Graphene: Preparations, Properties, Applications, and Prospects 0128195762, 9780128195765

Table of contents :
Graphene
Copyright
Preface
Acknowledgments
1 -
Introduction
1.1 What is graphene?
1.2 Fundamentals of materials science for carbon materials
1.2.1 Classification of carbon materials
1.2.2 Structure and nanotexture of carbon materials
1.2.3 Carbonization and graphitization
1.2.4 Carbon materials
1.2.4.1 Highly oriented graphite materials
1.2.4.2 Synthetic graphite materials
1.2.4.3 Fibrous carbon materials
1.2.4.4 Nanoporous carbons
1.2.4.5 Spherical carbon materials
1.2.4.6 Glass-like carbons
1.3 Construction and purposes of the current book
References
2 -
Preparation of graphene
2.1 Chemical vapor deposition
2.1.1 Synthesis of graphene films
2.1.1.1 On platinum
2.1.1.2 On nickel
2.1.1.3 On copper
2.1.1.4 On other metals
2.1.1.5 On silicon carbide
2.1.1.6 On other metal carbides
2.1.1.7 On others
2.1.1.8 Transfer of graphene films
2.1.1.9 Structure analysis of graphene films
2.1.2 Synthesis of graphene flakes
2.1.3 Synthesis of single-walled carbon nanohorns
2.1.4 Substitutional doping of heteroatoms
2.1.5 Graphene foams
2.2 Cleavage (peeling)
2.2.1 Mechanical cleavage
2.2.2 Cleavage in solution
2.2.3 Cleavage via intercalation compounds
2.3 Exfoliation via graphene oxide
2.3.1 Synthesis of graphene oxide
2.3.2 Exfoliation of graphene oxide
2.3.3 Reduction of graphene oxide
2.3.3.1 Thermal reduction
2.3.3.2 Chemical reduction
2.3.3.3 Hydrogen reduction
2.3.3.4 Hydrothermal reduction
2.3.3.5 Irradiation reduction
2.3.3.6 Electrochemical reduction
2.3.4 Fabrication of reduced graphene oxide foams (sponges)
2.3.5 Functionalization of reduced graphene oxide
2.3.6 Substitutional doping of heteroatoms
2.3.7 Fabrication of transparent reduced graphene oxide films
2.4 Other processes
2.4.1 Chemical synthesis
2.4.2 Synthesis via pyrolysis
2.4.3 Unzipping of carbon nanotubes
2.5 Concluding remarks
References
3 -
Electrical properties and applications
3.1 Fundamental electrical properties
3.1.1 Electronic structure of graphene
3.1.2 Effects of defects and edges
3.2 Applications to information technology
3.2.1 Transistor devices
3.2.2 Spintronics devices
3.2.3 Transparent electrode
3.3 Applications to social fields
3.3.1 Sensor devices
3.3.2 Photon detectors
3.3.3 Resistance standard
3.3.4 Electron field emission
3.4 Concluding remarks
References
4 -
Chemical properties and applications
4.1 Fundamental chemical properties
4.1.1 Hydrogenation
4.1.2 Oxygenation
4.1.3 Layer modification
4.2 Applications to energy storage and conversion
4.2.1 Lithium-ion batteries
4.2.2 Electrochemical capacitors
4.2.3 Lithium-ion capacitors
4.2.4 Lithium-sulfur batteries
4.2.5 Solar cells (photovoltaic cells)
4.2.5.1 Semiconductor solar cells (Schottky junction solar cells)
4.2.5.2 Polymer solar cells (dye-sensitized solar cells)
4.2.6 Fuel cells
4.2.7 Hydrogen storage
4.3 Applications to environment remediation
4.3.1 Adsorption of polluting molecules and ions
4.3.2 Sorption and recovery of oils
4.3.3 Capacitive deionization for water desalination
4.3.4 Catalysts
4.3.5 Chemical sensors
4.4 Concluding remarks
References
5 -
Mechanical properties and applications
5.1 Fundamental mechanical properties
5.2 Nanolubricants
5.3 Mechanical sensors
5.4 Mechanical reinforcement
5.4.1 Reinforcement of plastics
5.4.2 Reinforcement of ceramics
5.4.3 Reinforcement of metals
5.5 Reduced graphene oxide fibers
5.6 Concluding remarks
References
6 -
Thermal properties and applications
6.1 Fundamental thermal properties
6.2 Thermal interface materials
6.3 Nanofluids
6.4 Thermoelectric power
6.5 Thermal energy storage
6.6 Concluding remarks
References
7 -
Biomedical properties and applications
7.1 Biocompatibility
7.2 Cell management
7.2.1 Scaffolds for cell culturing
7.2.2 Stem cell differentiation
7.2.3 Cell imaging
7.2.4 Antibacterial activity
7.3 Drug delivery systems
7.4 Biosensors
7.5 Concluding remarks
References
8 -
Beyond graphene
8.1 Graphene derivatives
8.1.1 Graphane (hydrogenated graphene)
8.1.2 Fluorographene (fluorinated graphene)
8.1.3 Graphene oxide (oxidized graphene)
8.1.4 Graphyne and graphdiyne
8.2 Single-layer materials
8.2.1 Honeycomb layers of group IV elements
8.2.2 Honeycomb layers of group III–V compounds
8.2.3 Single layers of transition metal dichalcogenides
8.3 Layer-by-layer composites
8.4 Concluding remarks
References
9 -
Summary and prospects
9.1 Summary on graphene
9.2 Prospects
9.2.1 Importance of number of layers stacked
9.2.2 Two kinds of graphene materials
9.2.3 Field effect and zero bandgap
9.2.4 Extremely high thermal conductivity
9.2.5 Basics for molecular sensing
9.2.6 Basics for foreign atom doping
9.2.7 Importance of π–π interaction
9.2.8 Biomedical applications
9.2.9 New composite materials
9.2.10 Extension to organic chemistry
References
Index
A
B
C
D
E
F
G
H
I
K
L
M
N
O
P
Q
R
S
T
U
V
W
X
Y
Z

Citation preview

Graphene Preparations, Properties, Applications and Prospects Kazuyuki Takai Seiya Tsujimura Feiyu Kang Michio Inagaki

Elsevier Radarweg 29, PO Box 211, 1000 AE Amsterdam, Netherlands The Boulevard, Langford Lane, Kidlington, Oxford OX5 1GB, United Kingdom 50 Hampshire Street, 5th Floor, Cambridge, MA 02139, United States Copyright © 2020 Elsevier Inc. All rights reserved. No part of this publication may be reproduced or transmitted in any form or by any means, electronic or mechanical, including photocopying, recording, or any information storage and retrieval system, without permission in writing from the publisher. Details on how to seek permission, further information about the Publisher’s permissions policies and our arrangements with organizations such as the Copyright Clearance Center and the Copyright Licensing Agency, can be found at our website: www.elsevier.com/permissions. This book and the individual contributions contained in it are protected under copyright by the Publisher (other than as may be noted herein). Notices Knowledge and best practice in this field are constantly changing. As new research and experience broaden our understanding, changes in research methods, professional practices, or medical treatment may become necessary. Practitioners and researchers must always rely on their own experience and knowledge in evaluating and using any information, methods, compounds, or experiments described herein. In using such information or methods they should be mindful of their own safety and the safety of others, including parties for whom they have a professional responsibility. To the fullest extent of the law, neither the Publisher nor the authors, contributors, or editors, assume any liability for any injury and/or damage to persons or property as a matter of products liability, negligence or otherwise, or from any use or operation of any methods, products, instructions, or ideas contained in the material herein. Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library ISBN: 978-0-12-819576-5 For information on all Elsevier publications visit our website at https://www.elsevier.com/books-and-journals

Publisher: Matthew Deans Acquisition Editor: Glyn Jones Editorial Project Manager: Naomi Robertson Production Project Manager: Sruthi Satheesh Cover Designer: Mark Rogers Typeset by TNQ Technologies

Preface The Nobel Prize in Physics for 2010 was awarded to Profs. A. Geim and K. Novoselov of the University of Manchester for their groundbreaking experiments on graphene. The term “graphene” was proposed in 1986 in relation to the terminology used for graphite intercalation compounds. After the Nobel Prize, scientific and technological interest in graphene increased rapidly; as a consequence, a tremendous amount of literature declaring its research target to be graphene has been published. Unfortunately, this rapid growth in interest has caused serious confusion regarding the definition and terminology of graphene even in scientific journals. Two of the current authors, M. Inagaki and F. Kang, have authored three books in a series, entitled Materials Science and Engineering of Carbon: Fundamentals, Advanced Materials Science and Engineering of Carbon, and Materials Science and Engineering of Carbon: Characterization. It was anticipated that these books would provide a comprehensive understanding of a wide range of carbon materials (graphite, graphitized carbons, carbon blacks, activated carbons, pyrolytic carbons, glass-like carbons, porous carbons, carbon fibers, etc., in addition to diamond, fullerenes, carbon nanotubes, and graphene) through the detailed explanation and discussion of their structures, nanotextures, and fundamental properties. However, the occurrence of misperceptions and the flood-like increase in research papers on graphene-related (graphene-like) materials led these two authors to believe it was necessary to edit a book focusing on graphene-related materials, in addition to the three previously mentioned books, although graphene is a members of the carbon family. Therefore, they invited two authors, K. Takai and S. Tsujimura, who are specialists in the physical chemistry of graphene, and electrochemistry applications to biomedicals, respectively, to cover the widely spread applications of graphene in this book. In this book, the authors attempted to provide summaries and reviews on graphene and its related materials by differentiating materials with a high perfection of structure (graphene) from those that are highly defective, even with various functional groups attached (reduced graphene oxide). In addition, it is emphasized that the number of layers governs the properties of the flakes of graphene (the characteristics of graphene) which are quite different from those of graphite (many layers stacked) and possible to be obtained on the flake of only a few highly-crystalline layers stacked.

ix

Preface To understand graphene and its related materials, a wide range of fundamental knowledge on various carbon materials is essential: that is, knowledge such as carbonization, graphitization, intercalation, and so on, in addition to basic knowledge about chemistry, physics, biology, and others. For the readers’ convenience, it is recommended to consult the three books mentioned previously, which are published by Tsinghua University Press and Elsevier. These books will supply fundamental knowledge about carbon materials and provide an understanding of a broad range of topics in the current book. It will be truly pleasing to all of the authors if the content of this book delivers useful information to the readers and will lead readers to the correct understanding of graphene and related materials.

x

Acknowledgments The authors would like to express their sincere thanks to the people who kindly provided the data and figures for this book. The names and affiliations of the contributing persons are mentioned in the captions of figures. The authors also thank all of the people who took care of this book at Tsinghua University Press and also at Elsevier.

xi

CHAPTER 1

Introduction Chapter Outline 1.1 What is graphene? 1 1.2 Fundamentals of materials science for carbon materials 1.2.1 1.2.2 1.2.3 1.2.4

Classification of carbon materials 5 Structure and nanotexture of carbon materials Carbonization and graphitization 8 Carbon materials 12 1.2.4.1 Highly oriented graphite materials 12 1.2.4.2 Synthetic graphite materials 17 1.2.4.3 Fibrous carbon materials 20 1.2.4.4 Nanoporous carbons 23 1.2.4.5 Spherical carbon materials 26 1.2.4.6 Glass-like carbons 28

1.3 Construction and purposes of the current book References 31

5

6

29

1.1 What is graphene? The term “graphene” was firstly proposed in 1986 [1] and then recommended by the International Union of Pure and Applied Chemistry [2] as the name for a single twodimensional layer of carbon atoms bonding using sp2 hybrid orbitals, which occurs in graphite intercalation compounds. The single carbon layer occurring in the intercalation compounds was proposed to be called “graphene” which comes from the suffix “-ene” for polycyclic aromatic hydrocarbons such as naphthalene and anthracene and the prefix “graph-” from graphite. In the earliest literature [1], the authors noted that “it should be adopted for graphite intercalation compounds.” In the first-stage structure of the compounds, a two-dimensional (2D) carbon layer is sandwiched by two intercalate layers and isolated from other carbon layers, although more than two carbon layers are stacked in parallel with regularity in compounds with higher than a second stage, as schematically illustrated in Fig. 1.1. “Graphene” is commonly defined as an isolated single layer of carbon hexagons consisting of sp2 hybridized CeC bonding with p-electron clouds. From an engineering point of view, thin flakes consisting of a few layers of carbon atoms, including single-layer Graphene. https://doi.org/10.1016/B978-0-12-819576-5.00001-3 Copyright © 2020 Elsevier Inc. All rights reserved.

1

2 Chapter 1 Stage 1

Stage 2

Stage 3

A

A

A

A

A

A

B

B

Graphene A

Graphite stacking

A B

A

A

C

C C A

Carbon layer Intercalate layer

A

A

A

Figure 1.1 Stage structure of graphite intercalation compounds.

graphene, could be important because of their interesting structural, chemical, and physical characteristics. The promising potential applications of graphene and its related materials have been pointed out in technological fields. Numerous reports have been published, particularly after the Nobel Prize in Physics was awarded in 2010 to Profs. A. Geim and K. Novoselov of the University of Manchester for their groundbreaking experiments in graphene [3]. Sudden scientific and technological interest in graphene and its related materials has caused some confusion about the definition and terminology of graphene-related materials, even in scientific journals. A proposal regarding the nomenclature for two-dimensional carbon materials was been presented in the journal Carbon [4]. In much of the literature, however, the term “graphene” has not been used according to its strict definition, i.e., a single layer of carbon atoms consisting of sp2 hybridized bonds. Some authors do not pay enough attention to how many layers are stacked in their materials, although they have called them “graphene” Therefore, here, the terms “single-layer graphene” (or monolayer graphene), “doublelayered graphene” (or bilayered graphene), and “multilayered graphene” are used only for products that have confirmed numbers of stacked layers. Numerous reviews have been published from various viewpoints: providing an overview of graphene by pointing out what is fascinating about it [5e13], focusing on its production in relation to its structure [14e23], applications in electronics [24e30], energy storage applications [31e48], biological applications including sensors [49e68], functionalization including the formation of composites [69e80], mechanical applications [81,82], thermal properties [83,84], doping to improve its properties [85,86], and its toxicity [87,88].

Introduction 3 Before summarizing research on graphene-related materials and discussing the results, it is worthwhile to understand the flakes and/or films that determine how many numbers of layers can exhibit the unique functionalities characteristic of graphene, to differentiate it from graphite. Fig. 1.2 compares Raman spectra are for graphite and graphene flakes with different numbers of stacked layers [89]. As shown in Fig. 1.2A, highly oriented pyrolytic graphite (HOPG), for example, shows a sharp and strong G-band and G0 -band (2D band). The latter is presented as decomposing into two bands, G0 1 and G0 2, with heights of roughly ¼ and ½ of the G-band, respectively. On the other hand, single-layer graphene exhibits a sharp and single profile for both G- and G0 -bands and the G0 -band is much stronger than the G-band by almost four times. As shown in Fig. 1.2B, the G0 -band gradually shifts to a higher position, changing its profile to unsymmetrical and decreases its intensity relative to the G-band with an increasing number of stacked layers. A marked field effect on electrical conduction is a characteristic of graphene. Pioneering works were reported on few-layered graphene flakes [90]. Fig. 1.3A shows resistivity change with gate voltage Vg on few-layered graphene flakes as a function of temperature, revealing a marked field effect, a peak of several kU with rapid decay to about 100 U at high Vg, and a strong dependence on temperature. This field effect is shown by plotting the relative carrier concentration n/n0 (n0 is the concentration of carriers at 4 K) against temperature for three flakes, few-layered flakes, few-layered but much thicker flakes, and multilayered flakes in Fig. 1.3B. The results clearly reveal that the field effect depends on the number of stacked layers: the smaller the number of layers, the stronger the field effect. (A)

(B) G'-band G'-band (2D-band)

Graphite

Single-layer graphene

1500

2000

Graphite Intensity / a.u.

Intensity / a.u.

G-band

2500

Raman shift / cm-1

10-layered 5-layered 2-layered 1-layer

3000

2600

2700

2800

Raman shift / cm-1

Figure 1.2 Raman spectra measured by a laser with 514-nm wavelength for graphite to single-layer graphene [89]: (A) comparison of graphite and single-layer graphene, (B) Change in G0 -band with increasing number of stacked layers. 2D, two-dimensional; a.u., arbitrary units.

4 Chapter 1 (A)

(B) Few-layered (refer a)

6

5K

n0(T) / n0(4K)

8

ρ / KΩ

6

70 K

4

Few-layered (thicker)

4

2 2 300 K 0 –100

–50

Multi-layered

0

50

0

100

0

100

Vg / V

200

300

Temperature / K

Figure 1.3 Field effect for few-layered graphene flakes [90]. Graphene

Thermal conductivity / W/m·K

6000

Largest data Averaged value Theoretical prediction I Theoretical prediction II

5000 4000 3000

Highly-crystalline bulk graphite 2000 1000

Bulk graphite 0

1

2

3

4

5

6

7

8

9

Number of layers

Figure 1.4 Change in thermal conductivity at room temperature with number of layers stacked [91].

Extremely high thermal conductivity, more than 5000 W/m.K, is expected for single-layer graphene [91]. It depends on the number of layers stacked in the flake. In Fig. 1.4, thermal conductivities experimentally determined from the temperature dependence of the position of the micro-Raman G-band together with theoretical predictions are plotted against the number of stacked layers [91]. A marked dependence of room temperature thermal conductivity on the number of layers is clearly shown. Single-layer graphene has extremely high conductivity, but that of a flake composed of four layers is comparable to bulk graphite with high crystallinity, such as HOPG, although it is high compared with metals.

Introduction 5 These three experimental results suggest that we need a flake and/or film composed of less than four layers to obtain characteristics intrinsic to graphene.

1.2 Fundamentals of materials science for carbon materials Graphene is a member of carbon materials, which include graphite, carbon blacks, carbon nanotubes, carbon fibers, activated carbons, porous carbons, diamond, and fullerenes. To prepare graphene flakes and films, graphite materials such as natural graphite, HOPG, and kish graphite have been employed as starting materials. In addition to graphene, other carbon materials such as activated carbons and porous carbons including carbon foams and carbon nanotubes have been cited in references and/or competitive materials. Before discussing graphene, the fundamentals of carbon materials science will be briefly explained, emphasizing highly oriented graphite materials that have been used as raw materials of graphene. For a detailed explanation and discussion of carbon materials, readers of this book are suggested to refer to fundamental books written by the current authors (M.I. and F.K.) [92,93].

1.2.1 Classification of carbon materials Many kinds of carbon materials have been manufactured, synthesized, and widely used in various industries. These carbon materials are proposed to be classified on the basis of their chemical bonds of constituent carbon atoms using sp3, sp2, and sp hybrid orbitals. The sp2 hybrid bonding of carbon atoms results in two structures: flat layers composed of six-membered carbon rings, which had been represented by graphite but now include graphene; and curved layers, which introduce five-membered carbon rings into sixmembered rings, as occurs in fullerenes. Layers composed of sp2 orbitals, both flat and curved, are intrinsically anisotropic and have p-electron clouds at both sides of the layer, which creates broad diversity in the structure and properties of the materials. Carbon nanotubes can be placed between fullerene and graphene, because the tips of the tube include five-membered rings (fullerene-like) and its wall is composed of six-membered rings although it is rolled up (graphene-like). Fig. 1.5 shows the classification of carbon materials based on hybrid bonds, together with the diversity of each material. Because of the anisotropic nature and the presence of p-electron clouds in the carbon materials composed mainly of sp2 hybrid orbitals, the number of layers stacked in parallel has a strong influence on their properties. The importance of the number of stacked layers has been pointed out for carbon nanotubes and fullerenes, and now for graphene, as mentioned in the previous section. Large numbers of layers stacked with regularity have been called graphite, and various graphite-related materials have been produced in industries and used as important industrial materials, some of which are listed in Fig. 1.5.

6 Chapter 1

C 2s22p2

Crystalline Cubic Hexagonal

sp + 2π Carbyne

sp2 + π

sp3 Diamond

Planar

Curved

Graphene

Fullerenes

Non-crystalline Single-layer Diamond-like to carbon (DLC) Multi-layered

Carbon nanotubes Single wall Chirality Multi-walled

Graphite

Chain length, etc.

Single wall C60, C70, ··· Multi-walled

Crystalline Hexagonal Rhombohedral

Non-crystalline Turbostratic

P1=1.00 d002=0.3354 nm

P1=0.00 d002>0.344 nm

Graphitization degree Nanotexture

Oriented Planar orientation Highly-oriented graphite Synthetic (artificial) graphite Intercalation compounds Pyrolytic carbons

Axial orientation Carbon fibers C/C Composites

Not-oriented Point orientation Carbon blacks

Random orientation Activated carbons Glass-like carbons

Figure 1.5 Classifications and diversity of carbon materials.

1.2.2 Structure and nanotexture of carbon materials Most carbon materials composed of flat layers using sp2 hybrid orbitals consist of small units of layers stacked in parallel, as shown by the transmission electron microscopy (TEM) lattice fringe image and schematic illustration in Fig. 1.6. These units are called basic structural units (BSU) or crystallites. In a unit, two kinds of layers coexist with stacking regularity: random and regular stacking. The former is called turbostratic stacking and the latter is graphitic stacking (usually written as AB stacking). These BSUs are strongly anisotropic in the nature of their bond, with strong covalent bonding using sp2 hybrid orbital along the layer and weak van der Waals bonding across the layers. The aggregation of these anisotropic nanosized BSUs in the particles results in different textures owing to the different schema of preferred orientations of anisotropic BSUs (i.e., planar, axial, point, and random orientations), as illustrated with some representative carbon materials in Fig. 1.7. The aggregation of BSUs in different schema is called nanotextures. By the planar orientation scheme, films and platelets of highly oriented

Introduction 7

Figure 1.6 Basic structural unit of carbon materials: (A) lattice fringe image and (B) schematic illustration of a unit.

RANDOM NANOTEXTURE

ORIENTED NANOTEXTURE Planar orientation

Glass-like carbons

Random orientation

Highly-oriented graphites Cokes

Axial orientation Carbon nanotubes Carbon fibers

Point orientation

Fullerenes Carbon blacks

Figure 1.7 Nanotextures in carbon materials on the bases of preferred orientations of basic structural units.

graphite are produced and some coke particles are principally composed of planar orientation of BSUs. By the axial orientation scheme, fibrous carbon materials are produced from different precursors such as carbon nanotubes and vapor-grown carbon fibers with a coaxial mode of orientation and some mesophase pitch-based carbon fibers with a radial mode. By the point orientation scheme, various carbon spheres are produced, as well as various sized fullerene particles and different nanosized carbon blacks with concentric mode and mesophase spheres with a radial mode. The particles composed of these oriented nanotextures are still anisotropic. In addition, random aggregation of small BSUs occurs in so-called glass-like carbon (glassy carbon), the particles of which are isotropic in nature.

8 Chapter 1

1.2.3 Carbonization and graphitization Most carbon materials used in industry are produced from organic precursors such as pitches, biomasses, and organic polymers via heat treatment at high temperatures in an inert atmosphere during carbonization and graphitization processes. Fig. 1.8 summarizes changes in chemical and electronic band structures. Carbonization is performed after pyrolysis of the precursors from 800 to 2000 C, in which the BSUs are formed and their basic aggregation scheme (nanotexture) is established, accompanied by the emission of foreign atoms, oxygen, hydrogen, and nitrogen as gases and the polycondensation of six-membered carbon rings. Because the nanotexture of most carbon materials is established, this process is most important in the production of various carbon and graphite materials. The nanotexture governs the development of crystalline structures in carbon materials during graphitization. During this process, a large amount of shrinkage and the rapid emission of gas species occur that are associated with cracking of carbon particles that occur in many cases. Therefore, the process of carbonization is applied separately from that of graphitization. Above 2000 C, a change in crystalline structure (the development of a graphite structure) mainly occurs. The development of a graphite structure may be evaluated by different techniques, including x-ray diffraction (XRD), electromagnetic property measurements, Raman spectroscopy, and high-resolution TEM. In BSUs formed during carbonization,

Graphitization

Carbonization Pyrolysis

Electronic band structure

Graphite materials

Carbon materials

Conduction band Fermi level Valence band

r. t.

H2

CH4

CO2 & CO

High MW aromatics

Chemical and crystallographic structure

Low MW aliphatics

Cyclization Polycondensation Aromatization

Improvement in layer stacking Crystallite growth La

Lc

2000 1000 Heat treatment temperature / oC

3000

Figure 1.8 Schematic illustration of chemical, crystallographic, and electronic band structures in carbon materials with a planar orientation nanotexture. MW, molecular weight.

Introduction 9

Figure 1.9 Changes in x-ray diffraction parameters with heat treatment temperatures (HTT) for various carbon materials: (A) d002, (B) Lc measured from a 002 diffraction peak, and (C) La from a 110 peak.

turbostratic stacking with an interlayer spacing of about 0.342 nm is randomly changed to graphitic regular stacking with a spacing of 0.3354 nm (the spacing in the graphite crystal) with an increasing heat treatment temperature (HTT) of above 2000 C, which is measured as a decrease in the averaged interlayer spacing, d002, by XRD, associated with the growth of BSU sizes (crystallite size) along the a- and c-axes, La and Lc. The change in d002 with HTT depends on the materials after carbonization (carbon materials). Fig. 1.9 shows the changes in these parameters with HTT for various carbon materials. In needle-like coke with a planar orientation scheme, d002 decreases quickly, approaching the value of a graphite crystal (0.3354 nm), and Lc and La grow rapidly. In glass-like carbon with a random orientation scheme, in contrast, there is almost no decrease in d002 and no appreciable growth in Lc and La even after 3000 C treatment (i.e., no development of a graphite structure). Carbon blacks with a point orientation scheme have intermediate behaviors: the large-sized thermal black shows more improvement in structure than does the small-sized furnace black. Diffraction peaks of XRD for carbon materials are classified into three groups: 00l, hk0, and hkl, mainly owing to the strong anisotropy of BSUs and the coexistence of two interlayer spacings, as shown in Fig. 1.6B. Diffraction peaks with indices of 00l give averaged interlayer spacing, which decreases gradually from more than 0.344 nm for highly defective layers to 0.3354 nm for graphitic stacking and about 0.342 nm for turbostratic stacking with increasing HTT: in other words, with improving crystallinity of carbon, as shown in Fig. 1.9A. In turbostratic stacking of layers, there is no 3D regularity in stacking (i.e., no l index is defined), and so the diffraction peaks with hk0 indices for a graphitic structure are expressed as hk diffraction peaks for carbon materials mainly consisting of the turbostratic stacking of layers. Regular and random stackings, graphitic and turbostratic, are clearly demonstrated in diffraction profiles of

10 Chapter 1 (A) XRD

101 100 10

(B) TEM (002 lattice fringe image) d002=0.3345 nm Regular stacking (graphite structure) d002=0.342 nm Random stacking (turbostratic structure)

10

40 50 2θ / degree(CuKα)

d002>0.344 nm Minute BSUs

Figure 1.10 (A) Correspondence between x-ray diffraction (XRD) profile of hk line and (B) transmission electron microscopy (TEM) image. BSU, basic structural units.

hk0 and hk peaks, as shown for 101 and 10 peaks in Fig. 1.10. The profile of the 10 peak for a turbostratic structure shows a characteristic unsymmetrical profile, and that of a 100 peak for a graphitic structure is sharp and symmetrical and is associated with a 101 peak owing to the formation of 3D stacking regularity. With increasing HTT, an unsymmetrical 10 peak is modulated by the appearance of a 101 peak, together with its sharpened and improving symmetry. Peaks with indices of hkl are caused by the graphite structure, and so 112 peak is often selected as to indicate the formation of a graphite structure, because the 112 peak does not overlap with other peaks although the 101 peak overlaps with the 100 peak, as can be seen in Fig. 1.10. Because most carbon particles are more or less anisotropic, except for glass-like carbon with a random orientation scheme, their aggregation into a block gives texture on an larger scale, which may be called microtexture and macrotexture. The technique for evaluating these micro- and macrotextures has not yet been established. Figs. 1.11 and 1.12 demonstrate examples of these textures, which must be controlled for practical applications. In Fig. 1.11, scanning electron microscopy (SEM) images of cross-sections of commercially available isotropic high-density graphite are shown, demonstrating different pore structures [94]. Image processing of these micrographs suggests a relation between pore structure parameters (for example, the pore area) and the mechanical properties of these graphite blocks. To prepare carbon fiberereinforced plastics and carbon fiberecarbon composites, different macrotextures based on the orientation of carbon fibers are employed to obtain high strength and a high modulus of the composites, as shown in Fig. 1.12.

Introduction 11

Figure 1.11 SEM images of the cross-sections of isotropic high-density graphite blocks [94].

Chopped fibers Random

One-directional

Long fibers Three-dimensional Two-directional

Three-directional

Figure 1.12 Different scheme for reinforcing carbon fibers in composites.

12 Chapter 1

1.2.4 Carbon materials 1.2.4.1 Highly oriented graphite materials The blocks or platelets, which are composed of big carbon layers stacked in large amount with graphitic regularity, are called highly oriented graphite; the extreme case of this is a single crystal of graphite. In practice, however, it is difficult to obtain single large crystals and almost impossible to get those that are more than a few square millimeters. There are only two ways to find single crystals of graphite: in natural graphite ores and in so-called kish graphite. Resources of high-quality natural graphite are limited on earth and exist in Sri Lanka, Madagascar, and China. Even in these ores, the possibility of finding single crystals of a certain size is slim. Kish graphite is formed by the precipitation of supersaturated carbon from molten iron and cannot be large in size, but some have very high crystallinity and can be called single crystals. Alternatives to single-crystal graphite are HOPG and graphite films derived from some organic polymer precursors such as polyimides via high-temperature treatment. Natural graphite: Natural graphite is usually recovered as a powder from natural ores through milling and purification processes. After the final purification process, the powder has usually a purity of more than 99 wt%. Some of it consists of flaky particles with a highly crystalline structure and highly oriented nanotexture (flaky graphite), but some consists of aggregates of small crystals (called amorphous graphite) [95]. These two kinds of natural graphite can be differentiated by XRD, as shown in Fig. 1.13, which shows

Hexagonal 002 Rhombohedral 003

Flaky graphite Microcrystalline graphite Hexagonal 101 Rhombohedral Hexagonal 101 100

Rhombohedral 102

2θ / degree Hexagonal 004 Rhombohedral 006

2θ / degree

Figure 1.13 X-ray diffraction patterns of flaky and microcrystalline natural graphite [95].

Introduction 13 flaky graphite with sharp 100 and 101 peaks but a microcrystalline graphite broad 101 peak. Some flaky graphite contains metastable rhombohedral modification of the graphite structure, possibly as a result of shear stress during milling. Kish graphite: A part of carbons dissolved into molten iron at high temperatures is incorporated into the crystal lattice of iron to form alloys (different steels) and another part segregates as graphite. Graphite flakes segregated from molten iron are called kish graphite [96e98]. Relatively large amounts of kish graphite flakes are obtained as a byproduct during steel production; all of them do not always have high crystallinity because it depends on the segregation conditions. When they are produced at the temperature at which iron evaporates, kish graphite flakes have a single crystal nature. As shown in Fig. 1.14, flakes are thin with an irregular shape. Regular stacking of layers is observed by SEM and their single crystal nature is confirmed by a well-organized electron channeling pattern. The high crystallinity of kish graphite flakes was also confirmed by measurements of the resistivity ratio, r300 K/r4.2K, maximum transverse magnetoresistance, (Dr/r0)max, measured at 77 K, and Shubnikov-de-Haas oscillation in magnetoresistance. Graphite single crystals have been synthesized from molten iron by controlling the segregation process [99,100] and from Al4C3 by transporting decomposition gases [101,102]. Highly oriented pyrolytic graphite: Carbon deposited on a substrate by chemical vapor deposition (CVD) of hydrocarbon gases such as methane and propane at high temperatures can have a well-oriented texture and is called pyrolytic carbon. Pyrolytic carbons have extensively been studied to control their structure and texture by applying different deposition conditions, such as precursor hydrocarbon, the concentration and flow rate, the temperature of deposition, and the geometry of the furnace [103]. By hot-pressing at high temperatures under high pressure, the preferred orientation of graphite crystallites and their crystallinity can be markedly improved. Typical conditions are shown in Fig. 1.15 [104].

Figure 1.14 Kish graphite : (A) optical micrograph, (B) scanning electron microscopy image of the edge surface, and (C) electron channeling pattern on the basal plane. Courtesy of Prof. Y. Hishiyama of Tokyo City University.

14 Chapter 1 Pyrolytic carbon (PC)

deposited at 2100-2500 °C mosaic spread of 40-50°

Hot pressing 2800-3000 °C, 30-50 MPa Oriented pyrolytic graphite (OPG)

mosaic spread of less than 0.5 °C density of 2.226 g/cm3

Annealing under pressure 3400-3600 °C, 10 kg/cm2 Highly-oriented pyrolytic graphite (HOPG)

mosaic spread < 0.4°

Figure 1.15 Procedure of the production of highly oriented pyrolytic graphite.

The products of this process are called HOPG. To achieve high crystallinity and high orientation, the starting pyrolytic carbons need selected, as well as the hot-pressing conditions. As shown in Fig. 1.16A, HOPG has a highly oriented nanotexture along its plate, but electron channeling analysis suggests that HOPG consists of randomly oriented a-axes of crystalline domains, although their c-axes are almost perfectly oriented perpendicular to the plate. The channeling pattern on the surface of the plate (Fig. 1.16B) is somewhat distorted compared with that of kish graphite (refer to Fig. 1.14C) and the channeling contrast shows the random orientation of a-axes of the domains (Fig.1.16C). Graphite films derived from organic films: Highly crystallized graphite films were prepared from films of limited numbers of organic precursors, such as poly(p-phenylene vinylene) (PPV) [105], poly(p-phenylene-1,3,4-oxadiazole) (POD) [106,107], benzimidazobenzophenanthroline ladder polymer BBL [107e110], and the commercially available aromatic

Figure 1.16 Highly oriented pyrolytic graphite: (A) scanning electron microscopy image of the cross-section, (B) electron channeling pattern, and (C) electron channeling contrast. Courtesy of Prof. A. Yoshida of Tokyo City University.

Introduction 15 (A) PPV

(D) Polyimide Kapton

nn

(E) Polyimide PPT

(B) POD

n

(C) BBL

Figure 1.17 Molecular repeating units of polymers giving graphite films. PPV, poly(p-phenylene vinylene); POD, poly(p-phenylene-1,3,4-oxadiazole); BBL, benzimidazobenzo-phenanthroline ladder polymer; Kapton; PPT, commercially available aromatic polyimide films.

polyimides Kapton and PPT [111e113], the repeating units of which are shown in Fig. 1.17. Heat treatment of PPV-derived carbon films up to 3000 C gave a graphite film; as shown in its Raman spectrum in Fig. 1.18A, no D-band was detected [105]. The film of polyoxadiazole heat-treated at 2800 C gave sharp and strong 00l diffraction profiles and no hk0 diffraction lines by reflection mode and well-resolved hk0 and hkl lines and no 00l lines by transmission mode, as shown in Fig. 1.18B; this revealed the formation of a highly oriented, well-crystallized graphite film [106]. BBL polymer film was cast onto a glass substrate from its solution of trifluoromethanesulfonic acid (TFMSA) and then heat-treated at high temperatures to 3200 C by sandwiching it between two artificial graphite plates [110]. The resultant ultrathin films, which were less than 100 nm thick, showed an electron channeling pattern similar to that of HOPG and an orientation of graphite basal planes parallel to the film surface, as shown in Fig. 1.19A. By using methanesulfonic acid as solvent for the BBL polymer, film with an orientation of basal planes perpendicular to the film surface was obtained, as shown in Fig. 1.19B [109]. Two films derived from polyimide polytrimethylene terephthalate (PTT) and Kapton by heat treatment at 3200 C under a simple mechanical constraint (sandwiching between two

16 Chapter 1 (A)

(B) G Intensity / x103 cps

D

2000 °C

004

150

006

1 x 20

100

x 10

50 0 500

2500 °C

Intensity / cps

Intensity / a.u..

950 °C

Reflection mode

002

200

2750 °C 3000 °C 1000 1200 1400 1600 1800 Raman shift /

cm-1

101

400

Transmission mode 110

300

100 0 10

112

100

200

102 103 20

30

40

50

60

70

80

90

Diffraction angle 2θ / degree, CuKα

Figure 1.18 (A) Change in Raman spectrum with heat treatment temperature for poly(p-phenylene vinylene) film [105] and (B) x-ray diffraction patterns of polyoxadiazole film heat-treated at 2800 C [106]. a.u., arbitrary units; cps, counts per second.

Figure 1.19 002 lattice fringe images of the cross-section of graphite films prepared from a basic structural units polymer using trifluoromethanesulfonic acid (A) and methanesulfonic acid (B) as solvents and heat-treated at 2800 C [109].

Introduction 17 graphite plates) gave high crystallinity comparable to kish graphite and HOPG; rRT/r4.2K (one parameters to evaluate crystallinity of graphite) was 4.90 and 4.79 for PTT- and Kapton-derived film, respectively, although it was 4.7e5.5 for kish graphite and HOPG [112]. Exfoliated graphite and flexible graphite sheets: Because large flakes of natural graphite are not commonly available and natural graphite flakes cannot be formed directly into a sheet, a technique for preparing graphite sheets with no binding materials was developed by forming graphite oxides and via their thermal exfoliation and reduction, followed by compression into sheets. They are called flexible graphite sheets and have promoted the application of graphite [114]. Graphite sheets have characteristic advantages such as flexibility, resilience, and the ability to be formed into various shapes easily; in addition, graphite has intrinsic properties such as lubricity, chemical and thermal stability, and high electrical and thermal conductivity. Natural graphite flakes are chemically or electrochemically intercalated to form a covalent intercalation compound of graphite, graphite oxide, through the use of a mixture of concentrated sulfuric and nitric acids. After rinsing and drying, residue compounds are obtained in which the regular stage structure is lost but some sulfuric acid derivatives remain in the graphite gallery. These residue compounds are then rapidly heated to 900e1200 C, in which intercalates remaining in the graphite gallery are decomposed into gaseous products, yielding marked exfoliation of the pristine graphite. This exfoliated graphite consists of worm-like particles, as shown in Fig. 1.20. They are formed by preferential exfoliation perpendicular to the plane surface of the pristine flakes. Worm-like particles of exfoliated graphite are either molded or rolled into a sheet with no adhesives or binders. These are called flexible graphite sheets and are widely used as seals, gaskets, and packings. After rinsing and drying of intercalation compounds, the residue compounds are commercially available as expandable graphite and are used as the starting material for thin flakes of graphene oxide in some literature. 1.2.4.2 Synthetic graphite materials So-called graphite materials have been manufactured industrially into variously sized blocks, rods, and plates for different applications, such as large degree electrodes for electric arc furnaces, various jigs for processing metals, crucibles and heating elements to grow semiconductor crystals, brushes for electric motors, and neutron moderators for nuclear reactors. Graphite materials for these applications have mostly been produced using coke particles as fillers with pitches as the binder through carbonization and graphitization treatment at high temperatures. The production processes are summarized as a block diagram in Fig. 1.21. Briefly, pulverized coke particles are mixed and

18 Chapter 1

Figure 1.20 (A) Scanning electron microscopy images of worm-like particles of exfoliated graphite. (B) Appearance of the particles. (C) Cross-section of a particle. (D) Distribution of pore area on cross-section.

Filler Petroleum coke Pitch coke

Pulverization Coase grains

Sieving Fine grains

Binder

Mixing & kneading

Extrusion Molding

Forming

Pitch Resin

Isostatic pressing

Pitch, Resin

Re-calcination Halogen gases

Purification

Graphitization Machining

Impregnation

Calcination (Carbonization)

Carbon materials Graphite Electrode-grade Isotropic materials

high-density-grade

Figure 1.21 Production process of synthetic graphite materials.

Introduction 19 kneaded with binder pitch at a temperature slightly higher than the softening point of the pitch, in which the nanotexture (either needle-like or mosaic) and particle size of the filler coke and the mixing ratio of the filler to the binder have to be controlled according to the requirements of the applications. After the mixture is formed by extrusion, molding, or isostatic compression, it is heat-treated at 1400e1800 C (calcination or carbonization), followed by heat treatment at a higher temperature up to 2800e3000 C (graphitization) (see Fig. 1.8). The final products of this manufacturing process are called synthetic (artificial) graphite materials. All synthetic graphites are polycrystalline solids; of these, BSUs (crystallites) are much smaller than those in the highly oriented graphites described previously. To produce graphite materials for use in an electric arc furnace as an electrode, coarse-grained coke particles with a needle-like texture are usually employed to achieve high thermal-shock resistance. They are formed by extrusion or molding, and the slight orientation of coke particles cannot be avoided. To produce isotropic high-density graphite blocks, however, fine-grained coke particles have to be used to achieve a high density, and isostatic pressing must be employed for an isotropic nature in the final block. Microtextures of two graphite materials, electrode grade and isotropic grade, are compared in Fig. 1.22. Most synthetic graphite materials consist of filler coke and a carbonized binder (binder coke). The former is able to attain a higher degree of graphite structure development (graphitization degree) compared with the latter. The degree of graphitization achieved after final graphitization depends on the nanotexture of the starting coke particles, as shown in Fig. 1.23A. Needle-like coke has an extended flow pattern owing to the preferred orientation of crystallites, and so there is a higher degree of graphitization after hightemperature treatment compared with mosaic coke, which has a short range of orientation and is less ordered (Fig. 1.23B). However, the degree of graphitization reached by needlelike coke is far lower than that of highly oriented graphite materials.

(A)

(B)

100 µm

100 µm

Figure 1.22 Polarized light micrographs of cross-sections of synthetic graphites: (A) electrode grade and (B) isotropic grade.

20 Chapter 1

(A)

(B)

100 μm Figure 1.23 Polarized light micrographs of cross-section of particles: (A) needle-like and (B) mosaic coke.

1.2.4.3 Fibrous carbon materials A variety of fibrous carbon materials have been prepared by different techniques: electric arc-discharge, catalytic CVD, template-assisted CVD, melt-spinning, and electrospinning. Table 1.1 classifies these fibrous carbons according to whether the 002 lattice fringes of carbon layers are short or long and whether their diameters are on the micrometer or nanometer scale. Carbon layers constituting the wall of carbon nanotubes (CNTs) are long and oriented exactly parallel to the tube’s axis, but layers of carbon nanofibers are short and oriented in different modes, called tubular, herringbone, and platelet types, depending on the methods and conditions of synthesis. Carbon fibers with a diameter of about 7 mm are produced industrially by melt-spinning from different precursors, polyacrylonitrile (PAN), isotropic and anisotropic (mesophase) pitches, cellulose, and phenol as filament yarns, clothes, chops, webs, etc. Vapor-grown carbon fibers on the market are produced by catalytic CVD and have tubular nanotextures. By applying electrospinning, various organic polymers can be used as precursors to fabricate carbon nanofibers [115]. Table 1.1: Classification of fibrous carbon materials. Diameter of fiber Nanometer size

Micrometer size

Long carbon layers Carbon nanotubes

Single-walled Double-walled Multiwalled Graphite whisker

Short carbon layers Carbon nanofibers Carbon fibers

Tubular type Herringbone type Platelet type Polyacrylonitrile-based Pitch-based, etc. Vapor-grown carbon fibers

Introduction 21

Figure 1.24 Transmission electron microscopy images of CNTs: (A) by catalytic chemical vapor deposition [116] and (B) by electric arc discharge [119].

Carbon nanotubes: CNTs are synthesized by catalytic CVD during the initial stage of vapor-grown carbon fibers [116,117] and are found in carbon deposits on graphite anodes during arc discharge [118e120], as shown in Fig. 1.24. Under similar arc discharge conditions, a scroll of carbon layers were obtained in 1960, called graphite whiskers [121]. The disproportionation reaction of CO at high pressures and high temperatures with an iron catalyst produces single-walled CNTs (SWCNTs), which are almost free of amorphous carbon [122,123]. Double-walled CNTs (DWCNTs) are successfully obtained at a high yield by catalytic CVD using an Fe catalyst deposited on MgO [124,125]. CNTs aligned perpendicular to the substrate (called arrays) with well-defined patterns were synthesized by CVD of ethylene at 700 C on an Fe thin film (5 nm thick) on an Si substrate [126]. Accelerated growth of SWCNTs and DWCNTs is achieved by using an alcohol as the carbon precursor or by adding water vapor to the precursor gas [127e129]. Vertically aligned SWCNTs (called forest) were grown from ethylene on various metal catalysts in either Ar/H2 or He/H2 by adding a small amount of water vapor to atmospheric gas, as shown in Fig. 1.25. These SWCNT forests grew at a superior high rate [130,131]. Transparent thin sheets 3.4 cm wide, composed of well-aligned MWCNTs, were pulled from the side face of a CNT array [132]. Long yarns and ribbons were continuously spun from the SWCNT array at room temperature [133], passing through a volatile liquid during spinning, which was effective in making the spun yarns dense [134]. Carbon fibers: Commercially available carbon fibers have different nanotextures, particularly in their cross-sections, as illustrated in Fig. 1.26. Carbon fibers derived from isotropic pitch, phenol, and cellulose have a random orientation in cross-sections both parallel and perpendicular to the fiber axis. On the other hand, vapor-grown carbon fibers

22 Chapter 1

A match head SWNT forest (2.5 mm height)

Enlarged picture

Figure 1.25 Single-walled carbon nanotube (SWNT) forest [130].

Figure 1.26 Schematic illustration of nanotexture in carbon fibers. PAN, polyacrylonitrile.

(VGCFs) consist of well-oriented carbon layers that are parallel along the fiber axis and concentric in the perpendicular cross-section. Mesophase pitchebased carbon fibers are able to control their nanotextures in cross-section, such as straight radial, corrugate radial, and concentric, and consequently their properties. Depending on the nanotexture, structural changes with high temperature treatment are different for each carbon fiber. Fig. 1.27 compares changes in interlayer spacing d002 (one

Introduction 23

Interlayer spacing d002 / nm

PAN-based carbon fibers

Anthracene-derived coke VGCFs (10 μm diameter)

HTT / oC

Figure 1.27 Changes in interlayer spacing d002 with heat treatment temperature (HTT). PAN, polyacrylonitrile; VGCFs, vapor-grown carbon fibers. Courtesy of Dr. N. Iwashita.

parameters for evaluating the degree of graphitization) with HTT on PAN-based carbon fibers, VGCFs, and anthracene-derived coke. Graphitization was depressed on PAN-based carbon fibers with a random nanotexture, whereas it proceeded rapidly on VGCFs about 10 mm in diameter with a well-developed axial nanotexture, comparable to anthracenederived coke with a planar nanotexture. The diameter of the fibers has an influence on the development of the graphite structure, VGCFs with a 10-mm diameter have a much higher degree of graphitization than do those 2 mm in diameter. 1.2.4.4 Nanoporous carbons Pores commonly exist in carbon materials; they are categorized according to whether a specific gaseous molecule, usually nitrogen, can be adsorbed in open and closed pores. Open pores are classified by their widths: micropores less than 2 nm wide, mesopores of 2e50 nm, and macropores more than 50 nm. These pores are controlled by oxidation. The process and the resultant porous carbons are called activation and activated carbons, respectively. After the development of template-assisted carbonization processes, many template materials were proposed to control the pore structure. This process has the advantages of a high carbon yield because there is no scarification of matrix carbon by oxidation (gasification), a homogeneous pore size, easy control of the pore volume, high reproducibility of pore structure, etc. Activated carbons: The history of activated carbons goes back to the prehistoric era, when charcoals are known to have been used to purify water and as a medicine to stop diarrhea.

24 Chapter 1 (A) Granular activated carbon

Mesoopore Macropore

Micropore

(B) Activated carbon fiber

Micropore

Figure 1.28 Illustrations of the pore structure in (A) in granular activated carbon and (B) activated carbon fiber.

Granular activated carbons are now prepared from different precursors, including biomasses, and are used in a wide range of industries. Different processes of activation were employed, such as oxidation using diluted oxygen gas, air, water vapor, etc. (physical activation), and oxidation using ZnCl2 and KOH (chemical activation). Using KOH, an extremely high surface area of about 3600 m2/g was reported. In most carbon materials, macropores (>50 nm in size) and mesopores (2e50 nm) coexist with micropores, as shown schematically in Fig. 1.28A. Macropores and mesopores are influenced by the performance of their adsorption as pathways to micropores for adsorbates. Fibrous activated carbons, or activated carbon fibers (ACFs), have been prepared from different carbon fibers, such as PAN-based, isotropic pitchebased and phenol-based carbon fibers, the pore structure of which is schematically shown in Fig. 1.28B. ACFs have the advantage of a high adsorptionedesorption rate because almost all micropores are open at the surface of the fiber where adsorbate molecules can be adsorbed directly into micropores, in contrast to granular activated carbons, for which adsorbates have to pass through the macropore and mesopore to reach the micropore (Fig. 1.28A). Templated porous carbons: Template-assisted carbonization was successfully applied to prepare nanoporous carbons [135,136]. Table 1.2 summarizes templates classified as hard or soft, together with commonly used carbon precursors. The resultant carbons with pore structure characteristics (microporous or mesoporous, the representative BrunauereEmmetteTeller (BET) surface area, SBET, and the total pore volume, Vtotal) and the cycle performance of the templates. Ordered micropores are created by using zeolites. Ordered mesopores are made using mesoporous silicas and block copolymers, but MgO, colloidal silica, metal-organic frameworks (MOFs) etc. result in disordered mesopores, together with micropores. Representative SEM images of MgO-templated and mesoporous SBA 15etemplated carbons are shown in Fig. 1.29A and b, respectively. An MgO template has two

Introduction 25 Table 1.2: Pore control by template-assisted carbonization.

Template

Carbon precursors

Carbon synthesized

Template performance

Pore structure

SBET, Vtotal

Solvent

Cycle

Hard

Zeolites

Propylene, AN, FA

Microporous, ordered

3600 m2/g 0.6 mL/g

No

Hard

MgO (Mgacetate, citrate, Mg[OH]2) Mesoporous silica (MCM-48, MCM-41, SBA-1, SBA-15) Colloidal silica

PVA, pitches, PET, PIs, etc.

Mesoporous, disordered

1600 m2/g 1.7 mL/g

HF or conc NaOH Acetic or citric acid

sucrose, FA, PF, and MP

Mesoporous, ordered, or disordered

1520 m2/g 1.3 mL/g

HF or conc NaOH

No

RF, AN

2600 m2/g 1.4 mL/g

PF, etc.

HF or conc NaOH HCl, HF, etc. Sacrificial

No

Ni(OH)2, CaO, etc. Diblock and triblock copolymers MOFs (MOF-5, ZIF-8)

Microporous, mesoporous, disordered Mesoporous, disordered Mesoporous, ordered, or disordered Microporous and Mesoporous, disordered

Recycled for CaO No

Sacrificial

No

Soft

RF, EOA, PhF, AN FA, RF, PF

e 1354 m2/g 0.74 mL/g 2872 m2/g 2.06 mL/g

Easy recycle

AN, acrylonitrile; SBET, BrunauereEmmetteTeller surface area; EOA, triethyl orthoacetate; FA, furfuryl alcohol; MOF, metal-organic frameworks; MP, mesophase pitch; PET, poly(ethylene terephthalate); PF, phenol-formaldehyde; PhF, phloroglucinol-formaldehyde; PVA, poly(vinyl alcohol); RF, resorcinol-formaldehyde; Vtotal, total pore volume.

