “5-8-5” defects provided active sites for OER, while those coordinated with vacancy defects provided active sites for HER [109].

1.3 Controllable Synthesis of Three-Dimensional Graphene Materials The key to the controllable preparation of 3D graphene materials is the growth of local high-quality 2D graphene sheets, the construction of long-range 3D pore structures, and the seamless low-resistance interfacial connection. At present, the most common and efficient methods for synthesizing 2D graphene materials are

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mainly CVD on metal foils and oxidation–reduction strategy. Therefore, on this basis, CVD on 3D porous templates and pore-creating of GO assemblies are also the dominant strategies for the synthesis of 3D graphene materials. In addition, many templates have bifunctionality for both pore-forming and surface catalysis, thus leading to a new method via templated polymerization.

1.3.1 Porous Metal Templated Chemical Vapor Deposition CVD is a method developed in the 1960s to prepare high-purity, high-performance solid materials. In recent years, it has been widely used in the preparation of nanomaterials such as carbon nanotubes and graphene. The growth of graphene by CVD mainly involves carbon source, growth substrate, and growth conditions (e.g., gas, pressure, temperature, flow rate). According to the carbon solubility of different metals, the growth mechanism on flat metal surfaces is classified into surface catalytic nucleation (low carbon solubility, such as copper) and bulk dissolution-precipitation (high carbon solubility, such as nickel). The specific process is illustrated in Fig. 1.5, including the adsorption and decomposition of the carbon precursors, the dissolution and migration of carbon atoms, the precipitation during the cooling process, nucleation, and 2D reconstruction [13]. In 2011, Cheng et al. used 3D porous nickel foam instead of 2D copper foil or nickel foil as the template for CVD, and used methane as the carbon source to grow graphene at 1000 °C [116]. After removing the template under the support of polymethyl methacrylate, the resulting ultra-thin graphene can still maintain the foamlike 3D interconnected framework. As shown in Fig. 1.6a, this 3D graphene foam completely replicates the structure of the nickel foam template, exhibiting a fully

Fig. 1.5 Schematic diagram of the elementary steps for CVD growth of graphene. Reprinted from ref. [13], copyright 2018, with permission from American Chemical Society

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Fig. 1.6 a Scanning electron microscopy (SEM) and TEM images of obtained 3D graphene foam using nickel foam as the template. Reprinted from ref. [116], copyright 2011, with permission from Springer Nature. b Schematic of the synthesis of 3D porous metal template by powder metallurgy and then 3D graphene by templated CVD. Reprinted from ref. [130], copyright 2015, with permission from American Chemical Society

connected and self-supported monolithic structure constructed by ultra-thin graphene layers with a thickness of 1~3 layers. The obtained graphene material exhibits outstanding electrical conductivity, large specific surface area (about 850 m2 g−1 ), ultra-high porosity (about 99.7%), and extremely low density (about 5 mg cm−3 ). This unique 3D network structure fully demonstrates the intrinsically high electrical conductivity and mechanical flexibility of graphene in three dimensions and macro scales. The graphene foam/silicone rubber composite prepared by in situ polymerization, with a low mass content of graphene at 0.5 wt%, can achieve a high electrical conductivity of 10 S cm−1 which only changes slightly in the bending or stretching state. This novel structure provides a promising candidate for the development of flexible, foldable and stretchable devices.

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This innovative and pioneering work has brought great inspiration to the design and synthesis of 3D graphene materials, which was soon promoted to other templates and widely used in energy storage [71, 117–121] catalysis [122, 123], sensing [124, 125], electromagnetic shielding [126], and other fields. However, it is notable that the skeleton pores of commercial nickel foams are very large, resulting in 3D graphene foam with pores as large as 200 μm. Such skeleton is mechanically unstable which requires the use of flexible polymer support, and the bulk density is too low to meet the requirements for practical applications. Instead of commercial nickel foams, we can also prepare a family of monolithic porous templates by melting and annealing metal nanoparticles [127, 128] or metal hydrochloride [129] as precursors under high temperature. When used as the template for CVD, 3D graphene materials can also be obtained with significantly reduced pore size less than several microns. However, the pore structure is irregular and uncontrollable, and the overall scaffold is not stable enough with many interface defects. To combine the advantages of macroscopic metal foam and metal nanoparticles, Tour et al. prepared a self-supported 3D graphene foam by powder metallurgy templated method [130]. As shown in Fig. 1.6b, the metallic nickel powders (with a particle size of 2–3 μm) and solid-phase carbon source (sucrose) were uniformly dispersed in deionized water, and heated and stirred to uniformly coat the carbon source on the nickel surface. The dried mixture was then ground into powder and pressed into pellets by cold pressing at 1120 MPa for 5 min. 3D graphene foam will be formed during annealing at 1000 °C for 30 min under the protection of Ar/H2 atmosphere. Compared with the graphene material grown on nickel foams, this material exhibited a higher specific surface area (1080 m2 g−1 ), a better electrical conductivity (13.8 S cm−1 ), and a mechanically robust structure. By changing the composition of metal powders, pressing conditions, calcination temperature, and pore forming agents, a series of 3D porous templates and graphene materials can be prepared [131, 132]. For example, foam-like porous copper templates can be prepared by calcination at 800 °C for 45 min using metallic copper powders (with a particle size of 0.5–1.5 μm) as the precursor and MgCO3 powder as the pore-forming agent. Owing to the low carbon solubility of copper, the obtained 3D graphene material was dominantly formed by monolayer graphene, exhibiting a high electrochemically active surface area of 2500 m2 g−1 , which is almost close to the theoretical limitation of monolayer graphene [131]. Another kind of nanoporous metal foam can be fabricated by etching one metal from the bimetallic alloy precursors. The 3D graphene materials prepared by using this kind of templates will exhibit more regular pore structures, much smaller pore size, and 3D interconnected scaffold, which can further optimize the properties and performances of 3D graphene materials [133]. As shown in Fig. 1.7, Chen et al. used Ni30 Mn70 alloy foil with a thickness of 50 μm as the precursor. After the etching of Mn in 1.0 M (NH4 )2 SO4 aqueous solution, a 3D nanoporous nickel foam can be achieved with a regular pore size of 10–20 nm [134, 135]. Using it as the template and using benzene (or pyridine) as the carbon source, graphene or nitrogen-doped graphene with a 3D bicontinuous nanoporous structure can be prepared by CVD in the temperature range of 800–1000 °C. The obtained graphene exhibited a specific

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Fig. 1.7 Schematic of the synthesis of 3D graphene material using nanoporous metal foam as the template. Reprinted from ref. [135], copyright 2014, with permission from John Wiley and Sons

surface area of ~1000 m2 g−1 , and a pore size ranging from 100 nm to 2 μm, which can be facially adjusted by changing the reaction temperature and duration. Compared with the graphene foam prepared from macroporous nickel foam, the nanoporous nickel foam-derived material exhibited not only a much smaller bicontinuous pore size, but also a more robust and stable graphene framework. The reconstruction of templates during annealing at high temperature led to a periodic minimal surface, and thus retained the massless Dirac fermion properties of 2D graphene [136, 137]. Besides, Zhao et al. prepared 3D nanoporous copper foam by etching Cu40 Mn60 alloy in 0.05 M HCl solution, and then fabricated similar 3D bicontinuous porous graphene via CVD at 900 °C using acetylene as the carbon source [72, 138]. Since the metal substrate can catalyze the growth of high-quality graphene, the 3D graphene material prepared using porous metal as the template often possess

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high crystallinity and conductivity. The graphene thickness can also be adjusted by changing the reaction conditions and the metal composition. In addition, the morphology and structure of the porous metal template can be regulated by various preparation methods, thus rendering abundant opportunities to tune the porosities in obtained foam-like graphene materials. However, it is very challenging to achieve the replication growth of 3D graphene materials with pore sizes below 100 nm due to the coalescence of metal nanoparticles at high temperatures. Besides, this method is of high cost and low yield, which limit the scale up and practical applications.

1.3.2 Porous Metal Oxide Templated Chemical Vapor Deposition Analogous to carbon nanotubes, the prerequisite for the growth of graphene by CVD is the efficient adsorption and pyrolysis of carbon precursors. Therefore, the metal substrate is always regarded as indispensable for the preparation of carbon nanotubes and graphene due to its catalytic activity for pyrolysis. However, Rümmeli et al. found that some substrates that seem to have no catalytic activity, such as porous Al2 O3 , can indeed catalyze the growth of carbon nanotubes [139]. He systematically studied the graphitization behavior on a series of metal oxides (e.g., SiO2 , Al2 O3 , MgO, Ga2 O3 , ZrO). The results showed that graphene can be formed by CVD on the surface of metal oxide nanoparticles, such as MgO, and graphene followed the principle of self-limiting growth no matter how the reaction conditions changed [139, 140]. Nevertheless, there is no report of the growth of graphene or other carbon materials on the surface of single crystal oxide substrates. It suggests that the defect sites on the surface of nano-sized oxide particles may provide the essential catalytic activity for the cracking of the carbon source, and the nucleation and growth of graphene, which is verified by theoretical studies [141]. Therefore, metal oxide nanoparticles are expected to be used as a new kind of templates for the growth of 3D graphene materials, which may achieve unique nanostructures different from the metal templates. Because of its simple synthesis, variable structure, and moderate activity, MgO is currently one of the most widely used oxide templates for preparing 3D graphene materials [142, 143]. As shown in Fig. 1.8a, Wei et al. used the mesoporous MgO nanosheets obtained by calcining Mg(OH)2 nanosheets as the template. A meshlike graphene framework can be fabricated via CVD using methane as the carbon precursor, with a graphene yield of 30–50 mg per gram of MgO templates [144]. The pore size of obtained 3D porous graphene material was smaller than 10 nm, and the pore volume reached 2.35 cm3 g−1 with a high specific surface area of 1654 m2 g−1 . It indicated that the graphene thickness deposited on the surface of MgO was within two layers, verifying the high activity and self-limiting growth characteristics of MgO.

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Fig. 1.8 a TEM image (left) and optical image (right) of the obtained 3D porous graphene grown on MgO templates. Reprinted from ref. [144], copyright 2011, with permission from The Royal Society of Chemistry. b Schematic of the hierarchical structure of 3D porous graphene corresponding to the MgO templates. Reprinted from ref. [145], copyright 2015, with permission from Elsevier

Hu et al. used basic magnesium carbonate (4MgCO3 ·Mg[OH]2 ·5H2 O) as the template precursor, which will in situ decompose into MgO nanoparticles during CVD, and the benzene precursor will transform into 3D porous graphene nanocages [146]. The degree of graphitization, pore size, and specific surface area of obtained graphene will change accordingly as the reaction temperature increased. The average pore size was 7–15, 7–15, 10–25, 20–30 nm, and the specific surface area was 2053, 1854, 1633, and 312 m2 g−1 for samples prepared at 670, 700, 800, and 900 °C, respectively. In addition, the chemical composition and assembly structure of the obtained graphene can be adjusted by changing the gas composition and template structure (Fig. 1.8b) [37, 145, 147, 148, 149–151], showing great advantages of MgO in controllable synthesis of 3D graphene materials. Other kinds of porous mono-metal oxides, such as porous CaO obtained by calcination of CaCO3 [152, 153], porous anode Al2 O3 [154, 155], and ZnO tetrapod [156], can also catalyze the growth of

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3D graphene materials, but the resulting graphene materials exhibit relatively larger thickness and more defects. Layered double hydroxides (LDHs) is a family of 2D layered nanomaterials with a brucite-like structure, usually showing a regular micron-level hexagonal sheet morphology. The partial substitution of Mg2+ by metal ions with +3 valance (e.g., Al3+ , Fe3+ ) makes the laminates positively charged, which balances with the intercalated anions. During the calcination process, LDHs gradually undergo the removal of interlayer water molecules, anions, and laminate OH− , ultimately evolving into layered double oxides (LDOs). The structure transformation of hydroxides and the Kirkendall effect of bimetallic systems lead to abundant mesopores uniformly distributed in LDOs, which provide an ideal template for the growth of 3D mesoporous graphene materials. Wei et al. fabricated unstacked double-layer graphene using MgAl LDHs as the template precursor and methane as the carbon source at 950 °C [157]. As shown in Fig. 1.9, the MgAl LDHs transformed into mesoporous MgAl LDOs during annealing, and monolayer graphene was deposited on both sides. The as-obtained graphene was composed of two unstacked graphene layers separated by a large amount of meso-sized protuberances. The pillared graphene material exhibited a specific surface area of up to 1628 m2 g−1 , a pore volume of 2.0 cm3 g−1 , abundant mesoporous with a size of 2–7 nm, and a high electrical conductivity of up to 438 S cm−1 . This material has outstanding advantages in electrochemical applications, owing to the efficient conductive networks and smooth ion channels. Furthermore, by controlling the size and assembly of LDHs [46, 158], and regulating the pore-forming behaviors of LDHs (e.g., the introduction of volatile metal Zn [48] or catalytically active metal Fe [159]), the pore structure of 3D graphene materials obtained by CVD on MgAl LDO templates can be finely tuned.

Fig. 1.9 a Schematic of the synthesis of unstacked double-layer graphene on MgAl LDHs. b TEM images of unstacked double-layer graphene. Reprinted from ref. [157], copyright 2014, with permission from Springer Nature

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Some natural composite oxides, such as vermiculite [70] and diatomite [160], can also be used directly as templates for the deposition of 3D porous graphene. Specifically, as a very common and easily available raw material, NaCl nanocrystals can also exhibit excellent graphitization activity and template effect similar to metal oxides [161, 162]. But the obtained material has relatively poor crystallinity and thicker graphene layers. Compared with porous metal templates, metal oxide templates have huge advantages in the structure and morphology regulation of 3D graphene materials, especially the mesoporous structure. Besides, the purification process is relatively easier, and there is no metal residue, which is essential for highend electrochemical energy storage applications. Furthermore, the CVD process using powdered metal oxides can be industrially scaled up by the fluidized bed technology [44], and the template can be facilely recycled, suggesting a green, low-cost, and scalable strategy for the production of 3D graphene materials. However, the current understanding of the growth mechanism of graphene on metal oxide surfaces is still very preliminary, especially the influencing factors on graphene thickness. It largely restricts the design of optimal oxide templates and controllable synthesis of high-quality 3D graphene materials.

1.3.3 Post-Treatment of Graphene Oxides It is the most mature technical route for the preparation of graphene materials in industry by oxidation–reduction process using graphite microplatelets as the precursor. The obtained reduced graphene oxide (rGO) shows the morphology of large-sized 2D flakes in microscopic view, and processes abundant surface functional groups and defects. Therefore, it is well dispersed in liquid, and suitable for post-treatments, such as spinning coating, vacuum filtration, hydrothermal reaction, templated growth, layer-by-layer self-assembly, and activation and pore formation, leading to a variety of macroscopic assemblies fosr 3D porous graphene or GO [34–163]. Due to the hydrophilicity and surface functional groups of GO nanosheets, as well as the strong Van der Waals force generated after reduction, (r)GO hydrogels can be fabricated by reduction-derived self-assembly [164, 165] or crosslinkingassisted polymerization [166] during hydrothermal processes. For the first time, Shi et al. assembled 2D GO nanosheets into a 3D porous hydrogel through a simple hydrothermal treatment [51]. The water content was as high as 97.4%, while the graphene hydrogel can stably carry a weight of 100 g. As shown in Fig. 1.10, the rGO nanosheets were interconnected to form a skeleton, which supported micron-scale cavities and a 3D mechanically stable porous structure. By optimizing the drying method of hydrogels, the pore size and density of obtained aerogels can be adjusted to enhance the application performance. For example, Zhou et al. employed freezedrying to remove the water [167]. When deceasing the freeze-drying temperature, the pore size and wall thickness of the obtained 3D graphene material were greatly reduced by a variation of 80 times (from 800 μm to 10 μm) and 4000 times (from

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Fig. 1.10 a Optical images of GO before and after hydrothermal treatment. b Optical image of GO hydrogel. c SEM image of GO hydrogel. Reprinted from ref. [51], copyright 2010, with permission from American Chemical Society

Fig. 1.11 a Schematic illustration of the synthesis procedure of 3D nanoporous graphene foams. b Electron tomography reconstructed pore structure. Reprinted from ref. [171], copyright 2012, with permission from John Wiley and Sons

80 μm to 20 nm), respectively, while the Young’s modulus increased 15 times (from 13.7 kPa to 204.4 kPa). Yang et al. used evaporative drying instead of freeze-drying [168], during which the evaporation of water pulled graphene nanosheets together and induced the structure shrink. The obtained 3D graphene aerogel achieved a high density of 1.58 g cm−3 and large specific surface area of 720 m2 g−1 . Due to the large size of GO nanosheets, the pore size of 3D macroscopic assemblies obtained by the liquid phase self-assembly methods always reaches the micron level, which is expected to be regulated if various templates are used during self-assembly. For example, researchers employed a series of soft and hard templates to regulate the interface assembly of GO nanosheets, including activated alumina [169], lignocellulosic paper [73], polystyrene microspheres [170], silica nanospheres [69, 171], ice crystal particles [172], hexane droplets [173], and emulsion [174]. The pore size of resulting materials can vary from 20 nm to 2 μm. As shown in Fig. 1.11a, Zhao et al. mixed hollow silica microspheres with an average size of 28 nm and GO nanosheets in

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the liquid phase, which self-assemble to form a layered structure due to hydrophobic surfaces [171]. Silica microspheres were wrapped between GO nanosheets, and the subsequent removal of silica templates led to a 3D nanoporous graphene foam. In contrast to the liquid self-assembly, the pore structure in this material comes from the replication of hollow silicon microspheres, rather than the spontaneous construction of GO nanosheets. Therefore, the pore size can be precisely determined by the silica microspheres (Fig. 1.11b), resulting in an average pore diameter of 32.5 nm, a graphene wall thickness of 0.89 nm, a specific surface area of 851 m2 g−1 , and a pore volume of 4.3 cm3 g−1 . Additionally, the gas generated in confined space during the reduction of hierarchical GO assemblies can serve as a novel soft template for constructing 3D porous graphene materials [175]. In addition to constructing a 3D porous structure directly using GO nanosheets as the skeleton, we can introduce a new porous carbon layer on the surface of GO [176–94] or etch the GO nanosheets for in-plane holes [53, 178], thus enabling the construction of smaller-sized pore structures based on GO gels. On the one hand, templated pyrolysis (e.g., based on silica microspheres [179], molecular sieve SBA15 [180], and block copolymer micelles [181]) on both sides of GO nanosheets can form a 3D sandwich-like hierarchical porous structure. As shown in Fig. 1.12, Huang et al. prepared polypyrrole shells on both sides of GO nanosheets using silica microspheres as the template and pyrrole as the precursor, followed by carbonization at 850 °C and template removal to obtain a sandwich-like porous graphene material [179]. The thickness of a single sheet of porous graphene was between 10 and 25 nm, and the specific surface area was as high as 1558 m2 g−1 (1058 m2 g−1 contributed

Fig. 1.12 a Schematic illustration of the synthesis procedure and b SEM image of 3D sandwichlike hierarchical porous graphene. Reprinted from ref. [179], copyright 2014, with permission from John Wiley and Sons

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to micropores). On the other hand, etching the GO nanosheets with strong corrosive reagents such as KOH and H2 O2 can greatly increase the specific surface area and pore structure of the obtained graphene materials. For example, Ruoff et al. used KOH to etch delaminated GO nanosheets [53], and the obtained pore size was dominantly between 0.6 and 5 nm, rendering a high specific surface area of 3100 m2 g−1 and a large pore volume of 2.14 cm3 g−1 . Electron paramagnetic resonance results revealed a very low content of hydrogen and oxygen in such a graphene material with rich micropores and mesopores, indicating that the sp2 hybridized matrix of graphene was well retained with few edge sites. Therefore, the electrical conductivity of the obtained porous graphene material was still excellent (500 S m−1 ). Furthermore, by integrating self-assembly, surface polymerization, and etching, 3D hierarchical porous graphene materials with abundant “micropores-mesoporesmacropores” can be facilely constructed based on GO nanosheets [40, 178, 182]. The post-treatment of GO precursors provides a very convenient and efficient strategy for the design and fabrication of 3D graphene materials, which is also very versatile for heteroatom doping and hybridization during the synthesis process. However, it is notable that this method cannot precisely control the structural features of graphene, and the generated defect sites are too many to be fully recovered, which limit the properties and performances.