(B)

(A)

1 µm

100 nm

Figure 1.29 Scanning electron microscopy images of templated carbons: (A) MgO-templated and (B) mesoporous silica (SBA-15)-templated.

26 Chapter 1 advantages: it can be removed from the carbon matrix simply by washing with diluted acidic aqueous solutions, and it can be recycled, if needed, in contrast to other templates, zeolites, mesoporous silicas, and colloidal silicas, which need either HF or concentrated NaOH for removal, and block copolymers and MOFs, which are sacrificial. 1.2.4.5 Spherical carbon materials Various spherical carbon materials with a wide range of diameter have been developed that have different nanotextures based on their point orientation schema, whether radial or concentric (see Fig. 1.7). Fig. 1.30 shows the nanotextures of three representative carbon spheres. The concentric alignment of BSUs is found in carbon blacks (Fig. 1.30A), although the nanotexture becomes random at the center of the sphere. Radial alignment (Fig. 1.30C) is found in carbon spherules, which are formed from a mixture of poly(ethylene) and poly(vinyl chloride) by pressure carbonization [137]. An intermediate nanotexture, in which the alignment of BSUs is radial at the surface of the sphere but planar at the center, occurs in mesocarbon microbeads, as explained later. Carbon blacks: Carbon blacks are formed through incomplete combustion in the gas phase of either gaseous or mist-like hydrocarbons [138]. Carbon blacks are classified on the basis of the reaction process and the raw materials as furnace blacks, channel blacks, lamp blacks, thermal blacks, acetylene blacks and Ketjenblack. They are characterized by different-size spherical primary particles and their coalescence into aggregates (secondary particles). Carbon black has the smallest primary particle size with the most marked aggregation, as shown in TEM images in Fig. 1.31. Furnace blacks are produced from either the mist of creosote oil or natural gas consisting mainly of methane by their incomplete combustion at 1200e1400 C; they have marked aggregation of primary particles (called “structure” in the industry) (Fig. 1.31A). This is the main reason why they are used to reinforce rubber. Lamp blacks are manufactured by

Figure 1.30 Nanotextures in spherical carbons.

Introduction 27 (A) Furnace black

10 nm

(B) Thermal black

10 nm

(C) Acetylene black

(D) Ketjenblack

100 nm Figure 1.31 Transmission electron microscopy images of carbon blacks.

burning aromatic oils in shallow open pans with a limited air supply; they are mainly used for inks. Thermal blacks are produced by the thermal decomposition of natural gas; the size of the primary particles is usually large, as shown in Fig. 1.31B, and almost no aggregation is observed. Acetylene blacks are formed through the exothermic decomposition of acetylene gas at a relatively high temperature, 2400 C, in which the primary particles are highly aggregated (Fig.1.31C). Under conditions similar to those for furnace blacks, carbon blacks with high aggregation, compared with acetylene blacks, are produced and named Ketjenblack [139], as shown in Fig. 1.31D. Acetylene black and Ketjenblack are often used as electric conductive additives for electrodes in electrochemical devices because their well-developed aggregation and relatively high electrical conductivity.

28 Chapter 1 (A)

(B)

50 μm

20 μm

Figure 1.32 Mesophase spheres. (A) Polarized light micrograph of spheres formed in a pitch, and (B) scanning electron microscopy image of the spheres separated from matrix pitch (mesocarbon microbeads).

Mesocarbon microbeads: Anisotropic spherical particles are segregated from pitches during heating, as shown in Fig. 1.32A. They are called mesophase spheres and coalesce into so-called bulk mesophase to form coke [140]. The structure of mesophase spheres is close to a radial point orientation scheme, but in its center, the orientation of layers is planar, as illustrated in Fig. 1.30B [140,141]. Mesophase spheres formed in pitches are separated from an isotropic matrix using a solvent, either pyridine or quinoline [142], and are called mesocarbon microbeads (MCMB) (Fig. 1.32B). The graphitization behaviors were studied for two kinds of MCMBs. One type was separated from quinoline-soluble parts in coal tar pitch after heat treatment at 430 C for 120 min, and the other was from asphalt pitch after heat treatment at 430 C for 60 min [143]. After they were separated from the matrix pitch, MCMBs were graphitized by heat treatment at high temperatures. Their changes were almost the same as those of cokes prepared from the pristine pitches, as shown for coal tar and asphalt pitches in Fig. 1.33A and B, respectively. In other words, the graphitization behavior is governed by that of the mesophase spheres (MCMBs). 1.2.4.6 Glass-like carbons Glass-like carbons (glassy carbons) are produced by the carbonization of thermosetting resins, such as phenol-formaldehyde, poly(furfuryl alcohol), and cellulose, by very slow heating. They are characterized by an amorphous structure [144]. The heating rate during pyrolysis and carbonization of the precursor has to be slower than the rate of shrinkage to compensate for the evolution of the pores caused by the release of decomposition gases from the precursor block. Their properties are similar to inorganic glasses, such as high hardness, brittle conchoidal fracture, and gas impermeability. This is why they are called glass-like (or glassy) carbon, even though they are not transparent.

Introduction 29 (B)

(A)

100

100

d002 / nm

60 0.345

d002

40 0.340

0.335

20 1000

1400

1800 2200 HTT / oC

2600

3000

80

0.350

60 0.345 d002

Lc(004)

0.340

0.335

40

Lc(004), La / nm

80

d002 / nm

Lc(002)

0.350

La

Lc(002), La / nm

La

20

1000

1400

1800 2200 HTT / oC

2600

3000

Figure 1.33 Changes in x-ray diffraction parameters, d002, La, and Lc, with heat treatment temperature (HTT) for mesocarbon microbeads (open symbols) compared with coke prepared from pristine pitch (closed symbols) [143]: (A) prepared from quinoline solubles heated to 430 C for 120 min, and (B) from asphalt pitch heated to 430 C for 60 min.

In glass-like carbons, the BSUs do not grow with heat treatment above 3000 C. They have large d002 (about 0.344 nm) and very small Lc and La (about 3.5 nm for both), as shown in Fig. 1.9. The following experimental result shows how strongly graphitization of these carbons is depressed. When a rod of glass-like carbon was melted by directly passing an electrical current at a temperature above 3700 C under Ar gas pressure above 10 MPa, a graphite ball was obtained in a crater-like cavity at the middle of the rod, but the wall of the crater retained the characteristics of glass-like carbon, even though it had been heated to a temperature near the melting point of carbon [145]. For complete graphitization of these carbons, heat treatment under high pressure at 30 MPa at a temperature above 1600 C was needed [146,147].

1.3 Construction and purposes of the current book The purpose of this book is to provide a basic and through understanding of the preparation, structure, properties, and applications of graphene and its related materials. The book has nine chapters. In Chapter 1, Introduction, a short history of the term “graphene” is explained and the conditions for exhibiting its intrinsic characteristics are briefly discussed on the basis of the Raman spectrum, the field effect on electrical conduction, and extremely high thermal conductivity. Fundamental knowledge about the science and engineering of carbon materials is explained, followed by a description of some graphite materials that are frequently used as raw materials of graphene, such as natural graphite, HOPG, and kish

30 Chapter 1 graphite, and as the reference and competitive materials such as carbon nanotubes, activated carbons, and carbon blacks. In Chapter 2, Preparation, the procedure for preparation and conditions for CVD, cleavage (mechanical or chemical), and exfoliation through graphene oxide are summarized by focusing on structural perfection rather than the commonly used classification (i.e., top-down and bottom-up processes). The first two processes (CVD and cleavage) give less-defective layers in the resultant graphene compared with the last one (exfoliation). In the current book, flakes synthesized through CVD of organic precursors and prepared from graphite through mechanical and chemical cleavage are differentiated from those prepared from graphite through graphene oxide by oxidation, exfoliation, and reduction because they are defective, including some functional groups. In this book, therefore, the former is expressed as graphene, and the latter as reduced graphene oxide (rGO). From Chapters 3 to 7, the properties of graphene and rGO flakes are discussed in relation to their applications: electronic properties and applications in Chapter 3, chemical properties and applications in Chapter 4, mechanical properties and applications in Chapter 5, thermal properties and applications in Chapter 6, and biomedical properties and applications in Chapter 7. The authors expect that this may provide a better and easier understanding of both properties and applications, rather than independent explanations of each property and application in separate chapters. Chapter 3 discusses the electrical properties of graphene by referring to other carbon allotropes, carbon nanofibers, graphite, etc. Applications of graphene based on its electrical properties are summarized by dividing them into transistors, spintronic and sensor devices, photon detectors, etc. Chapter 4 summarizes various applications of graphene based on its chemical properties, including electrochemical properties, by classifying energy storage and environment remediation, after discussing the fundamental chemical properties of graphene. The section related to energy storage includes lithium-ion batteries, electrochemical capacitors, lithium-ion capacitors, solar cells, and fuel cells. The section on environment remediation includes adsorption, capacitive deionization for water desalination, catalysis, and chemical sensors. For electrochemical devices, batteries, capacitors, etc., the values of capacity and capacitance are not directly compared to avoid their misunderstanding, because the procedures and conditions of characterization are not unified internationally. In Chapter 5, the mechanical properties of graphene are summarized, followed by an explanation focusing on nanolubricants and mechanical sensors in comparison with other materials to demonstrate the advantages of graphene.

Introduction 31 Chapter 6 explains the thermal properties of graphene by focusing on its extremely high thermal conductivity and how important the number of layers stacked in the flake is for realizing this high conductivity. Applications for thermal interface materials, nanofluids, and thermal energy storage are also related to the high thermal conductivity of graphene. Chapter 7 explains biomedical properties demonstrating various applications, biocompatibility, cell management, delivery carriers for drugs, and biosensors. In Chapter 8, the materials coming after graphene (i.e., beyond graphene) are discussed by classifying graphene derivatives (graphane, fluorographene, graphene oxide, graphyne, etc.), single-layer compounds with layer structures (boron nitride nanotubes, transition metal sulfides, etc.) and their composites, and layer-by-layer composites. Chapter 9 presents the prospects for graphene by discussing what the materials science of graphene bring us: zero band gap, the scientific basis for the adsorption of molecules, extremely high thermal conductivity, biocompatibility, etc.

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CHAPTER 2

Preparation of graphene Chapter Outline 2.1 Chemical vapor deposition

41

2.1.1 Synthesis of graphene films 41 2.1.1.1 On platinum 42 2.1.1.2 On nickel 43 2.1.1.3 On copper 46 2.1.1.4 On other metals 48 2.1.1.5 On silicon carbide 49 2.1.1.6 On other metal carbides 51 2.1.1.7 On others 52 2.1.1.8 Transfer of graphene films 53 2.1.1.9 Structure analysis of graphene films 56 2.1.2 Synthesis of graphene flakes 60 2.1.3 Synthesis of single-walled carbon nanohorns 2.1.4 Substitutional doping of heteroatoms 69 2.1.5 Graphene foams 80

65

2.2 Cleavage (peeling) 81 2.2.1 Mechanical cleavage 81 2.2.2 Cleavage in solution 85 2.2.3 Cleavage via intercalation compounds

95

2.3 Exfoliation via graphene oxide 101 2.3.1 Synthesis of graphene oxide 102 2.3.2 Exfoliation of graphene oxide 105 2.3.3 Reduction of graphene oxide 107 2.3.3.1 Thermal reduction 107 2.3.3.2 Chemical reduction 111 2.3.3.3 Hydrogen reduction 115 2.3.3.4 Hydrothermal reduction 116 2.3.3.5 Irradiation reduction 117 2.3.3.6 Electrochemical reduction 121 2.3.4 Fabrication of reduced graphene oxide foams (sponges) 123 2.3.5 Functionalization of reduced graphene oxide 127 2.3.6 Substitutional doping of heteroatoms 133 2.3.7 Fabrication of transparent reduced graphene oxide films 135

Graphene. https://doi.org/10.1016/B978-0-12-819576-5.00002-5 Copyright © 2020 Elsevier Inc. All rights reserved.

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40 Chapter 2 2.4 Other processes 138 2.4.1 Chemical synthesis 138 2.4.2 Synthesis via pyrolysis 143 2.4.3 Unzipping of carbon nanotubes

2.5 Concluding remarks References 153

147

150

The preparation process has been classified into two methods: top down and bottom up. In top-down methods, pristine graphite flakes are separated into graphene flakes through the mechanical cleavage of crystalline graphite or the exfoliation of graphite oxide prepared from graphite through severe oxidation conditions followed by reduction. The former process (cleavage) has a much smaller possibility of introducing structural defects into the resultant graphene flakes (sheets), in contrast to the latter process (oxidation, exfoliation, and reduction), which introduces various structural defects associated with certain amounts of functional groups containing oxygen. In bottom-up methods, graphene films are produced by chemical vapor deposition (CVD) of various hydrocarbon gases on various substrates, which can result in high structural perfection, although it is difficult to produce a large quantity of graphene films. Because their properties, and consequently their applications, are influenced by structural defects in graphene layers, in this chapter the preparation of graphene is determined mainly by classifying processes that produce less defective graphene layers, including CVD and cleavage, and those that produce highly defective layers, including the exfoliation and reduction of graphite oxide. Fig. 2.1 schematically shows the fundamental concept of this chapter. Fig. 2.2 compares the structural perfection of single-layer graphene flake and reduced graphene oxide (rGO) flake with pristine graphite using Raman spectra. Graphene flake prepared by cleavage retains the high crystallinity of the pristine graphite, except for the slight appearance of a D-band, which is reasonably thought to result from a relative increase in the edge surface of the flakes. On the other hand, rGO loses the crystallinity of pristine graphite during the formation of graphene oxide (GO) and its exfoliation and reduction into rGO, as detected by the marked growth of the D-band, broadening of both the D- and G-bands, and disappearance of the G0 -band (two-dimensional [2D]-band). The spectrum for rGO is similar to that for pristine GO, which suggests that there is no structural improvement in the layer during reduction. GO flakes are known to have various oxygen-containing functional groups on the layers, which are mostly eliminated during reduction, but not completely, as explained in detail in Section 2.3. These structural and compositional differences influence the properties and application performance of the resultant graphene flakes and films. Therefore, in this book, graphene flakes and films synthesized by CVD and cleavage are differentiated from rGO flakes prepared through oxidation (the formation of GO), exfoliation, and reduction, as denoted by graphene and rGO, respectively.

Preparation of graphene 41 2.1 Chemical vapor deposition Films on various substrates, Flakes, Nanohorns, Foams, Doping. 2.4 Other processes

2.2 Cleavage

Chemical synthesis, Unzipping of CNTs.

Mechanical cleavage, Cleavage in solution, via intercalation compounds. Less defective layers

Graphene Highly defective layers

2.3 Exfoliation via graphene oxide Synthesis of graphene oxides, Exfoliation and reduction , rGO foams (sponges), Functionalization, Doping, Transparent films

Figure 2.1 Fundamental concept of this chapter. CNTs, carbon nanotubes; rGO, reduced graphene oxide.

Figure 2.2 Raman spectra: (A) Pristine graphite and single-layer graphene, and (B) pristine graphite, graphene oxide (GO), and reduced graphene oxide (rGO) via hydrazine reduction. 2D, two-dimensional.

2.1 Chemical vapor deposition 2.1.1 Synthesis of graphene films The synthesis of thin flakes of graphite through CVD onto the crystal surface of different metals and metal carbides is mostly performed in a chamber under high-level vacuum.

42 Chapter 2 Interactions have been observed between deposited carbon atoms and the substrate, such as the epitaxial growth of thin films and the precipitation of carbon dissolved into the substrate metal. Some films synthesized in this way were thin enough to be called graphene. The synthesis of thin graphite films on the surface of various metals was explored before the discovery of graphene [1e3]. 2.1.1.1 On platinum CVD onto the (111) plane of a Pt crystal (Pt [111]) was carried out in a scanning tunneling microscope (STM) by exposing it to ethylene (C2H4) gas under a pressure of 4.4  102 Pa at 300 K and then annealing at different temperatures in the microscope [4]. The formation of single-layer graphene was determined from the detailed analysis of STM images. A carbon layer was deposited on a clean Pt (111) surface by exposing it to approximately 1.3  104 Pa benzene (C6H6) at 1000 K [5]. A carbon layer with a (O19  O19)R23.4 degrees structure grew at the beginning of the exposure, and then islands of single-layer graphene with clear hexagonal shapes were formed. Hexagonal millimeter-sized single-crystal graphene films are grown on polycrystalline Pt foil by ambient-pressure CVD with a low concentration of methane (CH4) at 1040 C, as shown in Fig. 2.3 [6]. The graphene thus synthesized was separated from the substrate Pt by using it as the cathode of an electrolysis cell with an NaOH aqueous solution after being spin-coated with poly (methyl methacrylate) (PMMA) (bubbling method). The resultant graphene had high crystallinity with a low wrinkle height of 0.8 nm and a carrier mobility of greater than 7100 cm2/V. The collision of CH4 molecules with a Pt (111) surface under supersonic acceleration resulted in the formation of single-layer graphene

Figure 2.3 Hexagonal single-crystal graphene synthesized by ambient-pressure chemical vapor deposition for different periods at 1040 C under a flow of methane of 44 standard cubic centimeters per minute (sccm) with an H2 of 700 sccm [6].

Preparation of graphene 43 with no carbon by-products [7]. A Pt (111) surface after cleaning was exposed to ethylene up to saturation and then heated to about 1000 C under ultrahigh vacuum, which resulted in the decomposition of C2H4 and the formation of a single-layer graphene [8]. High-order commensurability of the resultant graphene lattice with the Pt lattice was detected by local conductivity variations observed through point-contact microscopy and atomic force microscopy (AFM). 2.1.1.2 On nickel

Resistance / kΩ

Two mechanisms for the growth of thin films composed of a few graphene layers were pointed out: carbon deposited by CVD from hydrocarbon gas on the substrate Ni (graphitic carbon) and carbon precipitated from the substrate Ni at a low temperature (carbidic carbon), because Ni metal can absorb a large amount of carbon into its interstitial sites at a high temperature [9]. CVD in a CH4/H2/Ar gas flow at 1000 C was performed on a thin foil of Ni (less than 300 nm thick) to make carbon precipitation as small as possible; this resulted in the formation of thin films [10]. After being transferred to SiO2 substrate, single-layer graphene had electron mobility greater than 3700 cm2/V at a carrier density of approximately 5  1012/cm2. The resultant films were reported to be conducting, transparent, stretchable, and flexible, and their crystalline quality was comparable to that of graphene flakes mechanically cleaved from graphite. Fig. 2.4 shows variations in resistance with bending for a film transferred to polymer substrate. Resistance

Curvature, κ / mm-1

Ry

Rx

Bending radius / mm

Figure 2.4 Variations in resistance perpendicular and parallel to bending directions, Rx and Ry, with bending radius [10]. Insets are the relation between anisotropy in resistance Ry/Rx and the curvature of the film and the appearance of the setup.

44 Chapter 2 parallel to the bending direction (Ry) increases with an increase in the bending radius, which is recovered by releasing the force, but that perpendicular to the bending direction (Rx) does not change. Thin carbon layers were obtained on an Ni(110) surface by reacting with CO at 600 K and 6.5  103 Pa [11]. A strong interaction between the precipitated carbon atoms and the Ni substrate was concluded from carbon near-edge electron-energy loss and surface extended energy-loss fine-structure analyses. Thin graphene films were deposited onto faceted Ni(755), which consisted of the terraces of the (111) face and the steps of a high-Miller index face [12]. Nanoribbons consisting of a few layers were synthesized by controlling the precipitation of carbon from Ni. An Ni layer about 500 nm thick was sputtered onto the substrate and then heated at 1300 C in a high vacuum (1.3  104 Pa), followed by slow cooling to 800 C (1 C/min). Upon slow cooling, the carbon atoms precipitated onto the surface of the Ni as thin films. High-quality transparent ribbons less than 8 nm thick, up to 350 nm wide, more than 3 mm long, and with a negligibly small D-band in the Raman spectrum were obtained. On the surface of polycrystalline Ni, thin carbon films consisting of 1e12 layers with a large area of about 1 cm2 were prepared by ambient-pressure CVD [13]. The films were composed of domains consisting of 1e2 layers with a lateral size of 1e20 mm. Few-layered graphene films 1  1 cm2 were synthesized on a polycrystalline Ni substrate by optimizing the temperature and time of the deposition and with a mixing ratio of acetylene of (C2H2)/H2 [14]. The graphene film obtained at 1000 C did not show the D-band in Raman spectrum and gave a G0 -band much stronger than the G-band. Both bands were very sharp, which suggested the formation of a high crystallinity graphene film with less than five layers. The mixing ratio of C2H2/H2 and deposition time influenced the crystallinity of graphene films, as shown in Fig. 2.5A and B, respectively, demonstrating that the optimum conditions are 1000 C for 5 min with a C2H2/H2 ratio of 2:45. Glassy graphene films (highly defective graphene films) were prepared on an Ni film by precipitating carbon dissolved into Ni [15]. A glassy carbon film was first formed on quartz substrate by carbonizing a spin-coated mixture of glucose and polyethylenimine at 1000 C, which was annealed at 850 C and 1000 C after depositing the Ni film onto the glassy carbon; this resulted in the formation of highly defective (glassy) graphene film and crystalline graphene, respectively. In Fig. 2.6AeC, the Raman spectrum and transmission electron microscopy (TEM) image with selected area electron diffraction (SAED) pattern are shown on the glassy carbon, the precipitated glassy graphene, and crystalline graphene, respectively. The glassy graphene film (Fig. 2.6B) demonstrates a strong D-band in Raman spectrum and a disordered structure in the TEM image and SAED pattern similar to the glassy carbon (Fig. 2.6B), which works as a carbon source for Ni substrate, whereas it provides a relatively strong G0 -band in Raman spectrum, similar to the highly crystalline graphene film, which is precipitated from the Ni substrate at a high temperature (Fig. 2.6C). The glassy graphene films thus prepared exhibited conductivity, transparency,

Preparation of graphene 45 (B) Intensity normalized by G-band/a.u.

Intensity normalized by G-band/a.u.

(A)

10/45 5/45 3/45 2/45

2600

HOPG 15 min 10 min 5 min

2800 2600 Raman shift / cm-1

2700 2800 Raman shift / cm-1

Figure 2.5 Two-dimensional Raman bands of graphene films deposited at 1000 C on a polycrystalline Ni substrate [14]: (A) with different mixing ratio of C2H2/H2 for 5 min, and (B) with different deposition times with a C2H2/H2 ratio of 2:45, compared with the band of highly oriented pyrolytic graphite (HOPG). a.u., arbitrary unit.

G

G’

Intensity / a.u.

(B’) G’

D+G

(C)

1000

D+G

D G

Intensity / a.u.

(B)

(A’)

D

Intensity / a.u.

(A)

G’

(C’)

G

3000 2000 Raman shift / cm-1

Figure 2.6 Glassy graphene film [15]: (A) Glassy carbon, (B) glassy graphene, and (C) graphene. (AeC) Raman spectra, and (A0 eC0 ) transmission electron microscopy images with selected area electron diffraction pattern insets. a.u., arbitrary unit.

46 Chapter 2 and flexibility comparable to those of graphene, as well as mechanical and chemical stability comparable to those of glassy carbon. Composite films of single-layer graphene and hexagonal boron nitride (BN) (h-BN) were prepared on an Ni(111) surface by CVD (i.e., BN/graphene/Ni and graphene/BN/Ni composite films); both graphene and h-BN had a relatively strong interaction with the Ni(111) surface [16e18]. Graphene was formed by heating the Ni/SiO2/Si substrate coated by Orange II (named Acid Orange 7, C16H11N2NaO4S) at 850 C for 5 min [19]. On the Ni layer, which was sputtered onto the SiO2/Si wafer heated at 500 C, the formation of single-layer graphene was confirmed from the Raman spectrum, showing a trace of the D-band and the strong G0 -band with I2D/IG of more than 5. This suggests an important role of the Ni layer for graphene growth. 2.1.1.3 On copper Single-layer graphene with a large area was grown on Cu foil by CVD of CH4 at up to 1000 C, as shown in Fig. 2.7A [20]. The films consisted of predominantly a single-layer, probably owing to the low carbon solubility of Cu (about 0.03%), in contrast to the large solubility of Ni (about 1.1%). A high-resolution scanning electron microscopy (SEM) image of graphene on Cu (Fig. 2.7B) shows that the grain boundaries and steps of Cu are associated with double- and triple-layered graphene flakes with wrinkles. Insets show TEM images of folded edges of single-layer graphene (1L) and two-layered graphene (2L). Wrinkles associated with the difference in the thermal expansion coefficients between Cu and graphene also cross the Cu grain boundaries, indicating that the

Figure 2.7 Graphene films formed by chemical vapor deposition of methane at 1000 C [20]: (A) scanning electron microscopy (SEM) image, and (B) high-resolution SEM image with transmission electron microscopy images of folded edges.

Preparation of graphene 47 graphene film is continuous. Polycrystalline Cu substrate was also used for graphene synthesis; the graphene grains show no definite epitaxial relation to the Cu substrate [21]. Using surface wave plasma CVD of CH4 mixed with Ar and/or H2 on a Cu or Al foil substrate at low-temperature such as 300e400 C, transparent conductive films with a large area (23  20 cm2) were synthesized consisting of a few graphene layers [22]. Patterned graphene film was grown on a patterned Cu layer on an SiO2/Si substrate by CVD of ethanol (C2H5OH) vapor at 700e800 C, which was applied to the electrodes for fieldeffect transistors [23]. Graphene was grown on the (100) face of a high-purity Cu single crystal by CVD of CH4, demonstrating the formation of defect-free graphene [24]. Graphene films grown on Cu foil from C2H5OH were reported to possess a lower defect density and better uniformity than those grown from pentane (C5H12) [25]. Films derived from C2H5OH have a sheet resistivity of 2700 U/sq and a Hall mobility of 110 cm2/V, compared with 5000 U/sq and 65 cm2/V, respectively, for C5H12-derived films. The high growth rate of large single-crystal graphene films by CVD of CH4 at 1000 C was possible by using Cu foil placed above an oxide substrate with a gap of about 15 mm [26]. It reached as high as 60 mm/s, in contrast to rates of less than 0.4 mm/s (0.03e0.36 mm/s) [27e31]. The acceleration of the growth of graphene by oxygen was explained by the significant increase in CH4 dissociation to CH3 radicals via the reaction with oxygen on the Cu(100) surface. Oxygen was supplied from the oxide substrate through an optimized gap [26]. Sequential oxygen passivation during CVD nucleation and growth made it possible to grow millimeter-size single-crystalline graphene. The rate of growth reached about 100 mm/min [32]. In a vacuum of about 1 Pa at 1040 C, where Cu foil was annealed, the first O2 passivation of about 0.1 standard cubic centimeters per minute (sccm) was performed, followed by the introduction of CH4/H2 with 1/300 for nucleation, and then the second O2 passivation followed by CH4/H2 with 1/5 was applied to grow the graphene layer. As compared in Fig. 2.8A, the application of O2 passivation before both nucleation and growth processes resulted in the formation of large graphene

Figure 2.8 Scanning electron microscopy images of graphene grown on Cu foil [32]: (A) with O2 during both nucleation and growth, (B) without O2 during nucleation, and (C) without O2 during growth.

48 Chapter 2 flakes as large as 2 mm, with fewer flakes. Without O2 at either nucleation or growth (Fig. 2.8B and C, respectively), the nucleation density of graphene flakes becomes much higher but they are much smaller, about 2 and 20 mm. Density functional theory (DFT) calculation predicted that this acceleration effect of O2 was caused by a significant reduction of the adsorption energy of carbon atoms on the oxygen-covered Cu surface. For the synthesis of carbon nanotube (CNT) arrays (or forests) via catalytic CVD, a pronounced acceleration of growth rate was reported by either selecting alcohols as the precursors or adding water vapor to the precursor gases [33,34]. Graphene films were deposited on a Cu foil by CVD at 1000 C CO2 mixed in H2 gas [35]. CO2 gas was activated using an Ni/Al2O3 catalyst at 200 C before mixing into the H2 flow with 200 sccm. The morphology and lateral size of the graphene film changed with the increase in the flow rate of CO2, from a hexagonal film 6 mm in size for a 5-sccm rate of CO2 to a round one 1 mm in size for 30 sccm. The maximum mobility for holes and electrons determined by field-effect transistors of the resultant graphene reached 3010 and 750 cm2/V, respectively. 2.1.1.4 On other metals Thin carbon layers were grown on an Ir (111) surface by CVD at a low pressure of C2H4 at 1120e1320 K [36]. The flake consisted of single-layer graphene and extended over the terraces and step edges of the Ir substrate. On the other hand, epitaxial growth of thin films was observed on an Ru (0001) surface by the precipitation of interstitial carbon in Ru metal through slow cooling from 1150 C to 825 C [37]. The first deposited graphene layer interacted strongly with the metal substrate, but the second layer completely detached from the first layer. It was a lens-shaped flake with a size of more than 100 mm and consisted of one to a few layers. As shown in Fig. 2.9, Raman spectrum demonstrated the characteristics of double-layered graphene and Raman maps for G- and G0 -bands suggesting the formation of double-layered graphene flakes with a large area, because the center position and width of these bands remained constant over large areas. Defective graphene layers were reported to deposit on (100) (111), and (110) faces of single-crystal Si at 900e930 C in a flow of a gas mixture of 180 sccm CH4 and 10 sccm H2, in which the Si substrate was mounted face down on the heating stage with a narrow gap between the Si and the heating stage to increase the local concentration of hydrocarbon radicals and C ions at the gap [38]. The thickness measured by AFM showed the formation of single-layered and double-layered deposits, but their Raman spectrum gave a strong D-band and weak G0 -band, which suggested a defective structure. Deposition of “graphitic” carbon and “carbidic” carbons on clean and potassium-covered Co(0001) surfaces was reported through the adsorption of C2H2 molecules followed by heating above 420 K [39]. Metallic alloys were also successfully used as substrates for CVD growth of highly crystalline graphene films [40,41]. A millimeter-sized single-crystalline, single-layer

Preparation of graphene 49 (A)

(B)

(C)

Intensity / a.u.

G’

1000

G

3000 2000 Raman shift / cm-1

Raman shift / cm-1

Raman shift / cm-1

Figure 2.9 MicroRaman characterization of double-layered graphene on Ru (0001) [37]: (A) Spectrum, (B) map by G-band, and (C) map by G0 -band for two adjacent double-layered graphene islands. a.u., arbitrary unit.

graphene film was synthesized on a substrate of SiPt alloy by CVD in the flow of CH4 with 44 sccm and H2 with 600 sccm for 20 min. The growth rate reached about 1.5 mm/s at 1155 C [40]. The alloy Pt silicide was prepared by the reaction between Pt deposited on SiO2 and tetraethyl orthosilicate gas at 700 C. On a foil of CueNi alloy prepared by electroplating Ni on Cu foil 20 mm thick, followed by annealing at 1050 C in an H2/Ar atmosphere, single-layer graphene film with a size of about 1.5 inches was successfully synthesized by CVD in CH4/Ar at 1050e1100 C within 2.5 h (growth rate of 4.2 mm/s) [41]. 2.1.1.5 On silicon carbide Graphene has been known to be grown epitaxially on different crystal faces of SiC single crystals and its structure and growth mechanism have been extensively studied [42e58]. Graphene formation on an SiC substrate was described as follows. The SiC surface is first reconstructed to a (O3  O3)R30 degrees structure and then to a (6O3  6O3)R30 degrees structure (carbon nanomesh) after annealing at 1400 C; finally, the graphene and/or graphite film forms on carbon nanomesh. By heating SiC single crystals at 2000 C at an ambient pressure of Ar, carbon film consisting of areas with different thicknesses were formed on the (0001) crystal face. Electron reflectivity spectra are shown in Fig. 2.10A for the four selected areas. They demonstrate that these four areas correspond to graphene films with thicknesses of one to four layers. The microphotoelectron spectra collected using a photon energy of 450 eV at different graphene thicknesses (Fig. 2.10B) show that the C (1s) spectrum consists of three components at binding energies of 283.4, 284.4, and 284.9 eV, corresponding to bulk carbide, graphene, and interface (carbon nanomesh with a [6O3  6O3]R30 degrees structure). With an increase

50 Chapter 2 (A)

(B) 4 layers

3 layers

2 layers

3 ML

1 ML Graphene 1 layer

0.5 ML Interface

3 5 7 1 Electron energy / eV

287

Carbide

285 283 Binding energy / eV

Figure 2.10 Graphene films formed on 4HeSiC(0001) [56]: (A) Electron reflectivity spectra for four areas with different thicknesses, and (B) C (1s) of microphotoelectron spectroscopy for three areas with different thicknesses.

in the thickness of graphene films, the contributions from both substrate SiC and interface to the C (1s) spectrum decrease markedly. Thin films were grown on the (0001) face of 6HeSiC [46] and also of 4H-SiC [58] by CVD and their electronic transport properties were determined. They were shown by low-energy electron diffraction (LEED) and STM measurements to grow epitaxially on the substrate, and by Auger spectroscopy to consist of a few layers. The properties of these films could be explained by the existence of a 2D electron cloud with large anisotropy, high mobility, and 2D localization. By vaporization of Si, thin films were epitaxially grown on the (0001) face of single-crystal SiC [48] and on the (0001) faces of both Si-terminated and C-terminated sides of 6H- and 4H-SiC crystals [42,50]. On the C-terminated (0001) face, the growth of CNTs caused by Si vaporization was also reported [59]. The thermal decomposition of a 6HeSiC wafer at increasing temperature from 1080 C to 1320 C led to the layer-by-layer growth of unconstrained, heteroepitaxial thin carbon films [45]. Epitaxial graphene grown on SiC substrate was studied by Raman spectroscopy [60]. Si-terminated 6HeSiC(0001) prepared by heating at 850 C under a silicon flux for 2 min in ultrahigh vacuum was annealed at 1300 C several times. Graphene synthesis on C-terminated 6HeSiC(0001) was performed by annealing under the same conditions without a silicon flux pretreatment. Significant blue-shifts of the Raman bands, D, G, and G0 , from the bands for both bulk graphite and micromechanically cleaved graphene

Preparation of graphene 51 (A) D

(B) G SiC

G’

G’

(C)

D

D G

SiC Double-layered G

Single-layered SiC substrate

Double-layered graphene

2550 2700 2850

Mechnically cleaved single-layered graphene

Graphite

SiC D

G SiC

G’ Larger layer number

on C-terminated-SiC G’ D+G Double-layered on Si-terminated-SiC

Single-layered SiC substrate

Figure 2.11 Raman spectra of epitaxial graphene films grown on 6HeSiC(0001) [60]. (A) Blue shifts of Raman bands caused by constraint of the substrate (B) Effect of the crystal face of the substrate SiC, either C- or Si-terminated. (C) Effect of the thickness of the graphene films grown. a.u., arbitrary unit.

were observed, as shown in Fig. 2.11A, which were attributed to the compressive strain caused by the SiC substrate. Double-layered graphene on C-terminated SiC has a stronger D-band and a little broader G-band, revealing a lower crystallinity than those on Si-terminated SiC, as shown in Fig. 2.11B. All Raman bands shift to lower frequencies with increasing thickness of the epitaxial graphene layer, as shown in Fig. 2.11C. This suggests that the effect of the substrate on epitaxial graphene becomes weaker with increasing thickness, and the graphene lattice relaxes. 2.1.1.6 On other metal carbides Single-layer graphene was synthesized on the face of TiC(111) by CVD of C2H4 under a pressure of 8  103 Pa at 1100 C [61] and at 1250 C [62,63]. The strength of the CeC bond in single-layer graphene was weakened because of orbital hybridization between the graphene layer and the substrate TiC(111) face [61]. On the TiC(111) face, a film consisting of two crystalline domains with a lateral size of about 300 nm was obtained [62]. Epitaxial growth of a graphite thin film on TiC(111) [64] and on TiC(755) [65] was also reported. On the TaC(111) face, single- and double-layered graphene films were formed by CVD of C2H4. The first layer was deposited at 1300 C to make the layer with good crystalline quality, and then the second layer was formed at 1000 C [66]. Layer-by-layer formation of the carbon film was observed by measuring the relative intensity ratio of the x-ray photoelectron spectroscopy (XPS) C (1s) peak as a function of exposure L (1 Langmuir ¼ 1  106 Torr s), as shown in Fig. 2.12. Compared with the formation of the first layer, the second one grew slowly and the growth rate of the third layer was much slower. The electronic states of the single-layer film were different from those of

52 Chapter 2

Figure 2.12 Change in intensity ratio of x-ray photoelectron spectroscopy C (1s) peak for graphitic carbon layer to that for substrate TaC with exposure to ethylene gas [66].

bulk graphite owing to hybridization of the p orbitals of the deposit with the d orbitals of the substrate. The lattice constant measured from the LEED pattern was 0.249 nm for the single-layer film and 0.247 nm for the double-layered film, which suggested that the interaction between the film and the substrate became weak upon the formation of the second layer. 2.1.1.7 On others Catalytic CVD of CH4 was carried out at 900 C on highly oriented pyrolytic graphite (HOPG) after the deposition of Fe by an electron beam sublimator at room temperature [67,68]. The same process was applied to heal structural defects in the topmost layer of graphite [69]. Diamond nanoparticles prepared by electrophoretic deposition onto the HOPG substrate could be converted to single-layer graphene by heating to 1600 C; the single-layer graphene had a lateral size of 10e15 nm and was placed epitaxially on the substrate [70]. Changes in structure and p-electron state during the gradual conversion from nanodiamond to single-layer graphene were discussed [71]. ZnS ribbons formed on Si were successfully used as a substrate (template) to synthesize graphene nanoribbons consisting mainly of 10 layers by CVD of CH4 at 750 C [72].

Preparation of graphene 53 2.1.1.8 Transfer of graphene films Transfer of patterned graphene films synthesized by CVD on different metals and metal carbide substrates to a desired substrate is an importance process for producing graphenebased devices because risks introducing additional structural defects into graphene films. Transfer methods that were reported involve stamping polydimethylsiloxane (PDMS) [10], PMMA [6,73e76], or polyvinyl alcohol (PVA) [77] film onto the graphene film, followed by etching of the underlying metallic substrate in different etchant solutions. Then, it is transferred to a polymer substrate, followed by thermal annealing for better adhesion of graphene to the new substrate. These polymer substrates are usually dissolved using different organic solvents such as acetone, leaving the graphene clinging to the desired substrate. Procedures using PDMS and buffered oxide etchant (BOE) are illustrated in Fig. 2.13A and B, respectively [10]. Using BOE or hydrogen fluoride (HF) solution, silicon dioxide layers can be removed, so the patterned graphene with the nickel substrate floats on the solution surface. After transfer to a desired substrate, further reaction with BOE or HF solution can completely remove the remaining Ni substrate, as illustrated in Fig. 2.13C.

Figure 2.13 Scheme for synthesizing, etching, and transferring patterned graphene film on Ni substrate [10]. BOE, buffered oxide etchant; HF, hydrogen fluoride; PDMS, polydimethylsiloxane; RT, room temperature.

54 Chapter 2

Figure 2.14 Scheme of transfer process via coating of poly(methyl methacrylate) (PMMA) film [73].

The quality of graphene films deposited on the Ni substrate was shown to be maintained by measuring the D-band in the Raman spectrum. The transfer procedure using PMMA is principally the same as for using PDMS, as shown in Fig. 2.14 [73]. On PMMA/graphene after etching the metal substrate, an appropriate amount of PMMA solution was dropped, in which the coated PMMA was partly or fully dissolved. This redissolution of the PMMA seeming to relax the underlying graphene film mechanically and to have better contact with the substrate [74]. The graphene films were detached from the Pt substrate after spin-coating of PMMA using H2 bubbles formed by the electrolysis of water in diluted NaOH electrolyte at the cathode [6]. The transparent graphene films after spin-coating of PMMA were directly used to construct a graphene/Si solar cell. The PMMA film on the graphene worked as an optically antireflective film in addition to transferring film for graphene [76]. To apply this composite film to Schottky junction solar cells, it is not necessary to remove polymer and or to add antireflection coatings (e.g., TiO2). A one-step dry, etching-free transfer method was also proposed, which consisted of drop-casting of an aqueous solution of PVA onto the surface of a CVD-grown graphene film on Cu, followed by heating at 100 C for about 15 min to evaporate water. This resulted in the formation of a solid film of PVA [77]. Finally, the PVA film was peeled from the Cu substrate together with the graphene film using a piece of pressure-sensitive tape attached to an edge of the upper surface of the PVA film. The Raman spectrum for the PVA film with graphene shown in Fig. 2.15 suggests that graphene transferred to PVA exhibited characteristics in the Raman spectrum of a strong G0 -band, much stronger than G-band, and sharp G- and G0 -bands; in other words, the graphene film synthesized by CVD maintains its high crystallinity even after it is transferred to PVA film. An efficient roll-to-roll transfer process of graphene film formed on a flexible Cu foil was developed, which consisted of three steps, as shown schematically in Fig. 2.16: adhesion of polymer supports, copper etching, and then dry transfer printing onto a target substrate [78].

Intensity / a.u.

Preparation of graphene 55

G’

G Graphene/PVA PVA

500

1000

1500

2000

2500

3000

3500

Raman shift / cm-1 Figure 2.15 Raman spectra for graphene/polyvinyl alcohol (PVA) transferred from Cu substrate and PVA [77]. a.u., arbitrary unit.

Graphene on polymer support

Polymer support

Graphene on Cu foil

Released polymer support

Target substrate Cu etchant

Graphene on target

Figure 2.16 Scheme for roll-to-roll transfer process of graphene formed on flexible Cu foil [78].

Graphene film formed by CVD on a flexible Ni substrate was also transferred to a flexible poly(ethylene terephthalate) (PET) substrate by the roll-to-roll process [79]. The graphene film had a sheet resistance of about 125 U/sq and an optical transmittance of 97.4%, but four-layered films prepared by repeating the roll-to-roll process four times on the same substrate had about 30 U/sq and about 90%. Direct transfer of graphene films from metal substrates to selected target substrates without using polymer-coating was proposed [80]. The process is illustrated in Fig. 2.17 using the example of transfer to a TEM metal grid and comparing it with the standard process using PMMA. This procedure avoids several wet chemical steps and introduces mechanical stresses and chemical contaminations, which may occur in the transfer process using PMMA and polyoxymethylene described earlier.

56 Chapter 2 Transfer by using PMMA

Direct transfer

FeCl3 tech FeCl3 tech

PMMA: polymethyl methacrylate IPA: isopropyl alcohol DI: deionized water

Figure 2.17 Direct transfer process compared with standard process using poly(methyl methacrylate) (PMMA) [80].

2.1.1.9 Structure analysis of graphene films Annular dark-field (ADF) imaging techniques in an aberration-corrected scanning transmission electron microscope (STEM) optimized for low-voltage operation at 80 kV were proposed to resolve and identify the chemical type of every atom in single-layer graphene and h-BN that contains substitutional defects [9,81e83]. An acceleration voltage of 80 kV for STEM measurements is essential so as not to cause damage to the graphene structure. In Fig. 2.18, ADF images observed on graphene films are shown with the overlay of the lattice [81,83]. From a single-layered to a double-layered region (Fig. 2.18A and B), the second layer (shown in blue) is stacked on the first layer (extended layer for an upper single-layer region, shown in red) by offsetting in the Bernal stacking (AB). At a boundary between two grains with a relative misorientation of 27 degrees (Fig. 2.18C and D) and at a defective region in single-layer graphene (Fig. 2.18E and F), the formation of pentagons, heptagons, and distorted hexagons is recognized.

Preparation of graphene 57

Figure 2.18 Annular dark-fieldescanning transmission electron microscope images of graphene. (A) Images from single-layered to double-layered regions [81], (C) at a grain boundary [83], and (E) at a defective region [81]. (B, D, and F) Corresponding image with an overlay of the lattice.

Grains in single-layer graphene were visualized using the dark-field technique of TEM [83]. Fig. 2.19 shows consequent applications of bright-field, electron diffraction, and dark-field techniques to explain how to visualize grains in single-layer graphene. Site-specific single-atom spectroscopy at a graphene edge was performed in a low-voltage scanning transmission electron microscope and the energy-loss near-edge fine structure of carbon at the edge of graphene was determined [84], which may reveal a new insight into the electronic and bonding structures at the edge of graphene and other carbon materials. The edge structure of graphene was reviewed in relation to its preparation process [85]. Electron beam irradiation introduces structural defects into the graphene layer, which could be detected by microRaman technique [86]. As shown on single-layer graphene in Fig. 2.20A, D- and D0 -bands, which are known to be caused by structural defects, are markedly developed with the irradiation of a 20-keV (medium-energy) electron beam for a short time (7 min); further irradiation introduces broadening of both G- and D-bands. By the first few minutes, the intensity ratio of ID/IG attains a maximum and then decreases gradually with an increase in irradiation time, as shown in Fig. 2.20B. The double-layered graphene behaved similarly to the single-layer graphene. These results suggest that the increase in structural defects with the increase in the irradiation dose of electron beam may

58 Chapter 2

Figure 2.19 Imaging of grains in single-layer graphene film using dark-field technique of transmission electron microscopy [83]. (A) Graphene looks uniform in bright-field image. (B) Its diffraction pattern reveals a polycrystalline nature. (C) Dark-field image corresponding to the aperture in diffraction pattern (B), showing the grains in real space. Using (D) several different aperture locations and color-coding, (E) Dark-field image overlay depicting the shapes and lattice orientations of several grains. (F and G) Images of regions where many grains emanate from a few points. Scale bars indicate 500 nm. 2.5

(A)

(B) 2.0

G

Single-layer

Intensity / a.u.

D 30 min

1.5 ID/IG

15 min

1.0 D’

Pristine

1200

1400 1600 Raman shift / cm-1

Double-layered

7 min

1800

0.5 0.0

0

10

40 30 20 Irradiation time / min

50

Figure 2.20 Irradiation of 20-keV electron beam on graphene [86]. (A) Change in Raman spectrum of single-layer graphene. (B) Dependence of ID/IG for single-layered and double-layered graphene films on irradiation time. a.u., arbitrary unit.

Preparation of graphene 59 (A)

4

(B)

1014

ID/IG

Intensity / a.u.

3 1013 D

G 1012

2

D’

1 1011 Pristine

1200

1400 1600 Raman shift / cm-1

0

0

5

10

LD / nm

15

20

25

Figure 2.21 Arþ bombardment of single-layer graphene on SiO2 substrate [88]; (A) change in Raman spectrum with increasing irradiation dose of 1011 to 1014 Arþ/cm2 for a single-layer graphene, and (B) change in ID/IG ratio with average distance LD between point defects for three different single-layer graphene samples. a.u., arbitrary unit.

make the structure of graphene amorphous. Irradiation of 500 eV Neþ and Heþ created defects in graphene on the SiO2 substrate and caused a significant change in the intensity of the D-band as well as a marked decrease in the mobility of charge carriers [87]. The bombardment of a low-energy (90 eV) Arþ ion beam created point defects in single-layer graphene D- and D0 -bands, increasing their intensities with increasing irradiation dose from 0 to 1013 Arþ/cm2. Then, it quickly broadened (amorphization) [88], as shown in Fig. 2.21A. The irradiation dose of 1011 Arþ/cm2 corresponds to one defect per 4  104 carbon atoms, and that of 1015 Arþ/cm2 leads to full disorder. The defect density s was determined by directly counting defects in the STM image, which was converted to the averaged distance between defects LD by using the formula LD ¼ s1/2, which was employed as a quantitative parameter to evaluate the defective structure of graphene after Arþ irradiation [88e91]. In Fig. 2.21B, ID/IG is shown as a function of LD, measured on three independent single-layered graphene films prepared by mechanical cleaving of an HOPG onto an SiO2/Si substrate. The observed relation was discussed using a model of the point defect, consisting of a structurally disordered region at the center of the defect and an activated region at the periphery. A simulated prediction was obtained, as shown by the solid line in Fig. 2.21B [88]. At a low defect density, i.e., LD of larger than 5 nm, the area contributing to Raman scattering is proportional to the number of defects, giving a simulation of ID/IG ¼ 102/L2D. Upon an increase in the defect density, i.e., LD of less than 5 nm, the activated regions of each defect start to overlap and eventually saturate, which leads to a decrease in D-band intensity after passing through a maximum. The effect of Arþ-ion irradiation was studied on graphene films prepared

60 Chapter 2 4

3

Single-layer

ID/IG

Double-layered Triple-layered

2

Multi-layered (HOPG) 1

0 1011

1012

1013

1014

1015

Ion dose / Ar+/cm2

Figure 2.22 Changes in ID/IG with Arþ-ion dose as a function of layer numbers [90]. HOPG, highly oriented pyrolytic graphite.

by mechanical cleavage of HOPG as functions of the irradiation dose and the number of stacked graphene layers [89,90]. As shown in Fig. 2.22, ID/IG increases with an increase in the irradiation dose up to about 1013 (corresponding to a decrease in LD to 4e5 nm in Fig. 2.22B) more markedly for single-layer graphene. In other words, more point defects are introduced to the layer of single-layer graphene. After passing the 1013 dose, ID/IG decreases to almost the same ID/IG values at the 1015 dose for all graphene films. To measure the Raman spectrum on single-layer graphene after Arþ irradiation, three laser lines, an activation voltage EL (corresponding wave length lL) of 1.58 eV (785 nm), 1.96 eV (632.8 nm), and 2.41 eV (514.5 nm), were employed [91]. Raman spectra measured by three laser lines on five single-layer graphene films showed similar changes with the irradiation dose (expressed by LD), and the same tendency to depend on ID/IG on LD was observed with different EL’s, as shown in Fig. 2.23.

2.1.2 Synthesis of graphene flakes In a microwave plasma-enhanced CVD (MPECVD) system with a quartz tube to generate the plasma, the growth of carbon layers is vertically aligned to sapphire substrate coated with the catalyst, such as NiFe and CoFe, 20e100 nm thick [92]. Typical SEM images of vertically grown carbon layers are shown in Fig. 2.24, called carbon nanowalls. The gases used were CH4 and H2 with a flow rate of 10 and 40 sccm, respectively, and the substrate coated by the catalyst layer was set as the anode under a direct current (DC) bias of e185 V. In early CVD of CH4, carbon nanoribbons started to grow across the grains of the catalyst and eventually developed into walllike structures

Preparation of graphene 61 (A)

(B) LD=2 nm

LD=5 nm LD=7 nm

LD=2 nm

Intensity / a.u.

Intensity / a.u.

Intensity / a.u.

LD=2 nm

1200

(C)

LD=5 nm

LD=7 nm

LD=5 nm LD=7 nm

LD=14 nm

LD=14 nm

LD=14 nm

LD=24 nm

LD=24 nm

LD=24 nm

1600 1400 Raman shift/cm-1

1200

1600 1400 Raman shift/cm-1

1200

1400 1600 Raman shift/cm-1

(D) 16

ID/IG

12 EL=1.58 eV (λL=785 nm)

8

EL=1.96 eV (λL=632.8 nm)

4 0 0

EL=2.41 eV (λL=514.5 nm)

10

20 LD / nm

30

40

Figure 2.23 Raman spectrum data of single-layer graphene irradiated by Arþ measured using different laser activation energies [91]: (A) EL ¼ 1.58 eV (lL ¼ 785 nm), (B) EL ¼ 1.96 eV (lL ¼ 632.8 nm), (C) EL ¼ 2.41 eV (lL ¼ 514.5 nm), and (D) ID/IG versus LD. a.u., arbitrary unit.

(A)

(B)

100 nm

100 nm

Figure 2.24 Scanning electron microscopy images of carbon nanowalls grown in microwave plasma-enhanced chemical vapor deposition of CH4/H2 mixed gas on a sapphire substrate coated by NiFe (40 nm thick) [92].

62 Chapter 2

Figure 2.25 Graphene nanoflakes grown vertically to the substrate [94].

when all of the ribbons met with one another. The top edges of the nanowalls were either folded double-layered or unfolded single-layer graphene, which suggests that most nanowalls were hollow with a nanometer-size space. Graphene nanoflakes grown almost vertically to various substrates, such as Si, Ni, and graphite, were synthesized by radio-frequency (RF) plasma-enhanced CVD of CH4 diluted in H2 with a substrate temperature Ts of 600e900 C [93e95]. SEM images of the nanoflakes are shown in Fig. 2.25. At the beginning, a 1- to 1.5-nm-thick carbon layer was deposited parallel to the substrate surface; then, the leading edge of the top layer curled up, becoming vertical [94]. An increase in the substrate temperature caused the growth rate of nanoflakes to increase and the vertical growth of nanoflakes to be disturbed. A low turn-on field of 4.7 V/mm for electron field emission was reported on these vertically grown graphene nanoflakes [93]. An electric double-layer capacitor with these nanoflakes grown directly on metal current collectors exhibited efficient filtering of 120 Hz current [95]. Few-layered graphene flakes were synthesized by RF catalytic CVD on the FeeCo/MgO catalyst using C2H2 at 1000 C, although optimization of the reaction conditions was required to minimize the formation of CNTs and amorphous carbon particles [96]. The as-produced graphene flakes were composed of 1e5 layers overlaid side by side, with a lateral size of 100e110 nm, as shown in Fig. 2.26. CVD synthesis of thin flakes was carried out with no substrate by passing liquid ethanol droplets into microwave-assisted Ar plasma, as shown in Fig. 2.27A [97]. The powder collected on the filter was easily dispersed in methanol by sonication, in which the presence of single-layered and double-layered graphene flakes was confirmed by TEM and electron energy-loss spectroscopy (EELS), but most flakes were strongly rippled (Fig. 2.27C). The Raman spectrum obtained from the synthesized flakes exhibited a symmetrical, sharp G0 -band at around 2670 cm1, revealing that the analyzed region of the flakes consisted of single-layered graphene, and the sharp G-, D0 -, and D-bands suggested high crystallinity of the flakes.

Preparation of graphene 63

Figure 2.26 Few-layered graphene nanoflakes synthesized by radio-frequency catalytic chemical vapor deposition [96].

Figure 2.27 Synthesis of graphene flakes by microwave-assisted Ar plasma [97]: (A) scheme of synthetic equipment, (B) graphene flakes dispersed in methanol, and (C) transmission electron microscopy image of the synthesized flakes.

64 Chapter 2

2 μm

500 nm

Figure 2.28 Scanning electron microscopy images of graphene flowers grown by chemical vapor deposition under high pressure of Ar. Courtesy of Mr. K. Muramatstsu of Incubation Alliance, Inc., Japan.

Via CVD of the decomposition gases of polymers under high-pressure Ar gas, the aggregates of graphene flakes were synthesized and were called “Graphene Flowers” (trade name) because of their appearance, as shown in Fig. 2.28 [98,99]. Petals of the flower consisted of a few graphene layers; their lateral size was 1e20 mm. Their growth rate was very high so as to be synthesized on a large scale and to be on the market as a dispersion in a solvent. DC arc discharge between graphite electrodes in H2/He yielded graphene flakes consisting of 2e4 layers, together with multiwalled CNTs (MWCNTs) in the wall of the chamber [100]. The discharge current was in the range 100e150 A with a maximum open circuit voltage of 60 V. The arc was maintained by translating the cathode to keep a constant distance of 2 mm from the anode. Arc discharge in He and NH3 mixed gas at 0.1 MPa pressure resulted in the formation of N-doped graphene flakes [101]. The flakes obtained (Fig. 2.29) were thought to be composed of 2e6 layers by high-resolution TEM

Figure 2.29 N-doped graphene flakes synthesized by direct current arc discharge in He/NH3 [101].

Preparation of graphene 65 observations and the N content was about 1%. Few-layered graphene flakes were synthesized by external magnetic fieldeassisted electric arc discharge, carbon vapor sublimated from a graphite anode in an Ar atmosphere deposited on a graphite cathode, and deposits consisting mainly of graphene flakes and CNTs [102]. Discharge was performed by 22 V arc voltage, 170 A arc current, and an Ar gas pressure of 0.07 MPa. By selecting the gap between the electrodes to be 60 mm, graphene flakes were obtained with few CNTs and carbon nanoparticles.

2.1.3 Synthesis of single-walled carbon nanohorns Characteristic dahlia flowerlike e aggregates of single-layered graphene were synthesized by either CO2 laser ablation or arc discharge under various atmospheres and called single-walled carbon nanohorns (SWCNHs) [103e106]. Two morphologies of the aggregates were differentiated: so-called dahlia type and bud type, as shown in Fig. 2.30A and B. By CO2 laser ablation of graphite substrate (purity of 99.99%) with no catalyst in the 0.1-MPa buffer gas of Ar, dahlia-type SWCNHs with a uniform size 80 nm in diameter were synthesized [103]. The aggregate consisted of horns of single-layer graphene with a diameter of 2e4 nm and length of 30e50 nm, and closed tips with an average cone angle of 20 degrees, as shown in Fig. 2.31A. With increasing laser beam intensity, SWCNHs became shorter in horn length and fewer in number, and then nanotubelike structures about 2 nm in diameter were densely packed inside the particles. Although the purity of as-prepared SWCNHs was already more than 95%, further purification of flowerlike SWCNH aggregates was possible by centrifugal separation using a colloidal suspension of SWCNHs in ethanol (1 mg of aggregates per 30 mL of ethanol) [104]. The purified SWCNHs showed

Figure 2.30 Two modes for aggregation of single-walled carbon nanohorns [106]: (A) transmission electron microscopy image of dahlia-type aggregates, and (B) that of bud-type ones with enlarged images as the insets.

66 Chapter 2

80 60

(A)

(B) Ar, dahlia-type Ar, bud-type

N2, bud-type

40 He, bud-type

20 0

800 200 600 400 Buffer gas pressure / Torr

Average diameter of SWNH aggregates / nm

Yield of SWNHs / %

100

100

Ar, dahlia-type

80 Ar, bud-type

60 40

N2, bud-type

He, bud-type

20 0

800 200 600 400 Buffer gas pressure / Torr

Figure 2.31 Yield (A) and average size (B) of single-walled carbon nanohorn (SWCNH) aggregates formed under different pressures of different buffer gases with aggregates: either dahlia or bud type [106].

a broad 002 diffraction peak, which was explained using a model associated with two graphene sheets apart from 0.40 nm belonging to neighboring horns in their bundles. On the other hand, SWCNH aggregates produced in He and N2 had the shape of a flower bud, as shown in Fig. 2.31B. The primary SWCNHs in the bud-type aggregates were shorter than those in the dahlia type [105,106]. The type, yield, and size of SWCNH aggregates depended strongly on the pressure of the buffer gas, as shown in Fig. 2.31. Dahlia-type aggregates were obtained in Ar at 0.1 MPa with a high yield of about 95%, but in Ar at pressures less than 0.08 MPa and in He and N2 at 0.1 MPa, bud-type aggregates formed at a lower yield, at 70e80% [106]. In Kr at 0.1 MPa, the SWCNH aggregates were reported to be dahlia petal type, but in Xe, only thin graphene flakes were obtained [107]. SWCNHs with a purity higher than 90% were synthesized by a simple pulsed DC arc discharge between carbon rods at the atmospheric pressure of air, He, and Ar [108]. DC arc discharge with a repetition of 1 Hz (duration of 0.5 s) was generated in a gap of about 1 mm between two carbon electrodes with a diameter of 5 mm. The arc current and voltage between electrodes was set to 120 A and about 15 V, respectively, and the arcing period was 30 s. The sooty products were deposited onto the surface of the chamber, which consisted of mainly SWCNH aggregates with a small amount of amorphous carbon. The particle size of SWCNH aggregates was homogeneous and smaller with a narrower size distribution than SWCNHs prepared by CO2 laser ablation [103], as shown in Fig. 2.32. Heating of the resultant sooty product in dry air at 500 C was effective in eliminating amorphous carbon, which was detected as a weak additional band at around 1510 cm1 between D- and G-bands owing to SWCNHs in Raman spectrum.

Preparation of graphene 67 (A)

(B)

30

Pulsed arc discharge FWHM=24 nm

Counts

CO2 laser ablation FWHM=64 nm

20

10

0

0

50

150 100 Particle size / nm

200

Figure 2.32 Single-walled carbon nanohorn synthesized by pulsed arc discharge in air [108]. (A) Transmission electron microscopy images, and (B) particle size distribution compared with that by CO2 laser ablation. FWHM, full width at half maximum.

SWCNHs were also obtained by arc discharging in He and Ar, although the content of amorphous carbon was slightly larger than in air. Preheating the carbon rod at 1000 C just before igniting the arc was also effective in reducing the content of amorphous carbon in the product. DC arc discharge in Ar at 0.2 MPa resulted in dahlia-type SWCNH aggregates about 100 nm in diameter with a yield of more than 95%, but at 0.3 MPa in bud-type SWCNHs aggregates [109]. N-doped SWCNHs were fabricated by a flowing nitrogeneassisted arc discharge method at atmospheric pressure [110]. SWCNHs mixed with single-walled CNTs were produced by DC arc discharge by using the catalyst powders Ni and yttria (Y2O3) mixed with graphite, which were packed into a hole at center of the graphite anode [111]. The formation process of SWCNH aggregates was discussed on the basis of experiments done in air, CO2, and CO atmospheres [112]. During DC arc discharge in air at about 0.05 MPa, only a small number of sooty products were formed at first. After several minutes, there was a large number associated with fast consumption of anode graphite and the pressure in the chamber increased by about 20%. In CO2 at the same pressure, formation of the products was almost the same as in air, except the chamber pressure became almost two times higher. In CO, on the other hand, generation of soot occurred immediately after arc ignition at the same pressure. The formation processes in air, CO2, and CO were explained by the reactions of atmospheric gases, O2 in air, CO2, and CO, with carbon atoms in the graphite electrode. SWCNHs in the sooty products formed spherical aggregates with a diameter of 80e120 nm and were dahlia type in air but bud type in CO2 and CO. The sooty products were mixtures of amorphous carbon, SWCNHs aggregate, and graphitelike particles. The first component could be burned out below 400 C,

68 Chapter 2

0.4

0.4

0.2 0.2

0.0

0.0

-0.2 200 400 600 800 Temperature / oC

1000

0.6 0.4 0.2

0.8 0.6 0.4 0.2

0.0 -0.2

0

200 400 600 800 Temperature / oC

1000

Weight derivative / %/oC

0.6

SWNHs

0.6

0.8

Graphitic particles

0.8

1.0

1.0 Weight / %

0.8

Weight derivative / %/oC

Weight / %

1.0

0

(B)

1.0

Burn-off amorphous carbon

(A)

0.0

Figure 2.33 Thermogravimetryederivative thermogravimetry curves of sooty products [112]: products in (A) air and (B) CO. SWCNH, single-walled carbon nanohorn.

the second at around 600 C, and the third at 700 C, as shown in thermogravimetry (TG)ederivative thermogravimetry curves in Fig. 2.33. By DC arc discharge in liquid nitrogen, SWCNHs were obtained from graphite electrodes as a powdery dispersion in liquid nitrogen [113,114]. Arc discharge in distilled water between a graphite cathode with a hole and a graphite anode containing 0.8 mol% Gd gave SWCNH aggregates including Gd nanoparticles with a size of 2e30 nm [115]. The hole (diameter and depth of 7 and 20 mm, respectively) of the cathode, where the arc discharge was generated, was filled with N2 gas by introduction at a rate of 2 L/min through two narrow holes (diameter of 2 mm). SWCNHs were recovered from the powdery products floating on the water. SWCNHs were also synthesized by alternative current (AC) arc discharge between two graphite electrodes (purity of 99.999%) in air [116]. To achieve stable arc discharging, two graphite electrodes were axially rotated and one electrode was advanced to the other electrode at a constant speed. During discharging, two electrodes were equally heated and the sooty products were collected by the stainless-steel collector surrounded the two electrodes, in contrast to DC arc discharging, in which carbon sublimed mostly at the anode and the sooty products were mostly deposited on the cathode. The product of AC arc discharge at a frequency of 50 Hz with an applied voltage of 24e30 V was a mixture mainly of bud-type SWCNHs and few-layered graphene flakes, as shown in Fig. 2.34A and B. The rate of consumption of the electrodes increased with an increase in the voltage applied, together with the yield of SWCNH aggregates, as shown in Fig. 2.34C. Using a welding arc torch with counterrelectrode of graphite containing Ni/Y catalysts in open air, the formation of SWCNHs together with CNTs was confirmed in the powdery products [117,118].

Preparation of graphene 69

Figure 2.34 Single-walled carbon nanohorn (SWCNH) aggregates synthesized by alternative current arc discharge [116]. (A) Transmission electron microscopy image of the sooty product. (B) Transmission electron microscopy image of bud-type SWCNH aggregates. (C) Changes in the rate of electrode consumption and that of soot production with voltage applied.

2.1.4 Substitutional doping of heteroatoms Substitutional doping of heteroatoms to graphene has attracted attention as a possible technique for improving the electronic and chemical properties of graphene, although the possible atoms are limited. Three types of N atoms doped in graphene layer were identified: substitutionary doped N atoms bonded to 3 C atoms in a sp2 configuration (graphitic), and where N is bonded with two coordinated C atoms as a member of a hexagon (pyridinic) and a member of a pentagon (pyrrolic), as shown schematically in Fig. 2.35. The configuration of N atoms doped substitutionary into single-layer graphene

Figure 2.35 Three types of N atoms doped in a graphene layer: pyridinic, pyrrolic, and graphitic.

70 Chapter 2 (A)

(B)

390

Pyrrolic

394

Pyridinic

Intensity / a.u.

Intensity / a.u.

Graphitic

Pyridinic

398 402 406 Binding energy / eV

410

390

Pyrrolic

Graphitic

395 400 Binding energy / eV

405

Figure 2.36 X-ray photoelectron spectroscopy N (1s) spectrum for N-doped graphene. (A) The graphene film synthesized by chemical vapor deposition of CH4 with NH3 [120]. (B) The graphene film synthesized by segregation processes [121]. a.u., arbitrary unit.

was discussed on the basis of measurements by nitrogen core-level x-ray absorption spectroscopy and x-ray emission spectroscopy combined with theoretical calculations based on DFT to obtain a specific picture of doped N atoms [119]. N-doped graphene films were synthesized by CVD of CH4 with NH3 at 800 C on a Cu substrate, most of which consisted of a few layers and occasionally a single layer. They exhibited n-type electrical behavior [120]. N atoms in the resultant graphene films mainly had a graphitic configuration, as shown by the XPS N (1s) spectrum in Fig. 2.36A. The N atom with a graphitic configuration in a single-layer graphene film synthesized by CVD was visualized as an STM image, as shown in a simulation by DFT calculations in Fig. 2.37 [122]. N-doped few-layered graphene films were synthesized by segregating carbon from Ni metal and nitrogen from B on an SiO2/Si substrate during annealing at 800e1100 C under vacuum (103e104 Pa), as illustrated in Fig. 2.38 [121]. N atoms in the resultant graphene films had either a pyridinic or pyrrolic configuration, as shown in Fig. 2.36B, the concentration of which reached 2.9 atm%. N-doped multilayered graphene flakes synthesized by DC arc discharge between graphite electrodes in an NH3 atmosphere contained mainly pyridinic N atoms, the N content of which was about 11 atm% [101]. High-power electrical annealing in NH3 could form CeN bonds mostly at the edges of graphene ribbons [123]. N-doping of the as-grown single-layer graphene occurred by exposure to a beam of low-energy nitrogen ions (25e150 eV), followed by annealing at 900 K. The resultant graphene contained pyridinic N atoms accompanied by vacancies and graphitic N atoms [124]. Single-layer graphene films synthesized on Cu foils by CVD at 900 C in a flow of C2H4/H2 with NH3 diluted in He contained only pyridinic N atoms, as shown in Fig. 2.39 [125].

Preparation of graphene 71

Figure 2.37 Scanning tunneling microscope image of N atom doped in single-layer graphene (A) and its simulation by density functional theory calculations (B) [122].

Figure 2.38 Schematic illustration of the segregation technique for growing nitrogen-doped graphene [121].

(A)

(B) Pyridinic 399.3 eV

Intensity / a.u.

Intensity / a.u.

G’-band

G-band D-band

1200

1600 2000 2400 Raman shift / cm-1

2800

394

400 398 396 Binding energy / eV

402

Figure 2.39 N-doped single-layer graphene [125]. (A) Raman spectrum on SiO2/Si substrate. (B) X-ray photoelectron spectroscopy N (1s) spectrum for N-doped graphene synthesized with NH3/He. a.u., arbitrary unit.

72 Chapter 2

Figure 2.40 N-atom configurations in graphene with simulated images [126]: four pyridinic N atoms associated with two vacancies (A) and with three vacancies (B).

Doped N atoms were shown to distribute heterogeneously in the graphene layer by Raman spectrum measurements. By adjusting the flow rate of NH3, the atomic ratio of N/C was controlled from 0 to 16. N-doping of as-grown single-layer graphene on a Cu substrate at 900 C was performed by introducing NH3 gas at 850 C [126]. In the resultant N-doped graphene films, different configurations of doped N atoms with the formation of associated vacancies were observed under STM analyses with ab initio calculations, as shown in Fig. 2.40. Nitrogen-doped SWCNHs were synthesized by an arc discharge method in an N2 flow that contained mainly pyridinic and pyrrolic N atoms at the defect sites and the edges of graphene layers [110]. N-doped single-layer graphene films were synthesized from the films of the mixture of melamine (C3N6H6) with PMMA, which were spin-coated on Cu foil deposited onto SiO2/Si substrate, by heat treatment at 1000 C in a flow of H2/Ar, N doped located mainly in a pyridinic configuration [127]. B-doped graphene films were grown on polycrystalline Cu foils using a mixture of CH4, H2, and B2H6 gases at 1000 C in a quartz tube furnace [128]. The quality of the B-doped graphene films was far more sensitive to the B2H6 gas concentration and flow rate compared with N-doped graphene films using CH4, H2, and NH3 mixed gases. Prior exposure of the quartz tube to B was sufficient to dope the graphene grown subsequently. Like nitrogen, B incorporates into the carbon lattice primarily in a graphitic configuration and contributes about 0.5 carriers into the graphene film. DFT calculations suggested that B interacts strongly with the underlying Cu substrate whereas N does not, as illustrated in Fig. 2.41. The distribution of B atoms was observed to be completely random, whereas N atoms displayed clustering. Structurally, N-doped graphene films were relatively defect-free whereas B-doped graphene contained a large number of defects. B-doped graphene was synthesized by CVD at 1000 C in a flow of Ar gas after bubbling in a triethylboraneehexane solution [129]. On the resultant film, many defects were accompanied by B dopants with a croissantlike configuration (Fig. 2.42A).

Preparation of graphene 73 (A)

(B)

B on Cu B-Cu: 0.239 nm C-Cu: 0.278 nm

N on Cu N-Cu: 0.335 nm C-Cu: 0.332 nm

Figure 2.41 Interaction of doped B and N with a Cu substrate [128]: (A) High-resolution scanning tunneling microscope images for single graphitic B and N and (B) structure and charge density differences in B- and N-doped graphene on Cu(111) substrate (B and N atom on top of a Cu atom). Red denotes excess charge and blue, a deficit of charge.

Figure 2.42 B-doped graphene film [129]: (A) scanning tunneling microscope (STM) image and (B) simulated STM image and corresponding ball-and-stick structural model of B3 dopant.

74 Chapter 2 This is shown by the presence of a B3 cluster in the ball-and-stick structural model in Fig. 2.42B. The resultant B-doped graphene films had unique sensing capabilities for toxic gases, such as NO2 and NH3. B-doped rGO flakes were prepared by mixing a GO suspension with a borane-tetrahydrofuran adduct under reflux in an oil bath, in which borane works as a source of boron and a reducing agent for GO [130]. Graphene flakes containing different amounts of N and B were prepared [131]. N-doped flakes were synthesized by arc discharge of graphite electrodes in an H2 and He atmosphere with either pyridine vapor or NH3, resulting in an N-content of 0.6 and 11 atm %, respectively. N-doped flakes with an N-content of 1.4 atm% were obtained by heating a nanodiamond in an atmosphere of He and pyridine vapor at 1650 C. B-doped graphene flakes were synthesized by arc discharge of graphite electrodes in an H2 and He atmosphere in the presence of diborane (B2H6) vapor or by using a boron-packed electrode; this resulted in a B-content of 1.2 and 3.1 atm%. Electrical conductivity was improved by doping, N-doping (electron doping), which was much more effective than B-doping (hole doping), as shown in Fig. 2.43. Graphene films with different contents of N and B were synthesized by CVD of the mixed vapors of polystyrene and dopant (either urea or boric acid) introduced by an H2/Ar gas flow at 1035 C on a Cu foil substrate [132]. By adjusting the concentration of dopant precursors, the N-content could be modulated from 0.9 to 4.8 atm% and the B-content from 0.7 to 4.3 atm%, as estimated by XPS measurement. After transfer of the graphene film by using PMMA onto an SiO2/Si substrate, their Raman spectra and field effects were measured. As shown by Raman spectra in Fig. 2.44A and B, the D-band and D0 -band appeared and increased in intensity, and the G0 -band decreased in intensity accompanied by broadening with an increase in the content of dopants, N and B. The position of the G-band for graphene changed to the negative side with an increase in the N-content, but to a positive side with B-doping, as shown in Fig. 2.44C.

log(σ / S/m)

2.4

0.6 at% N-doped

2.1 1.2 at% B-doped 1.8 Undoped 1.5

0

100 200 Temperature / K

300

Figure 2.43 Temperature dependence of electrical conductivity of N- and B-doped graphenes [131].

Preparation of graphene 75 (A)

(B)

(C)

4.8 at%N

Weight of boric acid mixed / mg

4.3 at%B

3.7 at%N B-doping

D

G D’

G’ 0.9 at%N

Pristine (0 at%)

1400 1600 2400 2600 2800 Raman shift / cm-1

4.3 3.3

1.9 at%B 1.3 at%B D

G D’

G’ 0.7 at%B

Shift of the G-band / cm-1

1.4 at%N

Intensity / a.u.

Intensity / a.u.

3.3 at%B 2.1 at%N

1.9 1.3

B-content (at%)

0.7 0.9 1.4

N-content (at%) 2.1

N-doping

3.7 4.8

Pristine (0 at%)

1400 1600 2400 2600 2800 Raman shift / cm-1

5

15 20 25 10 Weight of urea mixed / mg

Figure 2.44 N- and B-doped graphenes [132]: changes in Raman spectrum with increasing content of N (A) and B (B), and shift of G-band (C) with increasing content of N and B. a.u., arbitrary unit.

Si is a common impurity in graphene layers grown by CVD caused by the presence of Si sources (SiO2/Si wafer, quartz tube, etc.) and the isovalence of Si and C. These Si impurities in graphene were directly observed by a combination of Z-contrast imaging and atomically resolved EELS on a scanning TEM [133,134]. Si atoms bond with either three or four carbon atoms in the graphene lattice. In other words, Si substitutes for a single C atom or for a CeC pair, as directly observed from the ADF image in Fig. 2.45A, and energy-loss near-edge fine structures (ELNES) for an Si L-edge clearly exhibit the difference in these two configurations, threefold and fourfold configurations, as shown in Fig. 2.45B. Si atom with a threefold configuration was concluded to displace outwardly from the graphene layer by adopting sp3 hybridization, whereas that with a fourfold configuration stayed practically in the same plane as the graphene layer by sp2d like hybridization owing to a significant contribution of Si 3d orbitals, as shown schematically in Fig. 2.45C. Substitutional doping of Si into a graphene lattice was performed by a bubbler-assisted ambient-pressure CVD on Cu foil by introducing the vapor of the mixture of methoxytrimethylsilane (C4H12OSi) and hexane as Si and C precursors, respectively, as shown in Fig. 2.46A [135]. As shown by Raman spectra in Fig. 2.46, Si-doped graphene exhibited an increased D-band and a decreased G0 -band, compared with pristine graphene and N-doped graphene, which might be attributed to the graphene lattice distortion caused by Si doping.

76 Chapter 2 (A)

(B)

(C)

Si-C3

0.2 nm Si-C4

Intensity / a.u.

Si-C3(sp3)

100

0.2 nm

Si-C3

Si-C4(sp2d)

105 110 115 Energy loss / eV

120

Si-C4

Figure 2.45 Threefold and fourfold configurations of Si impurities in single-layered graphene [133]: (A) Annular dark-field images with structure models, (B) energy-loss near-edge fine structure for an Si L-edge, and (C) structural models. a.u., arbitrary unit.

Intensity / a.u.

G’

1000

1500

2500 2000 Raman shift / cm-1

3000

Figure 2.46 Raman spectra of Si-doped graphene film in comparison with N-doped and pristine graphene films [135]. a.u., arbitrary unit.

Preparation of graphene 77

Figure 2.47 S-Doped graphene [136]. (A) High-resolution transmission electron microscopy image. (B) Elemental mapping of S for the image in (A). (C and D) Enlarged images in green and blue frames of the image in (A).

Doping of S into graphene was performed by CVD of the vapor from S-dissolved hexane at 950 C on Cu foil [136]. Randomly distributed particlelike regions 2e5 nm in size were observed on high-resolution TEM images, as shown by the red arrows in Fig. 2.47A, which are aggregated S atoms, as shown by elemental mapping in Fig. 2.47B. Most of the S atoms formed two kinds of nanodomains: relatively regular regions and strongly disturbed regions (Fig. 2.47C and D, respectively). S-doping gave characteristic changes in the Raman spectrum, marked increases in intensity of D- and D0 -bands, but a strong G0 -band, stronger than the G-band, as shown in Fig. 2.48A. XPS spectra for C (1s) and S (2p) (Fig. 2.48B and C, respectively) indicate the formation of CeS bonds, a S (2p) peak of 163.7 eV, and a C (1s) peak at 285.3 eV. The S (2p) peak at 164.5 eV was attributed to neutral S. Metallic impurities, Si and Fe, in graphene were analyzed using ADF-TEM imaging. Si was likely to have come from the glassware used to process GO and Fe from the etchant employed to remove metal substrates [137]. Si atoms are mostly segregated into a single

78 Chapter 2 (A)

(B)

(C)

D’

Pristine graphene

1000 1500 2000 2500 3000 Raman shift / cm-1

285.3 eV 287.8 eV

290 288 286 284 282 280 Binding energy / eV

Intensity / a.u.

D

Intensity / a.u.

Intensity / a.u.

S-doped graphene

G

163.7 eV 164.5 eV

284.5 eV

G’

162.5 eV

166 165 164 163 162 161 160 Binding energy / eV

Figure 2.48 S-Doped graphene [136]: (A) Raman spectra,(B) x-ray photoelectron spectroscopy C (1s) spectrum, and (C) S (2p) spectrum. a.u., arbitrary unit.

vacancy site, i.e., substitution of a single Si atom for a single C atom (no vacancy formation) and occasionally in a divacancy site, i.e., substitution of a single Si atom with a vacancy for 2 C atoms, whereas an Fe atom is substituted into a divacancy site but not in single vacancy site, as shown in Fig. 2.49A and C. These impurities showed no evidence

Figure 2.49 Location of impurity atoms in single-layered graphene [137]. (AeC) Annular dark-field transmission electron microscopy image after noise correction. (DeF) Calculated structure. (A) Perfect graphene area. (B) Si atom in a C-vacancy site. (C) An Fe atom in a C-divacancy site. Single impurity atom (Si and Fe) in (D) a single vacancy site, (E) a divacancy site, and (F) four vacancy sites.

Preparation of graphene 79 by EELS analyses of being oxidized and were stabilized by location at vacant sites. Some theoretically calculated sites are shown in Fig. 2.49DeF. Substitutional doping of hetero atoms (P, S, Si, and Al) was discussed by theoretical calculation [138,139]. Energy for doping was estimated to be the lowest for Si, at 0.2 eV, but the largest for Al, at close to 10 eV. Structural changes caused by substitution are significant for these atoms because they are too large to fill in the sp2 framework of graphene and produce an out-of-plane distortion, as illustrated in Fig. 2.50 for single-layered and double-layered graphenes, respectively, mainly because of larger bond lengths between the heteroatom and carbon than the CeC bond in the layer. Si-doped graphene is semimetallic and P-doped graphene has a magnetic moment and a band gap of 0.67 eV. Al-doped graphene is metallic but very unstable. Si-doping was theoretically predicted to be efficient in improving the functionalities of graphene for hydrogen storage capacity [140], gas sensing [141,142], and the catalytic activity of NO reduction [143]. A review was published focusing on N doping into graphene [144]. In some papers, surface modification using heteroatoms (surface transfer doping) of graphene was discussed that included substitutional doping. In this broad meaning, doping (called chemical doping) was reviewed as a technique for band gap tuning of graphene [145].

(A) Graphene: 4x4 single-layer Al-C: 0.1860 nm P-C: 0.1783 S-C: 0.1780 Si-C: 0.1767

(B) Graphene: 4x4 double-layered P-P: 0.2230 nm Si-Si: 0.2360 Al-Al: 0.2528

Figure 2.50 Distortion of graphene layers by substitutional doping of hetero atoms [138].

80 Chapter 2

2.1.5 Graphene foams Graphene foams were synthesized by CVD of CH4 at 1000 C under ambient pressure using Ni foam as a template [146]. Before etching the Ni template by a hot HCl solution, a thin layer of PMMA was deposited on the surface of the graphene films deposited on an Ni surface to prevent collapse during Ni etching. After the PMMA was carefully removed by hot acetone, a monolith of graphene 3D network (graphene foam) was obtained, as shown in Fig. 2.51. The graphene foams synthesized with a CH4 concentration of 0.7 vol % were confirmed from high-resolution TEM and Raman spectrum to consist of few-layered graphene flakes and to have a bulk density of about 5 mg/cm3, corresponding to a porosity of about 99.7%, and a relatively high BrunauereEmmetteTeller (BET) surface area (SBET) of about 850 m2/g. Its composite with PDMS delivered a high electrical conductivity of about 10 S/sm3, even with low loading of the foam at 0.5 wt%. Graphene foams were also prepared by CVD of C2H5OH using Ni foam as the template, of which the Raman spectrum suggested a high quality of graphene layers because of the negligibly weak D-band [147]. NiO was electrochemically deposited onto this graphene foam. The resultant NiOegraphene composite foam exhibiting high capacitance (about 816 F/g) and stable cycling performance as electrodes of electric double-layer capacitors. Highly elastic graphene foams were fabricated by CVD of CH4 at 900 C using g-Al2O3 (primary particles about 7 nm in size) as the template [148]. After dissolving the template Al2O3 using HF, followed by heat treatment at 1800 C, the resultant graphene foam had a high SBET at 1940 m2/g with a Vtotal of 2.79 cm3/g and a Vmeso of 2.14e2.74 cm3/g, which revealed that it contained principally mesopores of 5e7 nm in size, the walls of which consisted of one to two layers of graphene. It had ultrahigh elasticity and mechanical toughness. By compression under 500 MPa, the bulk density increased to 1.2 g/cm3 and

Figure 2.51 Graphene foam synthesized by chemical vapor deposition of CH4 using an Ni foam template [146]. (A) Photograph of a 170  220-mm2 free-standing foam. (B) Scanning electron microscopy image. (C) Transmission electron microscopy image.

Preparation of graphene 81 the pore volume to 0.33 cm3/g from initial values of 0.13 g/cm3 and 2.79 cm3/g, respectively. By releasing the pressure, both the bulk density and pore volume completely recovered. Graphene foams were prepared using porous ZnO as the template and toluene vapor as the precursor at 760 C [149]. It had very low bulk density, as low as 180 mg/cm3, with a conductivity of 0.2 S/m and a Young’s modulus of 15 kPa, and was called aerographite.

2.2 Cleavage (peeling) 2.2.1 Mechanical cleavage Thin flakes of graphite were obtained by applying micromechanical cleavage (peeling) to micropillars formed using oxygen plasma [150], as shown in Fig. 2.52A. By repeating this micromechanical cleavage technique on freshly cleaved HOPG, thin graphene flakes 3e100 nm thick and with a lateral size of about 2 mm were obtained [151,152]. Arrays of graphite micropillars (about 2  2  5 mm3) were fabricated on the HOPG surface using micropatterning followed by oxygen plasma etching (inset in Fig. 2.52A), and then a pillar was detached from the surface and transferred into AFM using a silicon cantilever to leave thin flakes on the SiO2/Si substrate (Fig. 2.52B), in which the cleavage process was controlled by tuning the normal force between the cantilever and the substrate [151]. A large flake shown in Fig. 2.53A was composed of regions with different thicknesses from 0.8 to 3 nm (Fig. 2.53B) and a relatively large area 0.8 nm thick that was thought to be a single-layered region. The electrical properties of these graphene flakes composed from a few layers were measured as functions of thickness of the flake and gate voltage Vg. The results are shown in Fig. 2.54A and B by plotting resistivity r, Hall coefficient RH [152], and normalized conductivity s/smin [151] as a function of gate voltage Vg,

Figure 2.52 Micromechanical cleavage process [151]. (A) Pillars formed on highly oriented pyrolytic graphite (HOPG) surface by oxygen plasma etching (inset: arrayed pillars). (B) Cleavage using atomic force microscopy (AFM) cantilever.

82 Chapter 2

Figure 2.53 Graphene thin flake prepared by micromechanical cleavage [152]. (A) Photograph of relatively large flake about 3 nm thick. (B) Atomic force microscopy image of single-layered region of the flake on an SiO2 substrate.

(A)

(B) 8

at 1.7 K

at 5 K

1.6

4

at 70 K at 300 K

2

RH / kΩ/Τ

0 0.5 0

-0.5 -100

Normalized conductance σ/σmin

ρ / kΩ

6

δE

at 5 K

EF

EF EF

50 -50 0 Gate voltage Vg / V

1.4

12 nm

1.2 42 nm 95 nm

1.0 100

-90

-60

-30 30 0 Gate voltage Vg / V

60

90

Figure 2.54 Field effect of mechanically cleaved graphene flakes. (A) Dependence of electrical resistivity r and Hall coefficient RH on gate voltage Vg with schemes of band structure at different Vg’s on few-layered graphene flakes [152]. (B) Dependence of normalized conductivity s/smin on Vg on flakes with different thicknesses [151].

which clearly reveal the field effect. In Fig. 2.54A, r shows a sharp peak at a value of several kU and decays to about 100 U at high Vg; RH exhibits a sharp reversal of its sign, which suggests a change in carriers from holes to electrons with an increase from negative to positive Vg, although both r and RH shift to the positive Vg side owing to adsorbed water. In Fig. 2.54B, electric fieldedependent conductance measurement of these flakes shows marked modulation as a function of Vg; more markedly on the thinner flakes and no effect on flakes with more than 95 nm thick.

Preparation of graphene 83 (B)

Resistivity ρ / Ω/cm

29 nm

43 nm

10-4 52 nm 59 nm 79 nm

10-5

95 nm

150

at 4.2 K

0

18 nm

-2 45 nm

200 100 Temperature / K

300

100

0

2 3 4 1 Magnetic field B / T

35 nm

23 nm

45 nm

50 18 nm

35 nm 23 nm

111 nm

10- 6 0

(C)

at 4.2 K

2

Magnetoresistance Δρ/ρ

(A) Hall coefficient RH / x106 cm3/C

10- 3

0 5

0

3 4 2 1 Magnetic field B / T

Figure 2.55 Electromagnetic properties of thin graphite flakes with various thicknesses [153,154]. (A) Change in r with temperature. (B) Change in RH with magnetic field B. (C) Magnetoresistance Dr/r with B.

Before these works, cleavage of highly crystalline kish graphite using double-sided adhesive tapes was repeated until the flake became transparent at a thickness of 18e108 nm [153,154]. The electrical resistivity r, Hall coefficient RH, and transverse magnetoresistance Dr/r of these thin flakes were measured at temperatures of 4.2e300 K, as shown in Fig. 2.55. The high crystallinity of the thin flakes thus prepared was confirmed by the presence of a Shubnikov-de Haas oscillation in RH (Fig. 2.55B). These results showed the marked dependence of electronic properties on the thickness of the flakes, i.e., the number of stacked layers. Analysis of these experimental results showed that the overlap energy between conduction and valence bands decreased and the relaxation rate caused by lattice defects increased with a decrease in the number of the stacked layers. Flakes obtained through repeated micromechanical cleavage of graphite were suspended in a liquid and then attached to a micrometer-sized metallic scaffold to identify singlelayered graphene under TEM [155]. Detailed observations by TEM techniques showed that the flakes exhibited random microscopic out-of-plane deformations. Single-layered graphene flake showed more marked deformations than did double-layered flakes, in addition to folding [156]. The edges of single-layered graphene prepared from natural graphite by micromechanical cleavage were shown to have either a zigzag or armchair structure through high-resolution STM [157]. Atomically flat single-layered graphene flakes were obtained from kish graphite by rubbing on the cleaved surface of mica (muscovite) and compared with those on the SiO2 substrate [158]. The results of AFM

5

84 Chapter 2 (A)

(B)

Frequency / a.u.

1.0

Y / nm

Z / nm

SiO2 substrate

Mica substrate

Graphene/SiO2

Graphene/mica

0.5

(C)

X / nm

Y / nm

Z / nm

0.0 -400

-200

0

Height / pm

200

400 X / nm

Figure 2.56 Comparison of surface roughness for single-layered graphene on mica and SiO2 [158]: (A) atomic force microscopy (AFM) topographic data, (B) AFM topographic image of graphene on mica, and (C) that of graphene on SiO2. a.u., arbitrary unit.

topographic observations in Fig. 2.56 show that the fluctuation in height on the graphene surface was calculated to be 24 pm for the graphene on mica, in contrast to 154 pm for the graphene on SiO2. This difference in height fluctuation was the result of the surface roughness of the substrate: 34 pm for mica and 168 pm for SiO2. Graphite bonded onto borosilicate glass was cleaved off, leaving several bonded areas of graphite on the glass surface, which were then peeled off using adhesive tape, leaving many transparent thin flakes with a large area consisting of a single or few layers on the substrate [159]. On the glass substrate, graphene flakes could be identified under a polarized light optical microscope. The contrast increased with the number of graphene layers. Single-layered graphene prepared by micromechanical cleavage on a glass microscope slide was successfully used as a transparent conductive electrode in liquid crystal photonic devices [160]. By rubbing a fresh surface of layered crystals, graphite, BN, MoS2, NbSe2, and Bi2Sr2CaCu2Ox against a solid surface, it was possible to be obtained thin flakes with different thicknesses, including single-layered flakes, [161]. The optical contrast observed on these thin flakes was discussed using a model based on the Fresnel law [162]. Single-layered graphene can be visible on an Si wafer with a 300 nm thickness of SiO2 (Fig. 2.57A), but on 200 nm SiO2 it is completely invisible (Fig. 2.57B) under white light illumination. Only flakes thicker than 10 layers can be detected on 200 nm SiO2 under white light. The visibility of graphene flakes depends on the wavelength of illuminated light. Under l of 560 nm, steplike changes in contrast are clearly visible for single-, double-, and triple-layered regions on 300 nm SiO2 (Fig. 2.57C).

Preparation of graphene 85 (A) 300 nm SiO2

(B) 200 nm SiO 2

(C)

Figure 2.57 Graphene sheets on SiO2eSi substrate [162]: (A) on SiO2 with a thickness of 300 nm, (B) on 200-nm-thick SiO2 under white light, and (C) on 300-nm-thick SiO2 under green light, where single-layered and triple-layered graphene flakes are recognized.

Graphene flakes cleaved onto hydrophilic SiO2/Si substrate were rolled up with the assistance of bubbles formed in an aqueous solution of NaHCO3 at 50 C by inserting the substrate vertically [163]. The graphene nanoscrolls obtained showed p-type electrical conductance, which was distinct from the electrical properties of flat graphene flakes.

2.2.2 Cleavage in solution Cleavage processes of graphite in solution for the fabrication of graphene (in some literature called liquid-phase exfoliation) were reviewed by categorization into two major classes: i.e., surfactant-free and surfactant-assisted processes [164]. Various solvents, such as N-methylpyrrolidone (NMP), N,N0 -dimethylacetamide, g-butyrolactone, and 1,3-dimethyl-2- imidazolidinone, could be used to disperse thin flakes. Excessive sonication may lead to destruction of the graphene. Dispersion of sieved graphite powder (lateral size of 50

2800

1400 1600 1800 2600 Raman shift / cm-1

2800

Figure 2.134 Raman spectra of graphene films formed on Cu at 1000 C [127]: (A) graphene films derived from poly(methyl methacrylate) (PMMA) with different H2 flow rates, and (B) those derived from different precursors. a.u., arbitrary unit.

146 Chapter 2 (B)

(A)

250

(C)

11

Intensity

200 150

22

100

33

50 Substrate 0 1.0

3.0 2.0 2θ / degree

4.0

Figure 2.135 Graphene flakes synthesized under solvothermal condition [371]: (A and B) Transmission electron microscopy images; (C) electron diffraction pattern of hk reflections.

Hexagonal single-crystalline graphene films were synthesized on Cu foils from humic acid by thermal annealing at 1100 C in a mixture of Ar/H2 for 1 h under atmospheric pressure [370]. Graphene flakes were synthesized from a mixture of sodium and ethanol (1:1 in mol) by heating at 230 C for 72 h under hydrothermal conditions in a Teflon-lined autoclave, followed by pyrolysis [371]. In Fig. 2.135, TEM images of the resultant graphene flakes are shown together with electron diffraction patterns for hk reflections. The resultant graphene flakes were proved by AFM analysis to be single-layered with a lateral size of 0.1e10 mm. N-doped graphene flakes were synthesized by the detonation reaction of a powder mixture of cyanuric chloride and trinitrophenol at 320 C in autoclave, in which a momentary pressure reached 60 MPa and an equilibrium pressure was 30 MPa [336]. TEM images and Raman spectra of the resultant flakes are shown in Fig. 2.136, revealing the formation of thin flakes of multilayered graphene (4e8 layers stacked). Three configurations of N atoms, graphitic, pyridinic, and pyrrolic, were identified in the XPS spectrum. The yield of the flakes was about 0.1 g per 1 g cyanuric chloride. (B)

(C) Intensity / a.u.

(A)

G

G’ D

1000 1500

2000 2500 Raman shift / cm-1

3000

Figure 2.136 Multilayered graphene flakes synthesized by detonation [336]: (A and B) Transmission electron microscopy images; (C) Raman spectrum. a.u., arbitrary unit.

Preparation of graphene 147

2.4.3 Unzipping of carbon nanotubes An alternative route to the preparing graphene nanoribbons with straight edges is the unzipping of CNTs using different techniques. Unzipping of MWCNTs was performed by oxidation in H2SO4eKMnO4 aqueous solutions [372]. By increasing the amount of KMnO4 from 100 to 200, 300, and finally 500 wt%, unzipping of MWCNT proceeded stepwise, as shown in Fig. 2.137. Pristine MWCNT with a wall consisting of 15e20 layers was finally converted into a thin graphene ribbon with a width of >100 nm and linear edges. Oxidation in mixed acids of H2SO4 and KNO3 led to longitudinal unzipping of MWCNTs [373]. Using plasma etching, graphene nanoribbons with smooth edges and a narrow width of 10e20 nm were successfully obtained [374,375]. The unzipping process is schematically shown in Fig. 2.138. CNTs were embedded in PMMA on an Si substrate; then, the resultant PMMAeCNT film was peeled from the substrate, in which a narrow strip of the CNT side wall was exposed to the atmosphere. This exposed strip was etched by 10 W Ar plasma. Single- and double-layered graphene ribbons were prepared from SWCNTs and double-walled CNTs, respectively. Starting from MWCNTs with a diameter of about 8 nm, nanoribbons consisting of few layers and with a width of 10e20 nm were obtained

Figure 2.137 Transmission electron microscopy images of stepwise opening of multiwalled carbon nanotubes using increasing amounts of KMnO4: (A) 100, (B) 200, (C) 300, and (D) 500 wt% [372].

148 Chapter 2

Figure 2.138 Unzipping of multiwalled carbon nanotube (MWCNT) by plasma etching [374]. CNT, carbon nanotube; GNR, graphene nanoribbon; PMMA, poly(methyl methacrylate).

in a yield of about 20%. AFM images of some ribbons are shown in Fig. 2.139A and B. Their width and thickness were determined by AFM. The obtained ribbons were uniform in width along their length and smooth on their edges. The intensity ratio ID/IG in Raman spectra increased from almost zero for the starting MWCNTs to 0.38e0.28 for the ribbons obtained, which was thought to be caused mainly by their open edges. The electric-field dependence of conductivity of these nanoribbons was measured; the ribbon was about

100 nm

(C)

-Ids / 10-6A

(B)

MWCNT

(A)

Vg / V

Figure 2.139 Graphene ribbons prepared by unzipping of multiwalled carbon nanotubes (MWCNTs) [374]: (A) atomic force microscopy images of the ribbons with a width of 8 nm and a height of 1.8 nm (left) and with 13 and 2.0 nm (right), (B) and with 15 and 0.9 nm and a drain-source current Ids versus a gate voltage Vg for a ribbon with a width of about 16 nm.

Preparation of graphene 149

Figure 2.140 Scanning electron microscopy images of unzipped multiwalled carbon nanotubes by N2 gas expansion [376]: (A) curved, and (B) flat graphene nanoribbons.

16 nm wide, as shown in Fig. 2.139C. Graphene ribbons prepared by unzipping MWCNTs gave a V-shaped Ids versus Vg curves with Dirac point at near Vg ¼ 0 V, similar to that of bulk graphene, which reflected symmetric hole and electron transports at negative and positive gate voltages, respectively. MWCNTs were unzipped by an abrupt N2 gas expansion within their hollow core by a thermal shock [376]. Graphene nanoribbons obtained by unzipped MWCNTs are shown in Fig. 2.140. Unzipping of SWCNT induced by hydrogenation was theoretically discussed with ab initio calculations [377]. Merging and cutting of carbon structures (in other words, zipping of graphene and unzipping of SWCNT), with the help of atomic hydrogen were discussed theoretically [378]. Electrochemical oxidation and reduction of MWCNTs in 0.5 M H2SO4 aqueous electrolyte was reported to be effective for unzipping to obtain graphene ribbons with fewer defects than chemical reduction [379]. Electrochemical treatment using a solidstate proton-conducting polymer electrolyte membrane (Nafion) and placing MWCNTs at both the oxidized and reduced sides was reported to be effective for their longitudinal unzipping [380]. Commercially available MWCNTs (40e70 nm in diameter) were transformed into Ndoped graphene nanoribbons in a single-step electrochemical process in a 0.1-M LiClO4 formamide solution at a potential of 1.5 V for different periods (2e24 h) [381]. Formamide worked as a solvent and a source of nitrogen. Nitrogen in the resultant nanoribbons resulted in pyridinic and pyrrolic configurations with a total content of about 44 atm%. In Fig. 2.141A and B, the TEM image of the MWCNT before unzipping is compared with that after complete unzipping. The resultant nanoribbons had a width of 80e150 nm and stacking of five to 20 layers. After unzipping, MWCNTs had a strong D-band, as shown in Fig. 2.141C. The unzipping, which occurred mostly at the

(A)

(B)

Intensity / a.u.

150 Chapter 2 (C)

D G

G’ D’

N-doped (24h) N-doped (12h) Pristine MWCNF

0

3000 1000 2000 Raman shift / cm-1

Figure 2.141 Unzipping and N-doping of multiwalled carbon nanotubes (MWCNTs) [381]. (A) Transmission electron microscopy image of the pristine MWCNT, (B) that of N-doped unzipped MWCNTs, and (C) their Raman spectra. a.u., arbitrary unit.

electrodeeelectrolyte interface, produced irregular morphologies, probably owing to the accompanying decomposition reaction, unlike the well-defined longitudinal unzipping of CNTs explained earlier.

2.5 Concluding remarks CVD processes have the highest possibility for synthesizing graphene films with fewer structural defects, transparency, and homogeneous thickness on different substrates, which can be successfully transferred onto a target substrate such as SiO2/Si, and PET without a noticeable increase in structural defects. The synthesis of thin films with a large area and a controlled thickness is possible using this method, although several exploratory efforts have been made to find optimal conditions, such as the kind of precursor, its concentration, a carrier gas, its flow rate, the deposition temperature, the substrate, and the cooling rate after the deposition. Moreover, a large initial investment is required for instruments for CVD deposition on an industrial scale. This process is suitable for the patterning of graphene films combined with lithography techniques, which are often required to produce electronic circuits. For use as a transparent conductive thin film, the formation of thin small-sized graphene flakes into thin films has been achieved: for example, using drop-casting or spin-coating of the suspension of thin flakes on a substrate. For this purpose, thin flakes synthesized by CVD using microwave-assisted plasma and arc discharge are possible after suspending in a solution, but heavy instrumentation can be a barrier to industrial application. Thin graphene flakes recovered as a supernatant of the suspension of flakes cleaved from thick graphite flakes in a solution are useful for preparing transparent graphene films. However, the yield of thin enough flakes as a supernatant is low, and suitable solvents and additives for suspending graphene nanoflakes are limited. The dispersion of GO flakes after exfoliation by sonication is much easier, but GO flakes have to be reduced to prepare graphene films and the resultant rGO is usually defective and some functional groups containing oxygen remain even after the reduction.

Preparation of graphene 151 In the case of using a transition metal substrate, such as Ni for a CVD process, the precipitation of carbon atoms dissolved into the substrate metal has to be considered. On thin films prepared using this process, however, interaction with the surface of the substrate crystal is important; it results in the epitaxial growth of graphene and the weakening of CeC bonds in layers close to the substrate (initially deposited layers). The interaction of deposited graphene layers with the substrate was reported to be markedly weakened with an increased number of layers stacked in the flake. For polycrystalline Cu and Pt substrates, however, single-grain (single-crystal) graphene was successfully grown, which suggested that there was no epitaxial relation between graphene and the substrate. For SiC substrate, only weak electronic coupling of deposited few-layered graphene with the substrate was suggested by microRaman scattering. The cleavage of graphite crystals is the best method in terms of the structural and electrical qualities of flakes obtained by selecting a graphite plate with high crystallinity, but it is challenging to bring this process to large-scale production. The thickness of the thin flakes is difficult to tune through this process and it is laborious to get single-layered graphene flakes. The size and shape of the resultant flakes are governed by the graphite used, as demonstrated by the pillaring of HOPG, and the crystallinity of the graphite can be retained in the resultant graphene nanoflakes. However, even the highest-grade commercially available HOPGs are not a single crystal, as shown in Fig. 2.142. They consist of grains with different orientations of a-axes, although with an extremely high c-axis orientation [382]. Some flakes of kish graphite have very high crystallinity, as shown by the high value of r300k/r4.2K and by Shubunikovede Haas oscillation in the Hall coefficient RH. Using micromechanical cleavage, the patterning of thin sheets can be performed on pristine graphite, but special care is needed concerning the crystal perfection of the prepared thin sheets, because the patterning of graphite is often carried out using oxygen plasma. Cleavage also has to be performed carefully, because some distortion may be induced in cleaved flakes, particularly when flexible adhesive tape is used.

(A)

(B)

100 μm

Figure 2.142 Electron channeling contrast (A) and a-axis orientations of each domain (B) of highly oriented pyrolytic graphite surface [382].