1.3.4 Templated Polymerization For the above two completely different synthesis strategies, the porous template is the key to the successful construction of 3D graphene materials. For the post-treatment of GOs, the template always only plays the role of physical space-occupying, which prefers easily obtained nanoparticles and microspheres. For the templated CVD, the template not only provides the expected porous scaffold for replication, but also offers the catalytic activity for carbon source cracking and graphene growth, therefore mainly requiring metal and metal oxides. In fact, these metal and metal oxide templates can also catalyze the cracking and polymerization of some complex, solid carbon-containing precursors, such as NiO template for phenolic resin [183], FeO(OH) nanorods for fluoroaniline [184], and MgO nanosheets for pitch [185], aromatic polyimides [150], sucrose [186], and starch [187], leading to few-layered, porous, and conductive 3D graphene materials. For example, Kaskel et al. fabricated 3D graphene materials using ZnO nanoparticles as the template and sucrose as carbon precursors, which exhibited a specific surface area as high as 3060 m2 g−1 and a pore volume of 3.45 cm3 g−1 (Fig. 1.13a, b) [188]. However, it should be noted that it is difficult for the solid-phase carbon source to fully penetrate into the porous framework of the template. Therefore, the template polymerization can always replicate the outer morphology of the template, but can hardly achieve the internal porous structure as the CVD method (Fig. 1.13c) [186]. In addition, because the template is mixed with the carbon source in advance, the catalytic cracking and polymerization can start at the same time from multiple

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Fig. 1.13 TEM images of 3D graphene materials obtained from sucrose using different templates. a, b ZnO nanoparticles. Reprinted from ref. [188], copyright 2015, with permission from John Wiley and Sons. c MgO nanoparticles. Reprinted from ref. [186], copyright 2017, with permission from The Royal Society of Chemistry

sites, resulting in less-controllable microporous and mesoporous structures and more defects. Monodispersed nanocrystalline particles can construct a superlattice structure through self-assembly, which has a very regular and controllable 3D structure. It is expected to prepare periodic 3D porous graphene materials using such superlattice templates. For example, Lee et al. prepared 3D porous graphene materials via template polymerization at 1000 °C using superlattice SiO2 nanospheres as the template and polyvinyl alcohol/FeCl3 as the carbon/catalyst precursors [171]. The achieved pore size was about 200 nm with a specific surface area of 1025 m2 g−1 and electrical conductivity of 52 S cm−1 . Dong et al. used superlattice Fe3 O4 nanocrystalline as the template and adsorbed oleic acid molecules as carbon precursors, which were first polymerized at 500 °C to form an ordered mesoporous carbon skeleton and then graphitized at 1000 °C to obtain a 3D cross-linked, highly ordered mesoporous graphene materials [189, 190]. The specific surface area and pore volume of obtained graphene materials reached 1500 m2 g−1 and 2.5 cm3 g−1 , respectively. The pores of about 10 nm originated from the occupation of Fe3 O4 nanocrystals, and the mesopores of 2–4 nm correspond to the skeletal gap of the superlattice structure. The graphene thickness can be adjusted between 2 and 6 layers by changing the chain length of the hydrocarbons adsorbed on the surface of the nanocrystals [191], and the macroscopic morphology (spheres or thin films) can be tuned by changing the assembly structure of superlattice templates [192, 193].

1.4 Research Approach and Methodology

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1.4 Research Approach and Methodology In spite of great advances in the preparation of 3D graphene materials, the improvement in their structure and properties, and excellent performance in the fields of energy storage and catalysis, the research is still by trial-and-error due to the lack of general and effective principles for materials design, fabrication, and application. Based on the existing preparation strategies, the characteristics of the materials obtained, and the target application fields, this Thesis aims to establish the goals, principles, and methodologies for constructing optimal 3D graphene materials. The goal of constructing an ideal 3D graphene material mainly involves four aspects: (1) 3D conductivity and defect distribution. By building a 3D interconnected sp2 -hybridized carbon skeleton, the point defects can be optimized with a lower concentration and uniform distribution, ensuring high conductivity in three dimensions. (2) Facile control of graphene thickness and specific surface area. The number of graphene layers has a huge impact on the specific surface area, which determines its performance in applications such as electrochemical energy storage. (3) Hierarchical porosity and mechanical strength. Micropores, mesopores and macropores play different roles in different applications. The matching of multi-level pores can maximize the mass transport and interface reactions, while ensuring the mechanical stability of the 3D structure. (4) Single crystal domain and macro size. The size of the single crystal domain determines the transfer properties of the graphene material, and the size of the structural units determines the abundance of interface. To achieve the above objectives, the construction principles of 3D graphene materials specifically include: a 3D covalently connected monolayer low-defect graphene, a self-supported structure with abundant mesopores, and 3D assembly with as few interfaces as possible, thus rendering the full maintenance and demonstration of the intrinsic physical properties of graphene in 3D macroscopic materials. In order to achieve the fine regulation and controllable preparation of the ideal 3D graphene materials, this Thesis employed porous metal oxides as the growth template, and conducted research in order from four levels of the 3D graphene material’s growth mechanism, synthetic methodology, custom application and system promotion (Fig. 1.14). First of all, we employed advanced characterization methods such as SEM, spherical aberration correction electron microscopy, and 3D tomography reconstruction, as well as knowledge of surface catalysis and reaction processes, to fully understand the growth mechanism and construction principles of 3D graphene materials on the surface of metal oxides. The acquired knowledge will serve as guides to optimize the preparation and regulation strategy of 3D graphene based on metal oxides (Chapter 2). Secondly, in view of the difficulty of controlling the structure, composition and performance of 3D graphene materials, we developed the versatile templated polymerization strategy based on metal oxides for efficient synthesis and modification.

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Fig. 1.14 Conceptual framework of the research

The obtained materials were used as excellent platforms to investigate the application performance and structure-activity relationship of 3D graphene materials in the application of energy electrocatalysis (Chapter 3). Furthermore, in order to fully exploit and make good use of the advantages of 3D porous graphene materials, we proposed a new concept of “nanoreactor” to efficiently prepare heterogeneous electrocatalysts. The 3D porous graphene based hybrid nanomaterials provided not only high-efficient catalysts for water splitting, but also an ideal platform for the study of intrinsic activity mechanisms (Chapter 4). Finally, considering the wide variety of 2D graphene-analogous materials with different properties, the concept of 3D nanostructure engineering is extended to other material systems, such as MoS2 . A 3D mesoporous van der Waals heterostructure with excellent trifunctional electrocatalytic activities was successfully synthesized (Chapter 5). The research outcomes are expected to deepen the understanding of the structureactivity relationship of graphene materials and the growth mechanism of 3D graphene materials, provide new ideas and guidance for the design, preparation and application of high-performance 3D graphene materials, and make inspiring explorations for the development of graphene and other 2D nanomaterials. It will have a significant impact in many fields such as nanomaterials and energy catalysis.

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Chapter 2

Growth Mechanism of 3D Graphene Materials Based on Chemical Vapor Deposition

2.1 Introduction Chemical vapor deposition (CVD) is one of the most effective means to prepare high-quality graphene materials. By changing the energy supply method, the composition and morphology of the substrate, the composition and flow rate of the gas, and the temperature and pressure of the reaction, we can effectively adjust the number of graphene layers, the crystal domain size, stacking form, defect density, and heteroatom doping, which can tune the intrinsic physical properties of obtained graphene materials for various application requirements. As the same “bottom-up” strategy to construct sp2 -hybridized graphene structural units, metal foil-templated CVD usually produces large-scale 2D graphene films or graphene single crystals, while porous oxide-templated CVD can obtain 3D replicated porous graphene materials. A lot of research work has shown that porous MgO powder is one of the most unique templates. On the one hand, we can fabricate various MgO templates with different mesoscopic structures based on different synthesis pathways, enabling the synthesis of 3D mesoporous graphene framework materials [1], graphene nanocages [2, 3], porous graphene fibers [4], and more. On the other hand, the surface of MgO has moderate alkalinity, adsorption characteristics and catalytic activity. These unique properties lead to graphene materials with ultra-thin thickness, huge specific surface areas, high graphitization degree, and excellent conductivity, which fully realize the excellent intrinsic physical properties of 2D graphene in the 3D structure. In addition, unlike metal foils, the MgO powder-templated CVD process powder can be facilely scaled up via the fluidized bed reactor technology, thus creating the possibility for the large-scale production and industrial application of high-quality 3D porous graphene materials [5]. Although the porous oxide-templated CVD method has many advantages in the synthesis of 3D graphene materials, we still know little about the growth mechanism of graphene on metal oxide surface. Why is the number of graphene layers deposited on the surface of MgO much smaller compared to those on other oxides? Why can © Tsinghua University Press 2021 C. Tang, Construction Principles and Controllable Fabrication of 3D Graphene Materials, Springer Theses, https://doi.org/10.1007/978-981-16-0356-3_2

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the number of graphene layers and specific surface area of the graphene prepared using different MgO templates vary significantly? How to elucidate the construction principles and mechanical stability of 3D porous graphene materials? How to effectively control the layer thickness and pore structure of 3D graphene materials? Are there other efficient oxide templates for the synthesis of 3D graphene materials? These problems make relevant research still dominated by the trial-and-error methodology, which greatly limits the controllable construction and synthesis of 3D graphene materials. In order to fully understand the growth mechanism and construction principles of 3D graphene materials prepared by metal oxide-templated CVD, this Chapter starts with the atomic characterization of nano-sized mesoporous structure and lattices of mesoporous quasi-single-crystal MgO using double spherical aberration correction TEM. The results will help understand the key elements for self-limiting nucleation and growth of graphene on the surface of MgO. Based on the acquired knowledge, we rationally deduced a new kind of metal oxide templates, porous CaO, and realized the construction and regulation of 3D graphene materials with respect to the number of graphene layers, specific surface area, and porosity. The results have deepened and improved the understanding of the growth mechanism of graphene on oxide surfaces, and provided important guidance for the rational design of porous templates for depositing 3D graphene materials and the efficient regulation of resulting structure and physical properties.

2.2 Templated CVD Based on Mesoporous MgO 2.2.1 Morphology Replication Based on Mesoporous MgO Templates In the oxide-templated CVD process, the gaseous carbon source catalytically cracks on the surface of the oxide, and then nucleates and grows at specific active sites to form a continuous graphene layer. The graphene completely replicates the 3D porous structure of oxide templates to form a 3D porous graphene scaffold. Therefore, the nanostructure of the obtained 3D graphene material dominantly depends on the structure of oxide templates, and the number of layers and crystallinity of the graphene are determined by the surface chemistry of oxide templates. In order to achieve an ideal oxide template, our group proposed the use of polyethylene glycol as an inducer to adjust the nucleation and growth behavior of MgO in the high-temperature and high-pressure hydrothermal process to obtain a micron-sized Mg(OH)2 precursor, followed by calcination in an inert atmosphere to obtain mesoporous quasi-singlecrystal MgO nanosheets. Using this porous MgO as the template and methane as the carbon source, 3D mesoporous graphene framework materials were obtained by CVD at 950 °C [6].

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As shown in Fig. 2.1a, the mesoporous quasi-single-crystal MgO template presents a hexagonal morphology with an average size of 1 μm and a thickness of 50–100 nm. No obvious through-holes can be observed on the surface. The TEM image shows that a large number of mesopores of 5–10 nm are uniformly distributed in the hexagonal MgO nanosheets, which are arranged regularly and interconnected (Fig. 2.1b). As shown in Fig. 2.1c, the 3D mesoporous graphene framework material prepared by CVD method still maintains the hexagonal morphology and size of MgO after removing the template. More importantly, on the microscopic level, the 3D mesoporous graphene framework material also presents a highly porous structure with an average pore size of ~10 nm (Fig. 2.1d), indicating that the graphene has perfectly replicated the mesoporous structure of quasi-single-crystal MgO templates during the growth process. Furthermore, we compared the 3D topological structure of the MgO template and the obtained graphene material on the nanometer scale by high-resolution TEM technique. As shown in Fig. 2.2a, the high-resolution TEM image of MgO has contrastive brightness in different regions, showing an alternating regular arrangement of light and dark squares with a size of 5–10 nm (highlighted by the dashed box). This originates from the diffraction contrast of the sample, so the dark part corresponds to the

Fig. 2.1 a SEM and b TEM images of mesoporous quasi-single-crystal MgO nanosheets. c SEM and d TEM images of 3D mesoporous graphene framework. Reprinted from ref. [6], copyright 2015, with permission from John Wiley and Sons

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Fig. 2.2 High-resolution TEM images of a mesoporous quasi-single-crystal MgO nanosheets and b 3D mesoporous graphene framework. Reprinted from ref. [6], copyright 2015, with permission from John Wiley and Sons

structural unit of MgO, and the light part is the mesoporous space in the template. Therefore, the hexagonal prism structure of mesoporous quasi-single-crystal MgO is self-assembled by MgO single-crystal nanocubes with a size of 5–10 nm in a staggered arrangement, and the voids between the nanocubes constitute a 3D interconnected porous space. This unique structure is determined by the phase transition mechanism of Mg(OH)2 during high temperature calcination. The TEM image in Fig. 2.2b shows that the obtained graphene is dominated by a single layer; however, unlike common monolayer graphene sheets or rGO nanosheets, the graphene sheets in this material show a nanocage morphology with obvious bending and wrapping (highlighted by the dashed box). Additionally, these curved graphene sheets and the mesoporous space inside support each other and extend outward, which replicate the hexagonal structure of MgO templates in the mesoscopic view (Fig. 2.1c, d).

2.2.2 Nanostructure Characterization of 3D Mesoporous Graphene Framework In recent years, the material characterization technologies have been rapidly developed. For the TEM technique, the electromagnetic lens cannot be absolutely perfect. The converging ability near the edge is stronger than the center, resulting in the generation of spherical aberration. Besides, the electromagnetic lens have only convex lens and no concave lens. Therefore, the spherical aberration becomes the most important factor that limits the resolution of TEM images. By introducing a multi-pole correction device into the electronic optical system, the focus center of the electromagnetic lens can be effectively adjusted and controlled, so as to realize the correction of spherical aberration and achieve sub-angstrom resolution. In addition, the combination of electron microscopy, electron diffraction, and computer image processing technology has developed a 3D TEM tomography technology, which can intuitively

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Fig. 2.3 Spherical aberration-corrected TEM images of mesoporous quasi-single-crystal MgO nanosheets. Inset of (b) is the corresponding FFT result

and accurately characterize the structure of nanoporous materials in 3D space. The development and commercialization of these technologies provide us with opportunities to further characterize the nanostructures of mesoporous quasi-single-crystal MgO templates and 3D mesoporous graphene framework materials, thus creating the possibility for the understanding of the growth mechanism. First, we characterized the nanostructure of the mesoporous quasi-single-crystal MgO by a double spherical aberration-corrected TEM (FEI Titan Cubed Themis G2 300). The operating voltage is 300 kV and the limit resolution can reach 0.07 nm. Figure 2.3a shows the typical mesoporous structure of MgO nanosheets. On the outside of the dashed frame, we can see obvious lattice fringes corresponding to the MgO single crystal structural unit; while on the inner side, the poor contrast proves the disruption of the crystal structure and the presence of mesoporous holes. Notably, a certain amount of sample signal can still be observed in the region within the dashed frame, suggesting the existence of MgO crystal below the mesopore, which means that the mesopore is not a through hole, but exists in the bulk phase of MgO. The TEM image in Fig. 2.3b shows the MgO lattices around a mesopore. The lattice fringes in the entire region are very complete with a fringe spacing of 0.24 nm, which corresponds to the MgO (111) crystal plane. The Fast Fourier Transform (FFT) result of Fig. 2.3b reveals a set of very sharp and integrated diffraction spots, in consistence with MgO (111), suggesting the single crystal nature of the material. However, we can also find that there is still an obvious contrast distribution with an alternating regular arrangement of light and dark squares of about 1 nm along the same direction. It reveals the presence of periodic structural defects even in the crystalline region. The double spherical aberration-corrected TEM images not only verify the mesoporous quasi-single-crystal characteristics of the MgO template, but also reveal the microscopic crystal information in the MgO structural unit, which is very important for explaining the growth mechanism of 3D mesoporous graphene framework on it. Furthermore, we used the JEOL 2100F TEM with a high voltage of 200 kV, equipped a 3D tomography reconstruction sample holder to record TEM images

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Fig. 2.4 3D tomography reconstruction of mesoporous graphene framework materials. a, b TEM images recorded at different tilting angles. c, d 3D reconstructed structures viewed from different directions

at tilting angles from –65o to +65o . Subsequent series alignments were performed in the IMOD33 software using 5-nm gold nanoparticles as fiducial markers. The synchronous iteration (SIRT) algorithm was selected for 3D reconstruction. The 3D TEM tomography results are shown in Fig. 2.4. During the tilting process, the morphology of graphene changes continuously (Fig. 2.4a, b), while the mesoporous structure is always present with a relatively uniform pore size between 5 and 10 nm. The electron tomography shows the topological structure of the 3D graphene material more concretely and vividly (Fig. 2.4c, d), confirming that the monolayer graphene serves as the scaffold to build a self-supported 3D interconnected mesoporous structure.

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2.2.3 Growth Mechanism of 3D Graphene on the Surface of Mesoporous MgO With a combination of the characterization results from SEM, TEM, spherical aberration-corrected TEM, and 3D electron tomography, we have a very specific and in-depth understanding of the nanostructure of mesoporous quasi-single-crystal MgO templates and 3D mesoporous graphene framework materials. Based on this, the growth mechanism and key impact factors of 3D graphene materials on the surface of mesoporous MgO are analyzed. In addition to the moderate surface alkalinity, mesoporous interconnected structure, excellent mechanical stability, and high resistance to calcination, the intrinsic advantages for MgO being an ideal template for 3D graphene deposition lies in: regular mesoporous structure, abundant lattice defects, and exposure of O-terminated polar MgO (111) crystal plane. The first feature is the regular mesoporous structure and abundant lattice defects of MgO. During the dehydration process of Mg(OH)2 to form MgO at high temperature, the crystal structure changes from the hexagonal closest packed to the cubic closest packed. According to the inhomogeneous mechanism proposed by Ball et al., the dehydration reaction occurs simultaneously throughout the bulk of Mg(OH)2 , which then develops “donor” and “acceptor” regions [7]. The Mg2+ ions are postulated to migrate from the donor to the acceptor regions with a counter migration of protons generated by ionization in the acceptor region. The protons will combine with the excess OH– ions in the donor region to form water molecules, which then escape from the donor region to form lattice defects and intercrystalline pores that are uniformly distributed in the bulk phase. In addition, when the size of Mg(OH)2 is larger than 60 nm, nanopores of about 5–10 nm will be formed due to contractional strain-derived cracking [7]. Therefore, in order to obtain a regular mesoporous structure instead of large through-holes, we employed polyethylene glycol as a morphology inducer to increase the thickness of Mg(OH)2 precursor (> 60 nm). Due to the synergistic effect of lattice dehydration and fracture stress, the obtained MgO maintains the overall morphology of Mg(OH)2 and generates a large number of regularly arranged interconnected mesopores (5–10 nm, Fig. 2.2a), which offer the diffusion path of the carbon source and the replica template of the 3D graphene. Besides, the lattice defects and intercrystalline nanopores (~1 nm, Fig. 2.3b) provide abundant step and edge sites that are highly active for the adsorption and cracking of carbon sources, as well as the nucleation and growth of graphene [8]. The second feature is the exposure of O-terminated polar MgO (111) crystal plane. The distance between two adjacent oxygen atoms in the MgO (111) crystal plane is 2.9 Å, while the lattice constant of graphene is 2.5 Å, which do not match with each other. Studies have shown that after physical vapor deposition of fewlayered graphene on single-crystal MgO (111), the surface atoms of MgO (111) will be rearranged towards shortened O–O bond length by 10% compared with the bulk phase, thus enabling the match with the graphene lattice [9]. It is because that the MgO (111) crystal plane is a polar terminated surface (O2– /Mg2+ ), and the surface energy is unstable, creating the possibility for lattice rearrangement and graphene

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growth [10]. In other non-polar crystal planes, such as MgO (100) surface, atoms cannot undergo lattice rearrangement, so long-range graphene sheets cannot grow [10]. In addition, theoretical calculations revealed that when monolayer graphene is loaded on the polar MgO (111) surface, the interaction between graphene and O-terminated surface is much greater than that of Mg-terminated surface; when two layers of graphene are loaded on MgO (111), the bottom graphene layer forms strong chemical bonds with the O-terminated surface while shield the top graphene layer from the reactive O atoms [9, 11]. As a result, the exposure of O-terminated polar MgO (111) crystal planes is the key to ensure the stable deposition and self-limiting growth of monolayer long-range graphene.

2.3 Templated CVD Based on Porous CaO The catalytic substrate is the most critical factor that determines the growth behavior and structural properties of graphene during CVD process. For the preparation of 2D graphene films or single crystals, the metal substrates used for CVD include transition metals (e.g., copper, nickel), precious metals (e.g., iridium, platinum, gold), and liquid metals (e.g., gallium, indium). The catalytic growth behavior of graphene on different substrates is different, which provides a lot of opportunities for the regulation of resulting graphene structures. Analogously, the exploration of other porous oxide substrates beyond MgO is very important for the precise construction and structural control of 3D graphene materials. Our research work of mesoporous MgO-templated CVD reveals that the nanostructure of templates determines the porous scaffold of the resulting 3D graphene materials, and the crystal structure determines the growth behavior of graphene (e.g., carbon adsorption, cracking, nucleation, growth, stacking). After comparing various oxides, we find many similarities between CaO and MgO: alkaline earth metal oxides, face centered cubic packed with similar crystal structures, and versatile synthesis towards porous structures. Besides, surface chemistry of CaO and MgO are different due to a larger atomic radius and different lattice constant of Ca. Therefore, CaO is also supposed to be an excellent substrate for 3D graphene deposition, and provides possibility for the structure adjustment compared to that grown on MgO.

2.3.1 Synthesis and Characterization of Porous CaO To study the regulation principles and strategies of 3D graphene, we prepared two kinds of porous CaO with different structures as CVD templates. Similar to the synthesis strategy of mesoporous quasi-single-crystal MgO, we first synthesized the Ca(OH)2 precursor and prepared porous CaO by high-temperature calcination. Specifically, 59.0 g of Ca(NO3 )2 ·4H2 O was dissolved in 750.0 mL of deionized water, and then 300.0 mL of NaOH solution (2.5 mol L−1 ) was dropwise

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Fig. 2.5 SEM images of a Ca(OH)2 nanosheets and b annealing-driven porous CaO nanosheets. SEM images of c commercial CaCO3 nanoparticles and d annealing-driven porous CaO nanoparticles. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

added into the solution with continuous stirring, followed by aging at 100 °C for 24.0 h. The resultant white slurry was then filtered, washed, and freeze-dried to obtain the Ca(OH)2 precursor. As shown in Fig. 2.5a, its morphology is similar to the Mg(OH)2 precursor prepared by hydrothermal process, showing a hexagonal prism structure with an average size of 1.5 μm and a thickness of 100–150 nm. The Ca(OH)2 precursor was then calcined at 650 °C in air for 4.0 h to obtain the CaO template. The latter basically maintains the size and morphology of the precursor (Fig. 2.5b), while has a large amount of large pores and through-holes generated (50–100 nm). The shape and distribution of the holes are irregular, which is in sharp contrast with the mesoporous quasi-single-crystal MgO. This may be ascribed to the different phase transformation process of two kinds of hydroxides. The water loss process of Mg(OH)2 is relatively controllable and moderate, while the phase transformation of Ca(OH)2 is more turbulent and uneven, resulting in significant difference in crystallinity and pore structure of obtained metal oxide templates. In addition, we directly used commercial CaCO3 reagent (purchased from Sinopharm Chemical Reagent Co., Ltd.) as the precursor, and calcined at 950 °C for 1.0 h to decompose into porous CaO particles. As shown in Fig. 2.5c, CaCO3 powder is an irregular solid particle with a size of about 20 μm. It decomposes at

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Fig. 2.6 XRD profiles of Ca(OH)2 nanosheets, commercial CaCO3 nanoparticles, and corresponding annealing-driven porous CaO. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons.

high temperature to produce CaO and CO2 , thereby generating uniform voids in the entire particle phase and forming a 3D labyrinth-like bicontinuous porous framework. The width of the voids and the CaO bones is between 500 nm and 1 μm, which is much larger than the pore size of CaO calcined from Ca(OH)2 nanosheets. X-ray diffraction (XRD) profiles verify that the precursors of Ca(OH)2 and CaCO3 are both completely converted into CaO after high temperature calcination (Fig. 2.6).