152 Chapter 2 The process using GICs (both donor-type compounds such as K-GICs and acceptor-type ones such as HNO3-GICs) is thought to give fewer structural defects to the resultant graphene nanoflakes, because interaction between the graphite layer and the intercalate is only charge transfer, and the sp2-bonding nature of carbon atoms in the layer is preserved. In other words, chemical oxidation and a change in bonding in the layers do not occur. The main product of this process is nanosized scrolls of graphene layers, which might present a new science and engineering of quasi-1D nanocarbons, although the reproducibility of the nanoscroll formation and the possibility of having nanoscrolls consisting of single-layered graphene need to be explored. The preparation of thin flakes using GO was reported in a number of papers. The process consists of many steps, including oxidation of the pristine graphite to GO, exfoliation to separated thin flakes, reduction of GO flakes, and annealing of the products (Fig. 2.78). Each step involves various chemical reactions, particularly oxidation and reduction, associated with the attachment of oxygen-containing functional groups to carbon atoms and the removal of these groups accompanied by vigorous evolution of CO and CO2, respectively. Some carbon atoms in the pristine graphite are lost as CO and CO2 during exfoliation and reduction. In addition, the bonding state of most carbon atoms has to be changed from sp2 in pristine graphite to sp3 in GO during oxidation, and vice versa during reduction. These drastic changes in composition and bonding nature are thought to give a number of structural defects to the carbon layers. The presence of carbon atoms bonded to oxygen atoms is detected on GO in XPS C (1s) and O (1s) spectra, as expected from its preparation processes. However, a marked tailing of the spectrum to the high binding energy side in C (1s) spectra and weakened-appearing O (1s) spectra are often observed even after reduction reactions (some oxygen-containing functional groups remain in rGO) (see Fig. 2.83, for example). A study on the reduction process of GO flakes under TEM demonstrated that the content of oxygen atoms detected by EELS decreased gradually with step-by-step Joule heating and became negligibly small only after increasing bias to about 20 V (thought to be about 2000 C) [383]. The electrical conductance of the flake increased to about 1.5  105 S/m, a six orders of magnitude rise, during this reduction process. To get flakes consisting of single-layered to few-layered graphene flakes, the supernatant of the dispersed solution or sol of GO particles has to be used. Therefore, the yield of graphene flakes from the pristine graphite seems to be low, less than a few percent, and reproducibility in the thickness distribution in the resultant flakes is poor. Most resultant thin flakes contain ripples and folds and are partly scrolled. However, this process through GO is thought to be applied as a large-scale industrial process to produce rGO nanoflakes (powder), and rGO nanoflakes are easily modified with various species (functionalization and heteroatom doping). These are certain advantages for many applications, such as biomedical applications, which have not been explored for conventional carbon materials.

Preparation of graphene 153 The chemical synthesis of graphene via organic processes may be possible for obtaining graphenelike sheets with a homogeneous size, but it might be difficult to synthesize large flakes on a large scale. The synthesis of nanoribbons via pyrolysis of organic precursors has the potential to expand the scale of synthesis, but most products consist of graphitic regular stacking. Unzipping of CNTs is a scientifically interesting process to control the number of stacked layers, but it is difficult to establish it as a practical fabrication process for graphene flakes. Microscopic corrugation, rippling, and partial scrolling at the edge were frequently observed on thin flakes prepared by each of these routes. The formation of ultraflat graphene was reported through micromechanical cleavage by selecting cleaved mica as an appropriate substrate. However, thin flakes prepared through this method were shown to scroll at their edges and fold when the flakes were detached from the scaffold [155]. Single-layered graphene is susceptible to structural distortions and the suspended graphene flakes showed spontaneous rippling of about 1 nm. Therefore, it is difficult to prepare single-layered graphene with no ripples. To avoid structural distortion and scrolling at the edges, the selection of an appropriate scaffold or substrate seems to be an important factor. This rippling was invoked to explain thermodynamic stability and distinctive electronic properties, because single-layered graphene does not have p-stacking that favors a flat structure. Substitutional doping of heteroatoms into a graphene lattice is an effective technique for improving and/or modifying the properties of graphene films and flakes. Substitutional doping of nitrogen and boron was investigated using various graphite materials, including CNTs and fibers. It seems to be more evident on graphene than on graphite to demonstrate the configuration of doped nitrogen and/or boron atoms in graphene lattice and the effects of doping on the properties. In addition, other heteroatoms such as Si, P, and S were reported to be able to be doped into a graphene lattice.

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CHAPTER 3

Electrical properties and applications Chapter Outline 3.1 Fundamental electrical properties 174 3.1.1 Electronic structure of graphene 174 3.1.2 Effects of defects and edges 181

3.2 Applications to information technology

189

3.2.1 Transistor devices 189 3.2.2 Spintronics devices 207 3.2.3 Transparent electrode 221

3.3 Applications to social fields 226 3.3.1 3.3.2 3.3.3 3.3.4

Sensor devices 226 Photon detectors 233 Resistance standard 237 Electron field emission 242

3.4 Concluding remarks References 244

244

For the next generations on our planet, establishing a sustainable and smart society is urgent and the goal of much current research and development. In view of sustainability, materials science and engineering plays an important role in solving problems such as environmental pollution, consumption of resources and energy, and human health. Novel features of graphene in chemical, mechanical, thermal, and biomedical properties greatly contribute to problems to achieve sustainability. On the other hand, the peculiar electronic properties of graphene are important mainly for development of the infrastructure for information technology (IT), which is core technology for a smart society for our next generations. Graphene is a promising key material to develop new devices mandatory in the information technology sector for faster processing and larger storage of information and improvement in the manemachine interfaces. In addition, the novel electronic properties of graphene would be useful for applications to other fields, especially sensor devices, which are a core technology in the emerging field of the Internet of things (IoT). Fig. 3.1 schematically shows the fundamental concept of this chapter. The electronic structure of graphene is greatly responsible for novel features of graphene in Thermal properties, Chemical properties, Biomedical properties and Mechanical properties. Especially, we need pay much attention to the electronic structure of graphene and its modulation by the presence Graphene. https://doi.org/10.1016/B978-0-12-819576-5.00003-7 Copyright © 2020 Elsevier Inc. All rights reserved.

173

174 Chapter 3 Graphene Single layer of mono-atomic thickness 3. Electrical Properties and Applications 2 Chemical properties

6 Thermal properties

3.1 Fundamental properties

High thermal conductivity 7 Biomedical properties

Electronic structure Effects of defects and edges

High biocompatibility

Extended p-electron clouds 5 Mechanical properties High strength, Flexible layer

3.2 Applications to information technology Transistor devices Spintronic devices Transparent electrodes

3.3 Applications to social fields Sensor devices, Photon devices, Resistance standard, Electron field 3mission

Figure 3.1 The concept of the present chapter.

of defects and edges in considering application of graphene to transistor and spintronics devices, transparent electrodes, sensor and photon devices, and Electron field emission.

3.1 Fundamental electrical properties 3.1.1 Electronic structure of graphene The electronic states of graphene as well as other three-dimensional crystalline solids are represented by innumerable periodic wave functions infinitely extended over the space having eigen energies within finite energy range (energy band). We usually recognize a phenomenon on the real space coordinate described by r ¼ (x, y, z) positions. However, for periodic wave functions, it is convenient to denote them as their Fourier transform product on the coordinate described by wave number (reciprocal lattice space or k-space). Thus, each electronic state described by the wave function Jk(r) within the energy band is characterized by wave number k (inverse of wave length) or that multiplied by 2p. Here the wave numbers kx, ky, kz along x, y, z directions, respectively, are described by the wave number vector k. This is called as energy band structure of electrons in crystalline solid. Fig. 3.2A shows part of a graphene honeycomb lattice located on the real space coordinate, where carbon atoms are described by open circles. Here, the unit vectors along x and y directions are x ¼ (1, 0), y ¼ (0, 1), respectively. The dashed lines denote a primitive cell of honeycomb lattice, containing two individual carbon atoms in crystallography viewpoint labeled as A and B. Namely, the graphene honeycomb lattice is composed by connecting two sublattices (triangle lattices) consisting of A or B carbon sites. The vectors
pffiffi 
pffiffi  3 1 3 1 a1 ¼ 2 ; 2 a; a2 ¼ 2 ; 2 a denote primitive vectors, where a ¼ ja1j ¼ ja2j ¼ 0.246 nm is the lattice constant. A position of an arbitrary lattice point in graphene honeycomb

Electrical properties and applications 175 (A)

(B)

y

ky

b1

K

a1 a2

τ2 A

B

x

Γ

M

kx

b2 Figure 3.2 (A) The honeycomb lattice of graphene. (B) The reciprocal lattice of honeycomb lattice shown in (A).

lattice is represented by a lattice vector R ¼ n1a2 þ n2a2(n1, n2 ¼ 0, 1, 2, 3 .), where the nearest neighbor carbon sites (A and B sites) are connected by the vector s2 ¼ 13 ða1 þ a2 Þ. It means that the network of carbon sites of the graphene honeycomb lattice is described by two triangle sublattices connected by the vector s2, where each lattice consists of only A and B carbon sites. The two sublattices structure of a honeycomb lattice is featured as “bipartite lattice” in terms of topology in math, significantly influencing novel electronic properties of graphene. Here, it should be noted that a primitive cell is not unique and one can choose an arbitrary primitive cell so as to make the following calculation simplest. Indeed, another primitive cell is chosen in the case of ribbon-shaped graphene, taking the total symmetry including the geometry of edges into consideration. Fig. 3.2B shows the reciprocal primitive vectors     p1ffiffi; 1 2p in the reciprocal lattice corresponding to the b1 ¼ p1ffiffi; 1 2p ; b ¼ 2 a a 3

3

primitive vectors of graphene in the real space shown in Fig. 3.2A. Here, the reciprocal lattice space and real space are correlated by the relations of b1$a1 ¼ b2$a2 ¼ 2p, b1$a2 ¼ b2$a1 ¼ 0. The G point is the origin of the reciprocal lattice. The primitive cell for the lattice represented by the reciprocal lattice vectors b1, b2 appears as a hexagon denoted by solid lines in Fig. 3.2B. A primitive cell in the reciprocal lattice is usually termed as the (first) Brillouin zone, where highly symmetric points are labeled as     p1ffiffi; 0 2p. Especially, K points are classified into two types: K K ¼ p1ffiffi; 13 2p ; M ¼ a a 3

3

and K0 points due to two crystallographically individual atoms in the primitive cell of graphene honeycomb lattice. Areas in the vicinity of K or K0 points in the reciprocal space are also termed as “valley.” In graphene, it is considered that 2s orbital and two 2p orbitals of each carbon atom are hybridized and form s-band fulfilled by 3/4 of valence electrons of carbon atoms. The rest

176 Chapter 3 2p orbital (pz) of carbon atoms that do not contribute the hybridization are orthogonal to the electronic states in s-band. The pz orbitals having higher energy levels than the sband form p-band partially filled by 1/4 of valence electrons of carbon atoms. Thus, the electronic states for filled electrons having the highest energy (Fermi energy) in graphene are well described by pz orbital of carbon atom interacting each other in honeycomb lattice. The calculation of the band structure of graphene around the Fermi energy was performed with tight-binding approximation by Wallace [1]. In the tight-binding model, the interactions only between the nearest neighbor carbon atoms are considered. Thus, taking into consideration that the overlap between wave functions for adjacent pz orbitals is negligible, the Hamiltonian for one electron in p band of graphene is described as 
0 H  ðkÞ (3.1) Hb ¼ HðkÞ 0 HðkÞ ¼ g0 gðkÞ (3.2) where the resonance integral g0 and g(k) are defined as following: Z Z 3 g0 ¼  fz ðrÞ½VðrÞ  V0 fz ðr  s2 Þdr þ S fz ðrÞ½VðrÞ  V0 fz ðrÞdr3   gðkÞ ¼ expðik$s2 Þ þ expðik$D3 s2 Þ þ exp ik$D1 3 s2

(3.3) (3.4)

where r, f(r), V(r), V0, and D3 are the position vector of an electron, the pz orbital of a carbon atom, the lattice potential of graphene, atomic potential of a carbon atom, and the 120 degree rotation operator, respectively. g0 is a parameter representing the electronic interaction between the nearest neighbor carbon atoms in graphene, and estimated as 3.15 eV [2]. To diagonalize Hb so as that the Schrodinger equation has nonzero solution, the energy of p-band E of graphene is obtained from the following secular equation: E g0 gðkÞ ¼0 (3.5) g g ðkÞ E 0 Here, k ¼ (kx, ky) is the wave number vector of the wave function representing an electronic state in p-band of graphene, which is connected to the eigen value of the momentum of electron in the state as p ¼ Zk. Finally, the energy of p-band of graphene is obtained as sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi


pffiffiffi  ky a ky a 3kx a 2 þ 4 cos EC;V ðkÞ ¼ g0 gðkÞ ¼ g0 1 þ 4 cos cos (3.6) 2 2 2 where positive and negative sign in Eq. (3.6) correspond to EC and EV, where C and V represent the conduction band (p band) and valence band (p* band), respectively [1]. Here, one electron per a carbon atom is accommodated into resulting p and p* bands, where p band is fulfilled and p* is not occupied. Thus, the Fermi energy, which is the

Electrical properties and applications 177 highest energy among filled electrons in solid, is located between p and p* bands. Hereafter, the origin of the electron energy is fixed at the Fermi energy. Generally, the electron energy E(k) as the function of the wave number vector of wave functions in crystalline solids is called as the “dispersion relation.” Due to the translational symmetry of the crystal structure, any electronic states of graphene are represented by wave functions having wave number vectors within the Brillouin zone in the reciprocal lattice space (k-space) shown in Fig. 3.2B. Fig. 3.3A shows a three dimensional plot of EC,V(k) represented by Eq. (3.6) [3]. A cross-section of Fig. 3.3A along the K, G, M points in the Brillouin zone is also shown as Fig. 3.3B [4]. The point marked by the circle in Fig. 3.3A, where the valence and conduction band touch each other, is called the “Dirac point.” The Dirac point is located at the K point kx and ky axes and the Fermi energy in the energy axis in the three-dimensional plot shown in Fig. 3.3A according to the cross-sectional plot shown in Fig. 3.3B. Similarly, the top and bottom of the p (valence) and p* (conduction) bands are located at the G point in the Brillouin zone, seen in Fig. 3.3B. As described by Eq. (3.6), the electronic energy band structure of graphene has the mirror symmetry with respect to the energy plane at the Fermi energy. Especially, the electronic structure in the vicinity of the Dirac point is isotropic in the reciprocal space. Namely, E(k) versus k plot is cone-shaped and independent of the direction of wave number vector k as shown in the magnified image in Fig. 3.3A. Changing the origin of the reciprocal lattice space from the G point to the K point, E(k) in the vicinity of the Dirac point approximated as pffiffiffi 3 EC;V ðkÞ ¼  g0 ak (3.7) 2

(B) 4

Ek

kx

ky

3 2 1 0 -1 -2

Energy / units of γ0

(A)

-3

Κ

Μ Γ Wave vector

Κ

Figure 3.3 (A) The p electronic band structure of graphene obtained by the nearest-neighbor tight-binding model. The inset shows the magnified image for the electronic structure in the vicinity of the Dirac point [3]. (B) The cross-sectional image of (A) along K / G / M / K points in the reciprocal lattice [4].

178 Chapter 3 where k ¼ jkj, the norm of the wave number vector k is the distance from the changed origin (K point) in the reciprocal lattice space and the origin of the energy is that at the Dirac point. The electron energy proportional to the wave number in graphene is quite peculiar in contrast to the dispersion relation for a typical two-dimensional electron system with a parabolic energy function of the wave number. The linear band (dispersion) characteristic is responsible for many novel features in the electronic properties of graphene. By using Eq. (3.7), the electron density of states in the vicinity of the Dirac point is derived as DðEÞ ¼

8jEj 3pg20 a2

(3.8)

Integrating the density of states D(E) by the energy E, the carrier density of graphene is obtained as n¼

4 E2 3pg20 a2

(3.9)

which corresponds the density of electrons and holes for positive and negative E, respectively. As indicated in Eq. (3.8), the density of states is proportional to the energy, and is zero at the Fermi energy. This means that the energy gap between p (valence) and p* (conduction) bands is zero, so graphene is often referred to as “zero-gap semiconductor” (Fig. 3.4A). The electronic structure of multilayer graphene, which is a stack of more than two graphene sheets, is calculated by introducing the interlayer interaction between graphene sheets. The most stable structure of multilayer graphene is equivalent to graphite, where each sheet is stacked in the manner so called as AB (Bernal) stacking as shown in (A)

(B) V

C

V

EF

Edge state

E

C

EF

E

Figure 3.4 (A) The energy dependence of density of states D(E) for ideal graphene. The Dirac point is located at the point where the conduction (C) and valence (V) bands touch each other. (B) The energy dependence of the density of states for graphene with edges having the edge states at the Dirac point. In both of the panels, the Fermi energy EF is downshifted below the Dirac point, and filled states are denoted by hatch in order to illustrate the tunability of EF of graphene according to the charge transfer with the external systems.

Electrical properties and applications 179

Figure 3.5 (A) AB (Bernal) stacking structure for multilayer graphene. (B) The stacking manner for oddnumbered (solid lines) and even-numbered layers (dotted lines) [5].

Fig. 3.5A [5]. As mentioned earlier, carbon atoms in graphene are classified into two sites (A and B). In AB-stacked structure, the second layer is parallelly placed below the first layer so as that the B sites in the first layer (B1) are superimposed into the A sites in the second layer (A2), where the A sites in the first layer (A1) are superimposed into the center of the hexagon of the honeycomb lattice of the second layer. The third layer is parallelly placed below the second layer with the same distance between the first and second layers so as that A and B sites in the third layer (A3 and B3) are superimposed into A and B sites in the first layer (A1 and B1), respectively. Similarly, odd-numbered layers are stacked so that A and B sites are superimposed into A1 and B1 sites in the first layer, respectively, and even-numbered layers are stacked so as to superimpose them in it into the second layer in the same manner (Fig. 3.5B). Thus, the primitive cell of AB-stacked multilayer graphene consists of two graphene sheets with the lattice constant perpendicular to the sheets c0 as shown in Fig. 3.5A. In AB-stacked structure, the resonance integrals between pz orbitals of the nearest neighbor carbon atoms on the adjacent graphene sheets are denoted by g1 w 0.4 eV [6,7] for the interlayer distance the same as that of graphite c0/2 ¼ 0.335 nm [8]. The interactions between the next-nearest neighbor carbon atoms on the adjacent layers are represented by g3. In the tight-binding approximation including only the nearest-neighbor interactions between layers, the electronic band structure of multilayer graphene in comparison with that of monolayer graphene is shown as Fig. 3.6 [9]. The electronic structure of bilayer graphene consists of four bands, including pairs of conduction and valence bands, in one pair of which the bands touch each other at the Dirac point (K point in the reciprocal lattice space). In the other pair, the bottom and top of the conduction and valence bands are placed at the K point, but they are shifted upper

180 Chapter 3 (A)

(B)

E

E C2

C

kF

C1

kF2 kF1

k

EF

2γ 1

k

EF

V

V1

V2

Figure 3.6 The energy band structure of monolayer graphene (A) and bilayer graphene (B), where C (C1, C2) and V (V1, V2) denote the conduction bands and valence bands, respectively. The Fermi energy EF is downshifted below the Dirac point to indicate how to measure the position of EF of graphene according to the charge transfer with the external systems, where kF, kF1, and kF2 denote the Fermi wave number corresponding to the wave numbers for the electronic states having EF.

and lower by g1 from the Dirac point along the vertical energy axis, respectively. So, bilayer graphene also exhibits a symmetric electronic structure with respect to the Dirac point, and a zero-band gap feature similar to that of monolayer graphene. However, the electronic energies for the four bands are parabolic to the wave number, where the linear dispersion relation disappears. The electronic structures of multilayer graphene more than three layers are well understood by a combination of those for monolayer graphene and bilayer graphene [5]. Consequently, multilayer graphene with even number 2i (i ¼ integer) of layers has 4i of the parabolic electronic bands similar to that for bilayer graphene. In graphene with odd number (2i þ 1) of layers, 4i of the parabolic bands (bilayer type) and two linear bands like that of monolayer graphene appear in the electronic structure (Fig. 3.7). Thus, the electronic properties of multilayer graphene depend on the parity of its layer number, which is understood by the electronic structure described by a sum of contributions by partial bands (monolayer or bilayer types). It should be noted that these results are derived with consideration of only pz orbitals, and valid in the vicinity of the Dirac point (typically 0.1e0.2 eV in the energy axis). When the electron energy deviates far from the Dirac point, the contribution of hybridized orbitals between s and p orbitals should be taken into account. Moreover, the electronic

Electrical properties and applications 181

Figure 3.7 Electronic band structures of multilayer graphenes with layer numbers N ¼ 2, 3, 4, 5 around K point along kx axis (left panel for each N). The magnified image of the left (right panel for each N) [5].

structure of graphene is significantly modified with a large degree of loss in the planarity of graphene sheet by some reasons through contributions of 2s, 2px, and 2py orbitals [3].

3.1.2 Effects of defects and edges As shown in the previous section, graphene having ideal structure exhibits novel electronic properties due to its topologically preserved robust electronic structure with a linear dispersion band and zero-band gap structure, where the carrier scattering is suppressed, being responsible for the extremely high carrier mobility. However, the electronic structures are modified by a structural modification in graphene at the large degree, for example, atomic vacancies in the honeycomb lattice, where carriers are significantly scattered. Edges that are considered as one-dimensionally aligned atomic vacancies in honeycomb lattice are also important defects in graphene. Especially, the presence of edges plays a significant role in practical graphene materials typically having finite crystalline size. In graphene materials with bulk crystalline size, the contributions of the edge part are negligible in its electronic structure and properties. However, as the crystalline size of graphene reduces, the geometrical ratio of the edge parts to the basal plane parts becomes larger as usually seen in nano-sized materials, where the directions and chemical structures of edges significantly influence the electronic structure of graphene. The effect of the edges in graphene has been theoretically developed mainly by Yamabe and Fujita since the 1990s [10,11]. In geometrical point of view, the edges in honeycomb

182 Chapter 3

Figure 3.8 Ribbon-shaped graphene having armchair edge (A) and zigzag edge (B), where the edges extending in the vertical direction are shown. Carbon atoms belonging to sublattices are classified by labeling with A or B. The width of the ribbon is denoted by N [11].

lattice of graphene are classified into two types of edges: zigzag edge (trans-polyacetylene type) and armchair edge (cis-polyacetylene type) as shown in Fig. 3.8. Arbitrary-shaped edges are described by a combination of those two representative types of edges in graphene. With consideration for a ribbon-shaped graphene with infinitely extended edges parallelly faced, Fujita et al. theoretically predicted that nonbonding p electronic states called as “edge state” emerge at the Dirac point in addition to valence p and conduction p* bands inherited from ideal graphene structure when zigzag edges are introduced in graphene honeycomb lattice [11,12]. The edge state is almost spatially localized at the edge part of graphene and dominates the electronic states of edge parts. In contrast to the zigzag edge, the presence of the armchair edge does not significantly influence the electronic structure around the Dirac point. Fig. 3.9A shows the spatial distribution of wave functions near the edge part in ribbonshaped graphene with zigzag edges, where k denotes the norm of the wave number vector (B)

E

(A)

k=π

k = 8π/9

k = 7π/9

k = 2π/3

Figure 3.9 (A) Spatial distribution of wave functions near the zigzag-edge part in ribbon-shaped graphene for various k. (B) The electronic energy band structure of ribbon-shaped graphene with zigzag edges. Open ellipsoid denotes the flat band nature of the edge states [11].

Electrical properties and applications 183 along the zigzag edge direction multiplied by the lattice constant. It should be noted that a rectangular primitive cell was applied for this calculation taking the symmetry of the ribbon shape into consideration, where the zigzag edges are directly parallel to the wave vector at the K point in the Brillouin zone of bulk ideal graphene shown in Fig. 3.9B. Thus, the horizontal axis of Fig. 3.9A corresponds from the G point toward K point direction in the Brillouin zone shown in Fig. 3.2B. In the case of k ¼ p, which corresponds to an electronic state having the energy at the Dirac point, the wave function is completely localized at the zigzag edge. As the wave number decreases, wave functions extend away from the zigzag edges, and eventually become delocalized completely over the whole graphene plane at k ¼ 2p/3. Here, the energy of electronic states remains the Dirac point energy. The huge number of degenerated electronic states between k ¼ p and 2p/3 as indicated by the ellipsoid in Fig. 3.9B results in the large density of states at the Dirac point for graphene with edges in addition to those for valence p and conduction p* bands in ideal graphene (Fig. 3.4B). The emerging the edge states for the zigzag edges is understood in view of topology of the honeycomb lattices having two triangle sublattices consisting of A and B sites (bipartite lattice). As shown in Fig. 3.8, the zigzag edge is terminated by only one of either of A or B sites. This means local breaking of the symmetry between A and B sites for the zigzag edges. Suggested by the one electron Hamiltonian shown in Eq. (3.1), the electronic structure of graphene is determined by the paths between wave functions for adjacent carbon atoms belonging to each A and B sites connected by the resonance integral. Here, the topology of the network of carbon atoms dominates the solution for the Schrodinger equation. As long as the symmetry of A and B site remains, similar electronic structures to ideal graphene are obtained as solutions of the equation, even if we consider nonideal graphene structures with defects. However, the electronic structure is significantly modified if the symmetry between A and B sites is broken by defects. The symmetry breaking for sublattices is an intrinsic reason for appearance of the edge states in the zigzag edges. On the other hand, the symmetry remains in the armchair edge. It is trivial that A and B sites are symmetrical for ideal bulk graphene with an infinite honeycomb lattice. Indeed, the electronic structure of graphene with the armchair edges shows similar electronic structure to that of ideal graphene. No significant influence appears on the electronic structure for defects remaining in the symmetry for sublattices. This is also confirmed by computer simulations for the band structure of graphene based on the density functional theory (DFT) [13]. Fig. 3.10 shows the electronic structure of ribbonshaped graphene with the zigzag edges obtained by the first-principles band calculation. As clearly seen in the feature of band structure around 0 eV in the energy of the plot, edge states having the flat band nature around the X point in the horizontal axis appear for ribbon-shaped graphene structures with width N of 2e6 as shown in Fig. 3.10A. It should be noted that the dangling bonds at the edges are fully terminated by hydrogen atoms. So,

184 Chapter 3

Figure 3.10 (A) Ribbon-shaped graphene with zigzag edges terminated by hydrogen having width of N ¼ 2, 4, 6. (B) The electronic band structures for ribbon-shaped graphene with zigzag edges having width of N ¼ 2, 4, 6 [13].

the edge states that appear here are not typical surface states concerning the unsaturated chemical bonds. In these viewpoints, we can consider the influence of the presence of general defects in nonideal graphene. The presence of an atomic vacancy in the honeycomb lattice could be considered as the simplest disorder in graphene [14e16]. In the tight-binding scheme, the presence of an atomic vacancy results in an additional electronic state at the Dirac point, which gives the large density of states (Fig. 3.11) [15]. The additional electronic states are specially localized in the vicinity of the vacancy in the honeycomb lattice. Here, the emerging additional electronic states have quite similar nature to that of the edge states in the zigzag edges. It is well understood by the modification in the topology of honeycomb lattice. An atomic vacancy locally breaks the symmetry of A and B sites similarly to the presence of the zigzag edges. Indeed, the influence of vacancy -induced localized states becomes obscure in the total density of states as shown in Fig. 3.11, when the bipartite nature of graphene lattice is lost as increasing in the interactions between the next-nearest neighbor carbon sites. DFT calculations also show the appearance of the vacancy-induced electronic states similar to the edge states, indicating that the effects of the next-nearest neighbor interactions are weak in actual graphene materials [14,16]. This justifies the results obtained in the above discussion in the framework of simple tight-binding model. On the other hand, no additional electronic states emerge in the case of vacancies on adjacent sites in graphene honeycomb lattice, where the additional vacancy is generated by covalent bonding of two hydrogen atoms to the carbon atom at the neighbor site of an atomic vacancy site (Fig. 3.12). As two vacancies in the adjacent sites retains the symmetry of A

Electrical properties and applications 185

Figure 3.11 (A) Local density of states in the vicinity of a vacancy site for graphene with a vacancy and the total density of states for ideal graphene. (B) The total density of states in the vicinity of the Dirac point with various vacancy densities for graphene with a vacancy [15].

and B sites, the electronic structure of the ideal bulk graphene appears as similar to that of graphene with armchair edges [16]. Experimentally, a direct observation of the edge states in graphene with edges has been achieved by using scanning tunneling microscopy (STM) and scanning tunneling spectroscopy (STS) [17,18]. Due to the high chemical activity of edges of graphene, edges are oxidized and terminated by oxygen-containing functional groups having geometrically bulky form. Thus, the structure of edge parts of graphene is hardly observed in ambient conditions. It is mandatory to prepare well-defined edge structure for in situ observation and evaluation of graphene edges by STM/STS. Annealing graphene up to 800 C followed by exposing atomic hydrogen in ultra-high vacuum (UHV) enables the thermal decomposition of terminating functional groups and hydrogen termination in graphene edges. According to STM observation in UHV for graphene with well-defined edges obtained through this process, the zigzag edges are defective and short in comparison with the armchair edges widely extended with semi-ideal structure (Fig. 3.13) [17]. Indeed, the armchair edges are frequently seen in the edge parts of graphene, whereas the zigzag

186 Chapter 3 (A)

(B) 4 3

E-EF

2 1 0

-1 -2 -3 -4 Γ

(C)

Κ Μ

Γ 0

(D) 3

E-EF

2

10

20

30

DOS / a.u.

4

1 0

-1 -2 -3 -4

Γ

Κ Μ

Γ0

10

20

30

DOS / a.u.

Figure 3.12 (A) Graphene with a single vacancy fully terminated by hydrogen. (B) The electronic band structure and the density of states for graphene shown in (A). (C) Graphene with vacancies on both (A) and (B) sites obtained by bonding to additional hydrogen to the structure shown in (A). (D) The electronic band structure and the density of states for graphene shown in (C) [16].

Figure 3.13 (A) STM image of the armchair edge in a step edge on graphite. (B) A zigzag edge embedded in armchair edges in a step edge on graphite [17].

Electrical properties and applications 187 (B)

200

dI/dVS / nA/V

150

dI/dV / nA/V

(A)

0 -0.5

0.0 VS / V

0.5

0 -0.5

0.0 VS / V

0.5

Figure 3.14 (A) STS spectrum near the armchair edge. (B) The STS spectrum near the zigzag edge [17].

edges are observed in short length as a defect structure embedded in the armchair edges on rare occasions. Those experimental results are in good agreement with the theoretical predictions of instability of the zigzag edges having the large density of states at the Fermi energy in view of electron energy in spite of the thermodynamically stable nature of the armchair edges. The presence of the edge states is directly proved by the STS spectra of which vertical axis dI/dV almost corresponds to the local density of states. As shown in Fig. 3.14A, the STS spectrum in the vicinity of the armchair edges in graphene shows linear p and p* bands touched each other at the Fermi energy, where the dI/dV (the density of states) is zero and proportional to the energy. The features of the electronic structure of graphene as a zero-gap semiconductor appear as similar to that of ideal graphene. On the other hand, the peak of dI/dV at the Fermi energy in the STS spectrum indicates the presence of the large density of states at the zigzag edges as shown in Fig. 3.14B These findings are well confirmed by the significant enhancement in the tunneling current (bright spots) at the zigzag edge parts in Fig. 3.13B in spite of the similar contrast at the armchair part as that of the basal plane part in the microscope images. Generally, the electronic properties of materials are mainly dominated by the electronic states around the Fermi energy. Thus, emerging of the localized states like edge states having the large population at the Dirac point where the Fermi energy is located in the electrically neutral condition as shown in Fig. 3.9B greatly modifies the electronic properties of graphene with edges from that of ideal graphene. The large density of states at the Fermi energy works as an electron reservoir in phenomena involving the interactions between guest chemical species and graphene. This explains that defect sites in graphene such as vacancy, edges (crystalline grain boundary), and impurities are chemically active and play an important role in terms of graphene application toward electrodes, catalysts, adsorbents, and so on. Emerging spin magnetism is one of the most remarkable changes in the electronic properties induced by the presence of the edges in graphene. Ideal bulk graphene exhibits only diamagnetism by orbital momentum of electrons, where no spin magnetism appears.

188 Chapter 3 (A)

(B) N=10

m

0.8

m=0.19

1A 2A 0.4

2D Graphite 5A 0

Uc 4 Ult

8

Figure 3.15 (A) Local magnetic moment for ribbon-shaped graphene with zigzag edges having width of N ¼ 10, where numbers and alphabets denote the carbon sites defined in Fig. 3.8. “2D Graphite” means ideal bulk graphene. (B) Schematic view of spins in graphene with edges [11].

As shown in Fig. 3.9, the edge states appear as two flat bands in the wave number region of 2p/3 < k < p. Namely, a large number of degenerated states exist in the vicinity of the Fermi energy for graphene with edges, being responsible for the large density of states at the Fermi energy as shown in Fig. 3.4B. In the electrically neutral condition, the flat bands of edge states are half filled, and the Fermi energy is placed at the center of the edge states in the energy axis. In this case, even slight perturbations such as tiny electron correlation interactions could remove the degeneracy of the electronic states, because a huge energy gain is obtained by removing the degeneracy of the electronic states. Indeed, Fujita et al. show the energy gain for the spin-polarized states obtained by infinitesimally small Coulomb interactions is enough large to stabilize the ferromagnetic alignment of spins of electrons in graphene with zigzag edges [11,19]. It means that the parallel spin alignment is favored for the edge states in graphene with zigzag edges. By introducing edges, the spin magnetism is given to diamagnetic (nonmagnetic) graphene. Fig. 3.15A shows the magnetic momentum m that is given by the difference between densities of electrons having up and down spins as a function of the strength of the Coulomb interactions U between electrons in graphene with the zigzag edges. The local magnetic moment m at the topmost sites (1A shown in Fig.3.8B) in the zigzag edge in graphene rapidly increases with increases in the Coulomb interactions U and becomes almost 0.2 even at small U (～0.3 eV). It should be noted that the magnetic moment m rapidly decays as being away from the edge. This indicates that the edge states specially localized at the edge part of graphene are responsible for the derived spin magnetism here, which is schematically shown as Fig. 3.15B. On the other hand, the armchair edges without the edge states in graphene exhibit no magnetic moment in the case of the smaller strength of the Coulomb interactions, which is achievable in practical materials as similar to that for ideal graphene as shown in Fig. 3.15A. Since the origin and nature is same as those for the edge states as for the vacancy-induced localized states, we can obtain the spin magnetism of graphene by introducing vacancies.

Electrical properties and applications 189 (A)

(B) 0.8 Energy / eV

Energy / eV

0.8 0.4 0.0 π-bands

-0.4

0.0 π-bands

-0.4

-0.8 Γ

0.4

-0.8 Μ Κ Wavevector

Γ -20 0 20 DOS

Γ

σ-band

Μ Κ Wavevector

Γ -20 0 20 DOS

Figure 3.16 Spin-resolved electronic band structure and the density of states for graphene with a vacancy fully terminated by hydrogen (A) and with an unsaturated vacancy (B) [20].

Fig. 3.16A shows the electronic band structure for graphene with a vacancy where dangling bonds left after removing a carbon atom are fully terminated by hydrogen atoms [20]. Flat bands similar to the edge states appear around 0 eV in the region between M and K in the horizontal axis. Here, the energies for electrons with up spins and down spins are completely different from each other. In the electrically neutral condition, p electronic states having only either up or down spins are occupied, where electron spins align parallel in graphene, giving the spin magnetism. Similar nature appears also for the vacancy of which dangling bonds are not terminated as shown in Fig. 3.16B. In addition to the splitting of p states with respect to the vertical energy axis, s states are also spinpolarized as seen around 0.8 eV in the energy axis. Both of the typical surface states and vacancy-induced states originating from the dangling bonds and asymmetry in the A and B sublattices contribute the spin magnetism for graphene with an unsaturated vacancy. Here, the Coulomb interactions are mainly considered as the origin of the spin magnetism. However, the flat band nature with a large degree of degeneracy in the electronic states induced by nonideal local structures has a large advantage to induce spin magnetism in graphene, where many other perturbations such as spin-orbital coupling interactions are proposed as its origin as well as the Coulomb interactions [21].

3.2 Applications to information technology 3.2.1 Transistor devices The application of graphene to transistor devices was attempted just after the first report about the isolation of monolayer graphene due to its high affinity for field-effect transistor (FET) devices. The Si-MOSFET device, which is widely used as a fundamental infrastructure in the information technology field, has a stacked-layer structure like a semiconductor (Si) channel/insulator(SiO2)/gate electrode. In FET devices, the current

190 Chapter 3 between the source and drain electrodes attached to the semiconductor layer is switched through the charge accumulation at the interface with the insulator layer by applying voltage to the gate electrode. When the lengths between the source and drain electrodes and the width of gate electrode become shorter down to 10 nm through the miniaturization of semiconductor devices according to Moore’s law, the drain bias weakens the electric field by the gate electrode, resulting in the loss of “OFF” state by a large leakage current (short-channel effect) [22]. One of the strategies against this problem is to decrease the thickness of the charge accumulation layer in the semiconductor channel, where applying atomically thin materials for the semiconductor channel is an ultimate solution [23]. Moreover, the charge mobility achieved in the silicon material as practically used has almost reached its theoretical limit. To satisfy the demand for more switching frequency of the logic devices, a post-silicon material is required for a channel material in FET devices. In this viewpoint, extremely high charge mobility of graphene demonstrates a large advantage for achieving ultra-high speed switching of FET devices. As shown in Section 3.1, the charge mobility of graphene is expected to exceed that of silicon or other postsilicon semiconductor materials intrinsically. However, graphene is a zero-gap semiconductor as indicated in Section 3.1. So, a process to open the band gap is necessary to apply graphene as FET devices instead of Si. The strategy to open a band gap in graphene is classified into three approaches, although other methods also have been proposed and demonstrated, such as using uniaxial strain [24,25], graphene-substrate interactions [26]. The first approach is a lateral confinement of electrons in graphene [27e31]. When the size of materials is reduced into nano-dimensions, the discreteness of the energy levels in the electronic band structure becomes comparative with the energy scale of room temperature, which is a typical effect on the electronic structure of a nano-sized system, so called as “size effect.” This is easily understood for graphene as it is the large extreme of aromatic molecules, seamlessly connected to polycyclic aromatic hydrocarbons such as coronene and circumanthracene having large HOMO-LUMO gaps. Han et al. fabricated ribbon-shaped graphenes having various narrow widths by electron-beam lithography and oxygen plasma etching technique (Fig. 3.17AeC) [28]. In ribbon-shaped graphenes with nano-sized widths, electrons in graphene are confined into a quasi-one-dimensional (1D) system. The gate voltage Vg dependence of the conductance G of ribbon-shaped graphene as shown in Fig. 3.17DeF exhibit a region of depressed G with respect to Vg. This is in good agreement with bulk graphene having a minimum conductance in the gate voltage dependence (the charge neutral point) corresponding to the Dirac point, where the number of conduction carriers becomes minimal due to the zero density of states. However, ribbon-shaped graphene samples show a significant decrease in G as the temperature decreases in contrast to bulk graphene exhibiting quite less change in the conductance

Electrical properties and applications 191 (D)

G / μS

(A)

(B)

G / μS

(E)

(F)

G / μS

(C)

Figure 3.17 (AeC) Ribbon-shaped graphene samples attached with metal electrodes. (DeF) The gate voltage Vg dependence of the conductance of ribbon-shaped graphene of varying width W (W ¼ 24, 49, and 71 nm) [28].

30 mK and room temperature. Moreover, a gap region appears for 25 < Vg < 45 V for narrowest samples at 1.7 K, where the conductance is smaller than the detection limit (Fig. 3.17D). The large depression of the conductance around the charge neutral point with decreasing temperature strongly suggests the opening of an energy gap in ribbon-shaped graphene samples. The energy gap obtained from the differential conductance for ribbon-shaped graphene a samples shows the empirical scaling law, Eg ¼ ðWW  Þ , where W* ¼ 16 nm, a ¼ 0.2 eV nm, and W is the ribbon width in nm unit. The presence of the scaling law means we can tune the band gap energy of graphene by changing the width of a ribbonshaped sample. Interestingly, no systematic crystallographic directional dependence of the energy gap Eg is observed in this study. As mentioned in Section 3.1, the electronic structure of graphene having edges strongly depends on its edge direction in crystallographic viewpoint. Taking the six-fold symmetry of the honeycomb lattice of graphene into consideration, the energy gap Eg is expected to oscillate upon the crystallographic angle

g

E

Eg / meV

/ meV

192 Chapter 3

θ / degree

W / nm Figure 3.18 Energy gap for ribbon-shaped graphene samples with various widths [28].

with a periodicity of 60 degree. However, the relative orientation angle q dependence of Eg for ribbon-shaped graphene samples obtained from two graphene flakes exhibits no systematic angle dependence as shown in the inset of Fig. 3.18. Actual operation of FET based on this strategy is demonstrated by using chemically derived ribbon-shaped graphene as shown in Fig. 3.19 [29]. By centrifugation of a dispersion of expanded graphite in a 1,2-dichloroethane (DCE) solution of poly(m-phenylenevinylene-co-2,5-dioctoxy-p-phenylenevinylene) (PmPV), ribbon-shaped graphene samples with quite narrow width down to several nm are obtained after removing large pieces of exfoliated graphene and unexfoliated graphite. FET devices using ribbon-shaped graphene show poor transistor properties for the ribbon widths more than 10 nm, where their switching behavior is gradual upon applying the gate voltage and the current on/off ratio is less than 10. On the other hand, quite sharp switching curves are observed for FET devices using ribbon-shaped graphene with widths less than 10 nm as shown in Fig. 3.20. The on/off ratio as a switching device also achieves up to 106e107 at room temperature, which is worth it for practical uses. Here, it should be noted that oxygen-containing functional groups covalently bonded would remain in ribbon-shaped graphene samples due to oxidation process in this study. Thus, the observed energy band gap is not only caused by the electron confinement due to decreasing in size of samples but also band-gap opening originating from the chemical modification involving the formation of the covalent bonding as mentioned in the next section. As shown previously, ribbon-shaped graphene samples with nano-sized width successfully exhibit an energy band gap up to a few hundred meV, which enables a fabrication of FET devices with enough large on/off ratio at room temperature. However, it is still a

Electrical properties and applications 193

Figure 3.19 (A) Schematic view for the synthesis of ribbon-shaped graphene with nano-sized widths by oxidation of expanded graphite. (BeF) AFM (atomic force microscope) images for chemically derived nano-sized ribbon-shaped graphene samples with widths from 50 nm down to several nm [29].

(B)

Ion/Ioff

-Ids / mA

(A)

Vg / V

W / nm

Figure 3.20 (A) Gate voltage dependence of the drain-source current for FET device using chemically derived ribbon-shaped graphene with width of 5 nm. (B) The on/off ratios of FET devices using ribbonshaped graphene with various widths [29].

194 Chapter 3 challenging task to fabricate graphene samples with desirable nanometer-scale size for the electron confinement in well-controllable manner in a technical viewpoint for practical uses. In addition, the carrier mobility is usually quite low in nano-sized graphene in comparison with bulk [30]. This is one of the reasons why the nano-sized ribbon-shaped graphene samples do not show clear crystallographic angle dependence. The strong disorder supposed from the low carrier mobility would make features of the electronic structure of ideal graphene obscure ones [31]. The second approach to open the band gap of graphene is the chemical modification. As mentioned in Section 3.1, the electronic band structure of graphene with zero-gap feature around the Dirac point originates from the nearest neighbor interaction between p(p)orbitals of sp2-hybridized carbons on the honeycomb lattice. When a carbon atom on graphene lattice covalently bonds to other chemical species, the sp2-hybridization of the carbon atom is converted to sp3-hybridization accompanied with an energy gap between bonding states and antibonding states. The sp2-sp3 conversions of carbon atoms on graphene sheet by covalent bonding with other species are known in graphite oxide [32] and graphite fluoride [33] even from the 19th century. The idea is also applied to nanocarbon materials like single-wall carbon nanotubes, where sp2-sp3 conversions by fluorination are used for cutting a longer carbon nanotube into smaller nanotubes [34]. The band gap opening of nanographene induced by the covalent bonding with fluorine is also reported based on the change in the temperature dependence of the conductivity [35]. Hydrogenation of graphene is one of the promising routes to open the band gap [36], where adsorption of atomic hydrogen obtained by thermal cracking of gaseous hydrogen molecule enables a controllable hydrogenation. Atomic hydrogen adsorbing on graphene forms covalent bonding with carbon atoms from the front side and back side of the surface in alternative manner as shown in Fig. 3.21, where the sp2-sp3 conversions of carbon atoms are expected to result in opening the gap in the electronic structure of graphene.

Figure 3.21 Schematic view of the structure of the hexagon ring of carbon for graphene before (A) and after (B) hydrogenation [37].

Electrical properties and applications 195 (A)

(B)

10 5

(C)

(D)

400

0 100

200 0 6 4 2 0

(E)

(F)

ρxx / kΩ

ρ / kΩ

6 4 2 0

0 10 5

-50

-25

0

25

50 -50 Vg / V

-25

0

25

50

0

Figure 3.22 Gate voltage dependence of the resistance of graphene FET before hydrogenation under no magnetic field at 40, 80, and 160 K, which practically coincide (A). The gate voltage dependence of the resistance of graphene FET before hydrogenation at 4 K under the magnetic field B ¼ 14 T (B). The gate voltage dependence after hydrogenation under B ¼ 0 T at 4, 10, 20, 40, 80, and 160 K (C). The gate voltage dependence after hydrogenation at 4 K under B ¼ 14 T(D). The gate voltage dependence after annealing under B ¼ 0 T at 40, 80, and 160 K (E) and under B ¼ 14 T at 4 K (F) [37].

Elias et al. show hydrogenation of graphene samples obtained by micromechanical cleavage results in opening band gap [37]. Fig. 3.22A and C indicates significant change in the electronic structure of graphene before and after hydrogenation, respectively. The gate voltage dependence of FET with micromechanical cleavage shows a typical behavior for graphene FET, where the maximum resistance appears at Vg ¼ 0 corresponding to the charge neutral point of graphene FET. The behaviors are practically coincident at 40, 80, and 160 K. The weak temperature dependence of the resistance is in good agreement with the metallic nature of the electric conductance of graphene. However, the resistance of graphene exhibits strong temperature dependence after the hydrogenation as shown in Fig. 3.22C. As decreasing in the temperature, the resistance of graphene shows two orders of increasing. This strongly suggests a transition from a gapless metallic electronic structure to that with band gap for insulators. The change in the behavior of the resistance under the strong magnetic field (quantum Hall effect) before/ after hydrogenation also supports the opening of the gap in electronic structure of graphene after hydrogenation (Fig. 3.22B and D). Hydrogenation-induced opening of the band gap in the electronic structure of graphene is more clearly indicated in the temperature dependence of the resistance. Fig. 3.23A shows the temperature dependence of the resistance at the charge neutral point for graphene FET

196 Chapter 3 T/K

ρmax / Ω

(A)

1/T1/3 / K-1/3 Vg / V

T0 / K

(B)

n / 1012cm-2 Figure 3.23 (A): Temperature dependence of the resistance at the charge neutral point for graphene FET. The circles, squares, and triangles are for graphene before hydrogenation, after hydrogenation, and after annealing, respectively [37]. (B): Gate voltage dependence of the characteristic temperature T0.

(the maximum resistance in the gate voltage dependence), where the Fermi energy is located at the Dirac point. Before hydrogenation, the temperature dependence of the resistance of graphene is quite weak, indicating a metallic conductance for zero-gap semiconductor system as known for graphene. On the other hand, the resistance of graphene exponentially increases as temperature decreases, obeying the law of h 1=3 i exp TT0 , where T and T0 are temperature and the characteristic temperature. The h 1=3 i temperature dependence of the exp TT0 law indicates the variableerange hopping conduction for graphene after hydrogenation. This suggests that there are many localized states in the gap induced by hydrogenation. Interestingly, the characteristic temperature T0, which denotes the degree of localization and the density of the localized states, systematically changes with the density of carriers controlled by applying the gate voltage. Both the electron and hole carrier injections cause the significant reduction in T0, and the graphene gradually recovers metallic conduction with weak temperature dependence as

Electrical properties and applications 197 shown in Fig. 3.23B. The hydrogenated FET device exhibited the same behavior during repeated measurement, and was stable for many days as long as it was stored at room temperature. However, the metallic electronic structure of graphene was recovered after thermal annealing of the device. Fig. 3.22E and F shows the gate voltage dependence of the resistance under B ¼ 0 and 14 T, respectively, after annealing the FET device at 450 C in Ar atmosphere for 24 h. After the annealing, similar behaviors to those before hydrogenation corresponding to the electronic structure of graphene with zero-band gap appear under both of the zero magnetic field and the magnetic field of 14 T. The temperature dependence of the resistance is also recovered by the thermal annealing as shown in Fig. 3.23A. This indicates that opening the gap by hydrogenation is reversible in graphene. We can freely control the electronic structure of graphene as our demand. On the other hand, we can say that there is a problem of the thermal stability in the strategy of opening the band gap by hydrogenation. Moreover, it was reported that the thermal annealing at higher temperature damaged graphene irreversibly. The last method in the three main strategies to introduce the band gap into graphene toward its transistor device application is breaking the symmetry between layers in multilayer graphenes with AB (Bernal) stacking [38e40]. In the tight-binding model, which is simple but well describes the feature of the electronic structure of graphene as shown in Section 3.1, the introduction of a difference in the potential for electrons between stacked layers in AB-stacking manner for bilayer graphene results in opening band gap. Fig. 3.24A shows the band structure of bilayer graphene having potential difference D, where an extremely large D ¼ g1 (the resonance integrals between pz orbitals (A)

(B)

n / 1011cm-2

Figure 3.24 (A) Energy band structure for bilayer graphene with the potential difference D between two layers. (B) The injected carrier density into one of two layers versus the band gap for bilayer graphene [38].

198 Chapter 3 of the nearest neighbor carbon atoms on the adjacent graphene sheets as shown in Section 3.1) is used for illustrative purpose [38]. For the electronic band structure of bilayer graphene with equivalent layers, the features of the zero-gap semiconductor remain in spite of vanishing of the linear dispersion as indicated in Fig. 3.23. However, the potential difference between layers not only changes the parabolic bands into the oscillating bands but also induces a clear band gap. The energy gap between the valence and conduction bands at the K point where the Dirac point is located in monolayer graphene is equivalent e in the to the introduced potential difference. Here, the minimum band gap is defined as D figure. So, we can tune the energy gap of bilayer graphene by tuning the potential difference for electrons between layers. This effect is intrinsically absent in monolayer graphene. The potential difference between layers is experimentally tuned by carrier doping into each layer in asymmetric manner. Theoretically, the addition of carrier density e n w 1012 cm2 into one of two layers yields a gap Dw10 meV as shown in Fig. 3.24B. Opening the band gap by adding extra charge to one of two layers of bilayer graphene as schematically described as Fig. 3.25 is observed for the sample grown on SiC substrate by angle-resolved photoelectron spectroscopy (ARPES) [41]. As initially grown, bilayer graphene is slightly n-doped due to the difference in the electron negativity between graphene and SiC, where carriers from the dopant in SiC are only accumulated into the bottom layer directly faced to the substrate. The inhomogeneity in the doped carrier distribution among stacked graphene layers is well known for higher stage graphite intercalation compounds, where the charges are almost injected into only the layer directly adjacent to intercalants due to the extremely short charge screening length along the axis perpendicular to the graphene plane [42]. Thus, asymmetric doping in each layer is achieved for bilayer graphene due to the polar interface with the substrate or additional adsorption of dopant as the adsorbed chemical species on the surface of the top layer. ARPES measurement reveals the population of electrons in solid for each energy and momentum values individually. Thus, the spectroscopy directly gives the information (A)

π*

(B)

π*

(C)

π*

E=E0

π

π

π

Figure 3.25 (A) Monolayer graphene and its electronic band structure around the Dirac point in a threedimensional plot. (B) Bilayer graphene with symmetric charge population between layers and its band structure around the Dirac point. (C) Bilayer graphene with asymmetric charge distribution between layers and its band structure with an energy gap.

Electrical properties and applications 199

Binding energy / eV

0.005e0.2 (A) 0.0 -0.2 -0.4 -0.6 -0.8 -1.0 -1.2 -1.4

0.0350e-

0.0125e-

(B)

(C)

Momentum

Figure 3.26 Electronic band structure of bilayer graphene obtained by ARPES for as-prepared bilayer graphene (A), and bilayer graphene with progressive adsorption of potassium atoms (B, C). The numbers at the top of each panel denote the total doping electrons per unit cell in bilayer graphene. The solid lines in the left part of each panel are theoretically obtained band structure for bilayer graphene. The energy axis is measured from the Fermi energy for each plot [41].

about electronic band structure, namely, the electron energy as a function of its momentum. Fig. 3.26A shows the energies of electrons in bilayer graphene for each momentum. The splitting of the valence band of w0.4 eV for each plot is in good agreement with the theoretically predicted band structure for bilayer graphene. The Fermi level (E ¼ 0) crosses the conduction band of the as-prepared sample because it is n-doped from the SiC substrate. The number of doped electrons per unit cell is calculated by comparison with theoretically obtained band structure and the observed shift of the Fermi energy from the energy of crossing point between the valence band and conduction band. Most importantly, an energy gap appears between the valence and conduction bands in Fig. 3.26A. However, the band gap becomes closed after potassium atom adsorption on the surface of the top layer of the sample, although the total amount of the doped electron increases in the sample. This is explained by the recovery of the symmetry between layers in bilayer graphene caused by electron doping from the adsorbed potassium on the top layer. Charges distribute homogenously in bilayer graphene, where both of the top and bottom layers are equally e-doped from the adsorbates and the substrate, respectively in Fig.3.26B. However, further adsorption of potassium atoms on the surface results in breaking the symmetry between two layers in bilayer graphene. Thus, the band gap appears again in Fig. 3.26C accompanied with a further upshift of the Fermi energy due to more electron doping. The presence of band gap in multilayer graphene is also theoretically expected to influence the optical properties significantly, where the appearance of intense peaks in optical absorbance and their systematic shifts upon the degree of asymmetry in charge distribution among layers are predicted for multilayer graphene. Instead of ARPES measurement, which is complicated and time consuming, a

200 Chapter 3 simple optical absorption might be a promising tool to evaluate the band gap opening in multilayer graphene. The attempt of the band gap opening by introduction of breaking the symmetry layers stacked in the Bernal manner is initially performed by inhomogeneous charge doping among layers caused by chemical doping. However, with chemical doping it is difficult to control the amount of charge precisely. Moreover, reversible doping is impossible in most cases, where removing dopant usually involves damages on the graphene structure. Applying the external electric field perpendicular to bilayer graphene is a promising strategy to achieve well-tunable and reversible introduction of the band gap to graphene. Here, bilayer graphene can be regarded as a parallel plate capacitor with a dielectric medium of the vacuum between layers, where the electron potential between layers is continuously and reversibly controlled by the external electric field. The field effect transistor (FET) structure is suitable to apply the perpendicular electric field to graphene. However, applying gate voltage causes a charge injection to graphene as well as applying the perpendicular electric field to layers. Even if the band gap opens by the perpendicular electric field, the position of the Fermi energy is shifted away from the band gap, and the electronic states of graphene become metallic due to the injected carriers. Metallic conduction of a sample easily hides the figure print of the presence of band gap in simple conductivity or optical measurements. Wan et al. overcame this problem by using dual gate FET structure, which is equipped with top and bottom gate electrode as shown in Fig. 3.27 [43], by independently manipulating the voltages of the two gates with two important parameters: the degree of doping and the size of the band gap induced by the perpendicular electric field in bilayer graphene.

Figure 3.27 (A) Optical image for the top view of the dual gate FET structure. (B) Schematic side view of the dual gate FET. (C) Electric field perpendicular to bilayer graphene applied by the top and bottom gate electrodes [43].

Electrical properties and applications 201 (A)

(C)

Absorption difference / %

(B)

Energy / meV

Figure 3.28 (A) Allowed optical transitions between different sub-bands of bilayer graphene with an energy gap opened by electrical gating perpendicular to graphene. (B) Gate-induced infrared absorption spectra at the charge neutral point for various external perpendicular electric field. For clarity, the upper spectra were shifted by 2%, 4%, and 8%, respectively. Dashed lines are guides for eyes. (C) Infrared absorption spectra obtained by theory based on a tight-binding model with the band gap, which is taken as an adjustable parameter [43].

Infrared absorption spectra for bilayer graphene are shown in Fig. 3.28, where different electric fields perpendicular to graphene are applied by dual gate FET structure. The spectra were acquired at the charge neutral point, where the Fermi energy is located at the midpoint between the valence and conduction bands. The dual gate transistor structure enables applying the perpendicular electric field while retaining the electrical neutrality of graphene. Fig. 3.28A illustrates the allowed optical transition between the conduction band and valence band of bilayer graphene with the band gap. When opening the band gap in bilayer graphene, a direct transition I appears below 300 meV in typical experimental conditions, and its energy gradually shifts to the larger energy side as the size of the energy gap increases. Transitions II, III, IV, and V occur at and above the energy of the separation of sub-bands, which is within the energy scale of g1 w 400 meV as mentioned in Section 3.1. With increasing in the electric field perpendicular to bilayer graphene by dual gate electrodes, an infrared absorption peak corresponding to the direct transition I appears around 200 meV. The peak shifts to the larger energy side and achieves 300 meV with the electric field D ¼ 3.0 V nm1. This is direct evidence of tunable opening of the band gap by applying the electric field perpendicular to bilayer graphene. The observed maximum absorption reaches 5% even for a two-atom-layer-thick object, which is extremely strong and is highest among all known materials. The broad feature around 400 meV is caused by transitions between sub-bands II, III, IV, and V. Prominent asymmetric peak around 200 meV remains irrespective to the applied electric field. This is attributed to the Fano resonance of the zone-center G-mode phonon (1585 cm1) coupled with the continuum electronic transitions.

Bandgap / meV

202 Chapter 3

200

100

0

Experimental Self-consistent tight-binding DFT calculation Unscreened tight-binding

0

1

D / V/nm

2

3

Figure 3.29 Electric field dependence of tunable energy bandgap in graphene bilayer obtained by absorption peaks (closed square) and theoretical models (solid and dashed lines) [43].

Fig. 3.29 shows the gate-tunable band gap of bilayer graphene experimentally and theoretically derived as a function of applied electric field. Comparing theoretical predictions with the experimental results, self-consistent tight-binding calculation and ab initio DFT reproduce the observed energy gap. This indicates the importance of selfscreening of the electric field in bilayer graphene. The reason why DFT calculation shows a slightly smaller bandgap than that obtained by the self-consistent tight-binding model is partly owing to the different parameters for on-site interlayer coupling, which are 0.4 and 0.34 eV for the tight-binding and DFT calculations, respectively. An underestimation for energy gap in semiconductors is quite usual for DFT calculation with the local density approximation [44]. The energy gap is fixed by the crystalline structure of the material in ordinary semiconductors. However, the bandgap is variable and arbitrarily changeable from 0 eV up to 250 meV at room temperature by an external electric field in bi-layer graphene. An energy gap of 250 meV is narrower than in well-known semiconductors like silicon or gallium arsenide. This suggests that bilayer graphene is not only promising as a flexible tool for nanoscale electronic devices but also for new kinds of optoelectronic devices for generating, amplifying, and detecting infrared light. Indeed, an actual switching operation of FET device by the dual gate transistor structure was demonstrated soon after the first report of the band gap opening by perpendicular electric field for bilayer graphene [45,46]. Fig. 3.30 shows the resistance of bilayer graphene FET with dual gate structure as a function of top- and back gate voltages, where Al2O3 is grown as a gate insulator layer for

Electrical properties and applications 203 (A)

150

Resistance/Ω

(B)

104 106 108

108 Resistance / Ω

VBG / V

100 50 0

106

-50

-100

104

-150 -10

-5

0 VTG / V

5

10

-100

0 VBG / V

-100

Figure 3.30 (A) Resistance of bilayer graphene FET with dual gate structure as a function of top and back gate voltages at 300 mK. (B) The back gate voltage dependence of the resistance for bilayer graphene at the charge neutral point (solid envelope line) corresponding to the black solid line in (A). Pale-gray-colored lines denote cuts corresponding to the pale-gray-colored oblique lines in (A) [45].

top gate electrode on bilayer graphene fabricated on SiO2/Si substrate by the micromechanical-cleavage method. As shown in Fig. 3.30, an almost six orders of magnitude change in the resistance is achieved by manipulating the top- and back gate voltages. The switching operation of a bilayer graphene FET is enabled by the transition from the metallic conduction for zero-gap semiconductor to the semiconductor conduction with a band gap induced by the perpendicular electric field applied by the top- and back gate electrodes. The transition is well seen in the change of the source-drain voltage dependence of the channel current of the FET as shown in Fig. 3.31. There is a nearly 40

Current / nA

20

0 D (V/nm) -2.5 -1.6

-20

-40 -5

-0.9 1.8

-1.0 -2.5

0 VSD / mV

2.5

5

Figure 3.31 Source-drain voltage Vsd dependence of the channel current with the various electric fields perpendicular to bilayer graphene [45].

204 Chapter 3 linear dependence on applied bias voltage between the source and drain electrodes observed at lower electric fields (D ¼ 0.9, 0.1 V/nm), which features the metallic conduction for a zero-gap semiconductor. On the other hand, nonlinear behavior that is typically characteristic for a semiconductor with the band gap gradually becomes remarkable in the source-drain voltage dependence of the channel current as it increases in the electric field perpendicular to bilayer graphene. In Fig. 3.31, semiconductor behaviors appear with both polarities of the applied electric field perpendicular to bilayer graphene. The asymmetry between the potentials of electrons in each layer is the intrinsic origin of the band gap opening in bilayer graphene, where the direction of the gradient of the potential does not matter as seen in the ARPES study shown in Fig. 3.26. Here, it should be noted that the observed large on/off ratio up to 106 is achieved only at extremely low temperature (T ¼ 300 mK). As shown in Fig. 3.32, the inverse of the resistance (conductance) exhibits strong temperature dependence for the lower electric field in which condition a metallic conduction appears as “on” state of FET. Moreover, the temperature dependence of the logarithm of the conductance, which is proportional to the inverse of the temperature in typical semiconductor with the band gap, significantly deviates from the T-inverse law at the lower temperature region. This suggests that disorder plays an important role in the transport properties of the device. With the presence of disorder, the electronic states around the Fermi energy are localized even in metallic system or zero-gap semiconductor, where the resistance at the lower temperature region is dominated by a hopping conduction process [47]. Here, the temperature dependence of the inverse of the resistance r of the FET device

10-5 1/ρ / Ω-1

10-6 10-7

10-8

10-9 1.0

2.0 1/T / K-1

3.0

Figure 3.32 Inverse of the resistance as a function of the inverse temperature from 300 mK to 100 K. The black and gray curves are the fits to Eq. (3.10) with nearest neighbor hopping (NNH) and variable range hopping (VRH) tems, respectively (see text) [45].

Electrical properties and applications 205 is well described for the whole region from on state to off state with the band gap by the following equation. 1 ¼ r

1 1 1

þ þ E1 E2 X r1 exp r2 exp kB T kB T

(3.10)

   13 E3 0 where X ¼ r3 exp kB T for nearest neighbor hopping (NNH) or r3 exp TT0 for variable range hopping (VRH) [47], both of which are typical for the conductance in metallic     system with strong disorder. r1 exp kEB1T ; r2 exp kEB2T denote the activation-type conductance typical for semiconductors at the higher and the intermediate temperature region having the effective activation energies of E1 and E2, respectively, and kB is the Boltzmann constant. Electrons can hop between localized states in the energy gap even in the off state of FET with opening band gap. Thus, the effective band gap obtained by the resistance measurement becomes smaller and makes it difficult to achieve the off state of the FET at the higher temperature. Indeed, the estimated energy gap is about two orders of magnitude smaller than that expected from the optical measurement shown in Fig. 3.29. The presence of the localized states caused by disorder is also a serious disadvantage to achieve on state with enough lower resistance, which is important for developing lowpower consumption devices. To overcome the problem of the contribution of the electron conduction between localized states caused by disorder, the improvement of the growth process of the insulator layer for top gate electrode is attempted for the fabrication process of the bilayer graphene FET. The origin of disorder in graphene in FET device mainly comes from the insulator layers, which directly face graphene and induce defects on it. In typical fabrication process of a graphene FET with the dual gate structure, graphene is first transferred or directly prepared by micromechanical cleavage on a SiO2/Si substrate, followed by deposition of the insulator layer for the top electrode on graphene. The latter process is much more destructive for graphene, where usually metal oxide is deposited directly on graphene surface involving a thermal oxidation process or by sputtering. On the other hand, the former process causes much less damage, where graphene is fixed on the substrate by weak van der Waals interactions. Xia et al. applied a polymer coating and precise controlled deposition to fabricate an insulator layer for the top electrode to prevent inducing disorder in graphene. Fig. 3.33 shows a schematic view of their dual gate FET device with bilayer graphene and its cross-sectional view. As the insulator layer for the top gate electrode, they coated about 10 nm of an organic polymer layer made from a derivative of polyhydroxystyrene (the polymer NFC 1400-3 PC manufactured by JSR Micro, Inc.), where it interacts with graphene mainly by van der Waals force like that between the

206 Chapter 3 (A)

(B) Top gate ALD HfO2 NFC Bi-layer graphene Silicon oxide Silicon back gate

Figure 3.33 (A) Three-dimensional schematic view of the dual gate bilayer graphene FET. (B) The layer structure of (A). From top to bottom are: top metal gate electrode, HfO2 insulator layer for the top electrode, organic polymer NFC 1400-3CP layer, bilayer graphene, SiO2 as bottom gate oxide layer, and highly doped silicone substrate as the back gate electrode [46].

Drain current / μA

1

Δ(φbarrier) / meV

10-1

0 -20 -40 -60 -2.0

10-2 -2

-1

0

-1.0 0.0 Dave / V/nm

2 3 4 1 Top gate bias Vtg / V

5

6

Figure 3.34 Top gate voltage dependence of the source-drain current of the bilayer graphene FET with dual gate structure for various back gate voltages from 120 to 80 V at room temperature. The inset shows the perpendicular electric field dependence of the Schottky barrier height [46].

substrate and graphene. After that, 10 nm of HfO2 layer is deposited in a very precise manner by atomic layer deposition (ALD) based on the seed layer of the polymer, of which methyl and hydroxyl groups also help to grow high-quality and defect-free HfO2 layer. As shown in Fig. 3.34, the bilayer FET device with the dual gate electrodes fabricated by the polymer coating and ALD deposition of oxide for the insulator layer of the top gate electrode exhibits high on/off ratio even at room temperature. The on/off ratio reaches up to 102 at room temperature, which is important for practical application of graphene as the FET device. However, the temperature dependence of the on/off ratio exhibits significant increases as the temperature decreases as shown in Fig. 3.35. At 20 K, the on/off ratio achieves 2000, which indicates the significant influence of the localized states caused by

On/Off current ratio

Electrical properties and applications 207

103

ln(Ion/Ioff) – qφbarrier/kT

Dave ~ 1.3 0.1(V/nm)

102

101 0.00

0.01

0.02 0.03 1/T / K-1

0.04

0.05

Figure 3.35 Temperature dependence of on/off current ratio at the electric field perpendicular to graphene of 1.3 V nm1 for bilayer graphene fabricated with the polymer coating and ALD oxide deposition technique [46].

defects in graphene. The presence of disorder of graphene significantly suppresses the performance of the switching operation even in this device. To improve the gate electrode fabrication process for the graphene FET, further understanding of the graphene/insulator and the graphene/metal interfaces is necessary. Recently, many trials to for extracting quantitative information to evaluate the interface quality between graphene and electrodes have been performed to understand the origin of localized states in the band gap. Detailed analysis for the results of both quantum capacitance and transport measurements enables estimation of the density of states and the localization length of the localized states in the gap to give an estimate to improve the quality of the device. The improvement of gap states below w1011 eV1cm2 is considered as the first step for a bilayer graphene device to be a promising candidate for future nanoelectronic applications [48].

3.2.2 Spintronics devices To overcome the problems in the current logic devices for the information technology based on the FET devices typically using silicon as the channel material, innovations in processing technology aspects are also proposed as well as material engineering point of view. Existing information processing is performed by controlling the charge degree of freedom of electrons in materials. However, the density of bits describing the information in view of integration of devices has been approaching its intrinsic limit due to the interference among bits caused by Coulomb interactions between charges. In this viewpoint, information processing based on bits using spin degree of freedom of electrons has an advantage for more integration of devices. The difference between charges and spins regarding the integration limit comes from the origin of inter-bit interactions in principle. A Coulomb interaction between charges has a r2 dependence

208 Chapter 3 on the inter-charge distance, which are long-range interactions. To avoid unintentional interferences between bits in a logic circuit, enough larger separation between bits is necessary. On the other hand, the inter-spin interactions are mainly classified into two interactions: the exchange interactions and the magnetic dipolar interactions. As inter-spin distance increases, the exchange interactions exponentially decay, which are quite shortrange interactions. The magnetic dipolar interactions are longer-range interactions. However, this obeys a r3 law for the inter-spin distance, which are shorter-range interactions in comparison with a Coulomb interaction. Thus, an information bit based on the spin degree of freedom has more advantage in view of the integration of the logic circuit. Indeed, the information storage is mainly achieved by a hard disk drive device in the current information technology, where the magnetic polarity of tiny domains in a material is used as a bit to describe and record the information. Moreover, the spin-based bit is one of the most promising candidates for an information bit in actual physical devices for quantum computing (qubit). In quantum computing, superposition and/or entanglement of on and off states are used for calculation in contrast to a typical existing computer using two definite states [49]. Spins in a material are promising as real qubit in actual quantum computing devices, where quantum-mechanical superposition and/or entanglement of wave functions for spin up/down states is used for computer processing. “Spintronics” is an engineering field to extend exiting electronics technology by using spin degree of freedom instead of that for charge [50]. A logic device in the spintronics is a spin transistor, which is FET-type device equipped with ferromagnetic gate electrode and nonmagnetic channel material [51]. In spin transistor, the spin flowing in the channel is regulated by the electric potential and the magnetization of gate electrode. To achieve the long-range spin coherence, extremely lower spin scattering probability is required for a channel material of spin transistor. Intrinsic spin scattering by the spin-orbital (SO) interactions are an unavoidable problem even in ideal channel materials without defects and impurities as extrinsic scattering sources. The Hamiltonian for the SO interaction is described by consideration of the Dirac equation under a spherical potential as following, 1 1 vV l$s ¼ ll$s (3.11) 2m2 r vr where m, r, V, s, l, and l are electron mass, orbital radius, potential for electron, spin momentum operator, angular momentum operator, and the SO coupling constant, respectively. Assuming hydrogen-like atomic orbitals, l is calculated as HSO ¼

l¼

m0 Ze2 8pm2 r 3

(3.12)

Electrical properties and applications 209 where m0, Z, and e are the permeability of vacuum, the atomic number, and the elemental charge. The expectation value of r3 in Eq. (3.13) is obtained for hydrogen-like orbitals Jn,l,m as

 Z 1 Z3  3
 r J dsf ¼ J (3.13) n;l;m n;l;m 1 r3 3 n l l þ ðl þ 1Þ 2 Thus, the strength of SO interaction is proportional to the fourth power of the atomic number (fZ 4). This means that the spin scattering quite rapidly increases as increasing in atomic number of elements consisting of a channel material intrinsically. The scattering of spins is significant even in silicon having the atomic number of 14, and it reduces a lifetime of spin state describing an information bit. In this viewpoint, carbon with the atomic number of six is almost the lower limit for light elements to compose semiconductor compounds. So, graphene is also promising for a channel material of spin transistor devices due to its high mobility of spin flowing (spin current). Early studies on graphene toward spintronic application are mostly devoted to increase the efficiency of the spin injection to graphene from the ferromagnetic electrodes and to extend the relaxation length of spin transport in graphene. Fundamentally, an injection of spin-polarized electrons into less-conductive materials like semiconductors and graphene from ferromagnetic metal electrodes is quite inefficient in the diffusive transport regime because of the mismatch in the conductivity [52]. Electrons injected into the interface between ferromagnetic metal and less-conductive materials prefer to flow back into the metal electrode side. In this process, the spin orientation of electrons is significantly disordered, and spin polarization is lost. In order to prevent the back diffusion of spinpolarized electrons to ferromagnetic electrodes, a very thin insulating barrier layerelike metal oxide is inserted between electrodes and graphene to give a high-contact resistance. Tombros et al. demonstrated spin injection to monolayer graphene 3 years after the insolation of that by the Manchester group [53]. Magnetic cobalt electrodes (Co) with very thin Al2O3 barrier layer are attached on monolayer graphene on SiO2/Si substrate as shown Fig. 3.36, where the widths of electrodes vary so as to have different magnetic fields to switch the magnetization direction from each other. Here, the nonlocal fourelectrode method is applied to avoid problems like anisotropic magnetic resistance effect often observed in the local two-electrode method [52,54,55]. The principle of the nonlocal method is based on the fact that the charge current flows in anisotropic manner due to the applied electric field, while spin current diffuses isotopically into the material [56]. A current I is injected by a current source instrument from electrode 3 through the Al2O3 barrier into graphene and is extracted at the electrode 4, where the electrodes inject or extract only electrons having spin parallel to the their magnetization directions.

210 Chapter 3

Figure 3.36 (A) SEM image of a four-terminal monolayer spin valve device. (B) The nonlocal spin valve geometry. Schematic plots of spin injection and spin diffusion for electrodes having parallel magnetizations (C) and antiparallel magnetizations (D) [53].

The voltage difference is measured between electrodes 2 and 1, where electrodes detect the electric potential for only electrons having spin parallel to their magnetization directions. Here, the nonlocal resistance is defined as Rnonlocal ¼

Vþ  V I

(3.14)

In the case of parallel alignment of the magnetizations for all electrodes, the electrode 3 injects up spin electrons and electrode 4 extracts up spin electrons. Because bulk ideal graphene is nonmagnetic material (no spin polarization), the sum of up and down spin is zero in any position inside graphene. Thus, the densities of electrons having up and down spins inside graphene are schematically described as shown in Fig. 3.36B. Here, the Rnonlocal is positive in this case because electrodes 1 and 2 only detect the potential for electrons having up spin. Similarly, the densities of electrons having up and down spins are described as shown Fig. 3.36C in the case of antiparallel configuration of the magnetizations of electrodes, where Rnonlocal becomes negative. The devices as shown in Fig. 3.36 are called spin valves, which are typically used to perform spin transport experiments. Indeed, the nonlocal resistance changes depending on the configuration of the directions of magnetization for electrodes according to the magnetic field as shown in Fig. 3.37. The systematic change in the nonlocal resistance Rnonlocal proves successful injection of the spin polarized current into graphene. However, the spin injection efficiency is estimated as 10% in this result, according to the theoretical model described as the following formula [57]: 
P2 lsf L Rn ¼ exp  lsf 2Ws

(3.15)

Electrical properties and applications 211 18

Rnon-local / Ω

12 6 0 -6

-12 -18 -300

-200

-100

0 By / mT

100

200

300

Figure 3.37 Magnetic field dependence of the nonlocal spin valve signal Rnonlocal for graphene at 4.2 K. The sweep directions of the magnetic field are indicated by horizontal arrows. The magnetic configurations of the electrodes are illustrated for both sweep directions. The direction of the magnetization of each electrode is flipped at different magnitudes of the magnetic field depending on the width of each electrode, where the wider electrode remains in the direction of its magnetization until applying the larger magnetic field with opposite direction [53].

where P, lsf, W, s, and L are the spin polarization of the contact, the spin relaxation length, the conductivity, the width of graphene, and the distance between central electrodes, respectively. The gate voltage dependence of the nonlocal resistance Rnonlocal for parallel and antiparallel configurations for the magnetizations of electrodes is shown in Fig. 3.38A. (A) 8

Parallel

Rnon-local / Ω

4 0 -4

Spin signal

-8

Antiparallel

-12 Rgraphene / kΩ

(B)

16 12 8 4 0 -60

Dirac neutrality point

-40

-20

0 Vg / V

20

40

60

Figure 3.38 Gate voltage dependence of the nonlocal resistance for parallel and antiparallel configurations for the magnetizations of electrodes (A) and the resistivity (B) [53].

212 Chapter 3 In agreement with the principle of the nonlocal technique, Rnonlocal shows that the parallel configuration always yields a positive value and the antiparallel a negative one. In comparison with the nonlocal resistance Rnonlocal and conductivity, the spin current signal denoted by Rnonlocal quite weakly depends on the carrier density. There is only slight reduction in Rnonlocal around the Dirac point. The spin injection efficiency is lower also in the following reports [58,59] in the earlier studies. The lower spin injection efficiency is attributed to either the conductance mismatch problem originating from high conductance of the ferromagnetic metal electrodes and lower conductance of graphene [52], or other possible contact-related effects. To overcome the problem of the lower spin injection efficiency between ferromagnetic electrodes and graphene, there have been many attempts to improve the interface between graphene and ferromagnetic electrodes. One of the important factors to improve the quality of the interface between graphene and electrodes is the presence of pinholes in the barrier between electrode and graphene. The importance of the presence of pinholes in the barrier has been reported for the spin transport in magnetoresistance measurement for a magnetic tunnel junction device composed of a stack of (Al capping layer)/Fe/(Al2O3 barrier)/(superconducting Nb/Al bilayer) [60]. Popinciuc et al. fabricated two types of graphene spin valve devices with Al2O3 barrier layer deposited only underneath of the ferromagnetic electrode (type II) and whole part of graphene (type I) [61]. The type I devices shows higher contact resistances, and less gate voltage dependence in contrast to the lower contact resistance depending on the carrier density in graphene for type II devices as shown in Fig. 3.9AeF. It is suggested that a part of graphene becomes insulator, resulting in the higher contact resistance insensitive to the

15 10

10

0 30 VG / V

2 0.2 0.1 A / μm2

4

-20

0 20 VG / V

15 10

6

4

5 0

(G)

1

3 2 A / μm2

Device IIA (L=2 μm)

6

5 0

60

(F) Rc / kΩ

Rc -30

(D) Rc / kΩ

15

1

40

Device IB (L=1 μm)

5 0 -40 -20 0 20 IDC / μA

40

2

(H) Rc / kΩ

0 20 VG / V

(E)

Rc / kΩ

(C)

Device IIA (L=2 μm)

/k

2

2

1 -20

Rc / kΩ

(B)

Ω

Device IB (L=1 μm)

Rsq / kΩ

Rsq / kΩ

(A)

40

-30

0 30 VG / V

60

6 4 2 -40 -20 0 20 IDC / μA

40

Figure 3.39 Differential resistances of graphene (A) and (B) and contacts (CeH) for a type I (left column) and a type II (right column) device (see text), respectively [61].

Electrical properties and applications 213

Figure 3.40 AFM images of graphene flakes after aluminum deposition for fabrication of an insulator barrier layer, (A) height image and (B) phase image [61].

whole carrier density of graphene. Moreover, the contact resistance in the type I device is much enhanced at the zero-bias voltage as shown in Fig. 3.39G. This suggests the contribution of tunneling or the presence of pinholes in the barrier layer. AFM images shown in Fig. 3.40 strongly suggest the contribution of pinholes rather than tunneling. After the aluminum deposition, the oxide layer on graphene surface is granular in nature. The root-mean-square values 0.5e0.7 nm for the aluminum oxide layer on graphene are obtained by roughness analysis for several samples. The ferromagnetic electrodes in earlier studies might be accompanied with the imperfect barrier layer with pinholes. By using pinholes in barrier layer, whether intentionally or unintentionally introduced, the improvement in the spin injection efficiency by “mimic” tunneling mechanism though pinholes is achieved as a nonlocal resistance Rnonlocal of 10 U, which is proportional to the spin injection efficiency as shown in Eq. (3.15), corresponding to the efficiency up to 18%. However, using pinholes is uncontrollable and shows poor reproducibility, although relatively high spin injection efficiency was also achieved in several pinhole contact devices. The higher spin injection efficiencies were achieved by using TiO2-seeded MgO films as the tunnel barrier between ferromagnetic electrodes and graphene. By using TiO2-seed layer deposited on graphene in advance, as shown in Fig. 3.41, the uniformity of MgO overlayers is significantly improved and pinhole-less insulator thin film is fabricated, where 0.8 nm of MgO barrier layer was deposited in this study [62]. Fig. 3.42 indicates the magnetic field dependence of the nonlocal resistance for spin valve device made of monolayer graphene at room temperature. In both the magnetic field

214 Chapter 3

Figure 3.41 Schematic diagram of the fabrication of TiO2-seeded MgO films as the tunneling barrier between monolayer graphene and the ferromagnetic Co electrode. (A) Tilted angle evaporation of metals (B) Schematic drawing of the Co/MgO/TiO2/graphene structure for injection spin-polarized current [62].

RNL / Ω

T = 300 K

H / mT

Figure 3.42 Magnetic field dependence of the nonlocal resistance for spin valve device made of monolayer graphene at room temperature. The inset shows the electric circuit diagram for nonlocal spin transport measurement on this device [62].

sweeping down (black curve) and up (gray curve) side, the nonlocal resistance Rnonlocal of 130 U is obtained as indicated by the arrow. This means the spin injection efficiency is over 35%, which is larger than that observed in typical studies using a barrier with pinholes. The highest efficiency for spin injection into graphene is achieved by using fluorographene as the insulator barrier between permalloy (Ni80Fe20) ferromagnetic electrodes and graphene [63]. After the fabrication of typical FET device structure with Ti/Au electrodes as the source and drain for CVD graphene transferred on SiO2/Si substrate, the second

Electrical properties and applications 215

Figure 3.43 Optical microscope image of the graphene spin valve device with fluorographene layer as a barrier layer for spin injection [63].

layer of CVD graphene is transferred on the device, where the second layer completely overlaps the underlying graphene. The device is exposed to XeF2 gas for fluorination of the second layer of graphene. After the deposition of permalloy as ferromagnetic electrodes for spin injection, the whole device is exposed to XeF2 gas again, ensuring that no parallel conductive path exists between electrodes and graphene (Fig. 3.43).

Conductance / μS

Conductance / mS

Fig. 3.44 shows the conductances for underlying graphene channel and upper fluorographene as barrier layer for tunneling injection of spin polarized current, where by

Gate voltage / V

Figure 3.44 Gate voltage dependence of the conductances for graphene and fluorographene layers in graphene spin valve device with tunneling barrier between electrodes and graphene fabricated by using fluorination method [63].

ΔRNL / Ω

Calculated polarization efficiency / %

216 Chapter 3

Bias current / mA

Figure 3.45 Bias current dependence of the nonlocal resistance and corresponding spin injection efficiency [63].

using another sets of electrodes prepared for the second layer, the resistance of the second layer graphene is found to be 50 GU after fluorination. In addition, the fluorographene layer shows no gate voltage dependence without the charge neutral point in contrast to underlying graphene layer exhibiting the minimum conductance around the gate voltage of 70 V, which is attributed to influences of extrinsic hole doping by residue in the lithography process and charging in the SiO2 substrate. Insulator nature of the fluorographene layer is also supported by the fact that the top fluorographene layer is transparent in optical microscope image shown in Fig. 3.43. The spin injection efficiency is calculated form the results by using the observed nonlocal resistance as shown in Fig. 3.45. The spin injection efficiency monotonically increases with decreasing in the bias current and finally achieves up to 60% near no bias current. The nonlinear increase of the spin injection efficiency as the bias current approaches zero is consistent with other spin valve devices reported, although it is not entirely understood [64]. Most recently, a spin injection efficiency effectively achieving up to 100% is reported for a graphene spin valve device using two-dimensional boron nitride layers sandwiching graphene as the tunneling barrier, where it is also important factors to protect graphene from the supporting substrate for enhancing the total device performance [65]. The other challenge for applying graphene to spintronics devices is increasing spin lifetime. In an ideal case, graphene has intrinsically less spin scattering probability due to its small atomic number in comparison with typical semiconductors. The SO interaction is significant even in silicon, where the electronic bands around the Fermi energy are split into several bands because of SO interactions, resulting in the serious influence of spin degree of freedom in any electronic processes inside silicon. Theoretically, the spin lifetimes of around a microsecond are expected in graphene [66]. However it is largely reduced down to values ranging from tens of picoseconds to a few nanoseconds in

Electrical properties and applications 217 experimental reports [67e71]. Toward practical applications, spin lifetime of nanoseconds or longer are a mandatory requirement. The large discrepancy between the theoretical prediction and experimental results has been the most outstanding puzzle in graphene spintronics engineering. Currently, the source of spin relaxation originates from extrinsic factors such as impurities, defects, adsorbates, or static ripples. The lack of a fully comprehensive picture of all possible sources of spin relaxation in graphene prevents microscopic understanding of the reason for the short spin lifetime despite numerous theoretical studies [72]. For example, spin lifetimes of a few nanoseconds have been observed experimentally for graphene spin valve devices with high quality tunneling barrier for spin injections [73,74]. On the other hand, the lifetime is less than 1 ns in graphene spin devices with pinhole contacts between electrodes and graphene [61]. This strongly suggests significant spin relaxation induced by contacts between electrodes and graphene. Here, again the improvement in the quality of the interface between electrodes and graphene is necessary in order to extend spin lifetime in graphene. Almost theoretical models in the field of spintronics for metals and semiconductors are based on spin relaxation caused by an electron momentum scattering with the SO interaction. However, direct applications of these models to the spin transport in graphene with the smaller SO coupling have failed to explain the experimentally observed results, where two orders of discrepancy is common [66]. So, developing new mechanisms unique to graphene would be required to solve this problem. Interestingly, there are several reports indicating an order of longer spin lifetimes in bilayer in comparison with that in monolayer graphene [73,74]. To evaluate the spin lifetime in graphene, Hanle-type spin precession measurements by using spin valve devices are usually performed in the nonlocal method as shown in Fig. 3.46A. To measure a Hanle type spin precession, an external magnetic field B is applied perpendicular to

(A)

(B)

B⊥

10 RNL / Ω

S

0

-10 -100

100 0 B ⊥ / mT

Figure 3.46 (A) Schematic view of a Hanle-type spin precession measurements, where the magnetic field B is applied perpendicular to graphene and the ferromagnetic electrodes are magnetized in the parallel or antiparallel manner to each other. The right- and left-side ferromagnetic electrodes are spin injector and detector, respectively. (B) Typical Hanle curves in graphene measured by spin valve device [66].

218 Chapter 3 graphene. The ferromagnetic electrodes are magnetized in the parallel or antiparallel manner to each other as described by black arrows. The right- and left-side ferromagnetic electrodes are spin injector and detector, respectively. When spins parallel to the magnetization of the left electrode (injector) are injected into graphene, they flow toward the right side according to the bias voltage between electrodes with the precession motion as shown in the sketch due to the magnetic torque by the external magnetic field B. The rotation speed of the precession motion is fixed by the magnitude of the external field. Thus, the direction of spins at the left electrode (detector) depends on the distance between electrodes and the velocity of spins in graphene. The left electrode detects only the components of spins parallel to the magnetization of the electrode. Thus, the ratio of spin current between electrodes to bias voltage (¼Rnonlocal) oscillates according to the magnitude of external field B (the speed of the precession). Consequently, the nonlocal resistance Rnonlocal behaves as shown in Fig. 3.46B in the case of parallel (black curve) and antiparallel (gray curve) configuration of two ferromagnetic electrodes. If the contact resistance is much larger than the spin resistance of graphene, the nonlocal Hanle curve is described as


Z N 1 L2 t pffiffiffiffiffiffiffiffiffiffiffi exp  Rnonlocal f  dt (3.16) cosðuL tÞexp  ss 4Dt 4pDt 0 where L, D, ss, and uL are the distance between electrodes, the diffusion constant, the spin lifetime, and uL ¼ gmBB/h– , in which mB is the Bohr magneton and h– is the Planck constant [57,75]. Here it should be noted that the þ() signs are for the parallel (antiparallel) configuration of the magnetization of electrodes. By fitting the experimentally observed Hanle curve with Eq. (3.16), the spin lifetime ss and the diffusion constant D are obtained. The results of spin transport measurements by a spin valve device using monolayer graphene are summarized in Fig. 3.47 [74]. At low temperature, the spin lifetime corresponds to the change in the density of carrier by gate voltage, where the lower carrier density, the lower spin lifetime and diffusion constant. The spin lifetimes obtained by fitting of the observed Hanle curve are in the range of several 100 psec. This is one of the best spin lifetimes for monolayer graphene with high-quality tunneling barrier between the ferromagnetic electrodes and graphene. However, bilayer graphene exhibits quite different behavior from that of monolayer graphene as shown in Fig. 3.48. The temperature dependence of the spin lifetime for monolayer graphene is weak, and it becomes twice at the lowest temperature at maximum, where the spin lifetime at the Dirac point (black curve) shows almost no temperature dependence. On the other hand, the spin lifetime in bilayer graphene exhibits strong temperature dependence, where the temperature dependence is most significant at the Dirac point. Most remarkably, almost an order of larger spin lifetime is obtained for

Electrical properties and applications 219 (B)

H / mT

τs / ps

D / 10 -2 m 2 /s

(D)

D / 10 -2 m 2/s

H / mT

τ s / ps

(C)

R NL

R NL / Ω

/Ω

(A)

V g -V CNP / V

V g -V CNP / V

Figure 3.47 (A) Nonlocal resistance as a function of the magnetic field parallel to graphene plane for monolayer graphene. (B) The Hanle curve for the same device. The gate voltage dependence of the spin lifetime and the diffusion constant at 300 K (C) and 4 K (B) [74].

(B)

τs / ps

τs / ps

(A)

T/K

T/K

Figure 3.48 Temperature dependence of the spin lifetime with different gate voltages for spin valve devices using monolayer (A) and bilayer (D) graphene [74].

bilayer graphene in comparison with monolayer graphene. At the lowest temperature, the spin lifetime exceeds 6 ns in bilayer graphene. The significant difference in the spin lifetime in monolayer graphene and bilayer graphene would be related to the difference in the dominant spin relaxation mechanism. Indeed, the

220 Chapter 3

Vg / V

1

4000

0

2

-60

-30 0 30 Vg-VCNP / V

60

D / 10-2m2/s

T = 20 K

8000

0

Figure 3.49 Gate voltage dependence of the spin lifetime and the diffusion constant for bilayer graphene at 20 K [74].

gate voltage dependence of the spin lifetime in bilayer graphene is completely opposite to that in monolayer graphene as shown in Fig. 3.49. In bilayer graphene, the spin lifetime increases with decreasing carrier density as the Fermi energy approaches to the Dirac point. Interestingly, the behaviors of the diffusion constants are also different between bilayer and monolayer graphene. In monolayer graphene, the diffusion constant almost coincides with the behavior of the spin lifetime. However, the diffusion constant behaves irrespective to the spin lifetime in bilayer. Comparing the inverse of the spin lifetime 1/ss corresponding to the spin scattering probability to the diffusion constant, 1/ss monotonically increases with the diffusion constant increasing in bilayer graphene. This is reminiscent of a spin relaxation by DyakonovePerel (DP) model. In DP model, the more electron momentum scatters, the less spin relaxes in view of the motional narrowing concept. On the other hand, 1/ss decreases as the diffusion constant increases. This is attributed to the reasonably understandable ElliotteYafet spin relaxation mechanism, where the electron spin relaxes at some probability by scattering events of electron momentum through the SO interactions as schematically shown in Fig. 3.50. Also for spintronics devices, synthesis route of graphene becomes a problem when we consider practical applications in actual industry. Graphene samples obtained by the

Figure 3.50 Possible spin relaxation mechanism in graphene. The resonance scattering is proposed as the model specific to graphene [66].

Electrical properties and applications 221 (B)

(C) Resistance / MΩ

138

L

200 nm

1.2 0.8

137 0.4 136

500 nm

Device A: L = 2 μm

0

ΔR/R = MR / %

(A)

-1000 0 1000 Magnetic field / Oe

Figure 3.51 (A) and (B) Optical images of a large-scale spintronics device fabricated on epitaxial graphene on SiC wafer substrate. (C) The response of the spin valve operation using this device [76].

micromechanical cleavage method exhibit the highest quality in view of crystallinity. Especially in spin transport measurement, the highest quality sample has been used due to the extremely low spin lifetime in spite of the ideal values predicted by theories. However, the development of a spin valve device using graphene sample more compatible to the industrial application proceeds gradually, where epitaxial graphene as a large-area and comparatively high-quality graphene sample is widely used [76]. Fig. 3.51A and B shows the optical images as a large-scale spintronics device fabricated on epitaxial graphene on SiC wafer substrate, where several devices are mounted on the single wafer tip similar to the existing electronic devices on Si wafer tip such as memory and CPU. Fig. 3.51C shows the spin value response obtained by using this device. By analyzing the obtained data, enough high spin transport efficiency and the spin lifetime above nanosecond order are reported.

3.2.3 Transparent electrode Toward growth of the smart society based on information technology, the improvement in the man-machine interface is also an important problem as well as data processing and storage. Recently, input/output through touch screen devices has become quite a common manemachine interface. The key technology of the touch screen devices is a transparent electrode, which enables simultaneous detecting and displaying information on IT instruments. Atomically thin thickness of graphene has a large advantage to achieving higher transparency. Indeed, the first report on the isolation of monolayer graphene by a Manchester group is based on the sample obtained by micromechanical cleavage, where the number of graphene layers is evaluated by its degree of transparency [77]. Thus, graphene having both of high conductivity and transparency is quite suitable for the application toward transparent electrodes, which now depend on rare metals. Compared with existing transparent conductive electrodes like indium tin oxide (ITO), graphene has high mechanical strength, flexibility, and chemical stability [78e81]. However, a large-area graphene sample is necessary to evaluate simultaneously its conductivity and

222 Chapter 3 transparency. So, the application of graphene to transparent electrodes usually involves the challenge of preparing the large-scale graphene film. Watcharotone et al. reported their attempt to use graphene as transparent electrode material 3 years after the first isolation of monolayer graphene by micromechanical cleavage [78]. In that study, composites of silica and graphene oxide (GO) or reduced graphene oxide (rGO) were fabricated as transparent conductive film as shown in Fig. 3.52. The silica/rGO composite film is prepared by spin-coating of sols obtained by mixing GO synthesized by Hummers’ method [82] and dispersed in water and tetramethyl orthosilicate [83]. SEM images indicate a homogeneous morphology with slightly crumpled graphene sheets overlapping each other, which can act as a conduction path of carriers between silica nanoparticles.

Figure 3.52 SEM images of silica-GO composite film containing 6.6 wt% of GO as additive. (A) low magnification after high-temperature curing. High magnifications before (B) and after (C) curing [78].

Electrical properties and applications 223 0.99 0.98

10-1

0.97 10-2 0.96

Transmittance

Conductivity / S/cm

1

10-3 0.95 4 0 2 6 8 10 12 Graphene oxide concentration / wt%

Figure 3.53 GO concentration dependence of the electrical conductivity (circles) and the transmittance at 650 nm (squares) for silica-GO composite film [78].

As easily expected, the conductivity of the silica-GO composite film at room temperature increases with the increase in GO concentration in the film, and the transmittance exhibits an opposite trend for GO concentration as shown in Fig. 3.53. So, the transparencyconductivity trade-off is seen in the graphene-based transparent conductive electrodes as similar to other electrodes actually used. However, the increment in the conductivity shows an exponential dependence on the GO concentration in contrast to a linear dependence observed for the transmittance. The wavelength dependence of the transmittance (dispersion of the transmittance) of the pure silica film is essentially highly transparent over the 380e1000 nm wavelength range as shown in Fig. 3.54A. However, the transmittance becomes reduced as GO concentration increases. Especially, the reduction of the transmittance in the lower wavelength range is noticeable for the composite with the higher content of GO. The transmittance is also largely influenced by posttreatments. Fig. 3.54A shows the GO concentration dependence (B) Transmittance

Transmittance

(A)

Wavelength / nm

Graphene oxide concentration / wt%

Figure 3.54 (A) Wavelength dependence of the transmittance of the composite film after curing, containing different concentrations of rGO. (B) The GO concentration dependence of the transmittance at 650 nm for the composite film untreated and treated by various processes [78].

224 Chapter 3 of the transmittance at 650 nm for the composite film untreated and treated by various processes. For the composite film containing the large amount of GO or rGO, the transmittance is significantly reduced by posttreatments, which are usually performed to improve the conductivity of the films. Although the reduction of the transmittance is noticeable at the lower wavelength and the high GO/rGO loaded films, an acceptable transmittance is observed for these films. The highest conductivity of 0.45 S/cm is obtained for the highest GO content composite film (11 wt%), which is comparable to that reported for MWCNT/silica composite film (0.57 S/cm). However, the conductivity of ITO film in the literature is w1  104 S/cm [84,85], which is much higher than that of the GO/ silica composite film. The improvement in the conductivity of graphene-based transparent electrodes has been achieved by the synthesis of the large-scale, high-quality graphene film. By usual thermal CVD growth of graphene on Ni film with methane gas and subsequent etching of Nickell using FeCl3 aqueous, a high-quality large-area graphene film having a lateral size up to cm order with an electron mobility more than 1000 cm2 V1 s1 is obtained on arbitrary transparent and insulating substrates. Fig. 3.55A shows the wavelength dependence of the transmittance of the graphene film transferred on quartz substrate [81]. The transmittance is about 80% in the visible range, which is important for the application toward display devices. The transmittance can be increased up to 93% when the thickness of the graphene film is reduced by decreasing in the growth time and thickness of Ni film. However, thinner film is defective and smaller in lateral size. As shown in Fig. 3.55B, the graphene film exhibits the abnormal integer quantum Hall effect, which requires extremely high-quality sample to prevent any electron scattering as (A)

(B) Magnetoresistance / kΩ

Transmittance / %

90

80

70

60 400

600 800 1000 1200 Wavelength / nm

10 5 0 -5 -10 -15 -60 -40 -20

0 20 Vg / V

40

60

Figure 3.55 (A) Wavelength dependence of the transmittance of the graphene film prepared by CVD on Ni substrate and transferred on to SiO2/Si substrate with different conditions of UV treatment. The inset shows UV-irradiation time dependence of the transmittance Tr and sheet resistance Rs. (B) Gate voltage dependence of the Hall resistance and longitudinal resistance of the graphene film [81].

Electrical properties and applications 225 possible for observation of noticeable oscillation. This guarantees much better electron transport in the graphene film and is attributed to the small resistance of the film w280U per square, which is almost 30 times larger than the minimum resistance observed in an assemble graphene film [86]. Interestingly, UV irradiation treatment under atmospheric condition improves the transmittance as clearly seen in Fig. 3.55A and its insets, where the transmittance proportionally increases with increases in the irradiation time. Although the sheet resistance simultaneously increases with UV irradiation, its reduction is not significant and the resistance remains at the practical use level. The UV-irradiationinduced change is attributed to the reduction in the film thickness by etching process.

Anisotropy R y/ R x

Resistance / kΩ

One of the most interesting advantages of graphene-based transparent electrode is its flexibility. A graphene film has excellent mechanical properties for making flexible and stretchable electrodes [87]. Fig. 3.56 shows the change in the resistance by bending of graphene on flexible polydimethylsiloxane(PDMS)/polyethylene terephthalate (PET) substrate with a thickness of 0.3 mm. The resistance of the graphene film almost is maintained up to the bending radius of 2.3 mm, roughly corresponding to a tensile strain of 6.5%. Further bending up to ca. tensile strain of 18.7% causes slight increment in the resistance, but is completely recovered after releasing the bending. Interestingly, the resistance of the film exhibits strong anisotropy between resistances along the bending direction Ry and the perpendicular direction Rx as defined in the inset of Fig. 3.56. The Ry/ Rx reaches up to 102 at maximum. The advantage of graphene against commonly used ITO transparent electrode in view of flexibility is more clearly shown in Fig. 3.57 [88],

Curvature κ / mm -1

Bending radius / nm

Figure 3.56 Change in the resistance by bending of graphene on flexible PDMS/PET substrate with a thickness of 0.3 nm. The left-side inset shows the anisotropy between resistances along the bending direction Ry and the perpendicular direction Rx as defined in the right-side inset [81].

226 Chapter 3 200

150

ITO R2R Graphene

Tensile

ΔR/R0

3 2

Compressive

1

ΔR/R0

0

100

0

4 2 Strain / %

50

0 1

2

3 Strain / %

4

5

6

Figure 3.57 Stain dependence of the change in the resistance for ITO and graphene-based transparent electrode. The inset shows the tensile and compressive strain dependence of the change in the resistance for the graphene film [88].

where the graphene film is prepared by the role-to-role process compatible to the industrial application and is completely ready for large-scale production. When a strain is applied to ITO electrode by bending, its resistance rapidly increases, and becomes 200 times larger value at strain of 4%, which is out of the practical use [89]. On the other hand, the graphene transparent electrode shows almost no change in the resistance even up to a strain of 6%. Here, the graphene film is robust for both tensile and compressive strains as shown in the inset.

3.3 Applications to social fields 3.3.1 Sensor devices Toward smart society based on IoT, the use of sensors has explosively increased. Simultaneously the demand for devices cheaper, better in performance, that consume less power, and so on has rapidly emerged. To develop such new classes of sensor devices, innovative sensor materials are required as well as new concepts of sensing. Generally, sensors are classified into two groups: contact and noncontact sensors. On the former sensors, substances physically contact to the surface and cause stimulation, inducing a response signal, whereas the latter do not need contact directly with the environment. Chemical, pressure, and mechanical force sensors are representatives of contact sensors, whereas noncontact sensors include usually magnetic, electric field, radiation, and strain sensors. Two-dimensional materials like graphene are intrinsically quite sensitive to the surrounding environment because almost atoms are possessed at the surface. Thus, graphene is promising as a new class of materials for sensor devices, empirically for contact-type sensing devices.

Electrical properties and applications 227 (A)

(B)

Figure 3.58 (A) Change in the charge density obtained from the shift of the charge neutrality point upon adsorption of different concentration of NO2 gas. Upper inset is top view of the graphene FET device, and the lower inset shows the gate voltage dependence of the longitudinal and Hall resistance at 1 T. (B) Time evolution of the ratio of the change in the resistance for 1 ppm of NH3, CO, H2O, NO2-adsorbed monolayer graphene [90].