2.3.2 Templated Growth and Structure Regulation of 3D Porous Graphene The CaO templates prepared as above have different surface chemistry, crystal structures, and porous morphology compared with the mesoporous quasi-single-crystal MgO template, thus providing a material platform for a comprehensive understanding of the structure modulation and growth mechanism of 3D porous graphene materials. We first performed the CVD process using Ca(OH)2 -derived CaO as the template to verify the feasibility of graphene growth on CaO surface. The reaction conditions for CVD were the same as those for preparing 3D mesoporous graphene framework materials based on MgO templates. Specifically, the CaO powder was placed in the center of a quartz tube which was inserted into a horizontal tube furnace. Under the protection of Ar (200 mL min−1 ), the furnace was heated to 950 °C at a rate of 15 °C min−1 , and then CH4 (100 mL min−1 ) was introduced for the graphene deposition for 7.0 min. After that, the reactor was naturally cooled to room temperature. The obtained sample was then purified with hydrochloric acid, washed, filtered, and freeze-dried to achieve the expected 3D porous graphene material, which was denoted as Ca(OH)2 -G. SEM and TEM images confirm the successful growth of graphene with a 3D porous structure. As shown in Fig. 2.7a, the graphene nanosheets in Ca(OH)2 -G are

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Fig. 2.7 Structural characterization of 3D porous graphene material fabricated via CaO-templated CVD. The CaO template was obtained by annealing Ca(OH)2 nanosheets. a SEM image. b TEM image. c High-resolution TEM image. d Schematic illustration of the hierarchical porous structure. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

in a severely curved and wrinkled state, which are interconnected and assembled into a foam-like 3D porous structure. A large number of scaffold cavities, with an average size of 300 nm, are formed between the graphene nanosheets. They are much larger than the through-holes in the CaO templates, and much smaller than the gaps between the CaO nanosheets (Fig. 2.5b). Therefore, these graphene nanosheets and cavities are mainly derived from the replication of the surface morphology of CaO templates rather than the replication of the internal macropores; the porous structure has collapsed after the deposition of graphene and removal of templates. The porous graphene framework in Ca(OH)2 -G provides a 3D path for rapid electron transfer and interconnected macropores for smooth mass diffusion, which meet the electrode material requirements for various high-performance electrochemical energy storage devices. Besides, the graphene membrane itself is also wrinkled into a porous sheet with large pores ranging from 10 to 100 nm (Fig. 2.7b), which is ascribed to the replication of porous CaO templates. However, due to the large pore size and thin graphene, this porous structure is mechanically unstable and thus collapses obviously. The highresolution TEM image show that the thickness of obtained graphene nanosheets

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is within three layers, and there are many micropores and mesopores (Fig. 2.7c). Compared with the mesoporous quasi-single-crystal MgO, the Ca(OH)2 -derived CaO template has poorer crystallinity and structural regularity, which may lead to unique catalytic graphitization behaviors and generate more irregular pores and defects during catalytic cracking and graphene growth. Overall, few-layered graphene can also be grown on the surface of the CaO template, but the obtained graphene has more abundant defects and hierarchical porosity including micro-sized in-plane vacancies, meso-sized wrinkled pores, and macro-sized strutted cavities (Fig. 2.7d). Although a few reports have demonstrated the use of CaCO3 nanoparticles as templates to prepare porous carbon materials from solid carbon sources [13, 14], and the use of CaO nanoparticles as templates to prepare nitrogen-doped carbon materials via acetonitrile vapor deposition [15], this work realized the growth of few-layered graphene materials via CaO-templated CVD for the first time. On this basis, we continued to explore the modulation strategies and growth mechanisms of 3D porous graphene materials using CaO as the template. Using CaCO3 -derived porous CaO as the template, a 3D porous graphene material was prepared under the same reaction conditions, which is denoted as CaOC. As shown in Fig. 2.8a, the morphology of CaO-C is similar to Ca(OH)2 -G, showing a 3D porous framework structure, but the pore size and arrangement of the graphene nanosheets are more regular, mainly due to the difference of templates. The TEM image clearly shows the structural difference between CaO-C and Ca(OH)2 -G (Fig. 2.8b). The number of graphene layers in CaO-C is significantly increased, and the pores formed by wrinkles are much larger with a size around 200 nm, which is also smaller than that of the template (Fig. 2.5d), indicating a certain degree of structural collapse. In fact, the commercial CaCO3 nanoparticles can be decomposed into CaO and CO2 at about 825 °C, which is lower than the temperature for graphene deposition (950 °C). Therefore, we directly used CaCO3 nanoparticles as the template precursor, which will in situ convert into porous CaO templates during the heating and CVD process. After graphene growth and template removal, the 3D porous graphene material obtained is denoted as CaCO3 -C. As shown in Fig. 2.8c, CaCO3 C presents a completely different hierarchical structure from the other two samples, namely Ca(OH)2 -G and CaO-C. Compared with CaO-C, the CaCO3 -C replicates the topological structure of the porous CaO template obtained by the in situ decomposition of CaCO3 more perfectly, with respect to the size, shape, and arrangement of the strutted cavities (Fig. 2.5d). Despite that the size of the cavities reaches the micron level, the 3D porous scaffold of obtained graphene material remains intact after the removal of templates (Fig. 2.8c, d). In addition, we found that the inner wall of the quartz tube turned black immediately after the gaseous carbon source was introduced in the case of direct using CaCO3 , implying distinct graphitization behaviors. Figure 2.9 compares the 3D porous graphene prepared based on three different CaO templates. It can be seen that the graphene layers deposited using CaCO3 directly as the template are the thickest, up to nearly 10 layers; the graphene layers deposited on pre-annealed CaO templates from CaCO3 are relatively thinner, around 5 layers; the graphene layers grown on the surface of the Ca(OH)2 -derived CaO are

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Fig. 2.8 a SEM and b TEM images of 3D porous graphene fabricated via CaO-templated CVD. The CaO template was obtained by annealing commercial CaCO3 nanoparticles. c SEM and d TEM images of 3D porous graphene fabricated via CaCO3 -templated CVD. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

Fig. 2.9 Thickness comparison of graphene materials fabricated by different CaO templates. a CaCO3 -C. The template is commercial CaCO3 nanoparticles. b CaO-C. The template is porous CaO annealed from commercial CaCO3 nanoparticles. c Ca(OH)2 -G. The template is porous CaO annealed from Ca(OH)2 nanosheets. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

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Fig. 2.10 a Nitrogen adsorption-desorption isotherms and b pore size distribution of 3D porous graphene fabricated by different CaO templates. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

the thinnest, within 3 layers. The difference in graphene thickness indicates different catalytic activities and graphitization behaviors on different templates, and rationalizes the difference in porous structures. From a mechanical point of view, the thicker the graphene, the larger the cavity that can be stably supported, and the more regular the structure (such as CaCO3 -C). Instead, when the thickness of the graphene is not enough to balance its cavity size, the 3D structure will collapse as those of Ca(OH)2 -G and CaO-C. Due to the combined effect of template replication and stress matching, the obtained graphene materials thus exhibit different topological structures, which in turn affect the intrinsic physical properties and application performances. Generally speaking, the thinner the graphene layers, the larger the corresponding specific surface area. We quantitatively analyzed the specific surface area and pore structure of the obtained 3D porous graphene materials by nitrogen adsorption-desorption experiments at 77 K. Figure 2.10a shows the nitrogen adsorption-desorption isotherms for different samples. It reveals that the Ca(OH)2 -G with the thinnest graphene layers has the largest specific surface area (572 m2 g−1 ), but the specific surface area of CaCO3 -C with the thickest graphene layers (343 m2 g−1 ) is slightly higher than that of CaO-C (290 m2 g−1 ). Although the Ca(OH)2 -G sample exhibits the largest specific surface area and the highest amount of mesopores, it has the smallest pore volume (0.93 cm3 g−1 ), while the CaO-C sample with the smallest specific surface area has the largest pore volume (1.15 cm3 g−1 ). The opposite trend of specific surface area and pore volume indicates that the apparent physical properties of 3D porous graphene materials are determined by both the graphene thickness and porosity hierarchy. This is why the specific surface area and pore volume of graphene fabricated based on CaO templates, namely Ca(OH)2 -G, are much smaller than those based on MgO templates. Although the graphene thickness is similar (< 3 layers), the mesoporous structure of the latter is more regular and the crystallinity of graphene is higher, thus resulting in distinct properties.

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2.3.3 Growth Mechanism of 3D Graphene on the Surface of Porous CaO Based on the above characterization results, herein we will discuss the growth mechanism and construction principles of 3D graphene materials on porous CaO surface. The growth of graphene via CVD on specific substrates mainly involves four steps: (1) the adsorption of carbon sources on the surface of the substrate, (2) the catalytic cracking and dehydrogenation of the carbon sources, (3) the surface migration of carbon-containing species, and (4) the nucleation of graphene at surface active sites. In the case of conventional metal-catalyst-free substrates, there are few active sites for the adsorption, decomposition, and diffusion of carbon species, thereby leading to great challenges to grow large-area continuous graphene layers on most metal oxide substrates (e.g., SiO2 [16], SrTiO3 [17]). Owing to the regular mesoporous structure, abundant lattice defects, and exposure of O-terminated polar MgO (111) crystal plane, the 3D mesoporous quasi-single-crystal MgO template can provide a large number of active sites favorable for the adsorption and cracking of carbon sources, and the nucleation and growth of graphene with a self-limiting feature. Similarly, another kind of face centered cubic packed alkaline earth metal oxide, namely CaO, also exhibits moderate surface chemistry and self-limiting growth mechanism towards few-layered graphene. The catalytic activity of the CaO surface originates from the abundant lattice defects due to the porous structure, and the self-limiting growth behavior is ascribed to the encapsulation of templates by the fresh-deposited first layer of graphene. Although there is an obvious self-limiting growth effect of Ca(OH)2 -G on the surface of Ca(OH)2 -derived CaO template, the obtained graphene layer is much thicker in the case of CaCO3 -derived CaO template. Considering the same crystal structure and composition, the difference in graphitization behavior and graphene thickness is ascribed to the different porous structures of templates. As shown in Fig. 2.5, the pore size of CaCO3 -derived CaO is almost an order of magnitude higher than that derived from Ca(OH)2 , which will lead to different pathways for the precursor diffusion and distinct surface areas for carbon deposition. When CaCO3 is used directly as the template, the porous CaO generated in situ is supposed to exhibit the same porous structure as that annealed in advance (CaCO3 derived CaO). However, the obtained graphene layer is much thicker, which can be rationalized by the influence of CO2 produced by in situ decomposition. During the CVD process, the carbon source CH4 will be cracked and dehydrogenated to produce H2 (CH4 ↔ C + 2H2 ), which then undergoes a reverse water-gas shift reaction (CO2 + H2 ↔ CO + H2 O) with the in situ released CO2 . The total reaction is 2CO2 + CH4 ↔ C + 2CO + 2H2 O. According to Le Chatelier’s principle, the methane decomposition and sequential graphitization will be enhanced to produce a large amount of carbon-containing products, which can quickly deposit and turn black on the inner wall of the quartz tube, and also significantly increase the graphene thickness. To be more rigorous, the obtained carbon material should be described as porous graphitic carbon rather than 3D porous graphene materials.

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On the basis of the above characterization and discussion, porous CaO templates are demonstrated as another kind of promising templates for the synthesis of highquality 3D graphene materials. Both the nanostructure of employed templates and the gas composition during CVD can significantly affect the graphitization process, thus regulating the graphitization degree, hierarchical porosity, and intrinsic physical properties. Combining the results in MgO and CaO templates, we summarize the mechanism understanding of 3D graphene materials grown on porous metal oxides via CVD. Firstly, the crystal structure and surface chemistry of oxides (polar surface energy, lattice matching, lattice defects, and step sites) determine the graphitization activity and the self-limiting growth behavior. Secondly, the porous structure of templates affects the diffusion and deposition behavior of carbon sources, which can further regulate the growth and stacking characteristics of graphene. In addition, the growth conditions (e.g., temperature, pressure, gas composition, flow rate) will also have a non-negligible effect on the structure and physical properties of the obtained 3D graphene materials. These insights have enriched our understanding of the growth mechanism of graphene on oxide surfaces and the fabrication strategies of 3D graphene materials. It provides a solid foundation for the selection and design of different templates, the rational adjustment of reaction conditions, and the customization of 3D graphene materials in the future.

2.3.4 Application of 3D Porous Graphene in Lithium–Sulfur Batteries In order to explore the application prospects and structure-activity relationships of 3D porous graphene materials prepared via oxide-templated CVD process, we targeted at the application of lithium–sulfur (Li–S) batteries. Li–S batteries use sulfur as the cathode active material and metal lithium as the anode material, and the redox reactions between them are responsible for electrochemical energy conversion and storage. Li–S batteries have been claimed as the most promosing alternative for the next-generation high-performance energy storage devices due to the high theoretical specific capacity and specific energy, 1672 mAh g−1 and 2600 Wh kg−1 , respectively, which are much higher than those of current commercial lithium batteries. However, the practical application of Li–S batteries still suffers from great challenges including (1) the intrinsic insulation of sulfur and lithium sulfides, (2) the dissolution of polysulfides with a shuttle effect, (3) the huge volume change of cathode materials during operation, and (4) the safety issue due to lithium dendrites. Development of new energy materials, especially porous nanocarbon materials with tunable structural hierarchies and surface chemistries, can address these issues effectively. In the case of porous nanocarbon materials, the conductive network can greatly promote the rapid transfer of electrons in the cathode; the high specific surface area and

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porous structure enhance the adsorption and storage of polysulfides; the 3D mechanically flexible scaffold can buffer the volume change of electrode materials; the high specific surface area can render low current density favorable for inhibiting the growth of lithium dendrites. Consequently, the specific capacity, rate performance, coulomb efficiency and cycle stability of Li–S batteries will be significantly improved. The 3D porous graphene material proposed and prepared in this work exhibits high electrical conductivity, huge specific surface area, unique hierarchical porous structures, and excellent mechanical flexibility, thus rendering an ideal cathode scaffold for Li–S batteries. The 3D porous graphene materials fabricated with different CaO templates, namely Ca(OH)2 -G, CaO-C and CaCO3 -C, were employed as the cathode scaffolds for Li–S batteries to investigate their electrochemical applications and elucidate the relationship between the hierarchical porous architecture and Li–S battery performance. The sulfur was infiltrated into graphene samples via a facile melt-diffusion strategy with a C/S mass ratio of 3:7. Thermogravimetric analysis (TGA) under nitrogen reveals the resultant sulfur loading in as-fabricated C/S composites are 68, 56, and 52 wt% for Ca(OH)2 -G/S, CaO-C/S, and CaCO3 -C/S, respectively (Fig. 2.11a). Among them, the Ca(OH)2 -G/S sample has the highest sulfur content, almost soaking up all the sulfur precursor, and a higher weight-loss temperature

Fig. 2.11 Application of 3D porous graphene materials for Li–S battery cathodes. a TGA profiles of different C/S composites recorded under N2 atmosphere. b Galvanostatic discharge profiles (50th cycle). c Cycling performance for different samples at a current density of 0.5 C. The discharge capacity is specified based on the mass of whole cathodes. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

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resulted from sulfur volatilization, indicating a very strong interaction between Ca(OH)2 -G and sulfur. The large specific surface area and moderately defective surface of Ca(OH)2 -G provide abundant adsorption sites for sulfur, and the abundant micro/meso-sized in-plane pores enhance the entrapment and thermal stability of sulfur. The Li–S battery performance was evaluated in standard 2025 coin cells using the C/S composites as cathode materials. The cathode slurry was prepared by homogeneously mixing the C/S composites, the polyvinylidene fluoride binder, and the carbon nanotube conductive agents in N-methyl-pyrrolidone at a mass ratio of 8:1:1 by a magnetic stirrer for 24.0 h. The slurry was then coated onto a 15 μm aluminum foil by a routine doctor blade method, followed by a vacuum drying for 6.0 h at 60 °C. The areal sulfur loading amount is around 2.0 mg cm−2 . Disks of 13.0 mm were punched as the cathode. 1.0 M lithium bis(trifluoromethanesulfonyl) imide solution in 1:1 (v/v) 1,3-dioxolane/1,2-dimethoxyethane was selected as electrolyte. The polypropylene membrane from Celgard Inc. was used as the separator, and lithium metal foil was used as the anode. As shown in Fig. 2.11b, the charge-discharge curves of the assembled Li–S batteries exhibit two typical voltage plateaus near 2.3 V and 2.1 V. The high plateau corresponds to the rapid conversion of sulfur to soluble higher-order polysulfides. The reaction rate is fast and the plateau is shorter and less stable. The low plateau corresponds to the further lithiation into solid lithium sulfides, which is more sluggish and contributes dominantly to the discharge capacity. At a current density of 0.5 1 C (1.0 C = 1672 mA g− S ), the initial specific capacity of the CaCO3 -C/S cathode −1 can reach 1053 mAh gS with an effective utilization of sulfur (63%), which is much better than the Ca(OH)2 -G/S cathode. The specific capacity calculated based on the mass of active material (sulfur) can intuitively reveal the utilization efficiency of sulfur, but is greatly affected by the sulfur loading in the electrodes. To make the comparison among different samples more rigorous and meaningful, the specific capacities are calculated based on the mass of whole cathodes hereinafter. The cycle stability of each sample was investigated at a current density of 0.5 C. As shown in Fig. 2.11c, the Ca(OH)2 -G/S cathode delivers an initial reversible 1 specific capacity of 434 mAh g− cathode , which is lower than those of CaCO3 -C/S and CaO-C/S cathodes. However, the specific capacity of the Ca(OH)2 -G/S cathode is more stable, and exceeds the other two after 10 cycles. The cyclic fading rate of the Ca(OH)2 -G/S cathode is as low as 0.11% for initial 150 cycles, which is significantly lower than that of CaCO3 -C/S (0.17%) and CaO-C/S (0.20%) cathodes. In addition, the coulombic efficiency of the Ca(OH)2 -G/S cathode is also superior. It can stable at 90% without lithium nitrate addition, while the other two quickly decrease to 80%. Therefore, considering the specific capacity, stability, and efficiency of the assembled batteries, the performance of the Ca(OH)2 -G/S cathode is more attractive. The abundant mesoporous structure and surface adsorption sites of Ca(OH)2 -G can effectively suppress the dissolution and diffusion of polysulfides, thereby alleviating the shuttle phenomenon and improving the utilization of active sulfur, cycle stability, and efficiency.

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Fig. 2.12 Application of 3D porous graphene materials for Li–S battery cathodes. a Rate performance and b corresponding galvanostatic discharge profiles at different current densities. The dotted lines present discharge profiles of Ca(OH)2 -G/S recovered after the discharging at a high rate of 5.0 C. The discharge capacity is specified based on the mass of whole cathodes. c EIS of Ca(OH)2 -G/S cell compared with other ones. d Illustration of the hierarchical porous graphene for the cathode scaffold. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

The specific discharge capacity of all batteries gradually decreases as the current density increases (Fig. 2.12a); however, the rate capability of the Ca(OH)2 -G/S cathode is significantly improved compared to the other two samples. At a high current density of 5.0 C, the specific capacity of the Ca(OH)2 -G/S cathode can still 1 −1 reach 357 mAh g− cathode (656 mAh gS ), exhibiting a 74% retention of that at 0.1 C, while the specific capacity of the CaO-C/S cathode drops obviously to 130 mAh 1 g− cathode . It is notable that the at a high current density of 5.0 C, the low plateaus of the CaCO3 -C/S and CaO-C/S cathodes both disappear, while the Ca(OH)2 -G/S cathode exhibits two clearly distinguishable discharge plateaus. When the current density decreases again, the polarization is very small, and the discharge plateau can be recovered to the initial voltage (Fig. 2.12b). This indicates that under different current densities, all the 3D porous graphene materials can efficiently use and convert active sulfur (high plateau), while at high current densities only the Ca(OH)2 -G cathode can promote the further conversion of lithium polysulfides to lithium sulfides (low plateau). The electrochemical impedance spectroscopy (EIS) results show that the charge transfer resistance of the Ca(OH)2 -G/S cathode is the smallest, while that of CaO-C/S is the largest, in consistence with the rate capability. The high conductivity

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and low charge transfer resistance of the Ca(OH)2 -G cathode enable that electrons can still transfer quickly under high current densities, and promote the deep reduction of lithium polysulfides. Besides, the abundant in-plane mesopores of Ca(OH)2 -G for a short Li+ diffusion distance, and the interconnected cavities between graphene nanosheets for electrolyte reservoirs and a rapid intermediate diffusion can greatly improve the high rate discharge. These results demonstrate that Ca(OH)2 -G can be used as a promising cathode scaffold for Li–S batteries, which is contributed to its unique structural hierarchy as illustrated in Fig. 2.12d. First, the defective graphene layers with a high specific surface area and abundant in-plane micropores accommodate a high sulfur loading with intimate affinity. Besides, the small mesopores facilitate the entrapment of sulfur and polysulfides, thus mitigating the shuttle phenomenon and improving the stability and efficiency. Additionally, the interconnected large mesopores and macropores shorten the transport distance of Li+ ions and electrolyte, which is beneficial for rate capability. Furthermore, the integrated graphene scaffold with high electrical conductivity and mechanical robustness can also ensure the rapid electron transfer and sustain the huge volume variation during cycling. In summary, the 3D porous graphene materials are demonstrated to be excellent cathode scaffolds for highperformance, especially high-rate Li–S batteries. Better performances are expected with more meticulous design of the hierarchical structure and additional surface modification of 3D porous graphene materials.

2.4 Summary This Chapter first explored the growth mechanism and key impact factors of 3D graphene via CVD on oxide surfaces. Assisted by SEM, TEM, spherical aberrationcorrected TEM, and 3D electron tomography, the mesoporous quasi-single-crystal MgO template and the resulting 3D mesoporous graphene framework materials were characterized in detail, and the growth mechanism of 3D graphene on MgO surface was analyzed. On the one hand, the mesoporous quasi-single-crystal structure of MgO obtained from the high-temperature conversion of Mg(OH)2 provides an ideal template for the structural replication of 3D graphene. The nanopores ensure the efficient diffusion of the carbon source, and the lattice defects promote the carbon adsorption, cracking, and graphene nucleation. On the other hand, the exposure of O-terminated polar MgO (111) crystal plane ensures the stable self-limiting growth of monolayer graphene, during which the polar surface can realize atomic rearrangement and lattice matching. This provides effective guidance for the selection of oxide substrates and the preparation of 3D graphene materials in the further research. To achieve different kinds of 3D graphene materials, special attention should be paid to the pore structure, lattice defects, and exposed crystal faces of the used oxide templates.

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On this basis, this Chapter proposed a new kind of oxide templates, and successfully achieved the structural regulation of 3D graphene materials. Based on CaOtemplated CVD process, the thickness of obtained graphene can be adjusted between 1 and 10 layers, and the specific surface area can reach up to 572 m2 g−1 by tuning the reaction conditions. The research results show that the crystal structure and surface chemistry of oxides (e.g., polar surface energy, lattice matching, lattice defects, step sites), the porous structure of templates, and the growth conditions (e.g., temperature, pressure, gas composition, flow rate) comprehensively affect the structure and physical properties of the obtained 3D graphene materials. These insights have enriched our understanding of the growth mechanism, construction principles, and regulation strategies of 3D graphene materials via oxide-templated CVD process. Meanwhile, this Chapter also investigated the application of 3D porous graphene materials in Li–S batteries. When used as the cathode scaffold, the initial specific capacity of the Ca(OH)2 -G/S cathode at a current density of 0.5 C can reach 434 mAh 1 g− cathode , with a low cyclic fading rate of 0.11% for initial 150 cycles. The coulombic efficiency of the Ca(OH)2 -G/S cathode can stable at 90% without lithium nitrate addition. Besides, at a high current density of 5.0 C, the specific capacity of the Ca(OH)2 -G/S cathode can achieve a 74% retention of that at 0.1 C, exhibiting excellent utilization of active sulfur, cycle stability, and efficiency. Results revealed that the “microporous-mesoporous-macroporous” hierarchical porous structure plays a crucial role in the charging and discharging process of the battery, rendering 3D porous graphene materials as excellent cathode scaffolds for high-performance, especially high-rate Li–S batteries.

References 1. Ning GQ, Fan ZJ, Wang G, Gao JS, Qian WZ, Wei F (2011) Gram-scale synthesis of nanomesh graphene with high surface area and its application in supercapacitor electrodes. Chem Commun 47(21):5976–5978 2. Xie K, Qin X, Wang X, Wang Y, Tao H, Wu Q, Yang L, Hu Z (2012) Carbon nanocages as supercapacitor electrode materials. Adv Mater 24(3):347–352 3. Zhao J, Lai H, Lyu Z, Jiang Y, Xie K, Wang X, Wu Q, Yang L, Jin Z, Ma Y, Liu J, Hu Z (2015) Hydrophilic hierarchical nitrogen-doped carbon nanocages for ultrahigh supercapacitive performance. Adv Mater 27(23):3541–3545 4. Cui C, Qian W, Yu Y, Kong C, Yu B, Xiang L, Wei F (2014) Highly electroconductive mesoporous graphene nanofibers and their capacitance performance at 4 V. J Am Chem Soc 136(6):2256–2259 5. Tian GL, Zhang Q, Zhao MQ, Wang HF, Chen CM, Wei F (2015) Fluidized-bed CVD of unstacked double-layer templated graphene and its application in supercapacitors. AIChE J 61(3):747–755 6. Tang C, Wang HS, Wang HF, Zhang Q, Tian GL, Nie JQ, Wei F (2015) Spatially confined hybridization of nanometer-sized NiFe hydroxides into nitrogen-doped graphene frameworks leading to superior oxygen evolution reactivity. Adv Mater 27(30):4516–4522 7. Green J (1983) Calcination of precipitated Mg(OH)2 to active MgO in the production of refractory and chemical grade MgO. J Mater Sci 18(3):637–651

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8. Scott A, Dianat A, Borrnert F, Bachmatiuk A, Zhang SS, Warner JH, Borowiak-Palen E, Knupfer M, Buchner B, Cuniberti G, Rummeli MH (2011) The catalytic potential of high-κ dielectrics for graphene formation. Appl Phys Lett 98(7):073110 9. Kelber JA, Gaddam S, Vamala C, Eswaran S, Dowben PA (2011) Direct graphene growth on MgO(111) by physical vapor deposition: interfacial chemistry and band gap formation. Proc Spie 8100:81000Y 10. Rummeli MH, Bachmatiuk A, Scott A, Borrnert F, Warner JH, Hoffman V, Lin JH, Cuniberti G, Buchner B (2010) Direct low-temperature nanographene CVD synthesis over a dielectric insulator. ACS Nano 4(7):4206–4210 11. Ryou J, Hong S (2013) First-principles study of carbon atoms adsorbed on MgO(100) related to graphene growth. Curr Appl Phys 13(2):327–330 12. Tang C, Li BQ, Zhang Q, Zhu L, Wang HF, Shi JL, Wei F (2016) CaO-templated growth of hierarchical porous graphene for high-power lithium–sulfur battery applications. Adv Funct Mater 26(4):577–585 13. Zhao C, Wang W, Yu Z, Zhang H, Wang A, Yang Y (2010) Nano-CaCO3 as template for preparation of disordered large mesoporous carbon with hierarchical porosities. J Mater Chem 20(5):976–980 14. Xu B, Peng L, Wang G, Cao G, Wu F (2010) Easy synthesis of mesoporous carbon using nano-CaCO3 as template. Carbon 48(8):2377–2380 15. Shlyakhova EV, Bulusheva LG, Kanygin MA, Plyusnin PE, Kovalenko KA, Senkovskiy BV, Okotrub AV (2014) Synthesis of nitrogen-containing porous carbon using calcium oxide nanoparticles. Phys Status Solidi B 251(12):2607–2612 16. Chen JY, Guo YL, Jiang LL, Xu ZP, Huang LP, Xue YZ, Geng DC, Wu B, Hu WP, Yu G, Liu YQ (2014) Near-equilibrium chemical vapor deposition of high-quality single-crystal graphene directly on various dielectric substrates. Adv Mater 26(9):1348–1353 17. Sun JY, Gao T, Song XJ, Zhao YF, Lin YW, Wang HC, Ma DL, Chen YB, Xiang WF, Wing J, Zhang YF, Liu ZF (2014) Direct growth of high-quality graphene on high-κ dielectric SrTiO3 substrates. J Am Chem Soc 136(18):6574–6577

Chapter 3

Construction and Application of 3D Graphene Materials Based on Templated Polymerization

3.1 Introduction 3D graphene materials have excellent performance and broad application in highperformance batteries, supercapacitors, electrocatalysis, adsorption and separation and other fields. It is worth noting that although graphene exhibits a variety of unique physical properties, it is chemically inert. Most applications require not only high conductivity and large specific surface area, but also abundant active adsorption sites and hybridized active components, thus demanding the rational design and regulation of graphene materials. The oxide-templated CVD method can controllably synthesize 3D porous graphene materials, especially 3D graphene framework materials with regular porous structures, few layers and large specific surface area. However, the obtained structure dominantly depends on the structure of the oxide templates, making it difficult to tune the pore size. In addition, the CVD process is relatively high-cost, and there are many restrictions on heteroatom doping, modification, and hybridization. Therefore, it is urgent to develop a versatile method for controllable construction of 3D graphene materials with not only 3D porosity and large specific surface area, but also tunable doping, modification, and hybridization, thereby enriching the properties of graphene materials and improving the performance in applications. Inspired by the graphitization activity of oxide templates, this Chapter systematically studied the construction, regulation, and application of 3D graphene materials based on oxide-templated polymerization. Compared with the CVD method, this strategy is simple in process, low in cost, and easy to control. The obtained 3D graphene materials exhibit a hierarchical porous structure, a large specific surface area, a high heteroatom-doping content, and abundant active defect sites, leading to outstanding performance in electrocatalysis. In addition, this versatile synthesis method and the unique structure of obtained graphene provide us with a huge space for material regulation, such as the “defect engineering” strategy to in situ coordinate and disperse single-atom metal active sites.

© Tsinghua University Press 2021 C. Tang, Construction Principles and Controllable Fabrication of 3D Graphene Materials, Springer Theses, https://doi.org/10.1007/978-981-16-0356-3_3

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3.2 3D Porous Graphene Mesh 3.2.1 Synthesis of 3D Porous Graphene Mesh Herein, MgO was selected as the template for polymerization due to its ideal porous structure and catalytic graphitization properties. Specifically, the synthesis process of the 3D porous graphene mesh material based on the MgO-templated polymerization is shown in Fig. 3.1: (a)

(b)

(c)

Precursor preparation: Firstly, 30.0 g of sticky rice was washed several times and boiled in 1.0 L of deionized water at 100 °C for 4.0 h to gelatinize the rice. After removing the large-sized particles by filtration, a 900 mL of dilute dispersion was obtained. Then, 100.0 g of MgCl2 ·6H2 O was dissolved in the as-obtained dispersion, and 200.0 mL of NaOH solution (5.0 M) was dropwise added with continuous stirring, followed by 24.0 h aging at 100 °C to obtain a carbon source/template composite slurry. For nitrogen doping, the slurry was additionally mixed with melamine (15.0 g of melamine per 500.0 mL of slurry). The solid “carbon source/(nitrogen source)/template” composite was finally obtained after dry at room temperature for 4 days or at 40.0 °C for 24.0 h, for the synthesis of 3D porous graphene mesh (GM) and 3D porous N-doped graphene mesh (NGM). High-temperature polymerization: The powdery precursor was placed in a quartz boat, and put in the center of a horizontal tube furnace. The tube was heated to 950 °C at a heating rate of 15.0 °C min−1 under the protection of Ar, and then stabilized for 1.5 h for the decomposition, polymerization, and graphitization of solid precursors. After that, the sample was cooled to room temperature under the protection of Ar. Template purification: The obtained sample was collected and purified by 6.0 M hydrochloric acid at 80 °C for 24.0 h to completely remove the MgO templates. After filtering, washing, and freeze-drying, GM and NGM materials are obtained.

Fig. 3.1 Schematic of the fabrication of NGM materials. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

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Gelatinized sticky rice is a cheap and easily available carbon source, and can be uniformly mixed and strongly bonded with melamine and Mg(OH)2 during synthesis, leading to uniform composite precursors. In the case of GM synthesis, XRD spectra show that the as-obtained precursor is mainly composed of Mg(OH)2 and NaCl (Fig. 3.2a). NaCl is the by-product of the reaction between MgCl2 and NaOH, which itself can also serve as the polymerization template. The XRD profile of the sample obtained after high-temperature polymerization indicates that Mg(OH)2 has been completely converted to MgO, and the new broad peak around 12o reveals the formation of graphene. After acid treatment, the diffraction peaks of all templates (MgO and NaCl) disappear, indicating a purified graphene sample was obtained. Further, SEM images clearly show the changes in structure and morphology during the whole process. The Mg(OH)2 obtained by co-precipitation shows a hexagonal nanoflake morphology with a size of about 200 nm. Due to the coating of non-conductive carbon source, the edges of Mg(OH)2 nanoflakes are not very clear in SEM images (Fig. 3.2b). After high-temperature polymerization, the sample still maintains the morphology of hexagonal nanoflakes, but the structure is obviously more uniform

Fig. 3.2 a XRD spectra of (i) the naturally dried precursor, (ii) products obtained after hightemperature calcination, and (iii) purified graphene materials for GM. SEM images of b the naturally dried precursor, c products obtained after high-temperature calcination, and d purified graphene materials for GM. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

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and clear (Fig. 3.2c), indicating the transformation of carbon sources. As shown in Fig. 3.2d, the obtained graphene nanosheets after template removal self-assemble into a 3D porous framework. The size of the strutted cavities is determined by the size of graphene nanosheets, and ultimately the size of the MgO templates. The above results show that oxide-templated polymerization of solid carbon precursors can also produce high-quality few-layered graphene due to the unique catalytic graphitization activity of MgO. More importantly, this synthesis strategy offers more opportunities for efficient regulation of nanostructures and compositions of resulting graphene materials. Taking the most common nitrogen doping as an example, the CVD route often uses NH3 as the gaseous nitrogen source, and the doping content is very low; while in the case of templated polymerization, we can use solid nitrogen sources and mix them with the template and carbon sources uniformly in advance, thereby realizing in situ doping during the high-temperature polymerization with greatly increased contents. Melamine is a nitrogen-containing heterocyclic organic compound with a chemical formula of C3 H6 N6 and a nitrogen content of up to 66.7%. It is widely used in the synthesis of graphitic carbon nitride and is considered as an ideal nitrogen source for the synthesis of N-doped graphene. Here we employed melamine as the solid nitrogen source for nitrogen doping in MgO-templated polymerization. A certain proportion of melamine was added in the precursor slurry to prepare a “carbon source/nitrogen source/template” ternary precursor for the synthesis of NGM. The composition and morphology of samples during the synthesis process are similar to those for the synthesis of GM (Fig. 3.3). TGA was conducted under oxygen atmosphere for a series of samples obtained at different stages for NGM and GM. As shown in Fig. 3.4, the weight loss of the precursor between 260 and 350 °C is ascribed to the pyrolysis of gelatinized starch, and that between 380 and 400 °C is assigned to the water loss in the phase transition process from Mg(OH)2 to MgO. The weight loss of purified GM and NGM under oxygen atmosphere both occurs at about 550 °C, suggesting similar crystallinity and thermal stability of obtained graphene. It is worth noting that this temperature corresponding to graphene oxidation is slightly lower than that for graphene obtained by CVD (about 600 °C), which indicates relatively lower thermal stability and higher chemical activity of the graphene synthesized by polymerization. From TGA results, the yields of carbon for GM and NGM are determined to be 4.8% and 5.4% (mass ratio of graphene and templates), which is slightly higher than that of MgO-templated CVD method.

3.2.2 Nanostructure of 3D Porous Graphene Mesh The nanostructure of obtained 3D porous graphene mesh materials is further characterized by TEM technique. As shown in Fig. 3.5, ultra-thin graphene nanosheets with a lateral size of about 250 nm are cross-linked to form a 3D porous framework. The size of the strutted cavities is about 200 nm, which is consistent with SEM images (Fig. 3.5a, c). The high-resolution TEM images show that the graphene nanosheets

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Fig. 3.3 a XRD spectra of (i) the naturally dried precursor, (ii) products obtained after hightemperature calcination, and (iii) purified graphene materials for NGM. SEM images of b the naturally dried precursor, c products obtained after calcination, and d purified graphene for NGM. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons Fig. 3.4 TGA curves of those samples obtained at different stages for NGM and GM. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

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Fig. 3.5 TEM images of a, b NGM and c, d GM. High-resolution TEM images of e MgO and f nanoparticles in unpurified graphene mesh. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

obtained by templated polymerization are dominantly monolayer, similar to those synthesized by CVD on MgO. What is different is that a large number of nano-sized holes are randomly distributed in the graphene nanosheet (Fig. 3.5b, d), exhibiting a mesh-like morphology. The synthesis process of 3D porous graphene mesh materials can be described as follows. Mg(OH)2 nanosheets generated by co-precipitation in the precursor are

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converted into porous MgO nanosheets in situ during the high-temperature calcination. The MgO nanosheets, as well as the NaCl by-product, serve as the templates for graphene growth. The high-resolution TEM images of unpurified 3D porous graphene mesh materials confirm the presence of MgO and NaCl templates (Fig. 3.5e, f). The gaseous carbon sources generated by decomposing the gelatinized sticky rice at high temperature adsorb and nucleate on the surface of MgO to form continuous few-layered graphene nanosheets. On the one hand, due to the high catalytic graphitization activity of MgO and the characteristics of self-limiting growth, the obtained graphene is dominated by monolayer with high crystallinity. The electrical conductivity was tested to be 120 S m−1 and 580 S m−1 for NGM and GM, respectively, by the four-probe method. On the other hand, the in situ released water vapor and gas from the precursor decomposition and the random nucleation-growth behaviors from different sites on the substrate lead to a large amount of micro holes and edge sites, which can promote the efficient incorporation of nitrogen atoms. The porosity of obtained 3D porous graphene mesh material was evaluated by nitrogen adsorption-desorption experiments at 77 K. As shown in Fig. 3.6a, the nitrogen sorption isotherms of GM and NGM are both typical type-IV isotherms. In the low P/P0 range, the convex curve corresponds to the Langmuir monolayer adsorption of nitrogen molecules on the micro/mesoporous surface of the graphene nanosheets, contributing a small amount of specific adsorption capacity. In the medium P/P0 range, the obvious hysteresis loop indicates capillary condensation of nitrogen molecules originated from the presence of abundant mesopores in the graphene structure. When the relative pressure is close to 1, the nitrogen adsorption mainly occurs at the strutted macropores, resulting in further increase in specific adsorption capacity. The specific surface area was calculated by the multipoint Brunauer–Emmett–Teller (BET) method, and the pore-size distribution was calculated based on Quenched Solid Density Function Theory model using the adsorption branch. The specific surface areas are calculated to be 1655 m2 g−1 and 1100 m2

Fig. 3.6 N2 sorption isotherms of NGM and GM, inset figure illustrates the pore size distribution. b Raman spectra of NGM and GM. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

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g−1 for GM and NGM, respectively, indicating that the number of graphene layers is within two layers. Both NGM and GM contain a small amount of micropores and abundant mesopores (inset of Fig. 3.6a), with total pore volumes reaching 2.72 cm3 g−1 and 3.39 cm3 g−1 , respectively. The difference in the pore size may be ascribed to the nitrogen source affecting the high-temperature decomposition and polymerization behaviors, resulting in the size change of the nano-sized pores in NGM. Although the graphene grown on the surface of MgO has high crystallinity, there are a large number of structural defects in GM and NGM due to the abundant pores caused by the conversion of solid-phase carbon source, which can be analyzed by Raman spectroscopy. As shown in Fig. 3.6b, the Raman spectra of GM and NGM exhibit two distinct characteristic bands, namely D band near 1328 cm−1 and G band near 1590 cm−1 . The strong and wide D band indicates a large number of non-sixmembered ring structures. The intensity ratio of D and G bands (I D /I G ) can be used to compare the amount of defects in the graphene samples. The slightly higher I D /I G ratio of NGM with respect to GM (1.24 versus 1.11) is ascribed to the structural modification brought about by nitrogen doping.

3.2.3 Heteroatom Doping of 3D Porous Graphene Mesh As mentioned above, the templated polymerization provides convenience for the regulation of graphene structure and composition, especially heteroatom doping. In this work, we employed melamine as the nitrogen source to realize the highefficiency nitrogen doping in graphene mesh materials. The doping content can be facilely adjusted by changing the ratio of the nitrogen source in the precursor. As shown in Fig. 3.7, we conducted energy-filtered TEM to characterize the structure and element composition of NGM. The results show that nitrogen (Fig. 3.7b) and carbon (Fig. 3.7c) elements are uniformly distributed, in consistence with the porous structure of NGM (Fig. 3.7a). The electron energy loss spectrum (EELS) analysis can analyze the element composition, chemical bond, and electronic structure of the sample, and is especially suitable for low-atomic-number elements. The characteristic π* and σ* transitions in the carbon K-edge spectrum suggest the dominant sp2 hybridization state of carbon in NGM (Fig. 3.7d). The σ* transition near 401.5 eV in the nitrogen K-edge spectrum is ascribed to the substitutional dopants such as graphitic nitrogen, and the broad σ* band near 408.5 eV is usually assigned to pyridinic nitrogen (Fig. 3.7e) [2, 3]. It is worth noting that graphitic nitrogen and pyridinic nitrogen are widely regarded as the origin of electrocatalytic activity for ORR in carbon materials [4]. XPS survey spectra reveal that GM and NGM dominantly contain three elements: carbon, nitrogen, and oxygen (Fig. 3.8a, Table 3.1), with the atomic contents of 96.80 at.%, 0.41 at.%, 2.73 at.%, and 88.50 at.%, 7.60 at.%, 3.64 at.% for GM and NGM, respectively. The trace nitrogen doping of GM comes from trace components in sticky rice, while the high content of nitrogen doping of NGM comes

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Fig. 3.7 a TEM image of NGM and corresponding b N and c C energy-filtered TEM images. d Typical nitrogen and e carbon K-edge EELS spectra of NGM. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

Fig. 3.8 a The XPS surveys of NGM and GM. b The high-resolution N 1s XPS spectrum of NGM. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

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Table 3.1 XPS results and the comparison of N/O groups for NGM, GM and NG [1] Sample

NGM GM

Total content of carbon (at.%)

88.50 96.80 88.85

Total content of nitrogen (at.%)

7.60

0.41

7.48

Total content of oxygen (at.%)

3.84

2.73

3.64

37.1

40.0

53.5

Relative amounts of nitrogen species (%)

Pyridinic nitrogen (398.5 eV)

NG

Pyrrolic nitrogen (400.1 eV)

33.2

60.0

46.5

Quaternary nitrogen (401.2 eV)

15.1

N.A.