Soon after the isolation of graphene by micromechanical cleavage, the application toward gas sensing was attempted by the Manchester group due to its two-dimensional nature [90] wherein a graphene FET device itself can act as chemical sensor, where graphene channel is directly exposed the environment. Fig. 3.58A indicates that the change in the carrier density well corresponds to the NO gas concentration, which is mainly caused by charge transfer from graphene to NO molecules having strong electron negativity. This supports the detection of the NO concentration by monitoring the change in the carrier density of graphene by using conventional FET structure. Fig. 3.58B shows the time evolution of the ratio of the change in the resistance for 1 ppm of NH3, CO, H2O, and NO2-adsorbed monolayer graphene. Just after starting adsorption (region II), the resistivity rapidly changes for NO2 and NH3, while the changes for CO, H2O are moderate. Although the response speed depends on the gas species, the changes gradually become saturated after long time exposing (region II). This is attributed to the redistribution process due to inhomogeneous adsorption inside measurement instruments. In the case of evacuation after achievement of near-equilibrium state, the resistivity is recovered to the initial values, although it takes quite a longer time in comparison with the change upon adsorption for all chemical species. The difference in kinetics between adsorption and desorption suggests the positive adsorption heat and the presence of an energy barrier for adsorption on graphene. Taking the signal noise Dr/r of w104 in this measurement into consideration, the detection limit of the order of 1 ppb is estimated for the graphene chemical gas sensor. The Hall resistivity rxy, which theoretically diverges at the Dirac point, is more sensitive to the charge transfer with adsorbates in graphene. In extremely diluted condition, the Hall

228 Chapter 3 (A) 40

(B) Number of steps

Changes in ρxy / Ω

600 30 20 10 0

0

200

400 Time / s

600

400 200 0

-4

-2

0

δR / Ω

2

4

Figure 3.59 (A) Changes in hall resistance during adsorption and desorption of diluted NO2 gas. The curve with smaller change corresponds to exposing to inert He gas, which presumably shows no adsorption at room temperature with graphene. (B) Population of events versus size of changes in Hall resistance value [90].

resistivity rxy changes in quantized manner upon NO adsorption, where rxy decreases (increases) with stepwise changes for adsorption (desorption) process, as shown in Fig. 3.59A. Fig. 3.59B shows the population distribution histogram of the stepwise change in rxy, where two peaks appear in both increasing and decreasing sides. This suggests the presence of the minimum amount of the change in rxy. In fact, the changes in rxy for the peaks in the histogram correspond to the charge transfer of elementary charge into/from graphene. Thus, the observed steps in Fig. 3.59 correspond to adsorption/desorption of a single molecule of NO. So, graphene-based chemical sensor achieves the detection of the single molecular adsorption. Chemical sensor devices have also been fabricated by using graphene samples conventionally available, which is suitable for more practical application, although the detection of single molecular adsorption is achieved by using graphene sample prepared by micromechanical cleavage. Fig. 3.60 shows the time evolution of the source-drain currents upon NO2 and NH3 adsorption for typical FET devices using graphene prepared by bubbler-assisted CVD process with a TEB/hexane solution on Cu foil, where the boron content in the resultant graphene can be tuned by the ration of TEB in hexane [91]. By using nondoped graphene PG, NO2 detection is quite difficult in the range less than 8 ppb, where a signal-to-noise ratio (S/N) is obtained as 9.4. On the other hand, the chemical sensor can detect even 1 ppb of NO2 gas. The difference in the sensitivity is attributed to the higher S/N in boron-doped graphene (BG) device (N/S ¼ 31.5). Similar results were found for NH3 detection. Chemical sensors using BG show sensitivities of 90 and 60 ppb for NO2 and NH3, respectively. These are 27 and 105 times larger sensitivities than those for nondoped graphene for NO2 and NH3, respectively.

Electrical properties and applications 229 (B)

(D)

Time / s

Ids / μA

Time / s

Ids / μA

(C)

Ids / μA

Ids / μA

(A)

Time / s

Time / s

Figure 3.60 Time evolution of the source-drain currents upon NO2 and NH3 adsorption for typical FET devices using boron-doped (BG) and nondoped (PG) graphene prepared by CVD process with a triethylborane (TEB)/hexane solution [91].

Higher sensitivity of chemical sensors using BG is explained by stronger interactions between molecule and graphene. First-principles calculations [92,93] predicted that NO2 and NH3 strongly adsorb on B-doped site in graphene in comparison with nondoped graphene, resulting in the dramatically increased sensitivities to NO2 molecules. This is explained by the difference in the charge distribution around B sites as shown in Fig. 3.61. Thus, further tuning in the materials side would still be important to achieve the higher performance of graphene chemical sensors. Indeed, the performances of gas sensing for NO2, NH3, and other toxic chemicals largely vary with devices, each of where different material processes were carried out as collected in some of the published data in Table 3.1. Development of chemical sensors has been also attempted by using rGO, which is more suitable for large-scale production and cost friendly. Robinson et al. prepared rGO film by

230 Chapter 3

Figure 3.61 Charge density mapping with respect to the isolated atoms for the boron-doped graphene with (A) NO2 and (B) NH3 molecules, respectively [91]. Table 3.1: Typical performance of graphene-based gas sensors [91,94]. Target gas NO2

NH3

HCN DNT CEE DMMP

Preparation process of sensing graphene CVD-synthesized graphene CVD-synthesized graphene-like film MPCVD graphene CVD-synthesized graphene B-doped CVD-synthesized graphene Epitaxially grown graphene from SiC Mechanically cleaved graphene rGO Sulfonated rGO CNT/rGO composite rGO on PET substrate CVD-synthesized graphene B-doped CVD-synthesized graphene CVD-synthesized graphene Mechanically cleaved graphene rGO rGO/polyaniline composite Pyrrole-reduced rGO rGO reduced by p-phenylenediamine

Minimum concentration tested 100 ppb 65 ppm 100 ppm 8 ppb 1 ppb 2.5 ppm 1 ppm 2 ppm 5 ppm 0.5 ppm 500 ppb 20 ppm 1 ppm 500 ppb 1 ppm 20 ppm 5 ppm 1 ppb 70 ppb 0.1 ppb 0.5 ppb 5 ppb

CEE, chloroethylethyl sulfide (a simulant for mustard gas); CVD, Chemical vapor deposition; DMMP, dimethyl methylphosphonate (a simulant for sarin); DNT, dinitrotoluene (explosive); MPCVD, Microwave plasma CVD; rGO, reduced graphene oxide.

hydrazine reduction of a film obtained by spin-coating of dispersion of GO synthesized by Hummers’ method with methanol and water mixture [94]. Fig. 3.62A shows change in the conductance of the film upon 5 s pulses of diluted acetone injections. The chemical sensor devices prepared by both rGO film and CNT exhibit fundamentally similar performance.

Electrical properties and applications 231 (A)

(B) CNT film

-4 -7

-0.5

-1.5

-13 -16

3h 6h 18 h 24 h

0.0

-1.0

Rapid reponse

-10

ΔG/G0 / a.u.

ΔG/G0 / 10-3

-1

Slow reponse

0

rGO film

100 200 Time / s

300

0

20

60 40 Time / s

80

Figure 3.62 (A) Change in the conductance of the film prepared by rGO and CNT upon 5 s pulses of diluted acetone injections. (B) Responses on the same experiments for rGO film with different time duration for the hydrazine reduction process [94].

Comparing with CNT film, rGO film has a larger irreversible component, which is attributed to the adsorption on the vacancies, structural defects, and oxygen functional groups. This is reasonable because rGO distinctly has more defective and oxidized parts than CNT. Indeed, with extension of the hydrazine reduction process that repairs graphene structure in GO, the irreversible slow components are gradually suppressed. By examination with several toxic or poisonous chemicals, rGO sensor is found to exhibit excellent performance, especially in detection of HCN (70 ppb for lower limit) in comparison with CNT-based sensors. The main detection principle of these chemical sensors using graphene is attributed to the change in the conductance of graphene by charge transfer between adsorbates and graphene. However, except for ideal conditions like adsorption of alkaline metal in ultrahigh vacuum chamber, direct charge transfer between adsorbed molecules and graphene is not so probable, where the adsorption/desorption process occurs under nonideal conditions such as ambient atmosphere, low-level vacuum, and inert gas atmosphere containing residue impurities. Under such nonideal conditions, charge transfer proceeds through electrochemical processes involving surrounding chemical species like residue water, which makes the principle of chemical sensing on graphene a more complex one. Indeed, the oxygen adsorption on the electrically charged graphene dramatically changes in the adsorption rate as much as two orders. Fig. 3.63 shows the time evolution of the shift of the charge neutrality point VCNP from that for nonadsorbed graphene, where Vg,ad is applied to graphene during O2 adsorption [95]. As clearly seen, the charge transfer kinetic largely depends on Vg,ad corresponding to the shift of the Fermi energy from the Dirac point during adsorption of O2. Here, it should be noted that doping level graphene is maintained even after evacuation of O2 from the measurement chamber. Moreover, no

V shift / V

n ox / 10 12 /cm

2

232 Chapter 3

O2 exposure time / min

Figure 3.63 Time evolution of the shift of the charge neutrality point VCNP from that for nonadsorbed graphene, where Vg,ad is applied to graphene during O2 adsorption [95].

charge doping was observed even with long exposure to O2 at 77 K. These results strongly suggest the presence of an energy barrier in the charge transfer process. Taking into consideration the fact that the presence of water vapor greatly accelerates the doping kinetics, it would be concluded that the charge transfer between graphene and adsorbed molecules proceeds through an electrochemical process involving water. This is also consistent with a quite slow timescale of charge transfer in the range of seconds to minutes as a simple electron transfer process. Furthermore, the inverse of the carrier mobility corresponding to the scattering probability of carriers is almost proportional to the amount of the doped hole induced by oxygen adsorption, indicating a scattering by Coulomb potential (Fig. 3.64). The adsorbed oxygen

1/μ / Vs/m2

50

40

30

20 0

1

2 4 3 Doping density / 1012/cm2

5

Figure 3.64 Inverse of the carrier mobility as a function of hole density doped by oxygen adsorption [95].

Electrical properties and applications 233 molecules should act as a charged scattering source against carriers in graphene after charge transfer process. The conductance of graphene is also significantly modified by carrier-scattering adsorbed molecules. These influences should be carefully considered, as well as structural changes with irreversible adsorption, to understand the principle for chemical sensing by graphene.

3.3.2 Photon detectors Photo detection is also one of the important technologies to support modern society, especially in the field of imaging application such as cameras, remote controls, including televisions and blue-ray players, for the wide range of wavelength of light from g-ray to electromagnetic wave for communication for various purposes and environments. The zero-gap semiconductor nature of graphene with high conductivity seems a disadvantage for the photonics application. In fact, the electronic properties such as the absence of a band gap are not critical but perhaps even an advantage for photonics and optoelectronics. The high carrier mobility is an advantage in conversion of photons or plasmons excited by photons to electrical signals like currents and voltages in ultrafast speed. The twodimensional nature of graphene is also a large advantage in integration of photodetectors, which enables 100-nm-scale detection resolution and pixelization of imaging devices. Furthermore, the zero-gap band structure, large mobility of carriers, and two-dimensional features make graphene a promising candidate for high-gain (efficiency) photodetection due to its very high sensitivity to even slight electrostatic perturbation. Graphene has a peculiar feature also in view of optical properties. Indeed, graphene can absorb photons in the ultra-wide range of energy due to the linear band dispersion unlike any other semiconductors used for photodetection. Moreover, graphene has unique dispersionless absorption coefficient in theoretical models [96], which is actually observed in the visible region of wavelengths as shown in Fig. 3.65. The wavelength dependence of the transmittance for monolayer graphene is almost invariant. Theoretically optical 2 transmittance of graphene with linear band dispersion is calculated as accurately pe Zc , where Z, c, and e are Planck constant divided by 2p, speed of light, and elementary charge. So, the transmittance of graphene is determined by only universal physical constants. The inset of Fig. 3.65 shows the transmittance of graphene with different numbers of layers. With increasing number of layers, the transmittance is reduced 2 pe2 accurately by pe Zc Zc for each layer. These facts justify the basis of evaluation of the number of layers for graphene by optical microscope. Thus, graphene can generate electron-hole pair by all wavelengths in equal efficiency due to its linear electronic band dispersion. Graphene is more suitable for long-wavelength light detection, although it exhibits universal light absorbance for all range of wavelengths in theoretical approximation, which

234 Chapter 3 (B) Light transmittance / %

Light transmittance / %

(A)

Distance / μm

Wavelength / nm

Figure 3.65 (A) Optical microscope image of monolayer and bilayer graphene superimposed with the transmittance along the gray horizontal line. (B) Wavelength dependence of the transmittance for monolayer graphene in comparison with theoretical models. The inset shows the transmittance for graphene with different numbers of layers [96].

becomes invalid in the higher energy region as seen in the deviation for lower wavelengths in Fig. 3.65B. Liu et al. fabricated a large-area, transparent, and flexible photodetector for infrared wavelength region by using CVD-grown graphene [97]. The key for photodetection in this device is a PeN junction by chemical doping. After transferring large-area graphene grown by CVD and attaching electrodes, the half part of the graphene between the electrodes is masked by photoresistance. Here, CVD graphene is already heavily P-doped through the transferring process under ambient conditions. So, the part of graphene under the mask is P-dope area. Electron doping to the nonmasked part of graphene was carried out by using the compound 2-(2-methoxyphenyl)-1,3-dimethyl-2,3dihydro-1H-benzoimidazole (o-MeO-DMBI), a strong n-type dopant, which can be doped easily by vapor deposition without any influence in transparency of the sample. Fig. 3.66 Photoresist

Spin-coat photoresist

Expose and develop for electrode pattern Deposit Au

N

P

Deposit N-type dopant

Photolithograph with alignment

Figure 3.66 Schematic view of the process of fabrication of a photodetection device having PeN junction by chemical doping to CVD-grown graphene [97].

Electrical properties and applications 235 (A)

(B)

Figure 3.67 (A) Schematic drawing for the IR detection by using PeN junction in graphene. (B) Photoresponse of the current between electrodes upon IR illumination (>780 nm) [97].

shows a schematic drawing of the IR detection in this device, where the irradiated IR excites a hole-electron pair at the PeN junction boundary. The photoexcited carrier might flow toward opposite directions according to the bias voltage between electrodes. Thus, the IR irradiation would be detected as the increment in the source-drain current as a consequence of the charge separation of photocarriers. By irradiating IR to the surface of the device, the source-drain current abruptly increases and shows sudden decreasing upon the switching off of the light source (Fig. 3.67). This is a similar characteristic of conventional semiconductor photodetectors [98]. However, it should be noted that graphene PeN junctions are not the same as the conventional semiconductor diodes in all of their characteristics. PeN junctions in graphene do not exhibit the “rectification effect,” which is the most famous feature of the semiconductor diode, because of the lack of band gap and pseudospin conservation in graphene [99]. Thus, the PeN junction structure fabricated here cannot be used for photodiodes, but only as photodetectors. A photodetector with high gain (efficiency) has been developed by combination of semiconductor nanoparticle (quantum dot; QD) and graphene, the former and the latter of which act as a photocarrier generator and sensitizer, respectively [100]. The schematic setting of the photodetection device is shown in Fig. 3.68A. Thin film of PbS colloidal QD was fabricated from the solution by spin-coating on the surface of typical FET transistor based on graphene prepared by micromechanical cleavage on SiO2/Si substrate. Initially, graphene is p-doped by charge transfer from PbS QD films. The photogating is responsible for the detection mechanism in this device, where holes in the PbS QD excited by laser are transferred to graphene channel and drift toward the drain electrode [101]. However, photoexcited electrons remain in the QD, which act like the gate electrode through capacitive coupling with hole carriers in graphene during its lifetime (Fig. 3.69A). The photocurrent responds successfully upon laser irradiation with a rise time of 10 ms, while the decay time is quite a bit longer. This is attributed to the long lifetime of

(A)

Light

b) (B)

Quantum dots

II / μA 0.52

VBG

VDS

IDS

0.45

Figure 3.68 (A) Schematic view of the graphene-PbS quantum dot hybrid photodetector. (B) Spatial photocurrent profile recorded by the laser beam scanning across the surface of the detector. The inset shows an optical image of the graphene contacted by electrodes before being covered by QD film [100].

Figure 3.69 (A) Energy-level diagram of the graphene/QD interface, where photoexcited electrons that remain in QD induce hole carriers in graphene. (B) Photocurrent response having a rise time of 10 ms for a wavelength of 532 nm (top). The response of the photodetector after the laser is turned off with reset pulse of the gate voltage and the schematic view after the application of a reset pulse for 10 ms (bottom) [100].

Electrical properties and applications 237 (A)

(B)

1x108

ΔRes / Ω

R / A/W

103 101

10-1

10-16

10-14 Power / W

10-12

Small quantum dots

5x107 0 600

800 1000 1200 1400 1600 Wavelength / nm

Figure 3.70 (A) Photoinduced resistance change (DRes) as a function of optical illumination power. (B) Excitation laser wavelength dependence of the spectral responsivity of the photodetector [100].

photoexcited electrons remaining inside the Pb QD with capacitive coupling with hole carriers in graphene. The photocurrent decay can be accelerated by the application of a gate voltage pulse. The pulse generates an electric field that reduces the potential barrier, which keeps electrons trapped in the quantum dots at the grapheneequantum dot interface. The detectable area of the device is investigated by scanning the excitation laser beam across the surface of the detector. The spatial profile shown in the inset of Fig. 3.68B indicates the large detection area of this device. Most strikingly, the QD/graphene hybrid photodetection device exhibits extremely high sensitivity to weak photoirradiation. Fig. 3.70A shows the photoinduced resistance change (DRes) as a function of optical illumination power. The photodetector exhibits a responsivity of 5  107 A W1 for an optical power of even less than 10 fW. The higher sensitivity is caused by the synergy of the high optical efficiency of QD and extremely sensitive nature of graphene to the slight perturbation in the carrier density. Indeed, the wavelength dependence of the responsivity follows the absorption spectrum of Pb QD. This proves the role of QD as an electron-hole pair generator with high efficiency.

3.3.3 Resistance standard For the current science and engineering supporting modern society, a standard of physical quantity is mandatory and required as an invariant one. Such a strict requirement is guaranteed by the traceability system of measurement instruments based on the International System of Units (SI). An accuracy of measurement instruments is calibrated by standard instruments, which are calibrated by upper-classed standard instruments. The origin of the chain of the calibration of instruments is called the “primary standard.” The accuracy of the primary standard is continuously examined by mutual comparison among the primary standards in the world. The dc resistance primary standard is currently utilized by quantum Hall resistance (QHR) on two-dimensional electron systems [102]. The resistance of two-dimensional electron systems is quantitized in a perpendicular magnetic

238 Chapter 3 field at universal values Rk/i, where Rk (h/e2), h, e, and i are the von Klitzing constant, the Planck constant, the elementary charge, and an integer [103]. The QHR standards based on GaAs/AlGaAs heterostructure devices forming two-dimensional electron systems reproduce Rk/2 with a relative uncertainty down to 3  1011. Although the resistance is determined by only fundamental physical constants, it is not very straightforward to operate this standard because a very low and a strong magnetic field are necessary for measurements. To guarantee the accuracy of traceability system, the shorter chain of calibrations of measurement instruments is quite important as well as the accuracy of the primary standard. However, extremely higher costs and technical barriers for measurements at ultra-low temperature and high magnetic field, which are necessary for the current resistance standard, reduce the accessibility for the primary standard, resulting in the longer chain of the calibration path in practical use. Since graphene intrinsically having two-dimensional electron systems exhibits the quantum Hall effect even at room temperature, the metrologists recently considered a QHR standard using graphene, which possibly surpasses the existing GaAs-based one [104e107]. To use QHR as a resistance standard, a guideline is given for materials and shape of the standard devices [108,109]. According to the guideline, an ohmic contact resistance between electrodes and materials having a two-dimensional electron system should be lower than 100 U to obtain uncertainty of less than 1  109, which is the lower limit for practical use of a resistance standard. The contact resistances of current electrodes are in the range of 101e105 U for typical devices for the resistance standard. One of the technical challenges to apply graphene to a resistance standard is making ohmic contacts between graphene and metal electrodes. The contact resistances for electrodes in a graphene-based QHR standard device are shown in Table 3.2, where the resistance is measured in the center of the plateau of quantitized Hall resistance [110]. The resistance is in the range of 0.1e10 U. This is higher than those in typical QHR devices, but lower to get enough accuracy for a practical resistance standard device. In addition, a clear difference of the contact resistance is seen between Table 3.2: Contact resistances for electrodes in a graphene standard resistance device [110]. Contact No. 1 3 4 5 6 7 8

Rholes (kU) 5.6 0.95 0.03 1.4 0.3 1.0 0.3

Relectrons (kU) 1.25 6.3 2.7 4.8 1.1 5.5 0.8

Δρxy/ρxy / ppm

Electrical properties and applications 239

Figure 3.71 Deviations of the resistance obtained by a graphene QHR standard device from the theoretical quantum resistance RK/2 in ppm for various contact configurations. The square (Rpa) denotes the deviation for a poorly annealed device [110].

electron and hole carriers. This is explained by carrier doping by metal electrode attaching [111] and nonideal coupling between the gold contacts and the graphene [112]. Another group reported a contact resistance as low as about 20 U without any special treatment or effort on contacts of electrodes [113]. The observed contact resistance also should not prevent realizing an accurate measurement of the Hall resistance by graphene in this report. It should be noted that the asymmetry for hole and electron carriers is also suggested for the contact resistance on their devices. Deviations from theoretical quantum resistance RK/2 shown in Fig. 3.71 for a QHR standard device based on graphene are two orders of magnitude larger than those in a measurement at the same current by using GaAs/AlGaAs heterostructure devices. However, the averaged resistance value for different runs of measurement and configuration of electrodes is almost identical to those in conventional QHR standard device based on GaAs/AlGaAs heterostructure and Si metal-oxide-semiconductor devices within a deviation of 15  15 ppm. Here, it should be noted that an appropriate annealing of device is important for the precise measurement for graphene QHR devices. Indeed, a poorly annealed device shows a larger deviation of the resistance in comparison to a well-annealed device as shown in Fig. 3.71. Improvement of the breakdown current is also a challenging task to overcome barriers for practical application of graphene to a resistance standard. QHR device is known to lose its quantum Hall state above the critical current, which is called the breakdown current of the quantum Hall effect [114e116]. The presence of the breakdown current is phenomenologically understood by the inhomogeneity in the electron temperature caused by hot electrons induced by a large flowing current. However, the microscopic mechanism of the breakdown of quantum Hall has been not clarified even in conventional QHR device like GaAs/AlGaAs heterostructure devices [116]. In typical

240 Chapter 3 (B)

14 2 μA

12

10 μA

10 8

20 μA

-16

B = 16.5 Tesla T = 4.2 K v=2

-14 -12 -10 -8 Gate voltage / V

-6

Longitudinal resistance / kΩ

Hall resistance / 103kΩ

(A)

6000 5000

20 μA

4000 3000 10 μA

2000

2 μA

1000 0

-16

-14 -12 -10 -8 Gate voltage / V

-6

Figure 3.72 Hall resistance (A) and longitudinal resistance (B) for the first plateau of the quantum Hall effect in graphene for various currents of 2, 10, and 20 mA [113].

measurements for a resistance standard device, a test current of 50e100 mA should flow through the device to get enough accuracy. For practical application of graphene to a resistance standard, the breakdown phenomenon of quantum Hall effect is also being investigated for graphene. Bennaceur et al. have examined a graphene QHR device in view of the breakdown current [113]. Fig. 3.72 shows the gate voltage dependence of the Hall and longitudinal resistances of graphene for the first plateau of the quantum Hall effect in hole carrier regime with the current of 2, 10, and 20 mA. As clearly seen in Fig. 3.72, a quantitized Hall plateau corresponding to a resistance standard remains its finite width up to 10 mA in spite of the obscure plateau for 20 mA due to emerging of the breakdown. The longitudinal resistance also remains zero below 10 mA in the middle of the plateau, although it exhibits a finite resistance for 20 mA. Similar robustness of a graphene-based resistance against the large current flowing is also reported from other groups [110], where the importance of annealing devices is noted in view of cleaning of the surface of graphene. Fig 3.73 shows the source-drain current dependence of the longitudinal resistance of a graphene QHR device after successful annealing obtained by different combination of the electrodes on the device for hole and electron carriers. The longitudinal resistance almost remains zero up to 3.5 mA for both electron and hole carriers. Interestingly the result of the same measurement shows significant magnitude of a finite resistance for a poorly annealed device as shown in the inset of Fig. 3.73. These results suggest that graphene is promising as a new material as a resistance standard in view of metrology. With further improvement in contact resistances for both electrons and holes, and using wider graphene samples for higher breakdown currents, this would achieve a graphene QHR device with enough accuracy for practical use for a resistance standard. However, using graphene obtained by micromechanical cleavage is not

Electrical properties and applications 241

ρxx / Ω

(B)

ρxx / Ω

(A)

ρxx / Ω

Isd / μA

Isd / μA

Isd / μA

Figure 3.73 Source-drain current dependence of the zero longitudinal resistance at the quantum Hall plateau for hole and electron carriers [110].

conventional for a resistance standard widely used, where monolayer samples are smaller in size and quite rarely obtained with poor reproducibility. Thus, there have been many attempts for resistance standard devices using graphene samples more conventionally applied [117e124]. Lafont et al. fabricated a QHR standard device using a large-area graphene sample synthesized by propane/hydrogen CVD on SiC substrate [124]. Fig. 3.74A shows the deviations of the Hall resistance value from the theoretical quantized resistance RK/2 at 1.4, 2.2, and 4.2 K with a larger measurement current of 20 mA. Fig. 3.74B indicates the magnetic field dependence of the longitudinal and Hall resistances obtained by voltage measurement using (V1, V2) and (V2, V3) electrodes, respectively, where a current of 100 nA circulates between I1 and I2 electrodes. The Hall resistance shows a quite nice plateau corresponding to RK/2 over 9 T of the magnetic field B range (10 T < B < 19 T) with a 109 relative uncertainty even under the large current flowing condition (20 mA) as shown in Fig. 3.74A. The zero longitudinal resistance also completely coincides with the plateau of the Hall resistance as shown in Fig.3.74C, where the resistance values are lower than (30  20) mU. Compared with the results of a reference QHR standard device made of GaAs shown in Fig. 3.74B, the relative discrepancy between the quantized Hall resistances in the graphene-based device and in a reference GaAs one is equal to (2  4)  1010. It should be noted that the device reported here also exhibits quite a large critical current for the breakdown of the quantum Hall state. As shown in Fig. 3.75, the QHR standard device exhibits quite a nice plateau even with the current of 50 mA, where the resistance standard value is easily distinguishable. Thus, graphene might be able to replace GaAs used in existing QHR standard devices.

ΔRH/RH / 10-9

242 Chapter 3

(A)

Rxx / μΩ

R / kΩ

(B)

(C)

B / T Figure 3.74 Magnetic field dependence of (A) the deviation of the quantitized Hall resistance from the theoretical value RK/2 of 1.4, 2.2, and 4.2 K; (B) the longitudinal Rxx and Hall RH resistances; and (C) the zero longitudinal resistance at the center of the quantized Hall plateau for QHR device based on graphene, which is synthesized by CVD on SiC substrate. The results for a reference QHR device based on GaAs (LEP 514) is also shown in (B). The inset photo shows the top view of QHR device with electrodes [124].

3.3.4 Electron field emission Graphene is also promising for field emission devices, where thermally stable, lightweight, and abundant materials are required in replacing the existing field emission materials, which often depends on rare resources. The self-standing graphene structures on the cluster/film surface are attractive because a huge electron field emission property is expected due to the nanolevel-sharpened edge and the high electron conductivity. Graphene flower cloth (GFC) shown in Fig. 3.76A and B

R / kΩ

Electrical properties and applications 243

B/T

Figure 3.75 Magnetic field dependence of the Hall resistance and longitudinal resistance for a QHR device using graphene synthesized by CVD on SiC substrate with different source-drain current [124].

(A)

(B)

(C)

X-ray Be window and Ta target (200 nm) 4.0 mm

GFC

Ba getter High voltage GND

Glass tube GND

Figure 3.76 Graphene flower cloth (GFC) [125]: (A) photograph of a cloth, (B) SEM image of a single fiber with a partially enlarged image of its surface, and (C) schematic illustration of the circuit for X-ray tube using GFC field emitter.

was successfully applied to the cathode of an X-ray tube as shown in Fig. 3.76C [125]. The tube current increased exponentially above the tube voltage of 8.5 kV, reaching 500 mA at 11.5 kV, with very small fluctuation in voltage. The X-ray tube was successfully applied to X-ray fluorescence analysis of stainless steel. The hybrids of single-walled carbon nanohorns (SWCNHs) with single-walled carbon nanotubes (SWCNTs) (called NTNH) were prepared by growing SWCNTs on SWCNH aggregates using Fe catalysts [126]. A paste of the hybrid (NTNH) was screen-printed on an ITO substrate and then sintered at 500 C, which gave a turn-on electric field of 5 kV/cm, much lower than SWCNTs, SWCNHs, and their mechanical mixture (1:1) do, as shown in Fig. 3.77.

Emission current density / mA/cm2

244 Chapter 3 Hybrid (SWCNT/SWNH)

SWCNTs A mixture of SWCNTs and SWNHs

SWNHs

Electric field / kV/cm

Figure 3.77 Current density versus electric field for emitters of SWCNTs and SWCNHs [126].

SWCNHs were supposed to be effective to make SWCNTs dispersed with each other. The electron field emission was measured on flower-like aggregates of graphene sheets prepared by MPCVD of H2/CH4 with NieFeeCr catalysts on Al2O3 substrate [127].

3.4 Concluding remarks The novel electronic structure of graphene having the zero-gap semiconductor band structure with linear dispersion results in the peculiar electronic properties and functions. These electronic features, which are not available in existing conventional materials, is an important reason for applying graphene to various fields such as logic circuit devices for data processing, transparent electrodes for tach screen interfaces, chemical sensors for IoT technology, photodetection devices for imaging applications, and resistance standards for supporting science and engineering. Through these aspects, we are convinced that graphene is a promising material to support a sustainable and smart society for the benefit of our next generations.

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CHAPTER 4

Chemical properties and applications Chapter Outline 4.1 Fundamental chemical properties 252 4.1.1 Hydrogenation 253 4.1.2 Oxygenation 256 4.1.3 Layer modification 258

4.2 Applications to energy storage and conversion

259

4.2.1 4.2.2 4.2.3 4.2.4 4.2.5

Lithium-ion batteries 259 Electrochemical capacitors 268 Lithium-ion capacitors 278 Lithium-sulfur batteries 282 Solar cells (photovoltaic cells) 292 4.2.5.1 Semiconductor solar cells (Schottky junction solar cells) 4.2.5.2 Polymer solar cells (dye-sensitized solar cells) 295 4.2.6 Fuel cells 298 4.2.7 Hydrogen storage 306

4.3 Applications to environment remediation 4.3.1 4.3.2 4.3.3 4.3.4 4.3.5

292

311

Adsorption of polluting molecules and ions 311 Sorption and recovery of oils 321 Capacitive deionization for water desalination 325 Catalysts 332 Chemical sensors 341

4.4 Concluding remarks References 354

352

In this chapter, the applications based on chemical properties of graphene materials to the storage of electrical energies and environment remediation are discussed based on research results published in various journals. In the field of energy storage, lithium-ion rechargeable batteries (LIBs) and electric double-layer capacitors (EDLCs), as well as the devices based on the hybridization of capacitive functionalities with battery ones, lithiumion capacitors (LICs), are focused on. In addition to these electrochemical energy storage devices, graphene is expected to deliver high performance in solar cells and fuel cells. For hydrogen storage, graphene cannot yet satisfy the proposed target values, but many computational studies are predicting promising performances. The applications of graphene and its related materials to environmental remediation are discussed by classifying into adsorption of polluting gases and ions, sorption and recovery of oils, Graphene. https://doi.org/10.1016/B978-0-12-819576-5.00004-9 Copyright © 2020 Elsevier Inc. All rights reserved.

251

252 Chapter 4 Graphene Single layer of mono-atomic thickness 4. Chemical properties and applications 4.1 Fundamental properties 6 Thermal properties High thermal conductivity 7 Biomedical properties High biocompatibility

Large surface area, Giant aromatic layer, Extended π-electron clouds at both sides of the layer Functionalization

4.2 Energy storage Lithium-ion batteries (LIBs), Electric-double layer capacitors (EDLCs), Lithium-ion capacitors (LICs), Lithium-sulfur batteries, Solar cells, Fuel cells, Hydrogen storage,

3 Electrical properties High electrical conductivity 5 Mechanical properties

High strength, Flexible layer

4.3 Environment remediation Adsorption Oil sorption Capacitive de-ionization Catalysis Chemical sensors

Figure 4.1 Concept of this chapter.

capacitive deionization for water desalination, catalysts for environmentally friendly chemical processes, and carbon-based sensors for detection polluting and toxic chemicals. The concept of this chapter is summarized in Fig. 4.1. Applications of graphene and graphene-related materials in these fields have been reviewed from different points of view [1e9].

4.1 Fundamental chemical properties Theoretically, graphene is a two-dimensional (2D) flat layer with single atomic thickness consisting of sp2-hybridized carbon, in other words, it has aromatic nature possessing highly dense p-electron clouds both above and below the layer. Specific surface area of 2965 m2/g was theoretically calculated for infinite single-layer graphene [10] and a large specific surface area as 2600 m2/g has been cited in different literature. Practically, graphene films (sheets) synthesized by chemical vapor deposition (CVD) and cleavage methods are presenting the nanomaterials having almost ideal structure with a low density of structural defects, except the edges of the layer. In these graphene films (sheets), the lateral size and number of stacked layers can be controlled by the selection of the size of the substrate and the deposition/cleavage conditions, although certain

Chemical properties and applications 253 limitations in the controls of the size, structural perfection, and stacked layer number cannot be avoided, for example, whether a large enough sized single crystal of a substrate metal, whether single crystalline graphite to be cleaved is available, etc. Even by taking these practical limitations into consideration, however, the size of an aromatic plane is extraordinary large, in comparison to the conventional aromatic compounds, and the graphene layer creates extended p-electron clouds that can interact with p-electron clouds of other aromatic compounds and also with other cationic species. On the other hand, the so-called reduced graphene oxide (rGO) nanoflakes prepared from various kinds of graphite through their oxidation, exfoliation, and reduction contain different functional groups, mainly oxygen-containing groups, and many structural defects in the layers. On these rGO flakes, the extension of p-electron clouds is highly limited due to the limited size of the flake and also the presence of functional groups even inside of the layer. However, it has to be pointed out that the presence of functional groups makes the modification of flakes easier, for example, control of hydrophilicity, functionalization through various chemical bonds, etc. On both graphene and rGO, functionalization is possible, even on both above and below the layer.

4.1.1 Hydrogenation The chemical properties of monolayer graphene are expected to be largely different from bulk graphite, although graphite is known as a chemically inert material. The chemical stability of graphene in specific atmospheres is particularly important for its applications in nanodevices. Reversible hydrogenation was reported to be possible on thin graphene flakes [11]. Singlelayer graphene prepared by micromechanical cleavage could react reversibly with atomic hydrogen and transform its electronic state from highly conductive zero-gap semimetal to insulator [11]. The resultant hydrogenated graphene (graphane) was crystalline and retained its hexagonal lattice by making its lattice spacing shorter. Changes in Raman spectrum due to the hydrogenation, i.e., marked developments of D- and D0 -bands, and due to the annealing, i.e., marked intensity decrease in G0 -band, are shown in Fig. 4.2. Hydrogenation of graphene sheets mechanically cleaved from highly oriented pyrolytic graphite (HOPG) was also performed by electrolysis in distilled water by applying 10.0 V and 2  103 A [12]. Hydrogenated rGO flakes were prepared by heating GO in a hydrogen atmosphere under high pressure (60e150 bar) and temperature (200e500 C) [13]. Hydrogenation was performed by radio frequency hydrogen plasma (10 W, 1 Torr, and 1 min) on mechanically cleaved HOPG with different layer numbers from single to four layers on SiO2/Si substrate [14]. Changes in Raman spectrum with hydrogenation and annealing under vacuum are shown in Fig. 4.3. Appearance of D-, D0 -, and (D þ G)-bands related to the lattice defects in graphene are reasonably supposed to be introduced by the

254 Chapter 4 G’

Intensity / a.u.

G’

Pristine

Pristine

D

Hydrogenated

G

D

G

D+D’

D’

D+D’

Annealed

Annealed

1250

Hydrogenated

D’

1500 2500 2750 Raman shift / cm-1

1250

1500 2500 2750 Raman shift / cm-1

Figure 4.2 Raman spectra of pristine, hydrogenated, and annealed single-layer graphene [11]; (A) on SiO2 substrate and (B) free suspended.

(A)

(B) G

(C)

D

G

Intensity / a.u.

G D’

G’ 4 layers 3 layers

G’

D+G 4 layers

G’ 4 layers

3 layers

3 layers

2 layers

2 layers

2 layers 1 layer

1500 2000 2500 3000 Raman shift / cm-1

1 layer

1500 2000 2500 3000 Raman shift / cm-1

1 layer

1500 2000 2500 3000 Raman shift / cm-1

Figure 4.3 Changes in Raman spectrum of graphene with hydrogenation and annealing [14]; (A) pristine, (B) after hydrogenation under hydrogen plasma, and (C) after annealing.

hydrogen plasma treatment mainly due to the formation of sp3 bonds. The results demonstrate that the hydrogenation of graphene flakes of two to four layers are more feasible than single-layer graphene (Fig. 4.3B). Since the dominant species irradiated are supposed to be Hþ 3 ions and hydrogen radicals (H atoms) under the present plasma condition, the hydrogenation was reasonably supposed to occur on the top graphene layer and the intensity of the D-band was proportional to the hydrogen coverage on the top graphene layer. By vacuum annealing at 500 C for 30e45 min, the Raman bands related to the defects (D, D0 , and D þ G) induced by hydrogenation are eliminated (Fig. 4.3C).

Chemical properties and applications 255 (A)

(B)

400 1.0

3

200

8 4

100

Double-layered

x30

1

0.4

2 1

0

0 1200

0.6

1300 1400 1500 1600 Raman shift / cm-1

1700

0.2

0

Intensity of Si-H peak / a.u.

0.8 2

ID/IG ratio

Intensity / a.u.

300

Single-layer

e-beam dose (mC/cm2)

0.0 0

2 6 4 e-beam dose / mC/cm2

8

Figure 4.4 Hydrogenation of single-layer and double-layered graphene [15]; (A) Raman spectrum change of single-layer graphene with e-beam irradiation of hydrogen silsesquioxane (HSQ), and (B) changes in ID/IG ratio of single- and double-layered graphenes and intensity of SieH peak with e-beam dose.

Single-layer graphene on SiO2 substrate exhibits a prominent growth of D-band in its Raman spectrum after the hydrogenation by electron beam irradiation in the presence of hydrogen silsesquioxane (HSQ, (HSiO3/2)n), as shown in Fig. 4.4A [15]. Since both the pristine single-layer graphene and graphite irradiated by an electron beam without HSQ (just 30 keV electron dose) showed negligibly small D-band intensity, D-band growth is due to hydrogenation by reactive H atoms liberated from HSQ, which is confirmed rapid decrease in SieH peak in Raman spectrum, as shown in Fig. 4.4B. On a double-layered flake, the growth of D-band is very slow and ID/IG ratio increased slowly to about 0.05 with increasing e-beam dose up to 8 mC/cm2, although on a single-layer graphene, ID/IG ratio increased rapidly to about 3 after the same dose (Fig. 4.4B). If the front and reverse surfaces of graphene react with hydrogen atoms to create the same density of defects, the ID/IG ratio is expected to be larger for single-layer than for double-layered graphene flakes because of the presence of an intact lower layer for double-layered graphene, but the defect density of the former is supposed to be much higher than that of the latter, suggesting that atomic H can react with the carbon atoms in graphene basal plane. The hydrogenation rate of single-layer graphene was faster than that of double-layered one. By annealing at 100 C in air, the D-band of hydrogenated single-layer graphene decreased drastically in intensity relative to the G-band, ID/IG decreasing by more than a factor of six (dehydrogenation). After thermal annealing above 250 C, the dehydrogenated graphene was “activated”: it could react with molecular oxygen at room temperature, although the pristine graphene could not react.

256 Chapter 4 Graphene hydrogenated completely is named as “graphene,” which is classified as a graphene derivative and is discussed in more detail in Chapter 8 (Section 8.1.1). Graphane is a fully saturated hydrocarbon derived from a single-layer graphene with formula CH, where all of the carbon atoms are in sp3 hybridization forming a hexagonal network, and the hydrogen atoms are bonded to carbon on both sides of the plane in an alternating manner. This is a giant aliphatic hydrocarbon and its stability was confirmed on the basis of first-principles total-energy calculations [16]. Hydrogenation of just the carbon atoms at the edge of graphene layer leads to the formation of a giant molecule of aromatic hydrocarbon.

4.1.2 Oxygenation Oxidation of graphite in acidic solution has been carried out for the production of flexible graphite sheets through oxidation, thermal exfoliation and compression. Oxidation in oxygen gas has been studied as the process for gasification of carbon materials, particularly coal. Now Oxidation in acidic solution is commonly used as the preparation process for graphene oxide (GO) for the precursor of reduced graphene oxide (rGO). The oxidation products in acidic solutions have various functional groups, such as -OH, -COOH, etc., at the edges of graphene layers and at the defects in layers. The details of the preparation of GO were discussed in Chapter 2 as one of the routes to prepare graphene (rGO) (Section 2.3) and their properties and applications are discussed as one of graphene derivatives in Chapter 8 (Section 8.1.3). Oxygenation of graphene grown epitaxially on the Si face of SiC(0001), which was composed of 1e3 layers of graphene, was performed under ultrahigh vacuum (UHV) conditions by atomic oxygen produced by cracking O2 molecules on a tungsten filament at w1500 C [17] and by drop-casting Hummers’ oxidizing agents (H2SO4 þ NaNO3 þ KMnO4) [18]. The atomic oxygens are randomly chemisorbed on the basal plane of the graphene and are observed as bright protrusions in scanning tunneling microscopy (STM) images, as shown in Fig. 4.5A, the presence of oxygen atom on the surface being confirmed by Auger electron spectroscopy. The lateral (A)

(B) UHV oxidation at room temp. Annealing at 280 oC

Figure 4.5 STM images of graphene surface oxygenized under UHV (A) and annealed at 280 C [17].

Chemical properties and applications 257 diameter of the oxygen adatoms observed in the STM image was around 1.2 nm, approximately eight times greater than CeC bond length in graphene layer, of which density functional theory (DFT) analysis suggested that oxygen adatoms on the basal plane of graphene perturbed the planar pep interactions of the graphene lattice through sp3 bonding by forming epoxy structure. Although the chemisorbed oxygen was stable at room temperature, even under STM imaging conditions, it could be reversibly removed by annealing the oxidized surface at 260 C, as shown in Fig. 4.5B. When the reduced surface shown in Fig. 4.5B was again exposed to atomic oxygen, an oxidized surface comparable to that of Fig. 4.5A was reestablished, the UHV oxidation-reduction process being fully reversible [17]. In contrast, oxidation by drop-casting aqueous phase caused the patches of oxidized and nonoxidized areas, the former (oxidized areas) containing carbonyl and hydroxyl groups, in addition to the dominant epoxy groups, so that the remaining were partially oxidized even following annealing at 1000 C [18]. The changes in X-ray photoelectron spectroscopy (XPS) spectra for C(1s) and O(1s) on the oxidized area with annealing at different temperatures are shown in Fig. 4.6A and B, respectively. O(1s) spectrum of the as-oxidized area in Fig. 4.5B can be resolved into four components, where the main peak O1 is ascribed to the epoxy group, and the relatively small peaks O2, O3, and O4 are ascribed to ether, hydroxyl, and carbonyl groups, respectively. The peak O1 due to the oxygen in epoxy groups disappears by annealing at 300 C, but the other peaks O2 to O4 remain, these three components becoming clearly visible. By further annealing at 600 and 1000 C, complex reactions involving the C

(A)

(B) C(1s)

O(1s)

Intensity / a.u.

1000

600 oC

Intensity / a.u.

1000 oC oC

600 oC 300 oC

300 oC O1

As-oxidized O3 290

282 286 Binding energy / eV

536

As-oxidized O2 O4

532 528 Binding energy / eV

Figure 4.6 XPS spectra of as-oxygenized graphene and annealed at 300, 600 and 1000 C in UHV [18]; (A) C(1s) and (B) O(1s).

258 Chapter 4 atoms in graphene and the remaining oxygen functionalities are expected to occur, but XPS suggests the remaining of a minor amount of O on the surface of the graphene even after annealing at 1000 C. Oxygenation of thin flakes was reported to occur in O2/Ar gas flow at a temperature of 200e600 C [15]. Single-layer regions in a flake were oxygenized faster than triple-layered regions, of which the oxidation rate was comparable with bulk natural graphite and resulted in the random formation of etch pits on its surface.

4.1.3 Layer modification Ideal graphene layer has various features, high chemical stability, highly electrical and thermal conductivities, high mechanical properties with flexibility, etc. For practical applications of graphene and related materials in various fields, some modifications of these layers are often required by keeping these characteristics of graphene layers. On rGO-related materials, these modifications are mostly performed during the preparation processes by calling “functionalization,” and so have been described in previous chapter, Section 2.3.5. In this section, the modification of graphene layers is explained by focusing on the graphene films synthesized by CVD. Halogenation of graphene layers was studied by using epitaxially grown graphene films composed of 1e3 graphene layers, because halogenated graphene may be directly used as an electronic material and as an intermediate for further modification by organic radicals [19e21]. Fluorination was possible by exposing the graphene films on SiC substrate to XeF2 in Ar atmosphere at room temperature, irrespective of layer numbers stacked. C/F ratio at saturation coverage was evaluated to be 0.65 from XPS analysis. On the other hand, chlorination by using Cl2 gas occurred very selectively onto single layer under the irradiation of UV. A very small amount of bromine could be detected on the graphene film after the reaction with Br2 under UV light. Iodine could not react with the graphene at room temperature. The fluorinated and chlorinated graphene films were treated with alkyl Grignard reagents containing N and F in an Ar atmosphere at 60 C [19,20]. No reaction between fluorinated graphene and the Grignard reagent was detected by XPS, probably due to the strong bond of CeF. However, the chlorinated graphene could readily react with the Grignard reagent by replacing the chemisorbed Cl by alkyl groups in Grignard reagent. Chemisorbed Cl could be replaced by alkyl and aryl groups by using aryl bromide, bromopentafluorobenzene (C6F5Br), and alkyl bromide, (CHCl2Br) [21]. In Fig. 4.7, STM images of the epitaxially grown graphene on Si face of SiC(0001) are shown for the pristine graphene with low magnification (A) and high magnification (B), after chlorination by Cl2 gas under UV irradiation, showing chlorinated and nonchlorinated regions, (C) and after methylation by reacting with CH3MgBr in THF at 65 C.

Chemical properties and applications 259 (A)

(B)

Single-layer area

(C)

(D) Single-layer area

Double-layered area

Double-layered area

Figure 4.7 STM images of graphene, (A) and (B) pristine graphene, (C) after chlorination showing chlorinated (single-layer area) and nonchlorinated (double-layered area) regions, and (D) after methylation [19].

4.2 Applications to energy storage and conversion 4.2.1 Lithium-ion batteries Graphenes prepared through different processes and doped graphene have been tested as anode material for lithium-ion rechargeable batteries (LIBs). Graphene does also work as a support for electrochemically active materials, nanoparticles of metals, and metal oxides to enhance their performances as electrode materials of LIBs. Flakes of rGO were prepared by thermal exfoliation and reduction at 1050 C, supposed to be composed of few layers (approximately 2.1 nm thick), of which electrochemical behaviors were determined by using the coin-type cells with the counter and reference electrodes of Li metal in 1 M LiPF6/(EC þ DMC) electrolyte [22]. The working electrode was fabricated by mixing rGO nanoflakes with PVDF binder (1/9 in mass). As shown in Fig. 4.8A, discharge and charge capacity for the first cycle calculated at the range of 0.01e3.5 V is 2035 and 1264 mAh/g, respectively, giving a large irreversible capacity. After the second cycle, however, almost the same reversible capacity, around 1100 mAh/g, is obtained at least up to five cycles, which is much larger than the theoretically predicted capacity for graphite (372 mAh/g) and the capacity experimentally observed on various graphite materials. The discharge/charge curves for rGO are sloped, in contrast to the welldefined plateau for graphite (an inset in Fig. 4.8A), and are rather similar to disordered carbons. These characteristics in voltage curves are kept even with high current densities, as shown in Fig. 4.8B, although the reversible capacity decreases from 1100 mAh/g at the current density of 100 mA/g to 936, 718, and 445 mAh/g at the current density of 300,

260 Chapter 4 (A)

(B)

4

4 1000 mA/g 500 mA/g 300 mA/g

2

3

1st

1

3 Voltage / V

2nd~5th Voltage / V

Voltage / V

Charge, 1st~5th

Discharge

3

Charge

Discharge

2 1

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2

1

0 0

100 200 300 Capacity / mAh/g

Discharge, 2nd 0

0 0

500 1000 1500 Capacity / mAh/g

0

2000

200

600 400 Capacity / mAh/g

1000

800

Figure 4.8 Discharge/charge curves of rGO [22]; (A) the first to fifth cycles at the current density of 100 mA/g, in comparison with the pristine natural graphite (inset), and (B) the second cycle at different current densities.

500, and 1000 mA/g, respectively. The reversible capacity of rGO nanoflakes was experimentally shown to depend strongly on the process of reduction of GOs [23]. Cycle performances of rGOs prepared by thermal reduction at 300 and 600 C, by hydrazinereduction and by reduction under electron beam irradiation, are shown in Fig. 4.9A, in comparison with the original natural graphite and the pristine GO. Thermally reduced and electron beameirradiated rGOs show high reversible capacities as 800e1013 mAh/g in the

(A)

(B) 1: Original natural graphite 2: Pristine graphene oxide 3: Hydrazine-reduced 4: 300oC-reduced 5: 600oC-reduced 6: Electron-beam-reduced

Reversible capacity / mAh/g

1600 1400 1200 1000

D-band G-band 6

3

4 6 5

800 600 400

5 4

1 2 3

200 0

2

4

10 8 6 Cycle number

12

14

16

2 1600 1400 1200 Raman shift / cm-1

Figure 4.9 rGO nanoflakes prepared via different reducing processes [23]; (A) cycle performance as an anode of LIB and (B) Raman spectra.

Chemical properties and applications 261 voltage range of 0.005e3.0 V in 1 M LiPF6/(DMC þ DEC þ EC) electrolyte, but hydrazine-reduced rGO shows very low capacity, much smaller than the original natural graphite and the pristine GO. These high reversible capacities for rGOs were discussed in relation to Raman spectra (Fig. 4.9B); the rGOs showing high reversible capacities have high ID/IG values, in other words, highly defective, suggesting the presence of additional reversible storage sites, such as layer edges and other defects in the layer, in addition to the functional groups remained in graphene layers. From the supernatant of hydrazine-reduced GO dispersion, flower-like agglomeration of rGO nanoflakes were obtained, as shown in Fig. 4.10A, most of the petals of the flower consisting of two to three layers with the interlayer spacing of 0.37 nm and demonstrating corrugation and scrolling [24]. Discharge/charge curves in the potential range of 0.02e3.0 V at 1C rate for first and 100th cycles are shown in Fig. 4.10B, demonstrating clearly the presence of irreversible capacity due to the formation of solid electrolyte interphase (SEI) film in the first cycle. These graphene nanoflakes delivered reversible capacity of 650 mAh/g (irreversible capacity of about 300 mAh/g) in the first cycle and 460 mAh/g after 100 cycles with 1C rate. On rGO nanoflakes prepared by thermal exfoliation and reduction of GO and having the thickness of 2e5 nm (6e15 layers stacked), reversible capacity was reported to be 540 mAh/g with 1 M LiClO4/ (EC þ DMC) electrolyte and to increase up to 730 and 784 mAh/g by mixing of carbon nanotubes (CNTs) and C60, respectively [25]. rGO nanoflakes prepared from artificial graphite via its GO by rapid exfoliation and ultrasonication were reported to give the reversible capacity of 670 mAh/g with the characteristics common for all rGOs reported, i.e., sloped voltage curve and a large irreversible capacity at the first discharge/charge cycle [26]. On the other hand, very large irreversible capacity (about 680 mAh/g) and very small reversible capacity (84 mAh/g) were observed on flexible rGO films, which were

(B)

(A)

100th

3

1st

Voltage / V

Charge 2

1 Discharge 0

300 nm

0

100th 200 400 600 800 Capacity / mAh/g

1st 1000

Figure 4.10 rGO nanoflakes showing flower-like aggregation [24]; (A) SEM image and (B) discharge/charge curves.