N.A.

Oxidized nitrogen (402.9 eV)

6.7

N.A.

N.A.

Chemisorbed nitrogen (405.6 eV) 7.9 Relative amounts of oxygen species C = O (531.2 ± 0.3 eV) (%) C–O (533.1 ± 0.3 eV)

N.A.

N.A.

58.4

51.4

55.1

41.6

48.6

44.9

from melamine. The oxygen doping of both comes from the inevitable oxygen functional groups generated during high-temperature polymerization and acidic purification. The high-resolution N 1s XPS spectrum of NGM is deconvoluted into dominant pyridinic N (37.1%), pyrrolic N (33.2%), and graphitic N (15.1%) (Fig. 3.8b), in good agreement with the EELS results. Further, we compared the ability of regulating porous structure and composition by different synthesis methods including templated polymerization, CVD, and post-treatment of GO. Nitrogendoped graphene (NG-CVD) was synthesized by CVD method using mesoporous MgO as the template, and methane/ammonia as carbon/nitrogen sources at 950 °C. The nanostructure is similar to the un-doped mesoporous graphene framework as discussed in Chapter 2. Nitrogen-doped graphene oxide (NG-GO) was obtained by post-treatment of commercial graphene oxides with urea under hydrothermal reaction at 180 °C for 18 h. The SEM image (Fig. 3.9a) and TEM image (Fig. 3.9b) clearly show that NG-GO still maintains the structure of GO precursors after hydrothermal treatment, exhibiting the morphology of micron-scale flakes without in-plane pores. The nitrogen sorption isotherm of NG-GO exhibits a very low specific adsorption capacity in the low P/P0 range and a small hysteresis loop in the medium P/P0 range, suggesting much lower specific surface area and poor porosity compared with NGM. Accordingly, the specific surface area of NG-GO is calculated to be 241 m2 g−1 and the total pore volume is only 0.32 cm3 g−1 . Nevertheless, the hydrothermal treatment of GO precursors always lead to a higher doping content. When the mass ratio of GO to urea is 1:20, the nitrogen content of the obtained NG-GO is as high as 7.48 at.%, with a high content of oxygen-containing functional group (3.64 at.%). The highresolution N 1s XPS spectrum of NG-GO can be well fitted with pyridinic N (53.5%) and pyrrolic N (46.5%). The porosity and composition of graphene materials synthesized by different methods are compared in Table 3.2. Notably, compared with the post-treatment

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Fig. 3.9 a SEM and b TEM images of NG-GO. c N2 sorption isotherms of NG-GO, inset figure illustrates the pore size distribution. d The high-resolution N 1s XPS spectrum of NG-GO. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

Table 3.2 Comparison of porosity and composition of different graphene materials [1] Sample

Specific surface area (m2 g−1 )

Total pore volume (cm3 g−1 )

N content (at.%)

O content (at.%)

NGM

1100

2.72

7.60

3.84

GM

1655

3.39

0.41

2.73

241

0.32

7.48

3.64

1440

2.18

3.41

3.34

NG-GO NG-CVD

of GO precursors, the 3D graphene materials obtained by templated polymerization and CVD have a higher specific surface area and pore volume, suggesting much thinner graphene and better porosity hierarchy (especially microporous and mesoporous structure). Compared with the templated CVD method, the templated polymerization route is more efficient to modify the graphene with heteroatom doping while maintaining the high quality of graphene and hierarchical porous structure. Consequently, MgO-templated polymerization is an attractive strategy for constructing multifunctional 3D graphene materials with regulated surface chemistry and enhanced performances.

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3.2.4 Electrocatalytic Performance of 3D Porous Graphene Mesh In recent years, graphene materials, especially heteroatom-doped graphene materials (e.g., nitrogen, oxygen, boron, sulfur, phosphorus), have been widely reported as a family of efficient ORR electrocatalysts, which can be comparable with commercial Pt/C catalysts with better electrochemical stability and methanol tolerance [5–9]. In addition to heteroatom modification, a few research works have paid attention to the role of edge sites and topological defects in nanocarbon materials on resulting ORR activity [10–15]. Besides, some works also reported that nanocarbon materials can deliver considerable OER activity or ORR/OER bifunctional catalytic activity [16–18]. In this Chapter, we have prepared a new type of 3D porous graphene mesh material by an innovative synthesis method, which has a unique microporous-mesoporousmacroporous hierarchical structure that can enhance the mass transfer in the heterogeneous reaction process. Meanwhile, the abundant nitrogen doping, oxygen functional groups and topological defects near the edge of in the mesh materials are expected to provide efficient catalytic active sites. In addition, the highly conductive 3D graphene framework offers smooth electron conduction network and fully exposed active sites with a favorable gas-solid-liquid triple-phase boundary. Therefore, the 3D porous graphene mesh material is expected to be an excellent electrocatalyst for ORR and OER. The electrocatalytic performance of graphene materials was evaluated in a threeelectrode system (CHI 760D, CH Instrument, USA). Catalysts were loaded onto the working electrodes, including a rotating disk electrode (RDE) with a disk diameter of 5.0 mm, and a rotating ring-disk electrode (RRDE) with a disk electrode (glassy carbon, diameter: 5.0 mm) and a ring electrode (Pt electrode, inner diameter: 6.5 mm, outer diameter: 7.5 mm). Pt sheet was selected as the counter electrode and saturated calomel electrode (SCE) as the reference electrode. The electrolyte used 0.10 M KOH or HClO4 solution saturated with O2 or N2 . The fabrication of working electrode was carried out as follows: Taking NGM for example, 5.0 mg of catalyst was firstly dispersed in 0.95 mL of ethanol and 0.05 mL of Nafion solution (5.0 wt%), followed by 1.0 h sonication to form a homogeneous suspension. 10.0 μL of the suspension was pipetted onto the glassy-carbon disk electrode, which was mechanically polished and ultrasonically washed in advance. After solvent evaporation for 10.0 min in air, the working electrode was prepared for electrochemical measurements with a catalyst loading of 0.25 mg cm−2 . For the evaluation of electrocatalytic activity such as ORR and OER, the commonly used techniques include cyclic voltammetry (CV), linear sweep voltammetry (LSV), chronoamperometry test, and AC impedance test. The scan rate of CV tests was 100 mV s−1 , and that of LSV tests was 10 mV s−1 . The rotating rate of RDE and RRDE was 1600 rpm. All polarization curves were corrected with 95% iR-compensation. The overpotential was calculated by the following equation:

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Overpotential = measured value (versus SCE) + 0.0592 pH + 0.241 – 1.229 (versus RHE). Considering that NGM has high conductivity, high specific surface area, hierarchical porous structure, high nitrogen content, and abundant defective edge sites, we first evaluated the electrocatalytic activity of NGM. As shown in Fig. 3.10a, the current density under O2 saturation is significantly higher than that under N2 saturation with an obvious reduction peak at 0.65 V, indicating the reduction of O2 . LSV profiles further reveal the superior ORR activity of NGM with a limiting current density as high as 7.5 mA cm−2 in O2 -saturated 0.10 M KOH electrolyte (Fig. 3.10b). It is notable that the current in O2 -saturated electrolyte is also detectable, which is mainly assigned to the capacitive current due to the high specific surface area of NGM and other graphene materials. In order to evaluate the ORR electrocatalytic performance of such materials more rigorously, we subtracted the current under N2 saturation from that under O2 saturation, as shown by the solid line in Fig. 3.10b. Therefore, the limiting current density of NGM material in the O2 -saturated 0.10 M KOH electrolyte can reach 6.4 mA cm−2 , and that in the O2 -saturated 0.10 M HClO4 electrolyte can also reach above 6 mA cm−2 .

Fig. 3.10 a CV curves of NGM in O2 and N2 -saturated 0.10 M KOH. b LSV curves of NGM obtained in O2 and N2 -saturated 0.10 M KOH or 0.10 M HClO4 . c The disk (bottom) and ring current densities (top) recorded on an RRDE, and d corresponding electron transfer number (n). Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

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To further evaluate the ORR performances and pathways, RRDE measurements were carried out for different catalysts, including NGM, GM, NG prepared by the post-treatment of GO, and commercial Pt/C (20 wt%). From the disk current density curves in Fig. 3.10c, it can be observed that the onset potential (E onset ) and half-wave potential (E 1/2 ) of NGM (E onset ≈ 0.89, E 1/2 ≈ 0.77) are slightly lower than those of commercial Pt/C catalysts (E onset ≈ 0.94, E 1/2 ≈ 0.80), but a much higher limiting current density of 6.41 mA cm−2 is obtained compared with Pt/C (4.97 mA cm−2 ). For the GM, in spite of a much lower N content, a high limiting current density of 5.72 mA cm−2 can be detected, while the overpotential is much larger than NGM (E onset ≈ 0.83, E 1/2 ≈ 0.68). Notably, the N content of GM is extremely low. Interestingly, when an even higher N content is doped into the nanocarbon, the overpotential of NG is similar to that of NGM, but the current density is much lower than GM. If we normalize the current (at 0.47 V) with respect to the amount of N dopants, 1 the specific current density of GM catalyst (3.27 mA μg− N ) significantly surpasses −1 −1 that of NGM (0.26 mA μgN ) and GM (0.19 mA μgN ), and even higher than the best reported result (nitrogen-doped graphene/single-walled carbon nanotube hybrid, 1 2.60 mA μg− N ) [19]. Since GM and NGM share similar structures and porosity features, the large difference in electrocatalytic properties implies that the electrocatalytic activity does not only originate from the nitrogen dopants in 3D porous graphene mesh materials. Two obvious conclusions can be drawn from the above results. On the one hand, NGM and NG exhibit high and comparable contents of nitrogen doping, and thus similar onset potential; while for GM with a low N content, the overpotential is much larger. On the other hand, NGM and GM have similar hierarchical porous structures, especially abundant in-plane pores, and thus the current density at large overpotentials is much higher than that of NG regardless of the onset potential. In other words, the heteroatom doping in graphene materials, such as nitrogen doping, determines the onset potential of ORR, while the hierarchical porous structure has a non-negligible effect on the current density at large overpotentials. The hierarchical porous structure of the material not only ensures the full exposure of active sites and the efficient reaction in the triple-phase boundary, but also provides a large number of unstable dangling bonds and topological defects, which are most likely to be new active sites. The electron transfer number per O2 (n) is determined from the RRDE profiles (Fig. 3.10c) to study the reaction intermediate and pathway. The electron transfer number is calculated based on the disk (I d ) and ring current (I r ) as follows: n = 4I d /(I d + I r /N), where N is current collection efficiency of the Pt ring (0.26). As shown in Fig. 3.10d, NGM exhibits a similar n to Pt/C (around 3.8) over the whole potential range of 0.0–0.8 V, suggesting a dominant four-electron pathway of ORR. However, for both GM and NG, the n is much lower (2.5~3.2), indicating a different reaction pathway such as two-electron or “two + two” electron via the intermediate of OOH. The four-electron pathway on NGM clearly reveals the synergy effect between nitrogen doping and edge defects. Remarkably, the n of NGM obtained in 0.10 M HClO4 also approaches 4, further indicating the superb activity and selectivity of NGM in both acidic and alkaline conditions.

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In the study of the ORR electrocatalytic activity mechanism of carbon nanomaterials, there is a very important debate: Compared with heteroatom doping, do trace metal impurities (especially Fe) provide comparable catalytic activity? Since the conventional synthesis process of nanocarbon materials (e.g., carbon nanotubes, graphene, porous carbon) inevitably involves some metal catalysts, it is difficult to completely remove them in the purification process. In this work, the preparation of graphene uses the MgO template, which can be completely removed by acid treatment compared to metal nanoparticles, and the process does not involve any electrocatalytically active metals. So the ORR activity of the obtained graphene mesh materials should originate from the carbon material itself. In order to verify this idea, we conducted a poisoning experiment by adding KSCN into the electrolyte. If any Fe-based active sites existed in the catalyst, SCN– will be strongly coordinated with them, and the electrocatalytic performance will be significantly weakened. As shown in Fig. 3.11a, no activity loss of the NGM electrode is observed after the addition of KSCN, confirming the high intrinsic activity of this novel metal-free nanocarbon catalyst. The electrocatalytic activity is highly affected by the temperature, and the apparent activation energy of the reaction can be calculated by the LSV curves obtained at different temperatures [20, 21]. During the test, the electrolytic cell containing 0.10 M KOH was suspended in a thermostatic water bath with increasing temperature from 5 to 40 °C. The LSV tests were conducted at a scan rate of 10 mV s−1 . The potential was corrected for temperature as E versus RHE = E versus SCE + [0.241–6.61 × 10−4 (T – 298)] + 1.98 × 10−4 T pH. Then the Tafel slope was determined according to the Tafel equation η = b log(j/j0 ) based on the LSV curves, where η is the corrected overpotential, b is the Tafel slope, j is the current density, and j0 is the exchange current density. By plotting log j0 versus 1/T, the activation energy E a was calculated according to the Arrhenius equation ∂ (log j0 )/∂ (1/T ) = – E a /2.3R. As shown in Fig. 3.11b, the apparent activation energy of ORR catalyzed by NGM is

Fig. 3.11 a Chronoamperometric response of NGM in O2 -saturated 0.10 M KOH with the addition of KSCN. b Arrhenius plots of Pt/C and NGM, and the inset presents the calculated ORR apparent activation energies. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

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34.4 kJ mol−1 , which is very close to that of Pt/C catalyst (31.4 kJ mol−1 ), directly verifying the high electrocatalytic activity of NGM. Both GM and NGM exhibit superior ORR durability as revealed by the chronoamperometric responses (Fig. 3.12) at a constant potential of 0.60 V versus RHE. In O2 -saturated 0.10 M KOH electrolyte, the ORR current of the Pt/C catalyst is significantly reduced after 8.0 h tests, retaining only 82.3% of its initial current. In contrast, the currents of GM and NGM are relatively more stable. The ORR current of NGM can retain as high as 98.5% of the initial value, and the LSV curves almost overlap with each other before and after 8.0 h tests (Fig. 3.12b). In 1.0 M KOH electrolyte, the NGM can also deliver a nearly constant current for more than 10 h, which further confirms the excellent durability of NGM catalysts for ORR. Further, we investigated the OER catalytic performance of the obtained 3D graphene material. As shown in Fig. 3.13a, both NGM and GM have considerable

Fig. 3.12 a ORR chronoamperometric responses at a constant potential of 0.60 V versus RHE. b LSV curves for NGM obtained before and after 8.0 h chronoamperometric tests at a constant potential of 0.60 V versus RHE. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

Fig. 3.13 a The LSV curves recorded in the OER region in O2 -saturated 0.10 M KOH solution. b Comparison of the oxygen electrode activities of the recently reported highly active ORR/OER bifunctional catalysts. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

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OER activity, which is much higher than that of Pt/C catalyst, and is also comparable to commercial Ir/C (20 wt%) catalyst. The OER overpotential required for a current density of 10 mA cm−2 is about 410 mV. It is worth noting that, unlike ORR catalysis, the OER activity of GM is slightly higher than that of NGM, which is due to the different active sites required for ORR and OER. For ORR/OER bifunctional catalysts, the potential difference (E) between the potential corresponding to the OER current density of 10 mA cm−2 (E j=10 ) and the ORR half-wave potential (E 1/2 ) is one of the most important evaluation indicators. As summarized in Fig. 3.13b, the E of NGM is less than 0.9 V, which is one of the best carbon-based ORR/OER bifunctional catalysts reported in the literature, and is even better than many metal-based electrocatalysts. In summary, owing to the high nitrogen doping content and hierarchical porous structure, NGM exhibits excellent ORR and OER bifunctional catalytic activities. It demonstrates the superiority of the MgO-templated polymerization method to prepare functionalized 3D graphene materials. Additionally, the unique structure and tunability of NGM and GM also provide us with a very special and efficient platform for the study of electrocatalytic activity mechanisms, which can systematically explore the roles of nitrogen doping, edge sites, and topological defects. The acquired knowledge will in turn guide the rational design and controllable construction of more efficient graphene-based electrocatalysts.

3.2.5 Electrocatalytic Activity Mechanism of 3D Porous Graphene Mesh The above material characterization and catalytic activity evaluation show that in carbon-based electrocatalysis, the porous edge sites or topological defects in nanocarbon materials and heteroatom modification both play important roles on the resulting electrocatalytic activity. It may be the reason why the NGM sample exhibits outstanding electrocatalytic performance, and the ORR activity of GM is much higher than that of NG. In order to further elucidate our experimental observations and confirm the activity origins, density function theory (DFT) calculations were conducted based on the four-electron pathway mechanism for both ORR and OER [22]. Considering all the possible active sites derived from the nitrogen-doping, edge effects, and topological defects, a series of models are proposed systematically based on graphene nanoribbons (GNRs). As shown in Fig. 3.14, the studied models include pyrrolic N (PR), pyridinic N (PN), quaternary N on the edge (Q), quaternary N in the bulk phase (QN), five-carbon ring (C5), seven-carbon ring (C7), and five-carbon ring adjacent to seven-carbon ring (C5+7), compared with the pristine graphene nanoribbon (G). Specifically, the total energy of OER/ORR intermediates, graphene nanoribbons, and small molecules like H2 , H2 O, were all calculated by DFT calculations performed in Dmol3 package in Materials Studio of Accelrys Inc with

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Fig. 3.14 A schematic graphene nanoribbon with different kinds of active sites. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

Fig. 3.15 a Comparison of the overpotential values for different models for OER and ORR catalysis. Dot lines represent the overpotential for the pure graphene nanoribbon. b The edge effect of the nitrogen-doping for ORR catalysis. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

general gradient approximation (GGA) [23, 24] and Perdew-Burke-Ernzerhof (PBE) exchange-correlation functions [25]. An all-electron double numerical basic set with polarization functions (DNP basic set) was used during calculating. The convergence tolerance quality of geometry optimization was set to fine: 1.0 × 10−5 au for energy, 2.0 × 10−3 au Å−1 for maximum force and 5.0 × 10−3 Å for maximum displacement. The k-points for Brillioun zone were selected by Monkhorts-Pack method and set to 8 × 1×1 [26]. The overpotential for each active site was calculated to serve as the vital figureof-merit, as shown in Fig. 3.15. Different from the conclusions reported in most literatures [14], the N-doped moieties deliver a relatively higher overpotential compared to the pristine graphene nanoribbons, except for the pyrrolic N configuration, indicating that the nitrogen doping itself is not necessarily a good ORR active site. Among various nitrogen doping-induced sites, all the sites near the edge exhibit a

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much lower overpotential (Fig. 3.15b), indicating the importance of edge effects; besides, the QN configuration delivers the highest overpotential for ORR while the lowest overpotential for OER, suggesting that the favorable active sites for ORR and OER are distinct [14, 27], which is consistent with the above experimental results. When topological defects without heteroatoms are introduced into the graphene structure, the overpotentials of ORR and OER are greatly decreased in our DFT calculation. For example, the ORR overpotential of C5 is reduced by 0.33 V compared to the original graphene nanoribbon. It is notable that the most active site PR for the N-containing configurations is also a five-atom ring. When C5 and C7 are adjacent, the two will adjoin into a curved configuration C5+7 with a certain spatial curvature. The resulting ORR overpotential will be further reduced to 0.14 V, and the OER overpotential to 0.21 V, strongly indicating the high electrocatalytic activity of topological defective sites. In fact, according to the second law of thermodynamics, structural defects in crystalline materials are inevitable. For nanocarbon materials, a certain amount of topological defects (such as the most common Stone-Wales defect) can be easily introduced during in situ doping or subsequent processing [28]. Two typical volcano plots can be constructed when correlating the overpotential and descriptor (adsorption energy of OH*) for both ORR and OER (Fig. 3.16a). C5+7 is at the apex of two volcano plots, showing the best activity for both ORR and OER. The active sites on the left side exhibit two strong adsorption of OH*, while those on the right side too weak, thereby leading to higher overpotentials.

Fig. 3.16 a ORR and OER volcano plots of overpotential versus adsorption energy of OH*. b Optimized adsorption structure of C5+7 interacting with O, OH, and OOH species. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

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Fig. 3.17 Calculated free energy diagrams of ORR for a C5+7 defect and b PR dopant. Reprinted from ref. [1], copyright 2016, with permission from John Wiley and Sons

Topological defects can adjust the adsorption strength of graphene to intermediate species more effectively and moderately, and enables ideal ORR/OER bifunctional catalytic activities. Taking C5+7 as an example, the O and OOH species prefer to bind at the intersection of the five-carbon and seven-carbon rings (Fig. 3.16b). The O–O bond in OOH is thus elongated, making it easier to break and thus promoting the reduction of oxygen. The adjacent five-carbon and seven-carbon rings with different electron densities generate spatial curvatures and form a permanent dipole moment. The dipole moment is slightly weaker than that between C and N atoms, thus leading to more moderate adsorption strength and optimal catalytic activity. In order to further analyze the intrinsic reaction mechanism, we predicted the free energy diagrams of ORR sub-steps as shown in Fig. 3.17. The thermodynamic equilibrium potential of ORR is 1.23 V (versus RHE). However, the actual reaction is irreversible with a certain overpotential. As the overpotential increases (decreasing the U value in Fig. 3.17a), the desorption of OH– is the last step to become a nonendothermic reaction and downhill for C5+7. It is regarded to be the rate determining step. The overpotential is 0.14 V (= 1.23 – 1.09). For the PR configuration, the rate determining step is revealed to be the transformation of O* to OH* (Fig. 3.17b), corresponding to an overpotential of 0.78 V. The C5+7 and other active sites containing topological defects are thus demonstrated to optimize the adsorption behavior of the catalyst to intermediate species, make the adsorption and desorption more favorable, and change the rate determining step, ultimately achieving a significant improvement in catalytic performance [29, 30].