262 Chapter 4 prepared by vacuum filtration of aqueous dispersion of chemically reduced rGO [27]. Large irreversible capacity was discussed by relating to the residual oxygen-containing groups in the rGO. rGO foams prepared by heating the mixture of GO and ethylene glycol under hydrothermal condition at 180 C, followed by freeze drying, exhibited a high Li storage capacity, 172 mAh/g at 0.1 A/g without capacity loss after 3000 cycles [28]. N- and B-doping of rGOs were effective to increase the reversible capacity in 1 M LiPF6/ (EC þ DMC) electrolyte [29]. Doping of N and B was performed by heating rGO nanoflakes (less than three layers stacked) at 600 C in a flow of mixed gas of NH3/Ar (1/2 in volume) and at 800 C in a flow of BCl3/Ar (1/4 in volume), respectively. As shown in Fig. 4.11, the N- and B-doped rGOs exhibit reversible capacities of 1043 and 1549 mAh/g, respectively, in the first cycle, and 872 and 1227 mAh/g after 30 cycles, which are much higher than those of the pristine rGO (955 mAh/g in the first cycle and 638 mAh/g after 30 cycles). N-doped graphene composed of single or few layers was synthesized by CVD of acetonitrile (AN) at 950 C on Cu substrate, which delivered the reversible capacity of 0.05 mAh/cm2 with 5 A/cm2 rate between 0.02 and 3.2 V in 1 M LiPF6/(EC þ DMC) electrolyte [30]. N-doped rGO nanoflakes fabricated through the heat treatment of a thermally exfoliated and reduced rGO in NH3 atmosphere were reported to have high capacity and cyclability at least up to 550 cycles [31]. High capacity and rate capability of N-doped graphene was discussed based on in situ high-resolution TEM observations [32]. Co3O4/rGO composite (Co3O4-loaded rGO) was prepared by depositing Co(OH)2 nanoparticles on suspended rGO nanoflakes in isopropyl alcohol/water solution using Co(NO3)2 and NH4OH, followed by calcination at 450 C in air [33]. TEM images of the composite are shown in Fig. 4.12. Co3O4 nanoparticles with the size of 10e30 nm are

(B)

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600 49.0 %

400

Charge Discharge

200 0 0

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1600 1400 1200 1549 mAh/g 1000 800 55.6 % 600 400 200 0 0 5 10

1327 mAh/g

1227 mAh/g Charge Discharge

15

20

25

80

Coulombic efficiency / %

(A)

30

Cycle number

Figure 4.11 Cycle performance and Coulombic efficiency of doped rGOs at the current density of 50 mA/g and the voltage range of 0.01e3.0 V [29]; (A) N-doped and (B) B-doped rGOs.

Chemical properties and applications 263

Figure 4.12 TEM images of Co3O4/rGO composite with two different magnifications (A) and (B) [33]. The inset in (B) is the selected-area electron diffraction pattern of Co3O4 nanoparticle.

(A)

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400 600 800 1000 1200 Capacity / mAh/g

Figure 4.13 Discharge/charge curves for (A) Co3O4/rGO composite and (B) Co3O4 in 1 M LiPF6/ (EC þ DMC) electrolyte with a current density of 50 mA/g [33].

homogeneously anchored on thin layers of rGO. In Fig. 4.13A and B, discharge/charge curves in 1 M LiPF6/(EC þ DMC) electrolyte with the current density of 50 mA/g at the potential range of 0.01e3.0 V are shown for the composite and Co3O4 without graphene. The curves for Co3O4 show a long voltage plateau at around 1.1 V (typical characteristics for Co3O4), followed by a sloping curve down to the cutoff voltage of 0.01 V, although rGO shows sloping curves in this voltage range, as shown before (for example, Fig. 4.8). However, Co3O4/rGO composite exhibits much better cycle performance than Co3O4; the reversible capacity fades with the cycles for the latter, but it retains almost the same values for the former, even increasing slightly up to 935 mAh/g. In addition, Co3O4/rGO

264 Chapter 4 composite exhibited much better rate capability than Co3O4 at various rates between 50 and 500 mA/g, about 800 mAh/g for 50 mA/g and about 480 mAh/g for 500 mA/g. CuO/ graphene composite, which was prepared by mixing ethanol suspension of graphene synthesized by arc discharge method with aqueous solution of Cu(CH3COO)2 and ammonia, gave a discharge/charge curves similar to Fig. 4.13A with reversible capacity of about 600 mAh/g at 65 mA/g rate [34]. The composite exhibited marked rate capability by delivering reversible capacity of 150 mAh/g at 6.4 A/g rate, though graphene 48 mAh/g and CuO were nearly null. Mn3O4/rGO and Fe3O4/rGO composites were prepared by dip coating of Ni foam (180 mm thickness) into rGO and either Mn3O4 or FeOOH nanoparticles suspensions, alternative repetition of dipping by 10 times [35]. SEM and TEM images of the composite of Mn3O4 and rGO deposited on Ni foam are shown in Fig. 4.14. This composite delivered a reversible capacity of about 1000 mAh/g by exhibiting a shallow plateau at the voltage range of 1.3e0.4 V due to the formation of SEI layer on the surface of electrode and distinct plateau at 0.4 V due to the reversible reaction between Liþ and Mn3O4 particles. Rate capability at 0.1e5 A/g rate of Mn3O4/rGO and Fe3O4/rGO composites on Ni foam are shown in Fig. 4.15A and B, respectively. Mn3O4/rGO composite delivers reversible capacity of 550 mAh/g even at the rate of 5 A/g, though Mn3O4 on Ni foam gives almost null (Fig. 4.15A). Fe3O4/rGO composite gives higher reversible capacity as 1360 mAh/g at 0.1 A/g rate and 752 mAh/g at 5 A/g (Fig. 4.15B), the former retaining the capacity as 923 mAh/g and the latter as 1180 mAh/g after 100 cycles with 1 A/g rate. SnO2/rGO composites were prepared by reducing suspended GO and SnCl2 with NaBH4 at 120 C [36] and by dispersing rGO nanoflakes and SnO2 nanoparticles in ethylene glycol [37]. Crystalline SnO2 nanoparticles with 4e6 nm size were well dispersed on the rGO layers. The composite gives sloping voltage profile, as observed commonly on rGO, as

Figure 4.14 Mn3O4/rGO composite on Ni foam [35]; (A) and (B) SEM images of the composite on Ni foam, and (C) TEM image of the composite.

Chemical properties and applications 265 (B)

0.1 A/g

Mn3O4/rGO

100 mA/g

0.2 A/g 0.5 A/g 1 A/g 3 A/g 5 A/g

Capacity / mAh/g

Capacity / mAh/g

(A)

Mn3O4

Fe3O4/rGO

0.1 A/g

100 mA/g

0.2 A/g 0.5 A/g 1 A/g 3 A/g

5 A/g

Fe3O4

Cycle number

Cycle number

Figure 4.15 Rate capabilities for (A) Mn3O4/rGO and (B) Fe3O4/rGO composites, in comparison with Mn3O4 and Fe3O4 on Ni foam [35].

(B) 2nd 1st

Voltage / V

50th

Discharge capacity / mAh/g

(A)

SnO2/rGO

SnO2 nanoparticles

rGO Capacity / mAh/g

Cycle number

Figure 4.16 LIB performances of SnO2/rGO composite [36]; (A) discharge/charge curves and (B) cycle performance at 55 mA/g rate with the voltage range of 0.01e3.0 V in comparison with SnO2 nanoparticles and rGO.

shown in Fig. 4.16A, and a large irreversible capacity in the first cycle. However, the composite gives much improved cycle performance in comparison with the pristine rGO and higher reversible capacity as 520e570 mAh/g even after 100 cycles than SnO2 nanoparticles, as shown in Fig. 4.16B [37]. The composite of rGO nanoflakes with SnS2 nanoplatelets (diameter of about 100 nm) was fabricated via the reaction of SnO2 loaded on rGO with H2S gas, and was shown to have reversible capacity of about 650 mAh/g even after 30 cycles at 50 mA/g rate [38]. It delivered the reversible capacity of 230 mAh/g at a high discharge/charge rate of 6.4 A/g. SnO2/rGO composite particles were loaded on SiO2 spheres (about 100 nm diameter) under hydrothermal condition at

266 Chapter 4 150 C [39]. The composite delivered first discharge and charge capacities of 1548 and 820 mAh/g, respectively, and relatively stable reversible capacity of 715 mAh/g by retaining 580 mAh/g after 100 cycles with 100 mA/g rate. The composites of rGO with rutile-type and anatase-type TiO2 were prepared and the anodic performance in LIB was tested [40]. Carbon-coated Sn/rGO composite was prepared from SnO2/rGO composite by coating glucose and then carbonized at 500 C in Ar, followed by reduction of SnO2 to Sn in a flow of H2/Ar (1/19) at 550 C, and its LIB performance was studied in 1 M LiPF6/ (EC þ DEC) electrolyte [41,42]. When Sn content in the composite was less than 56.5 wt%, average particle size of Sn encapsulated between rGO nanoflakes was about 15 nm, but Sn became discoidal particles with the thickness of less than 50 nm when the Sn content was more than 56.5 wt%. The composite containing Sn nanoparticles exhibits a reversible capacity of 727 mAh/g for the first cycle and 645 mAh/g for the 100th cycle in the voltage range of 0.01e2.0 V, while a serious capacity fading occurs up to 30 cycles for the composite containing discoidal particles of Sn, probably due to the mild pulverization of the Sn discs. By enlarging the voltage range to 0.01e3.0 V, the capacity of the composite can reach as high as 777 mAh/g after 100 cycles, as shown in Fig. 4.17A. The rate performance shown in Fig. 4.17B demonstrates that the composite delivers a capacity of about 270 mAh/g even at a high current density as 3.2 A/g. Sn/rGO composite was also fabricated via the reduction of dispersed rGO nanoflakes in SnCl2 aqueous solution by NaBH4 at 0 C, which delivered the reversible capacity of 508 mAh/g after 100 cycles at 55 mA/g rate with the voltage range of 0.01e3.0 V in 1 M LiPF6/(EC þ DMC) electrolyte [43]. Sn nanoparticles (5e30 nm size) encapsulated by rGO shells loaded on rGO foam prepared by using NaCl crystalline particles as template exhibited a high rate performance in 1M LiPF6/(EC þ DMC þ DEC) electrolyte, 1022 mAh/g at 0.2 A/g, 780 mAh/g at (B) Discoidal Sn nanoparticles (0.01-2 V) Sn nanoparticles (0.01-2 V) Sn nanoparticles (0.01-3 V)

Capacity / mAh/g

1500 1200 900 600 300

Charge Discharge

1500 Capacity / mAh/g

(A)

1200 900 600 300 0

0 0

20

60 40 Cycle number

80

100

0

40 20 Cycle number

60

Figure 4.17 Carbon-coated Sn/rGO composite [41]; (A) discharge/charge curves and (B) rate capability with the voltage range of 0.01e2.0 V.

Chemical properties and applications 267 1 A/g, and 270 mAh/g at 10 A/g [44]. Sn nanoparticles were also loaded on N-doped rGO through hydrazine vapor reduction of the mixture of GO and SnO2 [45] and thermal reduction of the mixture at 500 C in H2/Ar atmosphere [46]. Sn/graphene composites were fabricated through deposition of Sn nanoparticles from its vapor onto a graphene foam prepared by the deposition on Ni foam from methane in hydrogen plasma, which delivered a reversible capacity of 466 mAh/g with 879 mA/g rate after over 4000 cycles and 794 mAh/g at 293 mA/g rate after 400 cycles [47]. Composites of metallic Si nanoparticles with graphene and rGO have been prepared by different techniques, a simple mechanical mixing [48], mixing suspended Si nanoparticles with rGO nanoflakes in water by adding NaOH and pyrenebutyric acid (C20H16O2) [49], combining electrophoretic deposition with radiofrequency magnetron deposition [50], anchoring of graphene-encapsulated Si on graphene trees using a microwave plasma enhanced CVD of CH4/H2 and SiH4/H2 mixed gases [51], ball-milling the mixture of coarse-grained Si, WC and graphite [52], and carbonizing the mixture of ball-milled Si and graphite with sucrose at 700 C [53]. Si/rGO composite was prepared by simple mechanical mixing of commercially available nanosized Si (about 40 nm size) and rGO nanoflakes in the weight ratio of 1/1, and its electrochemical performance was tested [48]. In Fig. 4.18A, discharge/charge curves for the Si/rGO composite at a current density of 100 mA/g at the voltage range of 0.02e1.2 V are shown; the initial discharge capacity is 2158 mAh/g, retaining 1168 mAh/g after 30 cycles. The cycle performance for the Si/rGO composite is compared to those for nano-sized Si powder and the pristine rGO in Fig. 4.18B, revealing that the composite shows an enhancement of the capacity retention. As can be seen from Fig. 4.18, it is interesting that Si/rGO composite does not give a large irreversible capacity, although both rGO and nano-sized Si give large irreversible capacities. Graphene-encapsulated Si nanoparticles anchored on graphene, vertically (A)

(B) Capacity / mAh/g

2.0

Voltage / V

1.5 30th

1st

1.0 0.5 0.0

30th 0

3000

2000 Si/rGO 1000 Pristine rGO

1st

1000 2000 Capacity / mAh/g

3000

0

0

Nano-sized Si

10 20 Cycle number

30

Figure 4.18 Si/rGO composite [48]; (A) discharge/charge curves and (B) cycle performances with a current density of 100 mA/g in the voltage range of 0.02e1.2 V.

268 Chapter 4 (A)

(B)

800

2.5 Capacity / mAh/g

2nd 1st Voltage / V

2.0 1.5

Charge

1.0 0.5

2nd

1st

Charge Discharge

600

0.1 A/g

400 200

Discharge

0.0 0

400 200 Capacity / mAh/g

600

0

0

20

40

60 80 100 120 140 160 Cycle number

Figure 4.19 LIB performance of carbon-coated Sn/rGO composite in 1 M NaClO4/(EC þ PC) [42]; (A) discharge/charge curves for first two cycles, and (B) rate performance.

aligned graphene trees, were fabricated, which exhibited high rate performance, reversible capacity of 1528 and 412 mAh/g at 0.15 and 8 A/g rate, respectively [51]. rGO nanoflakes were also tested as the cathode of LIBs. rGO functionalized by ethylene glycol was prepared by stirring GO suspension in ethylene glycol at 120 C, followed by precipitation in water, freeze-drying, and heat-treatment at 900 C [54], where ethylene glycol worked as both solvent and reducing agent. This functionalized rGO delivered discharge capacity of 280 mAh/g at 0.05 A/g without obvious capacity fading after 300 cycles in 1 M LiPF6/(EC þ DMC þ EMC) electrolyte as the cathode. Carbon-coated Sn/rGO composite had also high storage capacity for Naþ [42]. Discharge/ charge curves in 1 M NaClO4/(EC þ PC) electrolyte at the first two cycles with the current density of 0.1 A/g are shown in Fig. 4.19A, the first discharge curve exhibiting a short plateau at 0.5e0.9 V probably due to the formation of SEI films. Fig. 4.19B demonstrates its rate capability for Na storage; reversible capacity as high as 434 mAh/g at the current density of 0.1 A/g decreases with increasing current density, reaching 106 mAh/g at 3.2 A/g. However, it is recovered to 420 mAh/g by setting back the current density to 0.1 A/g even after the high current cycling. rGO-wrapped FeS2 composites delivered a high discharge capacity of around 700 mAh/g at 100 mA/g with a high retention as 609 mAh/g after 100 cycles in 1 M NaClO4/(EC þ PC) with 5 wt% fluoroethylene carbonate [55].

4.2.2 Electrochemical capacitors Graphene has been studied as electrode material of electrochemical capacitors in both aqueous and nonaqueous electrolytes. However, it is very regrettable to write here that gravimetric capacitance (often employed in the research works) of two-electrode cells

Chemical properties and applications 269 have to be calculated on the basis of the total weight of electroactive materials, including graphene, but in some papers (not rare) it has been calculated by using the weight of a single electrode, and also in some papers no clear description on the capacitance calculation is presented by just emphasizing that the large capacitance values they obtained. It is very risky to compare the capacitance values reported in different papers directly. Here, therefore, we avoid comparing the values themselves, except the comparison among the values measured in the same paper. The dispersed solution of GO was sonicated until it became clear, and then hydrazine monohydrate was added to reduce GO to rGO precipitates. Then the rGO particles were formed into films by using polytetrafluoroethylene (PTFE) binder to construct twoelectrode supercapacitor cells [56]. The cells exhibit nearly rectangular cyclic voltammograms both in aqueous and nonaqueous electrolytes, as shown in Fig. 4.20, indicating the formation of electrical double layers on the electrodes, and the capacitance in 5.5 M KOH being about 128 F/g and that in 1 M tetraethylammonium tetrafluoroborate (TEABF4)/AN about 95 F/g with the current density of 20 mA/g. rGO prepared by thermal exfoliation and reduction of GO at 1050 C, which was composed of three to six layers, delivered a capacitance of 117 F/g, and graphene nanoflakes of three to six layers synthesized from nanodiamond at 1650 C gave the capacitance of 35 F/g in 1 M H2SO4 with 100 mV/s rate [57]. In an ionic liquid, N-butyl-N-methyl-pyrrolidinium bis(trifluoromethanesulfonyl)imide (PYR14TFSI), as electrolyte, the capacitance of 75 F/g was obtained for thermally processed rGO flakes, but the energy density was as high as 31.9 Wh/kg because the working potential increased to 3.5 V, instead of 1 V in H2SO4. By using an ionic liquid, 1-ethyl-3-methylimidazolium tetrafluoroborate (EMIM.BF4), curved flakes of rGO, which was obtained by the injection of GO suspension into a stream of compressed air (a fluidized-bed situation), exhibited energy density of 85.6 Wh/kg at room temperature and 136 Wh/kg at 80 C at a current density of 1 A/g [58]. This high

(B)

200 100 0

-0.1

0.1

0.3

0.5 Volts

0.7

0.9

1.1

-100

-200

Specific capacitance / F/g

Specific capacitance / F/g

(A)

200 100 0

-0.1

0.3

0.7

1.1 1.5 Volts

1.9

2.3

2.7

-100

20 mV/s 40 mV/s

20 mV/s 40 mV/s

-200

Figure 4.20 Cyclic voltammograms with two different scan rates for rGO [56]; (A) in 5.5 M KOH and (B) 1 M TEABF4/PC.

270 Chapter 4 (A)

(B) Specific capacitance / F/g

1.2

Voltage / V

1.0 0.8 0.6 0.4 0.2 0.0

0

1600

3200 4800 Time / s

6400

8000

240 210 180 150 120 90 60 30 0 0

200

400 600 800 Cycle number

1000 1200

Figure 4.21 EDLC performances of rGO in 30 wt% KOH [59]; (A) galvanostatic charge/discharge curves with 100 mA/g and (B) cyclic performance with 500 mA/g.

performance was supposed to be due to curved rGO flakes to avoid their restacking and give a space for accepting large ionic liquid molecules, in addition to a high potential window with ionic liquid. rGO reduced by hydrazine vapor exhibits high EDLC performance in 30 wt% KOH aqueous electrolyte, a maximum capacitance of 205 F/g with a power density of 10 kW/kg at energy density of 28.5 Wh/kg and a long cycle life along with about 90% retention after 1200 cycles, as shown in Fig. 4.21 [59]. Exfoliation and reduction of GO was possible at low temperatures of 200e400 C under a high vacuum (700 nm TiO2, λ>450 nm

0.8

0.2

Composite, λ>600 nm

0.6

Mechanical mixture, λ>450 nm Composite, λ>450 nm

0.4

TiO2

0.0 200

1.0

C/C0

Absorbance / a.u.

(A)

TiO2 (P25)

300

400 500 600 Wavelength / nm

700

800

0.2

0

50

100 Time / min

150

200

Figure 4.105 TiO2-loaded rGOs (composites) with different TiO2/rGO ratios [326]; (A) reflectance spectra under the irradiation of a 125 W mercury lamp, and (B) photodegradation of MB by the composite (TiO2/rGO of 3/1) under the irradiation with different wavelengths, in comparison with TiO2, laboratory-made and commercially available TiO2 (P25), and mechanical mixture of rGO with P25.

In some literature dealing with photocatalytic activity of rGO-based composites, such as TiO2- and Au-loaded rGOs, the authors claimed that the composites are photoactive under visible light based on experimental results under the irradiation of the light with l > 400 nm. Fig. 4.105A and B demonstrate the reflectance spectra and MB degradation curves for the composite with TiO2/rGO ratio of 3/1 under the lights with different wavelengths [326]. The composite absorbs the UV-Vis light but cannot degrade MB under the light with >700 nm. If a high adsorptivity of rGO matrix for MB is taken into consideration, the photoactivity of this composite has to be much smaller than that expected from the observed change in MB concentration that remains (Fig. 4.105B). Therefore, the photoactivity under visible light of the rGO-based composites is strongly required to be studied in more detail based on more exact experiments by taking degradation and adsorption into consideration. Hydrogen evolution rate via water splitting by photocatalysis of TiO2-loaded rGO was markedly enhanced by codeposition of CuO with TiO2 on rGO surface [327]. The composites, CuOeTiO2/rGO and TiO2/rGO, were prepared by hydrothermal treatment of the mixture of GO, TiO2 (P25) with and without Cu(NO3)2 at 180 C. Hydrogen evolution by the composite was performed in 450 mL aqueous solution containing 10 vol% methanol under the irradiation of UV-Vis light with magnetic stirring. Hydrogen evolution was accelerated by adding CuO, as shown in Fig. 4.106A and B. CuO content in CuOeTiO2 was kept at 2 wt%, because hydrogen evolution rate did not change much even with adding more than 2 wt% CuO.

340 Chapter 4 (B) 3000

H2 evolution rate / μmol/g h

H2 evolution rate / μmol/g h

(A) 250

2500

200

2000

150

1500

100

1000

50 0

0

1

3 4 2 rGO content / wt%

500

5

0 0

1

3 4 2 rGO content / wt%

5

Figure 4.106 Hydrogen evolution rate of CuOeTiO2/rGO composites as a function of rGO content [327]; (A) TiO2/rGO and (B) CuOeTiO2/rGO.

H2 evolution / μmol/h

CdS-loaded rGO (CdS/rGO) and MWCNT (CdS/MWCNT) were prepared by hydrothermal treatment of the suspension of either GO or MWCNT with cadmium acetate and Na2S at 180 C [328]. Two composites, CdS/rGO and CdS/MWCNT (in 1/0.01 in mass), exhibit high evolution rate of H2 from the 100 mL aqueous solution containing 0.1 M Na2S and 0.05 M Na2SO3, using 0.10 g catalyst and light from Xe lamp (l > 420 nm), their activities being superior to CdS itself and mechanical mixtures of CdS with rGO, as shown in Fig. 4.107. Loading of MoS2 on N-doped rGO was prepared by

Catalyst

Figure 4.107 H2 evolution rate for CdS/rGO and CdS/MWCNT composites, in comparison with other photocatalysts [328].

Chemical properties and applications 341 hydrothermal treatment of the mixture of GO, (NH4)2MoS4, polypyrrole with hydrazine in DMF at 180 C, followed by the heat treatment at 600 C [329]. The composite exhibited efficient electrocatalytic activity toward hydrogen evolution reaction under acidic conditions with a low-onset potential of 112 mV. The experimental results as well as theoretical calculation by DFT demonstrated that the abundance of exposed active sulfur edge sites in the MoS2 and N-containing active functional moieties in N-doped rGO are synergistically responsible for the catalytic activity, whilst the distinguished and coherent interface between MoS2 nanoparticles and N-doped rGO surfaces facilitates the electron transfer during electrocatalysis. On TiO2/rGO composite, the photogenerated electrons by UV irradiation of TiO2 nanoparticles are transported across rGO to reduce Ag ions into nanoparticles at a location distinct from the TiO2-anchored site [330]. This result may pave the way for the development of new rGO-based catalysts, where selective anchoring of semiconductor and metal nanoparticles at separate sites on two-dimensional sheets of rGO act as a support. PteZnO/rGO composite was synthesized under hydrothermal conditions, where weakly basic Zn2þ ions were converted to ZnO, accompanying synchronous reduction of K2PtCl4 and GO without additional reducing agent [331]. PdeZnO/rGO composite was also obtained with the same procedure. ZnO deposited on the surface of rGO nanoflakes offers homogeneous nuclei to anchor metal particles, leading to the uniform dispersion of Pt and Pd nanoparticles. Average diameter of the deposited nanoparticles was 3.0 nm for Pt and 5.5 nm for Pd. These composites had much higher catalytic activity and stability for methanol oxidation reaction than commercially available catalysts.

4.3.5 Chemical sensors Graphene has been known to offer important advantages to the sensing devices: higher sensitivity, lower electrical noise at room temperature, and, as a consequence, capability of the detection in ppm level, even ppb level, of gas species at room temperature and atmospheric pressure, in comparison with other sensing materials, such as carbon nanotubes. In addition, graphene devices offer relatively easier electrical contact, easier manipulation, are more cost effective, and easier to be produced in mass. A number of reviews of these graphene-based sensors have been published [332e336]. The Hall resistivity of the single-layer graphene obtained by cleavage was measured during the adsorption and desorption of strongly diluted NO2 and is shown in Fig. 4.108 [337]. Each step in resistivity change corresponds to the addition and removal of a single electron, indicating that the response observed on graphene-based sensor is owing to adsorption and desorption of individual gas molecules.

342 Chapter 4

Adsorption Change in resistivity / Ω

30

20 Desorption 10 Reference 0

0

200

400 Time / s

600

Figure 4.108 Resistivity change observed near the neutrality point on a single-layer graphene with adsorption and desorption of NO2 molecules [337].

Adsorption of O2 molecules on double-layered graphene was shown to give strong influences on its field effect transistor (FET) performances [338]. O2 adsorption causes the hole doping, depending remarkably on the gate voltage, and the rate of the adsorption is determined by the Fermi level of the graphene. Micrometer-sized sensors made from graphene are capable to detect individual gas molecules that attach to or detach from the graphene surface with clear change in the local carrier concentration in graphene. As shown in Fig. 4.109A for NO2 gas, charge carrier density in single-layer graphene changes with the change in the concentration of NO2 [337]. The gas-induced changes in resistivity have different magnitudes for different gases, and the sign of the resistivity change indicates whether the gas is an electron acceptor (e.g., NO2, H2O, iodine) or an electron donor (e.g., NH3, CO, ethanol), as shown in Fig. 4.109B. Patterning of single-layer graphene grown by CVD of ethanol into nanomesh by using a combination of lithography and reactive ion etching was effective to enhance the sensitivity toward NO2 and NH3 at room temperature [339]. The array of the films at 10 mm wide in 10 mm gap prepared from hydrazine-reduced single-layer rGO by spin coating on Si/SiO2 substrate was studied on its sensing performance [340]. The performance of the array for 5 ppm NO2 is shown in Fig. 4.110, revealing the trend from larger and slower response at room temperature to smaller and faster response at higher temperature. rGO contains many functional groups and defects, and so can offer great potential for practical development in graphene-based sensors because it is easy and cheap to

Chemical properties and applications 343 (A)

(B)

50

Vacuum Exposure to gas Evacu- Annealingo with 1 ppm ation at 150 C

NH 3

20 2 10

Δ R/R0 / %

Charge carrier density / 1010/cm2

4

5

H2O

2

-2

1

-4 0.1

1 NO2 concentration / ppm

CO

0

NO 2 0

10

500 Time / s

1,000

Figure 4.109 Sensitivity of single-layer graphene for gas adsorption [337]; (A) change in charge carrier density in graphene with the concentration of NO2 and (B) changes in relative resistivity DR/R0 with exposure to different gases, evacuation, and annealing.

5 ppm NO2

149 oC

5 ppm NO2

1.0

Relative resistance R/R0

124 oC 102 oC

0.9

66 oC

0.8 21 oC

0.7 0

60

180 120 Time / min

240

300

Figure 4.110 Sensing performance of the array of single-layer rGO film for NO2 at different temperatures [340].

produce in large scale, readily functionalized, and able to tune its band gap, in comparison with graphene nanosheet synthesized by CVD process, which is not producible in large scale, with no functional groups and no band gap [334]. By using an inkjet printing technique, thin film of rGO reduced by using ascorbic acid was printed onto PET substrate, which exhibited high-sensing ability (sensitivity) for

344 Chapter 4 (A)

(C)

1.06

1.08

1 ppm 2 ppm

4 ppm

Ids / μA

Ids / μA

8 ppm

Pristine graphene NO2

1.04

20 ppm

1.02

1.07

Pristine graphene NH3

20 ppb 1 ppb

0

2 ppb

1000

1.00

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2000 Time / s

0

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(D)

0.88

20 ppb

0.84

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3000

4 ppm

1.12

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1.08

B-doped graphene NH3

8 ppb 1 ppb

2 ppm

2000 Time / s

8 ppm

B-doped graphene NO2

0.80

1000 1 ppm

1.16

Ids / μA

Ids / μA

(B)

4 ppb

4 ppb

2000 Time / s

1.04 3000

0

1000

2000 Time / s

3000

Figure 4.111 Sensor responses of pristine and B-doped graphene films for NO2 and NH3 at a fixed drain to source voltage (Vds) under UV illumination [342].

different gas molecules, for NO2 the concentration down to 500 ppb, in addition to the advantages of light weight, free-standing, and flexible [341]. B-doped graphene film synthesized by CVD of the mixture of triethylborane/hexane at 1000 C exhibits high sensitivity for toxic gases in comparison with pristine graphene film, as shown for NO2 and NH3 in Fig. 4.111 [342]. For NO2, B-doped graphene gives a clear signal with high signal-to-noise ratio (s/n) of 31 even at the concentration of 1 ppb, although pristine graphene can detect 8 ppb with s/n of 9.4. For NH3, B-doped graphene is much more sensitive than pristine graphene, the former detecting at 1 ppm with s/n of 50 but the latter at 20 ppm with s/n of 9.5. FET based on rGO film was successively used as chemical sensors for gas molecules. rGO film deposited on SiO2/Si substrate (back gate) from the suspension of mostly single-layer flakes was settled in a FET device, bridging the source- and drain-electrodes of Au as the conducting channel [343]. As shown in Fig. 4.112, the sensor exhibits an instantaneous decrease in resistance by flowing NH3 and fast recovery in resistance by stopping NH3 flow under a positive gate potential (40e10 V), far superior to the performance at

Chemical properties and applications 345 Gate potential:

Positive

Negative

60

0 -20

50

ON

-40 OFF 45

0

50

100

150

Gas flow

Resistance / kΩ

20 55

Gate potential Vg / V

40

200

Time / min

Figure 4.112 Performance of FET device based on rGO film under a flow of 1% NH3/air [343].

zero/negative gate potential. FET based on graphene grown epitaxially on 6HeSiC substrate exhibited high sensitivity for pH change because both hydroxyl (OH) and hydroxonium (H3Oþ) ions were able to modulate the channel conductance by doping holes and electrons, respectively [344]. The application of negative gate potential, which induces the accumulation of OH ions on the surface, produces a larger increase in conductivity, compared to the positive gate potential (accumulation of H3Oþ), for two graphene sheets composed of one to two and three to four layers, the relation between threshold voltage and pH being approximated to be linear with the slope of 98e99 mV/pH unit. Pyrrole-reduced rGO gave detectable response even for the 1 ppb NH3 gas [345]. The film of rGO was covered by porous conducting polymer poly(3,4ethylenedioxythiophene), PEDOT, (porous PEDOT/rGO composite) exhibited a highsensing performance for NH3 gas [346], as shown in Fig. 4.113, a high-response level even after repeating the exposure/recovery cycles and also a fast response and recovery even at ppb level of NH3 gas. This porous PEDOT/rGO composite showed high-sensing performance also for different gases, such as H2S, SO2, CH2Cl2, and CH3OH, much better than the pristine rGO film. The composites of rGO with poly(sodium 4-styrenesulfonate) (PSS) and poly(allylamine hydrochloride) (PAH), which were prepared from the mixed solutions of GO with PSS and PAH by hydrogen reduction at 100 C, and exhibited high sensitivity for NO2 gas with the concentration of 1e15 ppm [347]. Single-layer graphene film prepared on SiO2 substrate at 650 C by inductively coupled plasma CVD was patterned by laser interference lithography technique to the periodic array of graphene ribbons with the width of 200 nm at a pitch of 1 mm [348]. Hydrogen gas sensing performance of this graphene ribbon array was studied after the deposition of

346 Chapter 4

Figure 4.113 Porous-PEDOT/rGO composite film [346]; (A) repeated sensing for 5 ppm NH3 gas, and (B) sensitivities to different concentrations of NH3 gas.

Pd metal layer in 2 nm thickness. The array composed of 1000 graphene ribbons showed instant response to hydrogen, and quick and complete recovery in nitrogen atmosphere, although the pristine film showed a little slower response to hydrogen on and off, as shown in Fig. 4.114. In the array, the graphene ribbons work as a support for homogeneous deposition of Pd nanoparticles and an electrically conductive path with less electrical noise. Resistance change DR/R0 (sensitivity) for the array is much smaller than the pristine film because of the reduced number of Pd-nanoparticles deposited on the narrow ribbons. The sensors based on SnO2-loaded rGO have a sensitivity for H2 of a low concentration as 0.5%, but additional loading of Pt on SnO2/rGO composite by the reduction of H2PtCl6 in

(A)

(B) 30

H2 ON OFF

6

Δ R/R0 / %

Δ R/R0 / %

8

4 2

20

10

0

0 0

15

45 60 30 Time / min

75

90

0

15

30 45 60 Time / min

75

90

Figure 4.114 Hydrogen-sensing performances of single-layer graphene [348]; (A) Pd-decorated graphene ribbon array and (B) pristine graphene film.

Chemical properties and applications 347 1000

SnO2/rGO 0.5 %

1%

Resistance / kΩ

100

2%

3%

Pt-SnO2/rGO

10 0.5 % 1%

1

2%

3%

0.1 0

100

200

500 400 Time / s

300

600

700

800

Figure 4.115 Resistance changes by the pulses of hydrogen with different concentrations for the sensors based on SnO2/rGO and PteSnO2/rGO composites [349].

ethylene glycol at 150 C gives more enhanced sensitivity, as shown in Fig. 4.115 [349]. PteSnO2/rGO sensor was possible to respond in 3e7 s and to be recovered during 2e6 s, depending on the H2 concentration. Hydrazine-reduced rGO nanoflakes deposited on 36 YX LiTaO3 surface acoustic wave (SAW) transducers were prepared, of which gas-sensing performance was assessed by using H2 and CO in air at 25 and 40 C [350]. The dynamic responses of the composites are shown in Fig. 4.116 as the variation in operating frequency of oscillation due to the interaction with the gas. The responses for 1% H2 and 1000 ppm CO are approximately (A)

(B)

60 ppm 125 ppm 250 ppm 500 ppm

102.075

1000 ppm

at 25 oC 102.070 0.06 % 0.125 % 0.25 %

0.5 %

Frequency / MHz

Frequency / MHz

102.205 at 25 oC

1%

101.946 at 40 oC 101.945

102.200

at 40 oC

102.120

102.115 0

2000

Time / s

4000

0

2500

5000

7500

10000

Time / s

Figure 4.116 Dynamic response of rGO nanoflakes/LiTaO3 SAW sensor at 25 and 40 C [350]; (A) H2 and (B) CO with different concentrations.

348 Chapter 4 5.8 and 8.5 kHz at 25 C, but 1.7 and 7.0 kHz at 40 C, respectively. For the low concentration of CO (60 and 125 ppm), the frequency change is almost insignificant. The 90% response time is 116 and 400 s at 25 C, but 12 and 300 s at 40 C for 0.125% H2 and 250 ppm CO, respectively. The recovery time for CO was almost 30 and 20 min at 25 and 40 C, while that for H2 was less than 10 and 1 min at 25 and 40 C, respectively. A microelectrode sensor composed of the ODA-functionalized rGO film drop-casted onto a glass substrate was able to detect benzene, toluene, ethylbenzene, xylenes, and cyclohexane dissolved in water at low concentrations of 5e100 ppm [351]. For graphene films synthesized by CVD and cleavage, their functionalization was performed through p-p interaction. Fig. 4.117A and B shows the fabrication process of FET using flexible graphene aptasensor and interaction of Hg2þ in solution with aptamer on the surface of the graphene, respectively [352]. Single-layer graphene film was

Figure 4.117 Flexible graphene-based aptasensor on PEN film for Hg2þ sensing [352]; (A) fabrication process of the sensor through functionalization using p-p interaction and (B) capturing and detecting Hg2þ.

Chemical properties and applications 349 synthesized by CVD of CH4 at 1000 C, transferred onto a flexible polyethylene naphthalate (PEN) film and then settled between electrodes (source and drain) to construct FET. 1,5-Diaminonapthalene (DAN) was adsorbed on the graphene surface using p-p interaction, followed by immobilization of a DNA aptamer (30 -amine-TTC TTT CTT CCC CTT GTT TGT-C10 carboxylic acid-50 ) for capturing Hg2þ. This graphene aptasensor had excellent sensing performance for Hg2þ ions; detection limit 10 p.m., two to three orders of magnitude more sensitive than other mercury sensors, and high specific sensitivity to Hg2þ in mixed solutions. Si-doped single-layer and double-layer graphene films exhibit a significant grapheneenhanced Raman scattering effect for fluorescent organic molecules, such as crystal violet, rhodamine B, and methylene blue [353]. The effect is schematically illustrated in Fig. 4.118A, and Raman signals of a dye molecule, crystal violet, were shown as an example in Fig. 4.118B, revealing clear signals of crystal violet even at concentration as low as 5  105 M. Si-doped graphene can serve as a good substrate for enhancing the Raman scattering signals of organic molecules by quenching the fluorescence, much clearer than graphene without doping. N-doping is also effective to enhance the effect. N-doped graphene has high sensitivity for different molecules via graphene-enhancing Raman scattering [354]. The detection limits were reported to be 5  1011, 1  108, and 1  108 M for rhodamine 6G, protoporphyrin, and copper phthalocyanine, respectively. Oxygen functionalization was shown to be effective to increase the resistivity changes due to the adsorption of polar gas molecules by using double-layered graphene grown by CVD epitaxially on 6HeSiC(0001) substrate [355]. The functionalization of graphene was performed by using electron beamegenerated 5% O2/Ar plasma at room temperature.

(B)

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* Crystal violet G G’

Intensity / a.u.

Dye molecule

Si-doped graphene

N-doped graphene

Graphene without doping

1000

1400

1800 2200 2400 Raman shift / cm -1

2600

Figure 4.118 Enhanced Raman scattering effect of Si-doped graphene [353]; (A) schematic illustration of Si-doped graphene and different probe molecules, and (B) Raman signal of crystal violet on graphene films after immersing its solution with 5  105 M concentration.

350 Chapter 4 (A)

(B) NMP DMSO DMF DMAC

Δ R/R0 / %

DMSO DMF DMAC

AA H2O2

Δ R/R0 / %

NMP

AA EG

EG H2O2

Dipole moment / D

Dipole moment / D

Figure 4.119 Relations between DR/R0 and dipole moment of adsorbate gas molecules for double-layered graphene [355]; (A) nonfunctionalized (as-prepared) and (B) O2-functionalized graphenes.

Resistance change is observed more markedly on O2-functionalized graphene than on asprepared one, as shown in Fig. 4.119, the change in resistivity DR/R0 on O2-functionalized graphene being linearly related to dipole moment of adsorbate gas molecules; positive DR/R0 upon exposure to polar protic molecules of hydrogen peroxide (H2O2) and ethylene glycol (EG), and negative DR/R0 upon exposure to polar aprotic vapors of DMF, dimethylacetamide, n-methyl-2-pyrrolidone, and acetic anhydride. Although the average response time was about 10 s for both as-prepared and O2-functionalized graphene sensors, the average recovery time was about 100 s for the latter, but 1.5e2 h for the former. Polyaniline-loaded rGO exhibited high sensitivity for NH3 gas, much higher sensitivity than pristine rGO and polyaniline [356]. The composite was prepared by mixing MnO2-loaded GO with aniline monomers in solution, where MnO2 worked as the template and oxidant for aniline monomers during polymerization, followed by washing out of H2O2 by a diluted HCl after the polymerization of aniline. Carbon fiber electrode modified by rGO nanoflakes and rGO flowers was reported to work as sensor for the detection of ascorbic acid (AA), dopamine (DA), and uric acid (UA) [357]. As shown on DA in Fig. 4.120A, its oxidation and reduction peaks are clearly detected by rGO-flower-decorated carbon fiber electrode, much clearer than the pristine carbon fiber electrode and the reference glassy carbon electrode. Even in the mixed solution of AA, DA, and UA, three molecules are separately detected by using rGOflower-decorated carbon fibers, although other two electrodes, pristine carbon fibers and glassy carbon, give only overlapped broad oxidation peak, as shown in Fig. 4.120B. Metal-oxide-loaded graphene or rGO composites were proposed for highly sensitive, selective, and cost-effective gas sensing, which can operate at room temperature and apply to environmental protection by detecting toxic gases [334]. Graphene-based

Chemical properties and applications 351 (A)

(B) rGO-flower-decorated carbon fiber electrode

rGO-flower-decorated carbon fiber electrode

480 360

70

0 Carbon fiber electrode Glassy carbon electrode

-70

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Current / μA

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0.6 0.4 0.2 0.0 Potential / V vs. SCE

0.8

-0.2

0.6 0.4 0.2 0.0 Potential / V vs. SCE

0.8

Figure 4.120 Cyclic voltammograms of rGO-flower-decorated carbon fibers, pristine carbon fibers and glassy carbon [357]; in 0.1 M phosphate buffer solution of (A) 0.5 mM DA and (B) a mixture of AA, DA, and UA. Scanning rate was 50 mV/s.

composites with various metal oxides, ZnO, Cu2O, SnO2, WO3, Co3O4, NiO, In2O3, etc. exhibited high sensitivity for different gases [336]. ZnO-loaded rGO was prepared by simultaneous reduction of GO and Zn(AC)2 by diluted LiOH in methanol at room temperature, which showed enhanced sensing for CO gas [358]. The composite of rGO loaded by Cu2O nanowire mesocrystals was fabricated under hydrothermal conditions [359]. The mesocrystals were composed of highly anisotropic nanowires and have a distinct octahedral morphology. Owing to high specific surface area and improved conductivity, the Cu2O/rGO composites achieved high sensitivity toward NO2 at room temperature. The composites were fabricated from vertically aligned ZnO nanorods grown on ZnO conductive layer by covering graphene film grown by CVD of CH4 at 900e1000 C [360]. The composite demonstrates high sensitivity for ethanol vapor at 300 C, as shown in Fig. 4.121, possible to detect ethanol vapor of ppm-level concentration with a high sensitivity, the ratio of resistance in air to that in ethanol vapor being about 9 for 10 ppm ethanol. Vertically aligned SnO2 nanorods on both surfaces of CVD-grown graphene film were prepared by hydrothermal treatment, which exhibited improved sensing performances to various gases, particularly H2S [361]. The sensitivity for acetone vapor of rGO thin films was shown to depend on the degree of reduction by using hydrazine hydrate vapor at 100 C [362]. On this rGO sensor, accumulation of acetone vapor was observed during repetitive exposure to the vapor, which needed heating to complete recovery, although the sensor composed of SWCNTs showed almost no accumulation. Minimum detectable level of the rGO sensor was similar to that of SWCNT sensor for warfare chemicals (chloroethylethyl sulfide CEES and

352 Chapter 4 (A)

(B)

8

80 Sensitivity

100

Conductance / μS

10

6 4 2 0

60 40 20

0

50

100 Time / min

150

0

10

20 30 40 50 Ethanol concentration / ppm

Figure 4.121 ZnO-nanorod/graphene composite sensor [360]; (A) conductance change at different concentrations of ethanol vapor at 300 C and (B) change in sensitivity (ratio of resistance in air to that in ethanol) with ethanol concentration.

dimethyl methylphosphonate DMMP) and an explosive (dinitrotoluene DNT). For HCN gas, however, minimum detectable level of the rGO sensor was as small as 70 ppb but that of the SWCNT sensor was more than 4000 ppb. The sensor based on rGO reduced by p-phenylenediamine could achieve larger responses to DMMP with 30 and 10 ppm concentration, compared with the sensor based on hydrazine-reduced rGO [363].

4.4 Concluding remarks Possible applications for graphene and rGO based on their chemical properties are of a wide range as described in the previous sections and are divided into two fields, energy storage and conversion and environmental remediation. In each field, graphene is used as various morphologies, such as films, flakes, porous aggregates, etc. and with various modifications, such as doping of N and B, functionalization by organic radicals, forming as composites with other materials, etc. In many cases in applications, p-electron cloud extended on the graphene layer has a certain role for the development of the functionalities in each application. For example, the interaction of p-electron cloud of graphene layer with cation of the adsorbates (p-cation interaction) gives new possibility for the enhancement of the performances of electrochemical capacitors, capacitive deionization devices, and adsorption removal of cationic pollutants. In addition, the interaction of p-electron cloud with p-electron in the adsorbates (p-p interaction) is also important for the performance of adsorption removal of polluting organics, such as dyes. These functions of p-electron cloud have not been recognized clearly on graphite-based carbon materials, such as porous carbons, although all of them contain p-electrons in their basic structural units, crystallites, which are composed of stacks of multilayers of graphene. By stacking graphene layers, p-electron clouds of neighboring layers interact

Chemical properties and applications 353 Na+

Na+

Single-layer graphene

Na+

Graphene sheet Na+

Na+

Na+

Na+

Na+

Na+

Na+

Double-layered graphene Na+

Na+

Multi-layered graphene

Na+ Graphene layer

π-electron cloud

Na+

Na+

Na+

Na+

Na+

Na+

Graphene foam (sponge, 3D-architecture) Na+ Na+ Na+ Na+ Na+ Na+ Na+ Na+ Na+ Na+ Na+ Na+ + Na Na+ Na+ Na+ Na+ + Na+ + Na Na + Na Na+ Na+ Na+ Na+ Na+ Na+ + Na+ Na Na+ + + Na+ Na Na Na+ Na+

Figure 4.122 Schematic illustration of the assemblages of graphene layers by emphasizing p-electron clouds.

with each other to keep parallel stacking so that just the outermost p-electrons can interact with adsorbate cations, in other words, greatly diminishing the numbers of cations possible to be interacted (to be adsorbed), as illustrated by using Na þ ions in Fig. 4.122. The formation into a film tends to orient the graphene layers, some becoming few-layered graphene, and as a consequence the number of p-electrons possible to interact with adsorbate cations is reduced. In order to keep a large number of interactable p-electrons, random aggregations of single-layer graphene are required, as illustrated in Fig. 4.122, which are called foams. Fabrication of foams of carbon materials has been performed even in industrial scale and proposed various techniques, exfoliation of graphite flakes via graphite oxides, carbonization in an autoclave with rapid pressure release, carbonization of organic precursors with foaming agents, template carbonization, etc., as reviewed [364]. On graphene and rGO, many efforts to prepare graphene foams, as explained in Chapter 2, Sections 2.1.5 and 2.3.4, and their performances in various applications are also explained in the previous sections of this chapter. However, it must be mentioned here that their characterization is not yet established, as well as different naming, such as foam, sponge, and 3D architecture. For their characterization, various parameters of the structure and textures have to be clarified. Basic structural units of the foams have to be defined, because either single-layer, doublelayered, or multilayered graphene are important to understand the adsorption phenomena

354 Chapter 4 using p-electrons. Although the single-layer graphene may stack to form double-layered graphene in the preparation process of the foam, the single-layer graphene results in a foam with quite different performances from the double-layered graphene. Not only the number of layers stacked but also lateral size of graphene layers is also an important parameter, because lateral size limits the extension of p-electron cloud. Micro Raman spectroscopy focusing on the wall of the macropores may give important structural information. Foams are porous materials, including a large amount of macropores, where bulk density and porosity have been employed as conventional parameters for characterizing these materials, because macropores cannot be evaluated by gas adsorption measurements, such as N2 adsorption/desorption isotherms, which gives the quantitative information on nanopores in carbon materials, micropores, and mesopores [365]. For the quantitative evaluation of macropores, mercury porosimetry has often been employed. However, mercury porosimetry cannot be applied on carbon foams, because of high risk of collapse of fragile macropores in the foam during impregnation of mercury under pressure, and so a modification of pressurizing system by using U-shaped dilatometer has been proposed by application on exfoliated graphite [366]. Quantitative evaluation of macropores in exfoliated graphite has been proposed by applying the image processing techniques on the SEM images of the cross-sections [367]. For applying this technique on the graphene foams, statistical analyses are required on a certain number of cross-sections in different directions, because strongly anisotropic graphene layers tend to orient during their fabrication process. To understand pore structure on the walls, evaluation of nanopores, micropores, and mesopores, by using gas adsorption/desorption isotherms, is needed on some graphene foams, particularly on activated ones.

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CHAPTER 5

Mechanical properties and applications Chapter Outline 5.1 5.2 5.3 5.4

Fundamental mechanical properties Nanolubricants 383 Mechanical sensors 387 Mechanical reinforcement 393

374

5.4.1 Reinforcement of plastics 393 5.4.2 Reinforcement of ceramics 395 5.4.3 Reinforcement of metals 400

5.5 Reduced graphene oxide fibers 5.6 Concluding remarks 409 References 410

404

Polycrystalline graphite materials have unique mechanical properties. Their strength increases with an increase in temperature up to about 2500
C, roughly two times more at 2500
C than that at room temperature. The mechanical properties of bulk carbon materials are sensitive to their structure, such as the size reduced graphene oxide and stacking order of basic structural units (crystallites), and also to their texture, such as the orientation mode and degree of the basic structural units, apparent density, and pore sizes. Carbon fibers, which principally consist of crystallites with an axial orientation, although the degree of orientation depends on the precursors and temperature of heat treatment, give high strength and modulus with flexibility, and their industrial development has stimulated a wide range of applications. Graphene, which is an ultimate structural unit of graphite and carbon materials, has high mechanical properties, high strength, modulus, flexibility, and lubricity. However, bulk materials composed of graphene flakes have usually low density, even highly porous ones such as graphene foams, and are not expected to be used in applications requiring high mechanical properties. Graphenes are applied as nanolubricants, mechanical reinforcement additives, and components of mechanical sensors. They are competitors of graphite. Graphene fibers fabricated from reduced graphene oxide (rGO) were developed that have high mechanical properties; some of them have been reported to be comparable to carbon fibers and carbon nanotubes (CNTs), so they have stimulated fields of application similar to carbon fibers. Graphene. https://doi.org/10.1016/B978-0-12-819576-5.00005-0 Copyright © 2020 Elsevier Inc. All rights reserved.