3.3 Atomic Metal Sites Anchored in 3D Porous Graphene Owing to the high catalytic activity, excellent structural tunability, low cost and efficient synthesis methods, heteroatom-doped nanocarbon materials (one or more kinds of heteroatoms) are a family of very promising ORR/OER bifunctional catalysts, which can be used as the air electrode for next-generation energy storage devices

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such as zinc–air batteries. Due to the difference in electronegativity and electron affinity between heteroatoms and carbon atoms, the incorporation of heteroatoms can derive rearrangement of the charges [14] or spins [31, 32] of surrounding atoms, making the originally chemically inert sp2 carbon to be favorable adsorption and catalytic sites. As discussed in the previous section, more and more studies have shown that dopant-free topological defective sites in nanocarbon materials (e.g., seven-carbon ring, seven-carbon ring, combinations of five-seven-carbon rings) can also contribute remarkable electrocatalytic activities [10, 11]. In addition, the introduction of metal elements intentionally or unintentionally during the synthesis of nanocarbon materials, even if in a trace amount, can significantly improve the electrocatalytic performance of the resulting materials [33, 34]. The coordination of metal atoms with heteroatoms, especially the atomically dispersed active sites (or single-atom catalytic sites) such as Co/Fe–Nx –C, can effectively adjust the local electronic structure of the sp2 carbon matrix and optimize the adsorption behaviors of intermediates on the surface, thereby resulting in outstanding catalytic performance comparable to that of noble metal catalysts [35–38]. However, powerful synthetic methodologies are urgently required to precisely construct and tune the active sites of nanocarbon materials at the atomic scale, such as doping, edges, defects, and metal-nitrogen-carbon sites. Based on the achievements in the preparation, regulation, and electrocatalytic application of 3D porous graphene mesh materials, we further exploited the advantages of the templated polymerization strategy and the unique structure of resulting materials. We creatively proposed the “defect engineering” concept to fully use the abundant defective edges in 3D porous graphene mesh materials to efficiently construct atomically dispersed “metal-nitrogen-carbon” sites. This novel synthesis strategy and material structure are demonstrated to further improve the electrocatalytic performance of graphene materials, and expand the synthesis methods, nanostructures, compositions, and application scenarios of 3D graphene materials.

3.3.1 Material Synthesis and Characterization Similar to the preparation process of the 3D porous graphene mesh materials, the preparation of the atomic metal-coordinated 3D porous graphene material uses the “carbon source/nitrogen source/template/metal source” composite as the precursor for high-temperature polymerization. In order to simplify the experimental process, this work replaces sticky rice with amylopectin, which can omit the filtration step, and is easier to be promoted without concentration variation. Specifically, 10.0 g of amylopectin was boiled in 1.0 L of deionized water at 100 °C for 3.0 h to gelatinize. 50.0 g of MgCl2 ·6H2 O was dissolved in the as-obtained dispersion, then 100 mL of NaOH (5.0 M) solution was added dropwise with continuous stirring. After that, 30.0 g of melamine was added into the slurry with continuous stirring, followed by 24.0 h aging at 100 °C. The other processes and polymerization conditions were the same as those for the synthesis of NGM. The obtained 3D

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graphene material was denoted as NGM-a. If a certain amount of metal salt, such as 116.8 mg of Co(NO3 )2 ·6H2 O was added to the precursor slurry, a metal-coordinated 3D porous graphene material can be finally obtained under otherwise identical conditions, denoted as NGM-Co. Notably, we can facilely prepare 3D graphene materials with varied concentrations of different metal coordination by changing the type and ratio of the metal salts in this synthesis strategy, which is of high universality and adjustability. The preparation process is similar to that for the 3D porous graphene mesh materials, but the change of carbon sources and concentrations of each component can slightly affect the resulting morphology and nanostructure. Figure 3.18 shows the morphology of samples at different stages during the preparation of NGM-Co. The template is still in a morphology of hexagonal nanosheets (Fig. 3.18a), which is well maintained after high-temperature polymerization (Fig. 3.18b). After template removal by acid treatment, the self-assembled 3D porous graphene can be obtained. The size of the strutted cavities is also about 200 nm, similar to that of GM and NGM. The individual graphene nanosheets are ultra-thin and wrinkled with abundant in-plane holes (Fig. 3.18c, d).

Fig. 3.18 SEM images of the samples for NGM-Co a before carbonization and b after carbonization. c SEM and d TEM images of NGM-Co after purification. Reprinted from ref. [39], copyright 2016, with permission from John Wiley and Sons

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The structural porosity of NGM-a and NGM-Co was further characterized by nitrogen sorption experiments. As revealed in Fig. 3.19a, the nitrogen sorption isotherms are both typical type-IV isotherms, in consistence with those of GM and NGM. The pore size distribution is similar (Fig. 3.19b), showing a small amount of micropores and abundant mesopores. The specific surface area is calculated to be 272 m2 g−1 and 541 m2 g−1 with the total pore volume of 1.03 cm3 g−1 and 1.36 cm3 g−1 for NGM-a and NGM-Co, respectively. The similar morphology and pore size distribution illustrate the templated replication effect of MgO nanosheets. However, both the specific surface area and total pore volume are significantly lower than those of GM and NGM, which can be rationalized by the different interactions between different carbon sources and oxide templates, as well as the changed ratio of carbon sources to templates in the precursor. It is worth noting that under the same synthesis conditions except the addition of cobalt salts, the specific surface area of the obtained NGM-Co is twice that of NGMa, and the total pore volume is also 30% higher, indicating the significant role of metals on regulating the growth of graphene and the construction of porous structure. The pore size distribution diagram shows that NGM-Co and NGM-a have similar

Fig. 3.19 Pore structure analysis of NGM-a and NGM-Co. a The N2 sorption isotherms and b the corresponding pore size distribution. c The surface area histogram and d pore volume histogram versus pore size, indicating the distinct micropores and mesopores below 5.0 nm. Reprinted from ref. [39], copyright 2016, with permission from John Wiley and Sons

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Fig. 3.20 a Raman spectra for NGM-a and NGM-Co obtained at 633 nm. b XRD results for the unpurified NGM-Co and unpurified NGM-a, showing the presence of CoO, MgO, and NaCl. Reprinted from ref. [39], copyright 2016, with permission from John Wiley and Sons

porous structures in the range above 5 nm, while the mesoporous and microporous regions below 5 nm are significantly different (inset in Fig. 3.19b). Furthermore, Figs. 3.19c and d respectively show the contribution of pores of different sizes to the specific surface area and pore volume, from which we can clearly observe that the mesopores and micropores below 5 nm are obviously increased for NGM-Co. In addition, a relatively higher I D /I G ratio in the Raman spectrum is obtained for NGMCo (Fig. 3.20a), indicating more defective sites due to the increased micropores and mesopores. The difference in porosity and structural features can be ascribed to the catalytic and etching effect of cobalt/cobalt oxide nanoparticles during hightemperature carbonization. As shown in Fig. 3.20b, the XRD profiles identify that Co(NO3 )2 ·6H2 O in the NGM-Co precursor is converted into CoO nanoparticles after high-temperature calcination, which can provide templated catalytic effect similar to MgO nanosheets. Compared with MgO templates, the CaO nanoparticles are more active and may lead to carbon etching, thus significantly improving the mesoporous and microporous structures of NGM-Co below 5 nm.

3.3.2 Defect Engineering Towards Atomic Co–Nx –C Sites Compared with the 3D porous graphene mesh materials, the more abundant mesoporous and microporous structure of NGM-Co is expected to provide more topological defects as electrocatalytic active sites, and new opportunities for constructing atomically dispersed “metal-nitrogen-carbon” active sites. Theoretical studies have shown that the “metal-nitrogen-carbon” structure is more energetically feasible to form at the edge of graphene [40]. Here, the Co atoms are also expected to coordinate with the N dopants at defective edges or pores of graphene, leading to Co–Nx –C single-atom sites.

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After acid treatment, no obvious nanoparticles can be observed in the highresolution TEM image of the NGM-Co sample, as well as no signal of metal or metal oxides in the XRD spectrum (Fig. 3.21a), indicating that the metal or metal oxide nanoparticles have been completely removed if any. However, the energy dispersive spectroscopy (EDS) mapping of NGM-Co shows a weak but uniform distribution of Co in addition to N and C (Fig. 3.21b). The Co content was detected to be ca. 1.23 wt% by the inductively coupled plasma (ICP) optical emission spectrometer. Additionally, the dark field high-resolution TEM image shows a number of highly dispersed bright spots, which are most likely assigned to the highly dispersed Co atoms survived (Fig. 3.21c, d). Furthermore, XPS was performed to analysize the chemical features of NGMa and NGM-Co in detail. As shown in Fig. 3.22a, the XPS survey spectra reveal high amounts of nitrogen and oxygen (N: 4.42 at.% for NGM-a and 2.31 at.% for NGM-Co; O: 2.04 at.% for NGM-a and 3.22 at.% for NGM-Co), while no obvious Co 2p signal can be detected due to its ultralow content. When decomposing the N 1s XPS spectra, a new peak around 399.3 eV has to be included for NGM-Co in contrast to NGM-a (Fig. 3.22b, c), which can be assigned to the pyridinic N bound

Fig. 3.21 Characterization of NGM-Co. a High-resolution TEM image. The inset represents the XRD pattern without any characteristic peaks of cobalt species. b TEM image and corresponding EDS mapping. c Bright field and d dark field high-resolution TEM images. Reprinted from ref. [39], copyright 2016, with permission from John Wiley and Sons

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Fig. 3.22 a XPS survey spectra for NGM-Co and NGM-a. b N 1s XPS spectrum for NGM-a. c N 1s XPS spectrum for NGM-Co. d C 1s XPS spectrum for NGM-Co and NGM-a. Reprinted from ref. [39], copyright 2016, with permission from John Wiley and Sons

to Co (Co–Nx –C) with ca. 1 eV upshift from pristine pyridinic N (ca. 398.4 eV) [41, 42]. In addition, the C–N shoulder in the C 1s spectrum of NGM-Co obviously shifts to higher binding energy by 0.36 eV (Fig. 3.22d), which is ascribed to the strong electron-withdrawing effect of Co in the Co–Nx –C moieties. Consequently, the electron density of adjacent C atoms is decreased, and is expected to facilitate the adsorption of reaction intermediates and promote the electron transfer for an enhanced kinetics. The above characterization and analysis demonstrate that atomically dispersed metal sites coordinated in 3D porous graphene materials has been successfully realized by changing the precursor composition in templated polymerization methods. The abundant porous edges of graphene provide favorable sites to anchor and disperse metal atoms, which are coordinated with the nitrogen dopants at the edge to form atomic Co–Nx –C moieties. The “defect engineering” derived single-atom sites are expected to greatly improve the electrocatalytic performance of 3D graphene materials.

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3.3.3 Bifunctional ORR/OER Activities The electrocatalytic ORR/OER performance evaluation of NGM-Co and NGM-a samples is the same as above, and the results are shown in Fig. 3.23. As shown in Fig. 3.23a, distinct cathodic peaks are observed in the CV curves for NGM-Co and NGM-a in O2 -saturated electrolyte. Compared with NGM-a, the ORR peak current density for NGM-Co is obviously larger and the peak is positively shifted by ca. 60 mV. LSV curves further reveal the higher ORR activity of NGM-Co with the presence of Co species, exhibiting a much higher limiting current density and a more positive half-wave potential (Fig. 3.23b). Moreover, the Tafel slopes are determined to be 58, 119, and 59 mV dec−1 for NGM-Co, NGM-a, and Pt/C, respectively, indicating fast ORR kinetics on NGM-Co comparable to Pt/C (Fig. 3.23c). Compared with metal-free NGM-a sample, the NGM-Co catalyst exhibits similar morphology, nanostructures, topological defects, and nitrogen doping, merely except for the trace amount of Co–Nx –C species. As the widely recognized active origin for ORR, the nitrogen content of NGM-Co is nearly half that of NGM-a. However, the ORR activity of NGM-Co is significantly superior to NGM-a, highlighting the critical role of Co–Nx –C species on ORR catalytic activity.

Fig. 3.23 a CV curves of NGM-Co and NGM-a recorded in O2 -saturated 0.10 M KOH. b ORR LSV curves for different electrocatalysts and c corresponding Tafel plots. d OER LSV curves for different electrocatalysts. Reprinted from ref. [39], copyright 2016, with permission from John Wiley and Sons

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Additionally, the NGM-Co material also exhibits good OER catalytic activity, with a lower onset potential and much higher current density than Pt/C (Fig. 3.23d). The potential difference between the OER potential for 10.0 mA cm−2 and ORR potential at half-wave is 0.95 V. Notably, the OER LSV curves of NGM-Co and NGM-a are nearly overlapped, with only 20 mV difference on the overpotential for 10.0 mA cm−2 . In contrast to the distinct ORR activities, the similar OER performance of NGM-Co and NGM-a indicates that the optimal active sites for ORR and OER are different. For ORR, the Co–Nx –C species are demonstrated to be remarkably favorable, while for OER, the activity is dominantly determined by heteroatom doping, oxygen-containing functional groups, and topological defects. To elucidate the superior ORR activity and structure-activity relationship of NGM-Co, we conducted electrochemical impedance spectroscopy (EIS) and electrochemically active surface area (ECSA) tests. Taking ORR as an example, the test potential of EIS was set to 0.79 V (versus RHE), the AC frequency range was 0.1–105 Hz, and the amplitude was 5.0 mV. The ECSA was determined by measuring the capacitive current associated with double-layer charging from the scan-rate dependence of CV. This measurement was performed on the same working electrode in a potential window of 1.00–1.05 V versus RHE and scan rates ranging from 10.0 to 100.0 mV s−1 . Then linear fitting of the charging current density differences (j = ja − jc at a potential of 1.025 V versus RHE) against the scan rate was done. The slope is twice the double-layer capacitance C dl , which is used to represent ECSA. Compared with NGM-a, the substantially smaller semicircle in the medium-frequency region for NGM-Co clearly reveals a decreased charge transfer resistance favorable for interface reactions (Fig. 3.24a). In addition, the ECSA of NGM-Co is 16% higher than that of NGM-a (Fig. 3.24b), due to the larger specific surface area, enriched micro/mesopores, and improved hydrophilicity for NGM-Co (128o versus 141o ).

Fig. 3.24 a Nyquist plots of the impedance tested at a potential of 0.79 V versus RHE. b Charging current density differences plotted versus scan rate. Reprinted from ref. [39], copyright 2016, with permission from John Wiley and Sons

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3.4 Summary In order to efficiently regulate the composition, structure, and properties of 3D graphene materials, this Chapter first developed a synthesis strategy based on templated polymerization. 3D porous graphene mesh materials can be achieved by high-temperature polymerization at 950 °C using in situ prepared Mg(OH)2 /MgO nanosheets as the template, gelatinized starch as the carbon source, and melamine as the nitrogen source. The obtained graphene materials are dominantly thinner than double layers, exhibiting large specific surface area (GM: 1655 m2 g−1 ; NGM: 1100 m2 g−1 ), high electrical conductivity (NGM: 120 S m−1 ; GM: 580 S m−1 ), and high total pore volume (NGM: 2.72 cm3 g−1 ; GM: 3.39 cm3 g−1 ). The nitrogen doping content could be as high as 7.60 at.%. Compared with the post-treatment method of GO and templated CVD method, this strategy exhibits huge advantages in the regulation of the specific surface area, porous structure, and heteroatom modification of 3D graphene materials. Owing to the high nitrogen doping content and hierarchical porous structure, NGM exhibits excellent ORR and OER bifunctional catalytic activities. In O2 -saturated 0.10 M KOH, NGM delivers a high ORR limiting current density of 6.41 mA cm−2 , and the potential difference between the OER potential required for 10 mA cm−2 and the half-wave potential of ORR is less than 0.9 V. Additionally, the unique structure and tunability of 3D porous graphene mesh materials provide an efficient platform for the study of electrocatalytic activity mechanisms. Based on it, we systematically explored the roles of nitrogen doping, edge sites, and topological defects. Combining experimental and theoretical studies, the crucial role of topological defects in graphene materials on the ORR and OER bifunctional catalytic activities is revealed, which provides important guidance for the rational design and synthesis of high-performance graphene-based electrocatalysts. In addition to the elucidation of “defect chemistry” in the 3D porous graphene mesh materials, this Chapter further creatively proposed the concept of “defect engineering” to fully use the advantages of the templated polymerization strategy and the unique structure of resulting materials. With the addition of metal precursors during the preparation of precursors, the abundant porous edges of resulting graphene materials provide favorable sites to anchor and disperse metal atoms, enabling the formation of atomic Co–Nx –C moieties. This work further improves the electrocatalytic performance of graphene materials, and expands the synthesis methods, structural composition, and application scenarios of 3D graphene materials [43].

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Chapter 4

Nano-Confined Hybridization and Electrocatalytic Application Based on 3D Mesoporous Graphene Framework

4.1 Introduction Using porous metal oxides as the template, we can efficiently and controllably synthesize a series of 3D porous graphene materials via the CVD of gaseous carbon sources and the templated polymerization of solid carbon sources. The obtained material exhibits the advantages of high specific surface area, high conductivity, 3D interconnected porous scaffold, abundant and adjustable hierarchical pores, and flexible surface chemistry by heteroatom doping. Besides, such synthesis methods are simple, low-cost, and easy to scale up, thus playing an important role in promoting the design, production, and application of high-quality functional graphene materials. We have already conducted a systematic and in-depth study on the catalytic growth mechanism, porosity regulation strategies, and active site design of 3D graphene materials based on oxide templates. According to the structural characteristics of different graphene materials, we also investigated the application of 3D porous graphene materials in Li–S battery cathodes (Chapter 2) and nitrogen-doped/atomic metal coordinated graphene mesh materials in ORR/OER electrocatalysis (Chapter 3). However, considering the unique structural characteristics of 3D porous graphene materials, more application strategies and scenarios are required to be further developed beyond direct use as electrode materials and electrocatalysts. It is crucial to fully exploit the advantages of 3D graphene materials, and customize some unique application strategies, which can make the best use of the properties of graphene and improve its significance in both academic and industrial fields. In view of this, this Chapter takes the 3D mesoporous graphene framework as an example, and creatively proposes the concept of “nanoreactor” to efficiently prepare heterogeneous electrocatalysts. Under the combined effect of anchoring nucleation and confined growth, the obtained active phases were in quantum dot size, uniformly dispersed in the 3D graphene framework, and strongly coupled with graphene, resulting in greatly improved electrocatalytic activities for OER. Additionally, this highly efficient synthetic strategy and unique structural design made it possible to

© Tsinghua University Press 2021 C. Tang, Construction Principles and Controllable Fabrication of 3D Graphene Materials, Springer Theses, https://doi.org/10.1007/978-981-16-0356-3_4

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probe the fundamental electrocatalytic activity mechanisms, thereby ascertaining the host–guest chemistry of the dual-metal hydroxides in OER electrocatalysis.

4.2 Construction Principles of Hierarchical Graphene-Based Hybrid Electrocatalysts Compared with the direct use as electrode materials or electrocatalysts, the hybridization of 3D porous graphene materials with certain inorganic or organic nanomaterials can make full use of the structural characteristics and advantages. As summarized in Fig. 4.1, taking the application in electrocatalytic ORR/OER as an example, the graphene material itself can provide abundant highly active sites such as heteroatom doping, hierarchical structure, edge sites, intrinsic topological defects, or metalnitrogen-carbon moieties. When strongly coupled with other active components, the 3D porous graphene can serve as multifunctional substrates to (1) accelerate the 3D electron and mass transfer, (2) modify the nanoscale incorporation of active components, (3) endow inter-component charge transfer and manipulate the electronic structure of active sites, (4) bring about nano-confinement effects, and (5) assembly into self-supported electrodes, thereby resulting in significantly enhanced performances [1]. A large number of research works have demonstrated that heteroatom-doped or defect-rich graphene materials have high electrocatalytic activity for ORR which is even comparable to commercial Pt/C catalysts; however, the OER activity of graphene itself is far from satisfactory. In addition, nanocarbon materials are easily oxidized at the high potential for OER, resulting in poor durability. Therefore, it is

Fig. 4.1 Application and roles of 3D graphene materials in energy electrocatalysis. Reprinted from ref. [1], copyright 2017, with permission from Elsevier

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crucial to hybridize graphene with other metal compounds to develop highly active and stable OER electrocatalysts. Among various precious-metal-free OER catalysts, nickel-iron layered double hydroxide (NiFe LDH) is one of the most promising candidates [2–4]. At the beginning of the last century, scientists discovered that metallic nickel and its oxides have excellent electrocatalytic activity for OER, and Fe doping can further improve the performance by orders of magnitude [5, 6]. The incorporation of Fe regulates the physical structure and electronic structure of α-Ni(OH)2 (before reaction) and γNiOOH (during reaction), which can optimize the adsorption and desorption energy of intermediates, thereby rendering NiFe LDH highly active for OER. By optimizing the metal compositions and crystal structures, the OER activity of NiFe LDH-based materials has been significantly improved [7–10]. However, suitable nanocarbon materials or metal current collectors are still needed to ensure the electron transfer pathways due to the poor conductivity of NiFe LDH itself [11]. Due to the various types of nanocarbon materials and complex interaction between NiFe LDH and nanocarbon scaffold, the resulting hybrid structures can be classified according to the structural features of nanocarbon materials, including zerodimensional (0D) carbon quantum dots (CQDs, Fig. 4.2a), one-dimensional (1D) carbon nanotubes (CNTs, Fig. 4.2b), 2D graphene nanosheets, and 3D porous graphene frameworks (Fig. 4.2) [2]. The conventional 2D graphene nanosheets can

Fig. 4.2 Schematic of hierarchical NiFe LDH/nanocarbon hybrid electrocatalysts with different nanostructures classified according to the dimensionality of nanocarbon. a 0D CQDs. b 1D CNTs. c 2D graphene nanosheets. d 3D porous graphene materials. Reprinted from ref. [2], copyright 2016, with permission from John Wiley and Sons

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Fig. 4.3 a TEM image and b schematic illustration of the nNiFe LDH/NGF electrocatalyst. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

serve as the substrate for in situ growth of NiFe LDH nanosheets, or self-assemble with NiFe LDH nanosheets into layer-by-layer heterostructures via post-processing (Fig. 4.2c). In contrast, 3D porous graphene materials can interact with NiFe LDH nanosheets on “micro-meso-macro” scales to realize the full regulation and optimization in nano size, crystal structure, interface coupling, active site exposure, and 3D distribution, thus making the best use of the unique characteristics of graphene and the intrinsic activity of NiFe LDH. Based on the above discussion of the construction principles for 3D graphenebased hybrid electrocatalysts, we have proposed a “nano-confined hybridization” strategy to fully demonstrate the properties and advantages of 3D porous graphene materials. As shown in Fig. 4.3a, the 3D mesoporous graphene framework material prepared by MgO-templated CVD exhibits a huge specific surface area and abundant mesopores. If the growth of NiFe LDH nanosheets can be spatially confined in the mesopores (Fig. 4.3b), it is expected to obtain highly dispersed active phases in quantum dot size (matching the size of mesopores of graphene), which are stably confined in the framework of graphene and easily accessible. Meanwhile, the strong interface interaction between graphene and the active phases enables the fast charge transfer from the active sites in NiFe LDH to the graphene substrate. This application strategy allows the key characteristics of 3D graphene materials to be fully utilized in both the synthesis and application of nanocomposites, opening up a new direction for the application of 3D graphene materials.

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4.3 Nano-Confined Electrocatalysts Based on 3D Mesoporous Graphene Framework 4.3.1 Material Synthesis and Characterization The 3D graphene material prepared by the templated CVD method has a high degree of graphitization, a low content of defects, and hydrophobic surface; however, the synthesis of NiFe LDH is generally via liquid-phase co-precipitation or hydrothermal reaction. In order to ensure that the precursor of NiFe LDH nucleates and grows selectively on the graphene surface instead of in the liquid phase, we introduced ammonia gas during the CVD process to prepare a nitrogen-doped 3D mesoporous graphene framework material, denoted as NGF. As shown in Fig. 4.4a, nitrogen doping does not change the structure and morphology of the obtained graphene material, which still maintains a regular mesoporous framework. The significantly higher I D /I G ratio in the Raman spectrum (2.07) compared with NGM and NGM-Co samples obtained by the templated polymerization method indicates that there are more defect sites in the NGF structure due to the topological defects at the corners of the 3D mesopores, nitrogen doping, and surface oxygen-containing functional groups. XPS results show

Fig. 4.4 a TEM image of NGF. b–d TEM images of nNiFe LDH/NGF. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

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that the nitrogen content of NGF is 1.72 at.%, and the oxygen content is 4.72 at.%, which is beneficial for improving the surface hydrophilicity of graphene materials and promoting the nano-confined growth of NiFe LDH. Using the NGF as both the substrate and template, we scrupulously designed and facilely fabricated a novel composite based on graphene and NiFe LDH (nNiFe LDH/NGF) via a urea-assisted precipitation. Specifically, 25.0 mg of NGF was dispersed into 40.0 mL of N-methylpyrrolidone (NMP) by ultrasonic treatment for 30 min. 735.0 mg of Ni(NO3 )3 ·6H2 O, 340.0 mg of Fe(NO3 )3 ·9H2 O, and 9.00 g of urea were dissolved in 50.0 mL of deionized water. Then the as-obtained solution was mixed with the NGF/NMP dispersion and refluxed at 100 °C under continuous magnetic stirring for 6.0 h. The as-prepared nNiFe LDH/NGF was obtained after the product was filtered, washed and freeze-dried. bNiFe LDH was fabricated under otherwise identical conditions without NGF, and nNiFe LDH/GF was fabricated with un-doped 3D mesoporous graphene framework (GF) instead of NGF. The bNiFe LDH + NGF mixture was prepared for comparison, with the same mass ratio of nNiFe LDH/NGF, simply by mixing bNiFe LDH and NGF under sufficient grind. The TEM image reveals that abundant NiFe LDH nanocrystals are selectively and uniformly confined into the mesoporous graphene scaffold without any obvious aggregations (Fig. 4.4b, c). High-resolution TEM image in Fig. 4.4d shows that the NiFe LDH nanoplates exhibit a size of typically ca. 5 nm, which are divided from each other and encircled with graphene layers (highlighted by the dotted hexagon). In addition, a set of lattice fringes with a spacing of 0.25 nm can be clearly observed, which are assigned to the NiFe LDH (012) crystal plane. These results confirm the successful synthesis of quantum dot-sized NiFe LDH nanosheets with high crystalline confined in the mesopores of NGF. To the best of our knowledge, the in situ grown nano-sized NiFe LDH in nNiFe LDH/NGF exhibits the smallest size and best dispersion. The XRD profiles further confirm the well crystallized nNiFe LDH in the asobtained composites. As shown in Fig. 4.5a, nNiFe LDH/NGF and bNiFe LDH

Fig. 4.5 a XRD spectra of nNiFe LDH/NGF, bNiFe LDH, and physical mixture of bNiFe LDH and NGF. b Raman spectra of nNiFe LDH/NGF and control materials obtained at 633 nm. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

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Table 4.1 Summary of material compositions [12] Sample and method

bNiFe LDH

Relative amounts of different elements (at.%) Ni

Fe

N

ICP

39.18

14.75

–

2.66

XPS

23.13

6.69

8.99

3.46

EDS

39.00

15.76

–

2.47

TGA bNiFe LDH + NGF

–

–

–

–

2.41

6.20

7.58

3.05

12.58

4.00

2.62

39.06

16.22

XPS

18.89

EDS

32.94 –

–

–

–

2.95

37.02

XPS

14.85

5.04

5.66

2.95

EDS

27.26

10.53

2.54

2.59 –

XPS

12.57

–

ICP

TGA NGF

–

ICP

TGA nNiFe LDH/NGF

Atom ratio of Ni/Fe

–

–

–

C

O

N

4.72

1.72

93.56

Mass ratio of LDH/NGF

7.7

7.2

share the same diffraction peaks assigned to hydrotalcite materials, indicating that the presence of NGF substrate can regulate the nucleation growth and nanostructure of NiFe LDH, but not the crystal structures. Compared with the pristine graphene materials, the Raman spectra of the composites exhibit two new characteristic bands near 460 cm–1 and 545 cm–1 (Fig. 4.5b), which are assigned to the Fe3+ /Ni2+ -O-Ni2+ and Fe3+ -O-Fe3+ linkage bands, respectively [13]. The atomic ratio of Ni to Fe of nNiFe LDH/NGF is determined to be about 2.9 by ICP, XPS, and EDS (Table 4.1), which is close to the feed ratio and is also at the optimal Ni/Fe ratio range for OER catalysts [10, 13]. This ratio is slightly higher than the Ni/Fe atomic ratio of bNiFe LDH without graphene substrate, which may be due to the different adsorption capacity of oxygen-containing functional groups and defect sites on the surface of NGF to Ni2+ and Fe3+ ions.

4.3.2 Mechanism of Nano-Confined Hybridization Through the morphology comparison of the composites obtained by different synthesis methods and precursors, we try to explore the key factors and synthesis mechanisms of nano-confined hybridization based on 3D mesoporous graphene framework materials. If replacing NGF with GF as the substrate for NiFe LDH deposition, the morphology of obtained nNiFe LDH/GF is similar to that of nNiFe

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4 Nano-Confined Hybridization and Electrocatalytic …

LDH/NGF with a large number of NiFe LDH quantum dots highly dispersed in the graphene scaffold, but some large-sized NiFe LDH aggregations also appear on the surface (Fig. 4.6a, b). On the one hand, the mesoporous structure of the 3D mesoporous graphene framework material can provide an ideal space for the confined growth of NiFe LDH, thereby obtaining a highly dispersed and uniform active phase. On the other hand, the nitrogen doping and defect sites of the graphene template are very important for the adsorption, anchoring, and nucleation of the metal precursors, which can greatly hinder the nucleation and growth of NiFe LDH in the liquid phase. Assuming that the weight loss ratio for LDH is consistent in the bulk LDH and composites, the mass ratio of NiFe LDH and graphene can be calculated from the TGA results of different samples under the oxygen atmosphere. As listed in Table 4.1, the mass ratio of NiFe LDH to NGF of the nNiFe LDH/NGF sample is 7.2, and that of the bNiFe LDH + NGF sample is 7.7. Although the graphene materials used are the same and the metal ratios are very close for the two samples, the resulting morphologies are quite different. A large number of NiFe LDH particles with a size of hundreds of nanometers are attached on the surface of NGF (Fig. 4.6c), and almost no NiFe LDH nanosheets can be observed to be dispersed or confined in the graphene scaffold (Fig. 4.6d). It reveals that the NiFe LDH obtained by template-free liquid

Fig. 4.6 a SEM and b TEM images of nNiFe LDH/GF. c SEM and d TEM images of bNiFe LDH + NGF. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

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synthesis is in a much larger size, and the physical mixing cannot realize the effective hybridization between NiFe LDH and 3D mesoporous graphene materials. Furthermore, we evaluated the change of the porous structures for different composites. As shown in Fig. 4.7a, the nitrogen sorption isotherm of the NGF sample is a typical type-IV isotherm, and the larger hysteresis loop reveals abundant mesopores in this sample. For both in situ hybridization and physical mixing, the introduction of NiFe LDH with higher density and lower specific surface area significantly reduces the absolute specific adsorption capacity (Fig. 4.7a). In order to compare the changes of the porous structures before and after hybridization more clearly, we conducted a detailed analysis of the pore size distribution (Fig. 4.7b). The micropores and small-sized mesopores between 0 and 5 nm come from irregular wrinkles and defects in the NGF structure; the mesopores between 5 and 10 nm in NGF are generated by replication of the mesoporous MgO quasi-single-crystals during CVD; the mesopores larger than 10 nm are assigned to the piled pores between graphene nanosheets. As compared in Fig. 4.7b, the mesopores between 5 and 10 nm are obviously decreased for the in situ hybridized composites more than that for physical mixtures, which is rationalized by the occupation of nNiFe LDH in the in-plane pits of NGF. In contrast, the piled pores larger than 10 nm of bNiFe LDH + NGF are dramatically increased by 33.8%, which may originate from the LDH aggregation on and among graphene nanosheets. The above-mentioned characterizations suggest that the key to nano-confined hybridization based on the 3D mesoporous graphene framework material lies in the abundant adsorption sites on the graphene surface, the open mesoporous structure, and the in situ growth in liquid, all of which are indispensable. The highly active sites introduced by nitrogen doping enable metal ions to be selectively adsorbed on the surface of graphene and anchored to nucleation, reducing the ratio of liquid phase nucleation and growth. The open mesoporous structure promotes the continuous growth of NiFe LDH after its nucleation at the solid-liquid interface, and when it reaches the size of the mesopore, it stops growing and is stably confined inside. The in situ growth in liquid avoids the stacking of graphene and ensures the full dispersion

Fig. 4.7 a The N2 sorption isotherms of nNiFe LDH/NGF and control samples. b Pore size distribution comparison of NGF with in situ grown LDH and physically mixed LDH. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

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and uniform contact with the precursors of NiFe LDH, thus achieving the efficient hybridization of nano-confined active phases with in quantum dot size and conductive framework with high specific surface areas. This work provides new strategies for the design and synthesis of composite electrocatalysts and the customized application of 3D porous graphene materials.

4.3.3 Electrocatalytic OER Activity As mentioned above, the unique structure of nNiFe LDH/NGF fully demonstrates the respective advantages of the components and also affect each other, which is expected to bring about extraordinary electrocatalytic performance. Therefore, we studied its OER performance in O2 -saturated 0.10 M KOH electrolyte. The specific evaluation methods are the same as the previous chapters. Figure 4.8a presents the OER LSV curves for different samples. The redox peaks around 0.2 V overpotential

Fig. 4.8 a LSV curves obtained in O2 -saturated 0.10 M KOH for different electrocatalysts. b RRDE measurement of nNiFe LDH/NGF, with a constant overpotential of 350 mV on disk electrode and a CV scanning on ring. c LSV tests on the ring electrode. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

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99

are assigned to the Ni2+ /Ni3+ redox process. The nNiFe LDH/NGF hybrid delivers a similar onset potential but substantially higher current density than all samples, showing its outstanding activity. The overpotential required for a current density of 10 mA cm–2 , which corresponds to 10% efficient solar water-splitting devices, is adopted as a critical figure of merit for OER catalysts. As highlighted by the horizontal dotted line in Fig. 4.8a, the nNiFe LDH/NGF exhibits a low overpotential for 10 mA cm–2 at ~337 mV, which dramatically surpasses both components and the physical mixture, and also outperforms the commercial Ir/C catalyst by 73 mV. It is contributed from the synergistic and strong couple effects due to the novel nano-confined hybridization and unique structural features. Since the 3D porous graphene material is unstable at high potentials, and the oxidation of carbon may occur to contribute a certain amount of current, we investigated the origin of the LSV current for nNiFe LDH/NGF by RRDE technique. The catalyst was loaded on the disk electrode, and the potential was set at a constant value of 350 mV overpotential to catalyze OER reaction (current density is larger than 10 mA cm–2 ). CV or LSV tests were conducted on the ring electrode to identify the reaction products on the disk electrode. As shown in Fig. 4.8b, there is no oxidation peaks for any organic compounds observed in the CV, indicating the electrolyte was free of carbon-containing products from the oxidation of nNiFe LDH/NGF electrocatalysts. In the LSV curve recorded on the Pt ring (Fig. 4.8c), there is an obvious reduction current, indicating that the ORR reaction occurs. The O2 produced by OER on the disk electrode diffuse to the ring electrode and then is reduced by Pt to generate a current. Assuming that the disk current all comes from OER, the ratio of I ring /I disk will represent the current collection efficiency of the ring electrode. At a potential around 0.1 V (versus RHE), the I ring /I disk reaches 0.24, which is very close to the theoretical value of 0.26 for the RRDE electrode. It demonstrates that most of the disk current is assigned to OER on the nNiFe LDH/NGF electrocatalyst. Generally, Tafel plots can be derived from the LSV curves to evaluate the kinetic characteristics of the reaction and the activity of catalysts. However, due to the serious overlap of the OER current and the oxidation current of Ni2+ /Ni3+ in the Tafel zone, the Tafel slopes cannot be obtained directly from the LSV curves for NiFe LDH-based catalysts (Fig. 4.8). In order to decouple them, we conducted the chronoamperometry tests at different constant potentials (slightly higher than the Ni2+ /Ni3+ redox potential) to obtain the real OER current. As shown in Fig. 4.9a, the initial current value at different potentials is very large due to the influence of the Ni2+ /Ni3+ oxidation, but it drops rapidly and stabilizes within a few seconds. The stabilized value is used as the OER current at each potential to draw the corresponding Tafel plots (Fig. 4.9b). The nNiFe LDH/NGF electrocatalyst exhibits the lowest Tafel slope of 45 mV dec–1 , compared with those of bNiFe LDH + NGF (52 mV dec–1 ), nNiFe LDH/GF (58 mV dec–1 ), bNiFe LDH (62 mV dec–1 ), and Ir/C (54 mV dec–1 ) (Fig. 4.9c). Consequently, the lower Tafel slope of nNiFe LDH/NGF renders greatly boosted current densities at relatively higher overpotentials, such as 60 mA cm–2 at the overpotential of 480 mV, which is nearly four-fold that of Ir/C catalysts (Fig. 4.8a). Taking the overpotential required for 10 mA cm−2 and the Tafel slope as indicators of the catalytic activity and kinetic characteristics respectively, we comprehensively

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Fig. 4.9 a Current densities versus time obtained at specific potentials slightly above the Ni2+ /Ni3+ redox potential to decouple the redox and OER currents. b Tafel plot of nNiFe LDH/NGF measured from (a). c Tafel plots for different samples. d OER performance comparision of nNiFe LDH/NGF and other samples in 0.10 M KOH. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

compares the nNiFe LDH/NGF catalyst with other counterparts and references, as presented in Fig. 4.9d and Table 4.2. It is revealed that the novel nNiFe LDH/NGF hybrid exhibits a lower overpotential and Tafel slope than most reported graphene and non-precious metal-based materials in 0.10 M KOH, which is among the most active candidates for OER.

4.3.4 Structure-Property Relation and Activity Mechanism For heterogeneous electrocatalytic reactions, such as OER, the electrical conductivity, structural porosity, active sites, and interfacial coupling of electrocatalysts all play important roles on the resulting performances. In the case of composite electrocatalysts, such as nNiFe LDH/NGF, the intrinsic merits of each component and the strong interaction between each other both contribute greatly to the outstanding activity.

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Table 4.2 OER performance comparison of some graphene and non-precious metal-based materials in 0.10 M KOH [12] Samples

Overpotential @10 mA cm–2 (mV)

Tafel slope (mV dec–1 )

Reference This work

n-NiFe LDH/NGF

337

45

b-NiFe LDH + NGF

392

52

n-NiFe LDH/GF

372

58

NGF

615

302

b-NiFe LDH

459

62

Ir/C

410

54

C3 N4 -CNT

383

83

[14]

Ni-NG

~340

188.6

[15]

PNG-NiCo

~419

156

[16]

N-GSH

~400

83

[17] [18]

m-NiFe/CNx

360

59.1

G/Co3 O4

359

67

[19]

Ni@NiCoOH

460

65

[20]

Au@Co3 O4

~383

60

[21]

NG-CNT

~410

141

[22]

Co3 O4 -CNA

~290

70

[23]

MWCNT/Ni(OH)2

474

87

[24]

NiFe LDH/CNT

~300

35

[11]

CQDs/NiFe-LDH

305

35

[25]

On one hand, the NGF affords a highly conductive mesoporous scaffold with moderate active sites. As shown in Fig. 4.10a, the conductivity of NGF is determined to be 5400 S m–1 using the four-probe technique, while the bNiFe LDH itself is nearly non-conductive (~10–8 S m–1 ). After in situ hybridization, the obtained nNiFe LDH/NGF exhibits a considerable conductivity of 145 S m–1 , thus ensuring a smooth electron pathway. In addition, the porous graphene-based hybrid even after the decoration of LDH can deliver a considerable specific surface area of 240 m2 g–1 and a total pore volume of 0.30 cm3 g–1 , which accelerate the electrolyte ion insertion/extraction and gas diffusion (Fig. 4.7). Furthermore, the nitrogen incorporation into the graphene framework not only provides abundant adsorption and nucleation sites of metal precursors, but also tailors the electronic structure of the adjacent carbon atoms towards facilitated chemisorption of intermediates and additional activity for OER. As shown in Figs. 4.10b and c, the OER activity of nNiFe LDH/NGF and NGF is significantly improved compared to their un-doped counterparts. On the other hand, the introduction of NiFe LDH effectively moderates the surface characteristics of obtained catalysts, bringing about a large number of highly active sites. In spite of abundant nitrogen dopants and oxygen-containing functional groups, the hydrophilicity of NGF is not high enough, but it can be significantly improved

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Fig. 4.10 a Electrical conductivity comparison of NGF, nNiFe LDH/NGF and bNiFe LDH. b, c LSV curves of nNiFe LDH/NGF and NGF and their un-doped counterparts. d Contact angels of nNiFe LDH/NGF compared with NGF. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

after hybridization with NiFe LDH to deliver a contact angle of 16o (Fig. 4.10d). It is very beneficial for the affinity between solid catalysts and liquid electrolytes with enhanced permeation of electrolyte ions. Notably, the Tafel slope of bNiFe LDH itself is only 62 mV dec–1 , which is comparable to most of previously reported OER catalysts (Table 4.2), revealing the high intrinsic activity of NiFe LDH. However, the mechanism of the high intrinsic activity of NiFe LDH has not been thoroughly explored. Some works proposed that Fe3+ incorporation exerts a partial-chargetransfer activation effect on the surrounding Ni3+ sites, thereby altering their average oxidation state towards increased activities [13, 26]. Others hypothesize that the substitution of Fe3+ generates more favorable active sites themselves instead of the activation of Ni3+ sites [10, 27], owing to the near optimal adsorption of reaction intermediates induced by the shorter Fe–O bonds in the edge-sharing [MO6 ] octahedras [10]. Although NGF and NiFe LDH both have excellent properties and play important roles on the resulting high OER performance, the activity of nNiFe LDH/GF and bNiFe LDH + NGF with the same components is much inferior to that of nNiFe LDH/NGF. It suggests that the nNiFe LDH/NGF sample possesses unique structural effects and strong coupled interface, which may further enhance the catalytic performance. Therefore, we further investigated the in-depth influence and mechanism of the OER activity derived from the nano-confined hybridization. As shown

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103

in Fig. 4.11a, the Ni2+ /Ni3+ redox behaviors for different samples vary considerably, the integrated area of which is assigned to the extent of the Ni(OH)2 /NiOOH transformation before OER. When normalizing the integrated peak area on the basis of bNiFe LDHs, we find that the ratio of oxidized Ni is dramatically increased for nNiFe LDH/NGF hybrid by 20 times and also several times increased for other composites, indicating an enhanced Ni(OH)2 /NiOOH transformation of NiFe LDH. It has been widely demonstrated that the resulting NiOOH phase after Ni2+ /Ni3+ redox is crucial to the active sites of NiFe LDH for OER [10, 13], which is consistent with the negative correlation between the overpotential for 10 mA cm–2 and the extent of Ni2+ /Ni3+ redox (Figs. 4.8a and 4.11b). Besides, the high resolution Ni 2p spectrum of nNiFe LDH/NGF shifts to higher binding energy by 0.7 eV compared with bNiFe LDH (Fig. 4.11c), indicating a modified local electronic structure of Ni cations contributed from the strong binding between nNiFe LDH and NGF, i.e. C-N…Ni-O. Therefore, it can be concluded that the strong couple between NiFe LDH and the graphene framework regulates the electrochemical redox behaviors of Ni cations and the resulting active phases for OER with better catalytic activities. In addition, the strong interfacial bonding facilitates a rapid charge transfer between active phases and conductive substrates as evidenced by the lower Tafel slopes.

Fig. 4.11 a Enlargement of the Ni2+ /Ni3+ redox range in LSV curves. b The normalized transformation ratio of Ni2+ on the basis of bNiFe LDHs. c High-resolution Ni 2p spectra. d Charging current density differences plotted versus scan rate. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

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In fact, due to the nitrogen-doping-induced functional groups on graphene, the physical mixture bNiFe LDH + NGF also exhibits a considerable interfacial coupling, as confirmed by the shifted Ni 2p spectrum in Fig. 4.11c. However, the OER activity is obviously inferior to that of in situ hybridized nNiFe LDH/NGF, which can be elucidated by ECSA of the materials. As shown in Fig. 4.11d, the ECSA of nNiFe LDH/NGF is 60% higher than that of bNiFe LDH + NGF. The smaller size and improved dispersion of NiFe LDH in nNiFe LDH/NGF owing to spatial confinement offers fully exposed and easily accessible active sites. Furthermore, this unique nano-confined hybridization enables better catalytic stability. Figure 4.12a records the chronoamperometric responses of different samples at an overpotential of 350 mV. The current density of nNiFe LDH/NGF can be stabilized for more than 12000 s, and the initial increase may be ascribed to the improved active site utilization caused by gradual infiltration of electrolyte, while that of bNiFe LDH + NGF is gradually decreased to 77% of the initial value within 10000 s. The LSV curves of nNiFe LDH/NGF before and after the stability test almost overlap with each other (Fig. 4.12b). The TEM image and XPS spectrum of the catalyst after long-time cycle reveal that the spatially-confined morphology and physical structures are well

Fig. 4.12 a Chronoamperometric response at an overpotential of 350 mV. b LSV curves at the first test and after 12000 s chronoamperometric test at an overpotential of 350 mV. c TEM image showing the preservation of spatially-confined morphology after long-term cycle. d High resolution Ni 2p spectra with the same result as the fresh sample. Reprinted from ref. [12], copyright 2015, with permission from John Wiley and Sons

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105

preserved (Fig. 4.12c, d), suggesting a superior structure stability and thereby an enhanced catalytic durability. In summary, we employed the 3D mesoporous graphene framework material as the nanoreactor, and realized the efficient hybridization of active phases in quantum dot size and conductive scaffold with high specific surface area, leading to a novel NiFe LDH/graphene hybrid with excellent activity and stability for OER. This naonconfined design and fabrication strategy opens up new avenues and sheds light on a new family of advanced nano-architectured materials, and facilitates the smart hybridization of excellent components towards specific applications.

4.4 Intrinsic Activity Mechanism Study Based on Nano-Confined Electrocatalysts The nano-confined hybridization strategy based on 3D mesoporous graphene framework materials not only brings about a new design concept of nanocomposites and a series of high-performance electrocatalysts, but more importantly, it also provides a very unique and effective material platform for the investigation of intrinsic activity mechanisms. Taking nNiFe LDH/NGF as an example, the size of the active phase in this material is reduced to the quantum dot level due to defect-anchored nucleation and spatially confined growth, thus exposing the active sites to the greatest extent. In addition, the strong interfacial coupling between non-conductive active phases and highly conductive graphene substrates ensures the rapid electron transfer during reaction and realizes the full demonstration of the intrinsic activity. Furthermore, the nano-confined structural characteristic makes the structure and performance of the catalyst highly stable during long-term operation, and also provides opportunities to facilely tune the composition of active phases. In other words, the nano-confined hybrid based on 3D mesoporous graphene framework materials not only ensures the full exposure and utilization of the intrinsic active sites of the incorporated active phases, but also allows the composition to be adjusted with the same nanostructure, thereby making it ideal to study the intrinsic activity mechanisms of the active phases. Herein, we explore the catalytic mechanism of NiFe LDH based on a series of nNiFe LDH/NGF materials.

4.4.1 Nano-Confined Hybridization of Various Active Phases The above discussion has revealed that the key to nano-confined hybridization based on the 3D mesoporous graphene framework material lies in the abundant adsorption sites, the open mesoporous structure, and the in situ growth in liquid. The adsorption sites are mainly derived from nitrogen doping and topological defects. However, nitrogen doping itself will make a non-negligible contribution to the electrocatalytic

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Table 4.3 The compositions of G/NiFex materials based on XPS results [28] Sample Dosage for synthesis (mg)

G/Ni

G/NiFe0.30 G/NiFe0.56 G/Fe

Ni(NO3 )2 ·6H2 O 980.0

980.0

245.0

Fe(NO3 )3 ·9H2 O

340.0

1021.0

0

0 1361.0

Mole ratio of different elements C (at.%) O

30.10

35.17

35.33

52.93

44.74

38.50

44.22

34.77

Ni

25.16

18.42

9.02

0.00

Fe

0.00

7.91

11.43

12.31

0.00

0.30

0.56

1.00

Fe content (mole ratio of Fe/(Ni + Fe))

activity, and can also adjust the electronic structure and intrinsic activity of NiFe LDH via the strong interfacial coupling effect. To make the mechanism study more rational and accurate, we replaced the NGF with the mildly oxidized mesoporous graphene framework for in situ nano-confined hybridization, which can avoid the interference of nitrogen doping on the results. The mildly oxidized graphene was fabricated by a modified Hummers’ method. Specifically, 100.0 mg of the un-doped 3D mesoporous graphene framework (GF) was firstly dispersed in 25.0 mL of concentrated H2 SO4 . Then, the mixture was transferred into an ice-water bath to maintain its temperature under 10 °C and stirred for 4.0 h. After that, 25.0 mL of 10% dilute H2 SO4 was gradually added in 1.0 h, followed by slow addition of 75.0 mL of deionized water in another 2.0 h. To end the reaction, 2.0 mL of H2 O2 was added into the solution until no gas bubble was released. The mildly oxidized sample was finally obtained after filtering, washing, and freezedrying. The process for nano-confined hybridization was the same as described above, except the replacement of NGF with mildly oxidized GF. In order to synthesize NiFe LDH with different compositions, especially different Fe contents, the amount of Ni(NO3 )2 ·6H2 O and Fe(NO3 )3 ·9H2 O in the precursors is slightly changed as shown in Table 4.3. The resulting samples are denoted as G/NiFex , where x is the true molar ratio of Fe/(Ni + Fe) determined by XPS. The synthesis conditions and active phase composition of the G/NiFe0.30 sample are closest to those of nNiFe LDH/NGF, and the obtained nanostructure of them are almost the same. As expected, NiFe0.30 hydroxide nanosheets are uniformly dispersed and confined in the mesopores of the graphene framework (Fig. 4.13a– c), indicating that the mildly oxidized GF can also provide sufficient adsorption sites for nano-confined hybridization. For the G/Ni sample, however, much larger Ni hydroxide nanosheets grow outside the graphene with few confined in pores (Fig. 4.13d). With the incorporation of Fe, the resulting nanostructure and hydroxide decoration style become similar to that of G/NiFe0.30 , showing spatially confined hydroxide nanosheets in graphene with a size of ~5 nm and no obvious agglomerates (Fig. 4.13e, f). The results reveal that based on the mildly oxidized graphene substrates, we have successfully realized the nano-confined hybridization of NiFe hydroxide active phases. The obtained electrocatalysts have full exposure of active sites,

4.4 Intrinsic Activity Mechanism Study Based …

107

Fig. 4.13 a SEM image of G/NiFe0.30 . TEM images of b, c G/NiFe0.30 , d G/Ni, e G/NiFe0.56 , f G/Fe. Reprinted from ref. [28], copyright 2016, with permission from The Royal Society of Chemistry

enhanced electrical conductivity from graphene, and tuned compositions with similar nanostructure, rendering a promising material platform for the mechanism elucidation.

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4.4.2 Structure Transformation and Characterization of Confined Active Phases In order to compare the crystal structure and electronic structure of active phases in different samples, we systematically conducted XRD, TEM, XPS, and Fourier transform infrared spectroscopy (FT-IR) characterizations. As shown in Fig. 4.14a, G/Ni and G/NiFe0.30 share similar XRD patterns. The (003) peak is assigned to the intercalation of water and reactive ions in the α-Ni(OH)2 structure (the same as LDH), which is considered to be the most efficient phase of Ni(OH)2 for OER catalysis. The (012) diffraction peak at 2θ = 33.72o corresponds to a d-space of 0.26 nm, in consistence with the lattice fringes in the TEM images (Fig. 4.15a, b). The asymmetric nature of (012) peak suggests the formation of turbostratic αNi(OH)2 lattice. In the case of higher Fe contents, the XRD patterns are assigned to FeO(OH) rather than lamellar hydroxide structures (Fig. 4.14b). The hydroxide in the G/Fe sample matches well with a polymorph of γ-FeO(OH) and δ-FeO(OH), while that in G/NiFe0.56 corresponds to Fe0.67 Ni0.33 O(OH), suggesting the incorporation of Ni into the FeO(OH) framework. The weak intensity and high noise-signal ratio of the XRD patterns of G/NiFe0.56 and G/Fe indicate their relatively poor crystallization. The lattice fringes in the high-resolution TEM images for G/NiFe0.56 and G/Fe exhibit a spacing of 0.25 nm (Fig. 4.15c, d), which is assigned to the (100) lattice planes of FeO(OH). As summarized in Fig. 4.16, these results reveal that with the increase of Fe contents, the resulting NiFe phases are transformed from well-crystallized hydroxides (α-Ni(OH)2 and NiFe LDH) to amorphous oxyhydroxides (Ni substituted FeO(OH) or FeO(OH) polymorph, dominantly δ-FeO(OH)). This change in the crystal structure of active phases is believed to be the crucial factor of corresponding electrocatalytic activities.

Fig. 4.14 XRD spectra of different samples. a G/Ni and G/NiFe0.30 . b G/NiFe0.56 and G/Fe. Reprinted from ref. [28], copyright 2016, with permission from The Royal Society of Chemistry

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109

Fig. 4.15 TEM images of different samples. a G/Ni. b G/NiFe0.30 . c G/NiFe0.56 . d G/Fe. Reprinted from ref. [28], copyright 2016, with permission from The Royal Society of Chemistry

Fig. 4.16 The model structures showing the physical structure evolution. With the Fe content increased, the resultant materials are transformed from well-crystallized hydroxides: a γ-NiO(OH) derived from α-Ni(OH)2 , and b NiFe LDHs, to amorphous oxyhydroxides: c Ni substituted FeO(OH), and d FeO(OH) polymorph, dominantly δ-FeO(OH). Reprinted from ref. [28], copyright 2016, with permission from The Royal Society of Chemistry

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Fig. 4.17 a High resolution Ni 2p spectra of different samples. b High resolution Fe 2p spectra of different samples. Reprinted from ref. [28], copyright 2016, with permission from The Royal Society of Chemistry

The XPS results for different samples are listed in Table 4.3. No distinct spectrum of Fe species (or Ni species) can be identified from the G/Ni (or G/Fe) sample. As shown in Fig. 4.17, the Ni 2p and Fe 2p spectra of NiFe bimetallic (oxy)hydroxides are both slightly shifted compared with the monometallic hydroxides, indicating variable electronic structures altered by the Fe/Ni ratio. For the G/NiFe0.30 sample, Ni 2p1/2 and Ni 2p3/2 peaks are upshifted by ~0.9 eV after the Fe incorporation, while the Fe 2p1/2 and Fe 2p3/2 peaks are almost the same compared with G/Ni. For the G/NiFe0.56 sample, the upshift of Ni 2p peaks is decreased to ~0.4 eV, but an obvious downshift by ~0.4 eV for Fe 2p peaks is also detected. This difference in peak shift is believed to be contributed from the guest–host substitution and the bond length variation, which are regulated by the Fe contents. For a low Fe content, Fe substitutes into the γ-NiO(OH) lattice and expands the Ni–O bond in the γ-NiO(OH) host; while for a high Fe content, Ni substitutes into the FeO(OH) framework and contracts the Fe–O length in the FeO(OH) host. Furthermore, FT-IR was performed to characterize the chemical bonds and functional groups of different samples in detail. As shown in Fig. 4.18, several characteristic FT-IR peaks for hydroxides are detected in addition to the peaks for the oxidized graphene substrate. The peaks around 1380 and 1470 cm–1 are attributed to the interlayer CO3 2– and the O–H bend of lattice OH in metal (oxy)hydroxides [29, 30]. The wide band around 1630 cm–1 is ascribed to the O–H bend of interlayer or free H2 O. The enhanced absorption band near 3430 cm–1 is assigned to the O–H stretching vibration of interlayer or free H2 O [29]. It is noteworthy that a narrow peak appears at 3645 cm–1 only for G/Ni and G/NiFe0.30 samples, which is assigned to the O–H stretch of the brucite-like structure [31]. From the above characterizations in morphology, crystal structures, compositions, and functional groups, we can rationally conclude that the Fe/Ni ratio not only tunes the morphology and decoration style of active phases in the resulting composites, but also more importantly, regulates the physical and electronic structures. These diversities arisen from the substitution of guest metals into the host

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Fig. 4.18 FT-IR spectra of G/NiFe materials and the oxidized graphene substrate. Reprinted from ref. [28], copyright 2016, with permission from The Royal Society of Chemistry

(oxy)hydroxide framework will contribute to different properties and performances in OER electrocatalysis.

4.4.3 Electrocatalytic OER Activity The OER catalytic activity was evaluated in O2 -saturated 0.10 M KOH, and the experimental details are the same as those described in the previous chapters. The typical LSV curves show that the anodic current densities of all G/NiFex catalysts increase as the potential becomes more positive, but the Ni2+ /Ni3+ oxidation peaks of G/NiFe0.56 and G/Fe near 0.2 V are significantly smaller than those of G/Fe and G/NiFe0.30 , indicating the difference in the crystal structures of the active phase and the catalytic activity mechanisms. The moderate incorporation of Fe substantially decreases the onset overpotential from 361 mV of G/Ni to 303 mV of G/NiFe0.30 , while the further increase of Fe content gradually increases the onset overpotential to 430 mV of G/Fe. The overpotential required for 10.0 mA cm–2 are 567, 372, 390, and 482 mV, for G/Ni, G/NiFe0.30 , G/NiFe0.56 , and G/Fe, respectively. To evaluate the durability of different catalysts, we conducted a series of chronoamperometric tests at potentials for an initial OER current density of 1.0 mA cm–2 . As shown in Fig. 4.19b, although the operation overpotential (360 mV) is much higher than G/NiFe0.30 (305 mV) and G/NiFe0.56 (342 mV), the G/Ni sample exhibits the best stability with a nearly 100% current retention after 8000 s. In contrast, the G/Fe exhibits poor stability, which may be rationalized by the high overpotential (400 mV) and the amorphous crystal structure. It reveals the importance of the well-crystallized laminar structure on a better OER stability. Comprehensively, the G/NiFe0.30 hybrid exhibits the best OER performance considering both activity and stability.

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Fig. 4.19 OER performances of G/NiFe composites in 0.10 M KOH electrolyte. a LSV curves. b Chronoamperometric responses for different samples at a constant potential with an initial current density of 1.0 mA cm–2 . c Tafel plots. d OER performance as a function of Fe content in consideration of both kinetics (Tafel slope) and the activity (overpotential required to achieve 10.0 mA cm–2 ). Reprinted from ref. [28], copyright 2016, with permission from The Royal Society of Chemistry

It is notable that the LSV curves of these catalysts deliver some novel features at higher overpotentials. As shown in Fig. 4.19a, for those samples with higher Fe contents, the current density rises slowly at low overpotentials but increases quickly after certain potential. For example, the current density of G/NiFe0.56 at the overpotential of 500 mV even exceeds that of G/NiFe0.30 . In addition, the G/Fe sample with the highest Fe content exhibits the lowest Tafel slope of 60 mV dec–1 , even lower than that of G/NiFe0.30 by 16 mV dec–1 (Fig. 4.19c). This phenomenon indicates that in the case of low Fe contents, the active phase during OER process is Fe-substituted NiO(OH), which can decrease the overpotential. While in the case of a higher Fe content, the active phase transforms to Ni-substituted FeO(OH), which can accelerate the kinetics but requires a higher overpotential for activation. It is contributed from the different roles of Ni, Fe, and their oxyhydroxide frameworks on the OER activity, which will be discussed later. As shown in Fig. 4.19d, the OER performance of G/NiFex bimetallic (oxy)hydroxide hybrids exhibits a volcano-type relation with the Fe content, suggesting an optimized Fe content in the range of 25–60 at.%, in consistence with other reports [10, 13].

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4.4.4 Guest–Host Chemistry of Multi-Metal Electrocatalysts To elucidate the variation of OER activities determined by the Fe/Ni ratios, we firstly analyzed the Ni2+ /Ni3+ redox behaviors for different samples. The resulting NiO(OH) phase is believed to be crucial to the OER activity of NiFe (oxy)hydroxides whatever the real active sites [26], while the incorporation of Fe can significantly regulate the oxidation state of Ni2+ , leading to different redox behaviors and catalytic activities. As the scan rate for LSV curves is 5.0 mV s–1 , the overpotential can be converted to test time in Fig. 4.19a, and then the integration of the Ni2+ /Ni3+ oxidation peak describes the transferred charge which is assigned to the extent of the Ni(OH)2 /NiO(OH) transformation before OER. As shown in Fig. 4.20a, the Ni(OH)2 /NiO(OH) transformation is revealed to be suppressed with the Fe content increased, well consistent with previous observations in various NiFe systems [13] and CoFe (oxy)hydroxides [32]. Assuming that the activity contribution of the in situ transformed NiO(OH) phase for all samples is similar, we normalized the LSV curves based on the amount of redox active Ni (resultant Ni3+ ) in each electrocatalyst, as

Fig. 4.20 a The charge transferred during Ni2+ /Ni3+ transformation and the oxidized percentage from the total Ni2+ based on the integral redox peak area. b LSV curves normalized based on the amount of redox active Ni (resultant Ni3+ ) in each material. c Charging current density differences plotted versus scan rates. d TOF depicted based on Fe, Ni, or the total number of Fe and Ni of the working electrodes. Reprinted from ref. [28], copyright 2016, with permission from The Royal Society of Chemistry

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shown in Fig. 4.20b. It is notable that the normalized current density of G/NiFe0.30 is inferior to that of G/NiFe0.56 , which however exhibits an oxyhydroxide framework of FeO(OH) rather than NiO(OH). This result indicates that the intrinsic activities of the host-metal (Ni or Fe) oxyhydroxides are distinct, and the modulation effect derived from the guest-metal (Fe or Ni) substitution is also different. In addition to the active framework and modulation effect, the ECSA is another important factor that can affect catalytic activities. As shown in Fig. 4.20c, G/NiFe0.30 exhibits the highest ECSA, which is more than twice that of G/NiFe0.56 and G/Fe, owing to the nano-sized lamellar structure and uniform distribution of NiFe0.30 oxyhydroxides. The G/Ni sample delivers the lowest ECSA due to the aggregations of much larger Ni(OH)2 flakes. Given the same synthesis condition except Fe/Ni ratios for all samples, this difference in ECSA suggests that the Fe/Ni ratio can alter the physical structure and decoration style of in situ grown (oxy)hydroxides. In heterogeneous catalysis, the turnover frequency (TOF) is widely used to evaluate the intrinsic activity of catalysts, which is defined as the maximum number of chemical conversions of substrate molecules per second and per catalytic site. However, it is always challenging to accurately calculate the TOF value due to the controversy on real active sites and also the difficulty of determining the number of active sites that really take part in the reaction. Assuming only Ni or Fe or all of them are involved in OER, here, we systematically calculate the TOF values based on the number of Ni and Fe atoms, respectively or together (denoted as TOFNi , TOFFe , TOFtotal ). As plotted in Fig. 4.20d, both the TOFNi and TOFFe values increase linearly with the number of assumed active sites decreased. Based on this phenomenon, we cannot identify whether Ni or Fe atoms are responsible for the real activity sites in each sample, as the trends can be ascribed to the net effect rather than intrinsic origin. However, the TOFtotal of two bimetallic (oxy)hydroxides, namely G/NiFe0.30 and G/NiFe0.56 , rises almost tenfold compared with their monometallic counterparts. It unequivocally reveals that the synergistic role of guest and host metals is responsible for the significantly higher activity in the multi-metallic (oxy)hydroxides regardless of the host framework. The above characterization and discussion are mostly consistent with previous reports [10, 13, 32]; nevertheless, some discrepancies shed a novel light on the catalytic mechanism of the multimetal hydroxide/oxyhydroxide active phases. Previous works report that the NiO(OH) is the main active substrate in NiFe (oxy)hydroxide catalysts for OER, while the pure Ni(OH)2 /NiO(OH) itself has poor activity [26]. Besides, FeO(OH) is an unstable and inactive framework for OER [32]. The Fe incorporation, even at an ultra-low concentration (