373

374 Chapter 5 Graphene Single layer of mono-atomic thickness 5. Mechanical properties and applications

5.1 Fundamental mechanicalproperties 6 Thermal properties High thermal conductivity 7 Biomedical properties

High strength and modulus, Flexible layer, High lublicity,

3 Electrical properties High electrical conductivity 4 Chemical properties

High biocompatibility

Large surface area

5.2 Nanolubricants 5.4 Mechanical reinforcement 5.3 Mechanical sensors 5.5 Reduced graphene oxide fibers

Figure 5.1 Concept of the current chapter.

This chapter explores the fundamental mechanical properties of graphene describes and its applications in nanolubricants, mechanical sensors, and mechanical reinforcement. In the last section, the enhanced mechanical properties of graphene fibers are explained, including their fabrication processes. The concept of this chapter is summarized in Fig. 5.1.

5.1 Fundamental mechanical properties According to the classical theory of elasticity, the microscopic structure of a solid is ignored, in which the system is treated as a continuum body. In this treatment, deformation in solids is described by the continuous displacement field u(r); the displacement vector of the part of the solid that occupies position r in the equilibrium state [1]. As shown in 3.1, equilibrium positions r of the atoms of graphene are given by rðn; kÞ ¼ RðnÞ þ rk

(5.1) rk

where R ¼ n1 a2 þ n2 a2 ðn1 ; n2 ¼ 0; 1; 2; 3.Þ; n ¼ ðn1 ; n2 Þ; and ¼ 0 for k ¼ A; s2 ðk ¼ BÞ are the lattice vector, the vector specifying the primitive cells, and the vector denoting the position of A and B carbon atoms inside the primitive cell, respectively. Although the equilibrium positions of the A and B atoms are in a two-dimensional (2D) plane, 3D displacements uðn; kÞ with components of ui ðn; kÞði ¼ x; y; zÞ are allowed for each carbon atom here. Carbon atoms can also move perpendicular to the graphene sheet even in 2D graphene. The displacements are a function of time and are denoted by 6N coupled equations of motion for graphene consisting of N carbon atoms [2,3]: M

X d2 ui ðn; kÞ ¼ fij ðn; k; n0 ; kÞuj ðn; k0 Þ; 2 dt n0 ;k0 ;j

(5.2)

where fij ðn; k; n0 ; kÞ is the interatomic force constant matrix between carbon atom k0 in a primitive cell nʹ at site rk and carbon atom k in a primitive cell n; i, j are the Cartesian

Mechanical properties and applications 375 indices of the displacements; and M is the atomic mass of carbon. Here, the Harmonic approximation is applied in this description. As shown in Eq. (3.1), the Fourier transformed product in the reciprocal space well-describes deformations in the infinite graphene sheet extended in the 2D plane with translational symmetry. Here, the wave number vector in the Brillouin zone is usually denoted by q instead of k to distinguish the lattice system from the electron system. In addition, one chooses periodic time dependence with angular frequency u. Thus, the displacements and the equation of motions are written as 1 X k ui ðqÞeiq$rðn;kÞiut (5.3) ui ðn; kÞ ¼ pffiffiffiffiffiffiffiffi MN q X 0 k0 Dkk (5.4) u2 uki ðqÞ ¼ ij ðqÞuj ðqÞ; k0 ; j

0

where Dkk ij are the elements of the dynamical matrix D(q) described as 0

Dkk ij ðqÞ ¼

0 0 1 X fij ðn; k; n0 ; k0 Þeiq,½rðn ;k Þrðn;kÞ . M n0

(5.5)

The dynamical matrix accounts for the collective vibrational motion of the lattice at wave vector q. The eigenvalues u of the equations are obtained by solving the secular equation:  2  u I  DðqÞ ¼ 0; (5.6) where I is the identity matrix. The eigenvalues u are connected to phonons that describe the lattice vibration dynamics in quantum mechanical treatment, which is especially important to thermal properties of solids, in which the relation between eigenvalues u and wave vector q is known as the dispersion relation. Compared with experimentally observed phonon dispersions [4], one can determine actual values for the force constants that describe the dynamical matrix elements (Fig. 5.2). In terms of interatomic coupling parameters using Born’s long-wavelength approximation method [1], macroscopic elastic constants are derived from knowledge of dynamical matrix D(q) given by actual values of force constants obtained by fitting with experimentally observed phonon dispersion relations. At first, the dynamical matrix is expanded into the series in powers of small q (long wavelength): X 0 1X kk0 ð0Þ kk0 ð1Þ kk0 ð2Þ Dkk þ i qk Dij;k þ qk ql Dij;kl (5.7) ij ðqÞ ¼ Dij 2 k;l k The zeroth-order term in q on the right-hand side gives information about the optical phonons at the center of the Brillouin zone (G-point), which are the major phonon modes

376 Chapter 5 1600 1400 1200

ω / cm-1

1000 800 600 400 200 0

Κ

Γ

Μ

Γ

Figure 5.2 Phonon dispersion relations of graphene obtained by solving Eq. (5.6) [2], in which the force constants were determined to be in agreement with Mohr et al. [4].

observed in Raman spectroscopy for graphene. The second-order coefficients mainly contribute to the elastic constants cij,kl; cij;kl ¼ ½ik; jl þ ½ jk; il  ½ij; kl þ ðij; klÞ;

(5.8)

where the square brackets are defined by ½ij; kl ¼

M X kk0 ð2Þ D ; 2vc k;k0 ij;kl

(5.9)

and the round brackets are denoted as ðij; klÞ ¼ 

M XXX kk00 ð1Þ kk0 k0 k000 ð1Þ D Ghp Dpk;l ; vc k;k0 hp k00 k000 hi;j

(5.10)

where 0

Gkk hp ¼

X uðlÞ ðkhÞuðlÞ ðk0 pÞ l

u2l

:

(5.11)

pffiffiffi Here, vc, ul, and u(l) are the volume of the primitive unit cell given by a2 3 2 with a ¼ 2:46  A as the lattice constant of graphene, the optical eigenfrequencies and eigenvectors at the G point of the Brillouin zone. For graphene as a 2D crystal, we can denote the elastic constant as the dimension of tension coefficients. Thus, we can denote them as gij,kl using Voigt’s notation: g12

g11 ¼ ½11; 112D þ ð11; 11Þ2D ; ¼ 2½12; 122D  ½11; 222D þ ð11; 22Þ2D ; g66 ¼ ½11; 222D þ ð12; 12Þ2D ;

(5.12) (5.13) (5.14)

Mechanical properties and applications 377 Table 5.1: Tension coefficients of graphene in units of 104 dyne/cm [3]. [11, 11]2D 47.28

[11, 22]2D 22.38

[12, 12]2D 12.45

(11, 11)2D 6.73

g11

g12

40.55

9.24

g66 15.66

where the hexagonal symmetry results in (11, 11)2D ¼ (12, 12)2D and (11, 22)2D ¼ e(12, 12)2D. Numerical results for the tension coefficients obtained by the empirical force constants model [4] are summarized in Table 5.1. Here, the longitudinal pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi and transverse sound p velocities ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffifor in-plane displacements are given by cl ¼ g11 =r2D ¼ 23:08 km/s and ct ¼ g66 =r2D ¼ 14:34 km/s, where r2D is the surface mass density of graphene. These results were derived within the harmonic approximation. The importance of anharmonic effects was noted for both of in-plane and out-of-plane vibrational modes of the graphene lattice [5]. The most striking anharmonic effect is the presence of a linear term in the dispersion relation of the acoustic bending band of graphene at long wavelengths (q/0). This implies a reduction in the amplitude of out-of-plane oscillations compared with a flexural mode with a quadratic dependence on the long-wavelength limit. Moreover, nonlinear elasticity was observed in graphene, in which a phenomenological nonlinear relation was found between applied stress and the observed strain by atomic force microscopy (AFM) measurement [6]. Mechanical properties were measured on suspended single-layer graphene films prepared from highly oriented pyrolytic graphite by micromechanical cleavage [6]. Two cleaved graphene films were placed on two SiO2-Si substrates with an array of circular holes (of 1.5 and 1 mm in diameter and 500 nm deep) and the cantilever suspended on the hole indented the center of the film with a diamond tip (tip radii of 27.5 and 16.5 nm). From two graphene films with the holes of two different diameters, 67 values for stress and Young’s modulus data were obtained using cantilevers with two different tip radii. The Young’s modulus and elastic modulus measured on the 2D structure (single-layer graphene) were divided by interlayer spacing in graphite at 0.335 nm, to convert to parameters corresponding to a 3D structure for comparison with bulk graphite and related materials. Fig. 5.3A shows the distribution of Young’s modulus converted to a 3D structure value on two graphene films, which gave a Young’s modulus for single-layer graphene of 1.0  0.1 TPa. Fig. 5.3B shows fracture behavior measured on films of different sizes using different tip radii. Results revealed no sign of slippage or other irreversible deformation before catastrophic fracture; fracture occurred at large deflections above 100 nm and forces above 1.8 and 2.9 mN for the smaller and larger indenter tips, respectively. These forces were large enough to break standard Si AFM tips, but the

378 Chapter 5 (B)

(A) 15

Load / mN

Counts

3 10

2 Fracture

5

1

0 0.7 0.8 0.9 1.0 1.1 1.2 1.3 1.4 Effective Young’s modulus / TPa

0

0

100 150 50 Indentation depth / nm

200

Figure 5.3 Suspended single-layer graphene [6]: (A) histogram of effective Young’s modulus, and (B) fracture behaviors of four sheets with different tip radii (R) and film diameters (2a). Dashed line in (A) represents Gaussian fit.

diamond tips used in this study were not damaged. The graphene film still hung around the edge of the hole, which suggested that the fracture started at the indentation point. The breaking strength was calculated to be 130  10 GPa. These determined that the Young’s modulus and breaking strength were almost comparable with those reported for bulk graphite and multiwalled CNTs (MWCNTs): for bulk graphite, 1.02 TPa and 118e121 GPa along the in-plane direction [7], and for MWCNTs, 0.27e1.47 TPa and 3.6e63 GPa, respectively [8]. Spring constants were measured on stacks of graphene layers (less than five layers) suspended over a hole of SiO2 substrate [9]. Graphene flakes with different thicknesses were prepared by rubbing kish graphite across an SiO2 substrate (mechanical cleavage). The ones with thicknesses between 2 and 8 nm were selected and suspended on the holes (trenches) of the substrate. Fig. 5.4A illustrates a scheme of an AFM tip pushing down on a suspended graphene flake. The force applied to the graphene flake was calculated from the deflection of tip h and the effective spring constant k of the suspended graphene flake was obtained from the relation between force and the deflection of flake Dz (for example, Fig. 5.4B). The model of a doubly clamped beam in equilibrium with a static force and under axial tension gives the relation k ¼ 16:23 Ewðt=LÞ3 þ 4:93T=L;

(5.15)

where w, t, and L are the width, thickness, and length of the flake, respectively, E is the Young’s modulus and T is tension. Therefore, the values of the spring constant measured for graphene flakes with different thicknesses were plotted against the flake constants, w(t/L)3, for each flake in Fig. 5.4C. The k values measured for thick flakes more than 5 nm

Mechanical properties and applications 379 (A)

(C) 6 Deflection η

Force / nN

(B) 20 15 10 5

5 Spring constant k / N/m

Deflection Δz

4 3 2 1 0

0 -5-5

-1 0 10 15 5 Δz of graphene / nm

0.0

0.1

0.2 0.3 0.4 w(t/L)3 / pm

0.5

0.6

Figure 5.4 Scheme of (A) measurement using atomic force microscopy tip, (B) relation between force applied and deflection of graphene flake Dz, and (C) a plot of spring constant measured against w(t/L)3 [9].

fit with the theory, which suggested that Young’s modulus E was about 0.5 TPa, significantly below the bulk value of 1 TPa for graphite. Graphene sheets, including single-layer graphene, suspended over a rectangular pit of SiO2 substrate as graphene-sealed microchambers forming a sealed square drumhead with a width W of 4.75 mm were used to measure the elastic constants and mass of single-layer graphene by changing the difference in gas pressure inside (pint) and outside (pext) the pit (Dp ¼ pint e pext) [10]. With a significantly long time (typically 24 h), the internal and external pressure equilibrated through leakage of the graphene membrane, in which the leakage rate depended on the gas species but depended less on the thickness of graphene films. Air and argon had similar leakage rates but helium exhibited two orders of magnitude faster rates. However, the graphene membrane had enough impermeability for a short period, which enabled it to apply graphene as a well-defined force uniformly distributed across the entire surface of the membrane by pressure difference Dp. The product of elastic constants of graphene Et/(1-v), in which E, t, and v are the Young’s modulus, the thickness of the membrane, and the Poisson’s ratio, respectively, was estimated using the pressure difference and deflection measured by AFM as 390  20 N/m, agrees well with that for bulk graphite. Pressure difference Dp was also used to control the strain in graphene, which emerged as a change in the resonance frequency f [10]. Fig. 5.5 shows the external pressure pext dependence of resonance frequency f, in which vibrations of the membrane are actuated

380 Chapter 5 (A)

(B)

90

30

Frequency3 / 105MHz3

Amolitude

Frequency / MHz

120

15 0 40

60

80

Frequency/MHz

60 30 0

0

20 10 Pressure / kPa

30

23.0 15.0 7.5 0.0 0

90 30 60 Pressure / kPa

120

Figure 5.5 External pressure dependence of resonance frequency of single-layer graphene membrane. (A) The inset shows a resonance frequency curve at pext ¼ 27 Pa. The cubic of the frequency is proportional to the pressure except for the vicinity of the minimum value of the resonance frequency. (B) Solid curve shows the fit with the model (see text) [10].

and measured using an optical method. The frequency depends on the pressure, pext, with a cubic law, which is well-explained by a simple model for the membrane as an elastic body described as sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi c2 Et 3 f ¼ Dp ; (5.16) 2048m3 W 4 ð1  vÞ in which m, E, t, and v are the mass of graphene per area, the Young’s modulus, the membrane thickness, and the Poisson’s ratio, respectively, and c2 ¼ (0.8 þ 0.062v)3. By fitting the data with Eq. (5.16) as shown in Fig. 5.5B, the mass of graphene per area is obtained as m ¼ ð9:6  0:6Þ  107 kg/m2, which is 30% higher than the theoretical value of 7:4  107 kg/m2. The extra mass was attributed to adsorbates, which is significant for the single atom-thick material. Uniaxial strain ranging from 0% to e0.8% caused marked splitting and redshift of the G-band, and also redshift of the G0 -band in Raman spectra of single-layer graphene [11,12]. Single-layer graphene films prepared by microcleavage were deposited onto two flexible substrates, polyethylene terephthalate and acrylic resin, and strained by either twopoint or four-point bending [11]. As shown in Fig. 5.6A, G-band shifted to the low Raman shift side (redshift) with increasing uniaxial strain and split into two, G- and Gþ-bands, above 0.5% strain, but the G0 -band did not split although it also had redshift. The 2D0 band exhibited the same redshift with strain. Shifts of these G-, Gþ-, G0 -, and 2D0 -bands with increasing strain were approximated by linear relations, the slopes of which were e31.7, e10.8, e64, and e35 cm1/% strain, respectively, in agreement with firstprinciples calculation. The redshift of the G0 -band was most sensitive to small strain ranging from 0% to e0.8%. Raman scattering from the Gþ- and G- bands showed

Mechanical properties and applications 381 G-

G+

G'

Uniaxial strain (%)

0.77

Uniaxial strain (%) 0.77

0.74

0.74

0.66 0.61 0.50 0.37 0.29 0.11 0.00

G 1560 1575 1590 Raman shift / cm-1

Intensity / a.u.

Intensity / a.u.

0.80

0.66 0.61 0.50 0.37 0.29 0.11 0.00

2700 2650 2600 Raman shift / cm-1

Figure 5.6 Changes in Raman spectra of single-layer graphene with increasing uniaxial strain [11]: (A) G-band, and (B) two-dimensional band. a.u., arbitrary units.

dependence on distinctive polarization that reflected the angle between the axis of the stress and the graphene crystal axes. For graphene flakes on an SiO2-Si substrate, which consisted of domains of single to four layers, Raman spectra were measured before and after covering an SiO2 layer 5 nm thick of by pulse laser deposition (PLD) [13]. In the single-layer graphene domain, a marked change in Raman spectrum was observed after SiO2 covering, as shown in Fig. 5.7A. There was a marked intensity increase in the D-band and a clear appearance of the D0 -band, which suggested the formation of structural defects from the SiO2 coating. By increasing the thickness of the domain to four layers, the effect of the SiO2 coating decreased gradually, and finally no effect was observed on the bulk graphite (multilayered stacking), as shown in Fig. 5.7B. With annealing of these SiO2-coated graphene flakes, the bands in Raman spectra showed blueshifts, as shown on the G-band as an example in Fig. 5.8A, owing to the compressive stress induced, which was estimated to be as high as 2.1 GPa, as shown in Fig. 5.8B. Fig. 5.7C suggests that different amounts of defects were introduced by different coating methods. Radio-frequency (RF) sputtering, PLD, and electron-beam (e-beam) evaporation, and also HfO2 deposition by PLD had the same effect, but poly(methyl methacrylate) (PMMA) deposition by spin coating had no effect on single-layer graphene.

382 Chapter 5 (A)

(B)

(C) D

G

D Intensity / a.u.

D

G'

G D’

Graphite

G'

G D’

4 layers

After SiO2 deposition

SiO2 by RF

G'

SiO2 by PLD

3 layers SiO2 by E-beam 2 layers Pristine graphene

PMMA by spin coating 1 layer

1200

2400 1800 Raman shift / cm-1

3000

1200

HfO2 by PLD

3000

2400 1800 Raman shift / cm-1

1200

1800 2400 Raman shift / cm-1

3000

Figure 5.7 Raman spectra of graphene flakes coated by SiO2 5 nm thick [13]. (A) Single-layer graphene before and after SiO2 coating by pulse laser deposition (PLD). (B) Graphene domains with different thicknesses as well as the bulk graphite after SiO2 coating by PLD. (C) After SiO2 deposition by radio-frequency (RF) sputtering, PLD, electron-beam (e-beam) evaporation, as well as poly(methyl methacrylate) (PMMA) coating by spin coating and HfO2 coating by PLD. a.u., arbitrary units.

(B) 1 layer 2 layers 3 layers 4 layers

1595 1590 1585 1580 0

100 200 300 400 500 Annealing temperature / oC

Compressive stress / GPa

G-band frequency / cm-1

(A)

2.5 2.0 1.5 1.0 0.5 0.0

-0.5

0

100 200 300 400 500 Annealing temperature / oC

Figure 5.8 Graphene flakes coated by an SiO2 layer 5 nm thick [13]. (A) Change in G-band frequency with annealing temperature for domains with different thicknesses. (B) Compressive stress controlled by annealing temperature on single-layer graphene domain.

The elastic properties of a graphene flake on the hole of an SiO2 substrate were measured using AFM under ambient conditions [14]. The graphite flake cleaved by an adhesive tape was pressed against the SiO2-Si wafer to obtain flakes of various sizes and thicknesses. By selecting more than 50 flakes with thicknesses varying from 2.4 to 33 nm (8e100 layers) completely covering the hole of the substrate, the bending rigidity, D, and tension, T, were calculated from induced deflection of the flakes. Parameters D and T depended on the thickness of the flake, as shown in Fig. 5.9A and B, respectively. Both parameters increased quickly with an increase in thickness, which was associated with fluctuations in

Mechanical properties and applications 383 (B)

10-11

Frequency / GHz

10-12

10

1.0

10-13

0.1

10-14 0

(C) 102

Tension T / N/m

Bending rigidity D / Nm

(A)

20 10 30 Thickness / nm

40

0

20 10 30 Thickness / nm

40

R

10 R

1 0.1 0

20 10 30 Thickness / nm

40

Figure 5.9 Thickness dependences of (A) bending rigidity D, (B) tension T, and (C) the frequency of the fundamental mode calculated [14]. The inset in (C) shows the displacement profile of this mode.

the observed data. For flakes with a thickness less than 10 nm, the data points of D are close to the theoretically calculated line in Fig. 5.9A, which suggests the absence of stacking faults in thin flakes. Tension T varied from flake to flake and depended on the thickness. It tended to saturate at around 20 N/m. The resonance frequency calculated from measured D and T increased with increasing thickness, depending on the size of the holes of the SiO2 substrate, as shown in Fig. 5.9C. The frequencies were slightly below 1 GHz for the hole 540 nm in diameter, whereas the frequency could be over 10 GHz for the small hole 84 nm in diameter.

5.2 Nanolubricants Solid lubricants to reduce friction in the interfaces of materials have been a major application of graphite as a precursor material of graphene. Thus, it is no wonder that graphene has been investigated by many research groups as a solid lubricant and an additive for lubricants, as well as graphite. The composite in which rGO flakes were covalently assembled onto an Si wafer was fabricated by placing a single layer of 3-aminopropyl triethoxysilane (APTES) in the interface between rGO and Si [15]. Conversion to graphene oxide (GO) from graphite and thermal exfoliation were repeated twice to achieve the high dispersion of GO flakes in water. The surface of Si wafer was coated by a single layer of APTES by immersing the wafer into an acetoneewater solution of APTES; then it was coated by rGO by immersion into a GO aqueous suspension at 80
C for 12 h. The resultant GO-APTES-Si composite was heated at 200
C to reduce GO to rGO. Fig. 5.10 shows the change in friction coefficient with time, measured by using a ball-on-plate tribometer, for the composite before reduction, i.e., GO-APTES-Si and rGO on an Si wafer without APTES (rGO-Si) together with the resultant composite, rGO-APTES-Si. The composite exhibited a low

384 Chapter 5

0.6 The composite GO/APTES/Si

0.4 0.2 0

(C) Friction coefficient

Friction coefficient

(B) 0.6

400

1200 800 Time / s

The composite rGO/Si

1000

3000 2000 Time / s

2000

4000

4000

6000 8000 10000 12000 Time / s

0.6 The composite rGO/APTES/Si 0.4 0.2 0

0

0

(D)

0.2 0

0.2

1600

0.6 0.4

The composite rGO/APTES/Si

0.4

0 0

Friction coefficient

Friction coefficient

(A)

0

100

200 Time / s

300

400

Figure 5.10 Changes in friction coefficients with time for different lubricants [15]: (A) composite graphene oxide (GO)/3-aminopropyl triethoxysilane (APTES)/Si, (B) reduced graphene oxide (rGO)/ APTES/Si, (C) rGO/Si at an applied load of 0.1 N, and (D) the rGO/APTES/Si composite at 0.2 N.

friction coefficient as 0.24 and an antiwear life of more than 10,800 s (Fig. 5.10B), although GO/APTES/Si had almost the same friction coefficient as 0.25 but a much short antiwear life of 1600 s (Fig. 5.10A). The composite rGO-Si without APTES had a similar friction coefficient of 0.23, but life was much shorter at 4300 s (Fig. 5.10C), probably owing to failure of the rGO film. These experimental results suggest that rGO flakes possess good self-lubricating properties and that covalent bonding of rGO to the silicon substrate mediated by APTES makes the antiwear life longer. Under a high load of 0.2 N, however, even the composite rGO/APTES/Si had a short lubricating life of about 300 s (Fig. 5.10D), which suggests that lubricating rGO is more effective at relatively low applied loads. Oleic acidemodified single-layer rGO flakes were prepared by mixing an rGO suspension (0.5 mg/mL) with oleic acid at 80
C [16]. rGO was produced by oxidation of thermally expanded graphite according to Hummers’ method and then reduced with hydrazine in Ar at 80
C, the Raman spectra of which suggested the formation of single-layer rGO flakes. The resultant oleic acidemodified rGO flakes were mixed into a gear oil, polyalphaolefin9, with a density of 0.83 g/cm3 and viscosity of 9 cSt at 100
C, in 0.01e5 wt%, for which the tribological properties were measured using a four-ball tribometer. The friction coefficient and wear scar diameter are shown as a function of the rGO content in the oil in

Mechanical properties and applications 385

0.050

Friction coefficient

Friction coefficient

0.055

0.060 0.055 0.050 0.045 0.040

0.01 0.1 1 Graphene content / wt%

5

0.045 0.040

0

0.02 0.04 0.06 0.08 Graphene content / wt%

0.1

0.70 0.65 0.60 0.55

Wear scar diameter/mm

(B)

0.060

Wear scar diameter / mm

(A)

0.78 0.63 0.56 0.49 0.01 0.1 1 5 Graphene content / wt%

0.50 0.45 0

0.02 0.04 0.06 0.08 Graphene content / wt%

0.1

Figure 5.11 Gear oil mixed with oleic acidemodified reduced graphene oxide flakes [16]: (A) friction coefficient, and (B) wear scar diameter as a function of the graphene content in the oil.

Fig. 5.11A and B, respectively. Both the friction coefficient and wear scar diameter decreased by adding rGO. The former was at a minimum with a graphene content of 0.02 wt%, and the latter was at a minimum at 0.06 wt% and then increased gradually with an increase in the rGO content. Improvement in tribological properties by adding rGO flakes to a gear oil was thought to be caused by the excellent dispersibility of graphene flakes because oleic acid modified and formed a protective rGO layer on the surface of each steel ball at rGO contents below 0.06 wt%. Modification of rGO flakes was performed using a mixture of stearic and oleic acids (3:5 in mass) in cyclohexane solution [17]. The addition of modified rGO was effective in improving the maximum nonseizure load, friction coefficient, and wear rate of the lubricating oil. Nanocomposites of single-layer rGO with buckminsterfullerene C60 were synthesized by mixing octylamine (OA)-intercalated GO with C60 in toluene, followed by heating at 600
C under a high vacuum for 80 min to remove OA and excess C60 [18]. The resultant composites were thought to have layer-by-layer texture. Their lubricant wear protection properties (wear loss in volume and wear scar diameter) in a grease were compared with other additives using a four-ball tribosystem. The results are shown in Fig. 5.12, which reveals that the composite C60/rGO exhibited much less wear loss and gave a smaller wear scar than other additives, graphite, C60, and even MoS2. rGO nanoflakes prepared by hydrazine reduction at 50
C were mixed with an ionic liquid (IL), 1-butyl-3-methylimidazolium hexafluorophosphate, at different concentrations of 0.025e0.100 mg/mL with sonication. The resultant mixtures (nanofluids) were spin coated onto a diamond-like carbon (DLC) deposited onto a steel substrate by magnetron sputtering in an Ar/methane atmosphere [19]. The thickness of DLC deposited on steel was approximately 2 mm and the thicknesses of the coated nanofluid were in a wide range of 0.5e2.0 mm. The composites thus prepared, rGO/IL/DLC, were subjected to friction and wear tests under high vacuum and radiation environments (imaging a space

12 0.6

10 8

0.4

6 4

0.2

Wear scar diameter / mm

0.8

14

C60/rGO

MoS2

C60

0

Graphite

2 Pristine grease

Wear volume loss / 10-3mm3

386 Chapter 5

0

Additives (1.0 wt%)

Figure 5.12 Wear volume loss and wear scar diameter of test ball for greases with various additives [18]. rGO, reduced graphene oxide.

Figure 5.13 Space tribology on steel for the composite reduced graphene oxide (rGO)/ionic liquid (IL)/diamond-like carbon (DLC) compared with rGO/IL [19]: (A) friction coefficient, and (B) wear loss.

environment) using a ball-on-disk tribometer. In Fig. 5.13A and B, friction coefficient and wear loss measured on an rGO/IL/DLC composite were compared with a composite without a DLC layer, i.e., rGO/IL/steel, revealing that the obvious advantages of the friction coefficient and wear loss of the composite rGO/IL/DLC. The optimal content of rGO in IL was 0.075 mg/mL; both friction coefficient and wear loss became minimal. The composite rGO/IL/DLC with the optimal composition was also compared with a composite containing 0.075 mg/mL MWCNTs instead of rGO [20]. Changes in friction coefficient with sliding time are shown in Fig. 5.14 for two composites and the ILs

Mechanical properties and applications 387

Figure 5.14 Friction coefficients of composites reduced graphene oxide (rGO)/ionic liquid (IL)/diamond-like carbon (DLC) and multiwalled carbon nanotube (MWCNT)/IL/DLC compared with ILs [20]: (A) under the load of 10 N and (B) under 30 N.

(without carbon additives). The composite CNT/IL/DLC exhibited the lowest friction coefficient (0.047) at an applied load of 10 N owing to the rolling of MWCNTs, whereas rGO had a negligible effect on the friction-reducing property of ILs. Under 30 N, however, the situation was the opposite: the effect of rGO on the frictionreducing property of ILs was significant and a minimum friction coefficient (0.057) was obtained. Different tribological behaviors of rGO and MWCNTs in ILs during friction were discussed. At a low applied load, long CNTs were shortened and rGO flakes were broken into small pieces, but at a high applied load, rGO flakes stacked to form thick sheets, which worked as a tribofilm that separated the two contact surfaces, DLC and steel ball. Aqueous dispersion of 1 wt% of single-layer GO flakes exhibited a small friction coefficient (approximately 0.05) using a sintered tungsten carbide ball and stainless steel flat plate, much smaller than purified water (over 0.4) and slightly smaller than the oil lubricant, poly-alphaolefin (approximately 0.08) [21]. No obvious surface wear on the ball and plate after 60,000 cycles of friction testing was observed when aqueous GO dispersion was used. GO flakes modified by alkyl imidazolium ILs were effective as additives for multialkylated cyclopentanes to improve stability and tribological properties, in which the GO flakes were partly reduced during modification by ILs [22]. Graphene-based nanofluids and nanolubricants were reviewed by focusing on developments [23].

5.3 Mechanical sensors To develop novel mechanical sensors, graphene with excellent properties for both mechanical and electrical performance is an ideal candidate for sensing materials, in

388 Chapter 5 which mechanical stimulation in the surrounding environment should be rapidly and sensitively converted into an electrical signal with high mechanical durability and reliability. The electrical conductivity of single-layer graphene, which was synthesized by chemical vapor deposition (CVD) and suspended with no substrate, was sensitive to strain, but its piezoresistive sensitivity represented by the gauge factor (GF) was not high, about 1.9 with an applied strain less than 3% [24]. Single-layer graphene prepared by mechanical peeling also gave a GF of 1.9 [25]. Graphene sheets synthesized by CVD of methane on an Ni-coated SiO2-Si substrate and transferred to a polydimethylsiloxane (PDMS) substrate, which was 95% transparent and composed of one to three layers, increased in resistance from 492 to 522 kU with an applied strain up to 1%; GF reached about 6.1 [26]. Graphene films were grown on fluorophlogopite mica by plasma-enhanced CVD of CH4 with no catalyst; they were composed of nanograins of one to few graphene layers and delivered high sensitivity to strain [27]. The IeV relations were linear under strains ranging from 0.29% to 0.37% and the resistance change DR/R0 with strain ε was approximated to be linear, giving a GF of 37, as shown in Fig. 5.15A. The GF increased with an increase in film resistance R0, and the maximum GF exceeded 300, as shown in Fig. 5.15B. A sensor was fabricated by photolithographic patterning of graphene film on a flexible PDMS substrate (see insets in Fig 5.16B) [28]. A single-layer graphene film was synthesized by CVD of CH4 in H2 gas at 1000
C on a Cu foil and transferred using PMMA onto PDMS. The softness and flexibility of the PDMS substrate allowed the sensor to be used under stretching (Fig. 5.16A), bending, and even torsion. Before use, however, 1.4e3.2% stretching was needed to release inherent wrinkles in the film formed during CVD growth. In Fig. 5.16B, resistance change with stretching along two orthogonal

(A)

(B) GF=37

350 300

0.10

250

0.05

200

GF

ΔR/R0

0.15 a)

150

0.00

-0.05

100

-0.10

50

-0.15 -0.4

0 -0.2

0.0 0.2 Strain ε / %

0.4

104 105 106 107 108 109 Film resistance R0 / Ω

Figure 5.15 Graphene films prepared by a plasma-enhanced chemical vapor deposition [27]: (A) change in DR/R0 with strain ε, and (B) change in gauge factor GF with a film resistance of R0.

Mechanical properties and applications 389 16

(A)

(B)

ΔR/R0

12

Stretching along y-axis

8 along x-axis

4 0 0 20

(C)

26

0

R / kΩ

10 5

-5 0

2

4 6 8 Time / s

10

Strain / %

15

30

22

3.0 Resistivity / 10-2Ωm

34

4

12 8 Strain / %

16

20

200

400 600 Cycle

800 1000

(D)

2.4 1.8 1.2 0

Figure 5.16 Graphene strain sensor [28]: (A) photograph under stretching, (B) resistance change DR/R0 with strains along two orthogonal directions, (C) resistance R response with repeated application of 5% strain, and (D) cyclability with 10% strain.

directions is shown, revealing that the strain resolution was 1% (corresponding to a length increase of 0.7 mm in the sensor) and the GF for the x- and y-directions reached 42 and 71, respectively. Resistance changes with 5% application of cyclic tensile strain along the electrode direction and 10% strain up to 1000 times are shown in Fig. 5.16C and D, respectively, suggesting the high stability of the sensor. This sensor could detect strains induced via bending and torsion as well as tensile. Graphene woven fabrics synthesized by CVD of CH4 on a Cu mesh transferred onto a PDMS sheet (inset in Fig.5.17A) exhibited a high GF for a wide range of strain [29]. The sensor responded to different strain modes, static and dynamic, as shown by the response for dynamic strain with different frequencies in Fig. 5.17A. The GFs of the sensors were calculated to be about 103 under low strains (2e6%) and to reach 106 under high strains (more than 7%), as shown in Fig. 5.17B. Percolative networks of graphene flakes were fabricated on flexible plastic substrates by spray-deposition of a graphene suspension and exhibited high sensitivity to strain [30]. Graphene flakes were prepared from commercially available graphite intercalation compounds by microwave-assisted exfoliation, followed by mechanical cleavage by sonication. The relative change in resistance, DR/R0, of this percolative graphene film linearly increased with an increase in mechanical strain ε; it was much more sensitive than

390 Chapter 5 (A)

(B) 2

GF / 106

Current

0.2 Hz 0.1 Hz

GF / 103

6

1 Hz

4

1 0

2

0

1 2 3 Strain / %

4

2

4 Strain / %

6

0.02 Hz

0 0

Time

8

Figure 5.17 Strain sensor of graphene woven fabrics [29]: (A) current responses for 5% strain at different frequencies with the inset of the photograph of the fabric (under 20% strain), and (B) gauge factor (GF) of the sensor under different strains.

(A)

(B)

20

15 10

10

5

Residance / kΩ

Graphene sensor 20

GF

ΔR/R0 / %

30

Reference metallic sensor

0 0

0.5 1 Strain ε / %

1.5

0 0

114 110 106 102 2490 2492 2494 2495 Cycle number

1000

2498

2000 3000 Cycle number

2500

4000

Figure 5.18 Percolative graphene film sensor [30]: (A) change in relative resistance, DR/R0 with strain ε compared with reference metallic sensor, and (B) change in gauge factor (GF) with 4000 straining cycles.

the reference metallic sensor, and its GF remained over 4000 straining cycles, as shown in Fig. 5.18. The introduction of rippling in a graphene sheet using a rippled PDMS substrate was effective in widening the range of applicable strain to 30% by keeping the GF at about 2 [31]. Strain sensors were developed by the spray deposition of multilayer rGO nanoflakes suspended in 1-propanol onto polycarbonate substrate [32]. Multilayer rGO was prepared by thermal exfoliation of GO at 1150
C and sonication in 1-propanol. Quasistatic tests of this flexible sensor were performed by measuring piezoresistive response using a universal tensile machine with a linearly increasing load. The total resistance of the sensor increased by increasing strain ε above 0.015% and the GF reached about 200 at ε of 0.135%, as shown in Fig. 5.19. To demonstrate the potential of this sensor for in situ nondestructive applications, response analysis was performed using an array of eight

Mechanical properties and applications 391 (B) Observed Fitted

GF

Resistance R / kΩ

(A)

Strain ε / %

Strain ε / %

Figure 5.19 Graphene sensor for strain measurement [32]: changes in (A) total resistance and (B) gauge factor (GF) for the sensor with ε.

sensors deposited onto a carbon fiberereinforced carbon plate (550  550  2.63 mm3) with excitation by a mechanical vibration up to 100 kHz. A composite was prepared from a powder mixture of GO and polyvinylidene fluoride (PVDF) by the irradiation of focused solar light, in which the reduction of GO and melting of PVDF occurred simultaneously [33]. The composite containing 2.2 wt% rGO exhibited a GF of 12.1. A ZnO-nanorod/graphene composite [34] worked as a mechanical sensor as well as a chemical sensor, as described previously (4.3.5). It accommodated flexural deformation without mechanical or electrical failure for a bending radius below 0.8 cm repeated bending and releasing up to 100 times. Resonators composed of single-layer graphene exhibited high sensitivity to mass and strain by optimizing the gate voltage [35e38]. The device was fabricated by depositing commercially available graphene flakes on an SiO2-Si substrate, followed by mechanical peeling to obtain single-layer graphene, which was confirmed by Raman spectra. On this graphene, metal electrodes (Cr-Au) were patterned by electron-beam lithography and then an SiO2 layer of the substrate was etched out to make the graphene sheet suspended [36]. Typically, vibrations with fundamental resonant frequencies are in the megahertz range for graphene, which are actuated either optically by using laser incidence or electrically through capacitive coupling and detected optically by interferometry [35]. Even with no driving force, the graphene resonator still oscillated owing to thermal excitation by a root mean squared amount xth ¼ [kBT/keff]1/2, in which keff is the effective spring constant and kB is the Boltzmann constant. Thus, the ultimate limit on force sensitivity is determined by thermal fluctuations in the resonator: 1  4keff ðkB TÞ 2 f dF ¼ ; (5.17) Qu0

392 Chapter 5

8

(B)

6 4 2 0 5.5 5.0 4.5 Frequency / MHz

(C)

0.2 Amplitude / nm

10 Amplitude / nm

(A)

0.1

0.0 15.5 16.0 Frequency / MHz

16.5

Figure 5.20 Scanning electron microscopy image of graphene resonator bridging a trench structure on the Si substrate: (A) vibrational spectrum at the primary resonance frequency (B) before and (C) after annealing at 600
C [38].

where Q and u0 are the quality factor and the eigen-angular-frequency, respectively. At maximum, a force sensitivity of 0.9 fN/Hz1/2 was obtained [35]. The mass sensitivity of the device reached about 2  1021 g. Because graphene has a very large surface area, adsorbates on the surface significantly influence its mechanical properties. Fig. 5.20B and C show vibrational spectra for the graphene resonator shown before and after annealing at 600
C, respectively. After annealing, a significant increase in quality and the resonance frequency were observed. This is partially attributed to a decrease in weight and the damping effect induced by thermal desorption of adsorbates on the surface of graphene. However, the increase in strain for graphene from thermal shrinkage of the holding part of the device also greatly contributed to the change in mechanical properties of the resonator in a typical case [38]. Anyway, it is better to clean the graphene part by ohmic heating for measuring mass with a graphene resonator, as shown in Fig. 5.21. This demonstrates that most as-fabricated devices have a mass higher than graphene itself

Pentacene adsorption

Figure 5.21 Built-in strain versus normalized mass r/rgraphene for a single-layer graphene resonator [36].

Mechanical properties and applications 393 (rgraphene) and built-in strain, even after cleaning by ohmic heating. To recover normalized mass r/rgraphene caused by the adsorption of pentacene molecules, annealing under high vacuum was necessary. The sensitivity of this graphene-based device was somewhat larger than the CNT-based device [39], but the graphene-based one had the advantages of higher reproducibility and a larger surface area for capturing the incoming mass flux, i.e., pentacene molecules [36].

5.4 Mechanical reinforcement Fundamentally, materials with excellent mechanical properties such as an extremely high modulus and tensile strength, such as graphene, are promising in the field of structural materials for constructing buildings, roads, ports, and so on. Although, graphene costs too much as a material in such fields where huge amounts of low-cost materials are usually consumed, several attempts have been made to use graphene as a small amount of additive to enhance the mechanical properties of typical structural materials such as plastics, ceramics, and metals.

5.4.1 Reinforcement of plastics The reinforcing effect of rGO dispersion in the matrix resin was investigated on rGOeepoxy composites with different dispersion states of the flakes [40]. rGO flakes prepared through liquid oxidation, followed by thermal exfoliation at 1050
C, were mixed with epoxy resin in ethanol in a planetary ball mill. A high shear stress applied during milling broke up the agglomeration of rGO flakes and improved their dispersion in the epoxy matrix (a highly dispersed composite). For comparison, the blended solution of rGO and epoxy resin was sonicated to obtain less-dispersed composites. In the composite containing highly dispersed rGO by only 0.2 wt%, fracture toughness KIC was markedly improved, more than the composite containing poorly dispersed rGO, as shown in Fig. 5.22C. The tensile strength was pronouncedly lowered with an increase in the rGO (B)

60

Highly-dispersed

50 40 30

Poorly-dispersed

0.2 0.1 0.0 rGO content / wt%

(C) 3.3 Highly-dispersed

3.1 2.9 2.7

Poorly-dispersed

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0.8

KIC / MPa/m1/2

70

Tensile modulus / GPa

Tensile strength / MPa

(A)

Highly-dispersed

0.6 Poorly-dispersed

0.4

0.2 0.1 0.0 rGO content / wt%

Figure 5.22 Composite of rGOeepoxy with different dispersion states [40]: (A) tensile strength, (B) tensile modulus, and (C) fracture toughness KIC. rGO, reduced graphene oxide.

394 Chapter 5 content in the poorly dispersed composite although there was no change on the highly dispersed composite (Fig. 5.22A). The tensile modulus increased slightly with an increase in rGO content irrespective of the dispersion state, as shown in Fig. 5.22B. The glass transition temperature (Tg) increased from 146.3 to 157.4
C and electrical conductivity increased from the order of 1011 S/m to the order of 108 S/m, respectively, by adding highly dispersed 0.2 wt% rGO. To enhance the reinforcement effects in a resin, the chemical modification of GO was investigated for better interactions with polymer compounds. The functionalization and reduction of GO were simultaneously performed in an ethanol solution of 4hydrazinopyridine by dispersing GO flakes at 80
C (pyridine functionalized and reduced graphene oxide [Py-rGO]) [41]. Py-rGO was functionalized by adding N-containing functional groups. The Py-rGO flakes thus prepared reinforced the polyimide matrix, as shown in Fig. 5.23A. The addition of 0.5 wt% Py-rGO markedly improved the tensile strength and modulus to 527 MPa and 58.6 GPa, respectively, from 104 MPa and 5.7 GPa for the pristine polyimide. The addition of Py-rGO also enhanced the degree of imidization of the matrix poly(amic acid) and oxygen barrier properties of polyimide films. Chemical defect-healing of Py-rGO through intramolecular cross-dehydrogenative coupling with the FeCl3 catalyst (H-Py-rGO) enhanced the reinforcing effect on the polyimide matrix, as shown in Fig. 5.23B [42]. The addition of H-Py-rGO by 0.5 wt% reinforced the tensile strength and modulus up to 661 MPa and 25.2 GPa, respectively.

600

60

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40

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20

0

1 0.5 3 5 0 Content of Py-rGO / wt%

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(B) Tensile modulus / GPa

Tensile strength / MPa

(A)

1000

50

800

40

600

30

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0 0.5 3 1 5 Content of H-Py-rGO / wt%

Tensile modulus / GPa

Graphene was used to reinforce thin films of elastomer (elastic polymer) 20 mm thick, which resulted in over a 50% increase in elastic modulus at a very low loading of 0.1 wt% [43]. Water-dispersible polyurethane (wPU) and natural rubber latex (NRL) were blended with GO via aqueous dispersion, and then GO was reduced using ascorbic acid. Composite films of rGO-wPU and rGO-NRL were prepared by dip-coating glass mandrel

0

Figure 5.23 Changes in tensile strength and modulus with the addition of reduced graphene oxide (rGO): (A) pyridine functionalized and reduced graphene oxide (Py-rGO) [41] and (B) defect-healed Py-rGO (H-Py-rGO) [42].

Mechanical properties and applications 395 in the dispersions. The content of rGO was controlled at 0.05, 0.1, and 0.2 wt% in the composites with wPU and at 0.08 wt% in the composites with NRL. Because of the high transparency of thin films (thickness of 20e30 mm), the high and homogeneous dispersion of rGO flakes in the NRL matrix was confirmed with optical microscopy. Because both wPU and GO were anionic, GO flakes were only sparsely covered by wPU, indicating weak interaction. However, rGO flakes were observed to be covered entirely with wPU, indicating strong interaction. For the composites of rGO-wPU, the tensile strength and ultimate elongation increased to 17 MPa and 749%, respectively, by adding rGO at 0.1 wt% from 13 MPa to 745% for the pristine wPU film. For the composites of rGONRL, similar reinforcing effects were observed: the tensile strength and ultimate elongation of the composite added by 0.08 wt% rGO were 10.9 MPa and 1129%, respectively, although the NRL film resulted in 9.4 MPa and 1088%. The reinforcement of polyvinyl chloride (PVC) films was successfully performed by mixing graphene flakes in an N,N-dimethyl formamide (DMF) solution with PVC, followed by drop casting on a glass plate and annealing at 100
C to obtain composite films 5e7 mm thick [44]. Graphene flakes were prepared from natural graphite flakes by cleavage in cetyltrimethylammonium bromide solutions in glacial acetic acid and were confirmed to have an average thickness of 1.18 nm (one to three stacked layers). By adding graphene flakes, the tensile strength of the film increased markedly from 24 MPa for the neat PVC to 55 MPa for the film containing 2 wt% graphene, which was associated with a decrease in the strain to failure from 124% to about 40% (refer to Fig. 2.62).

5.4.2 Reinforcement of ceramics Reinforcement of Si3N4 ceramics was reported to be possible by adding graphene [45e49]. The mixture of 1 wt% graphene flakes into Si3N4 containing 4 wt% Al2O3 and 6 wt% Y2O3 prepared by milling in water with the surfactant polyethylene glycol was compressed under the isostatic pressure of 20 MPa at 1700
C [46,47]. By adding 1 wt% graphene flakes, the fracture toughness KIC improved to 7.8e9.9 MPa m1/2 from 6.6e6.9 MPa m1/2 for neat Si3N4, depending on several kinds of multilayered graphene flakes used, although hardness was retained at 14.6e16.4 GPa. Using either GO or commercial graphene flakes 1 nm thick and a lateral size of 200 nm, composites with Si3N4 containing 2 wt% Al2O3 and 5 wt% Y2O3 were prepared by spark plasma sintering at 1625
C under a uniaxial pressure of 50 MPa in vacuum [48]. Using rGO, more marked reinforcement of Si3N4 was obtained; KIC improved to 10.4 MPa m1/2 by adding 4 wt%, although the increase in flexure strength was small, as shown in Fig. 5.24. The addition of graphene flakes resulted in less reinforcement than the addition of rGO, as shown in Fig. 5.24. Graphene nanoribbons prepared by unzipping MWCNTs were successfully used as fillers for Si3N4 reinforcement [49]. MWCNTs synthesized by the floating catalyst CVD

396 Chapter 5 (A)

(B)

1100

Flexure strength / MPa

KIC / Mpa.m1/2

10 rGO/Si3N4 8 6 graphene/Si3N4 4 0

2 6 4 Graphene content / vol%

8

rGO/Si3N4 1000 graphene/Si3N4

900 800 700

0

2 1 3 4 Graphene content / vol%

5

Figure 5.24 Composites of Si3N4 with reduced graphene oxide (rGO) and graphene flakes prepared by milling [48]: changes in (A) fracture toughness KIC and (B) flexure strength with graphene content.

method were unzipped by heating to 1000
C in Ar after oxidization via a procedure similar to Hummers’ method. Unzipped graphene nanoribbons thus prepared were mixed with Si3N4 powder in ethanol and then sintered under spark plasma. The KIC of the composite containing 4.3 vol% graphene nanoribbon increased to 6.4 MPa/m2 from the neat Si3N4 of 2.9 MPa/m2. The random dispersion of graphene nanoribbons was also thought to inhibit the grain growth of the Si3N4. This toughening of Si3N4 ceramics is reasonably thought to result from to crack deflection and bridging along crack paths, as shown in Fig. 5.25. The crack is deflected from the original path (white arrow) owing to the presence of the nanoribbon (Fig. 5.25A) and two Si3N4 grains are bridged by the nanoribbon (Fig. 5.25B and C). The matrix Al2O3 was also successfully reinforced by adding rGO flakes [50]. The fine powder of Al2O3 with an average size of 70 nm was mixed with GO in an aqueous suspension, followed by the reduction of GO using hydrazine monohydrate at 60
C and then sintered under spark plasma at 1300
C with 50 MPa pressure in Ar. The resultant (A)

(B)

(C)

Nanoribbon

Figure 5.25 Scanning electron microscopy images of indentation cracks in Si3N4 with 4.3 vol% graphene nanoribbons [49]. See text for explanation.

Mechanical properties and applications 397 composites containing 2 wt% rGO flakes had a density of 3.81 g/cm3 and KIC of 5.21 MPa m1/2, compared with 3.85 g/cm3 and 3.40 MPa m1/2 for neat Al2O3. The electrical conductivity of the composite delivered 172 S/m at room temperature; it was 13 orders of magnitude higher than that of neat Al2O3. The presence of rGO flake was also thought to hinder the grain growth of Al2O3. Several attempts were reported to reinforce GO in a more practical composite with ceramic as the material for actual construction [51e53]. The most widely used component of concrete as the major construction material is ordinary portland cement (OPC). OPC essentially consists of hydraulic calcium silicates (nCaO$SiO2), usually containing one or more forms of calcium sulfate, and iron, aluminum, or magnesium oxides according to European Standard EN 197e1. OPC paste with no additive is a brittle material characterized by a weak tensile strength owing to the presence of relatively large pores that initiate macrocracks [54]. To investigate reinforcing OPC with GO, 0.03 wt% GO with a lateral dimension of 0.5 mm2 and a thickness of approximately 1 nm was incorporated into OPC (ASTM C150, Type I) with a waterecement ratio of 0.5 using a high-speed shear mixer (up to 12,000 rpm for 30 s twice) [55]. The resultant mold was sealed with polyethylene sheets to prevent the escape of moisture. The porosity for OPC with or without GO after 28 days of curing is shown in Table 5.2. With 0.03 wt% GO, the GO-OPC composite had a total porosity of 28.2%, which was 13.5% lower than that for plain cement [56]. The decrease in porosity was attributed to the improvement in the degree of hydration in the GO-OPC composite. Adding GO decreased the total porosity of the cement paste. Fig. 5.26A shows the distributions of pore sizes for the OPC-GO composite and plain cement. In the GO-OPC composite, the population of smaller pores was large in contrast to the plain cement with larger pores. Adding GO increased the failure stress as well as the failure strain. The compressive strengths of the GO-OPC composite and plain cement are shown in Fig. 5.26B. At all test ages, the GO-OPC composite remarkably increased in compressive strength. The improvement in mechanical properties in the GO-OPC composite is explained by the Table 5.2: Porosity, average pore diameter, and total pores of ordinary portland cement with or without graphene oxide (GO) [56]. Mixes Plain cement GO cement

Total porosity (%) 32.8 28.2

Gel pore*1 (mL/g) 0.022 0.046

Capillary pore*2 (mL/g) 0.173 0.125

*: Classified according to Zhang and Zhan [57] and *1: