Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials 0128176806, 9780128176801

Table of contents :
Cover
Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials
Copyright
Dedication
Contents
About the author
Preface
Acknowledgments
Introduction
1 The fundamental aspects of spintronics
1.1 Introduction
1.2 Fundamental aspects of spintronics
1.2.1 Spin polarization
1.2.2 Spin relaxation
1.2.3 Spin injection
1.2.4 Ohmic injection
1.2.5 Tunnel injection
1.2.6 Ballistic electron injection
1.2.7 Hot electron injection
1.2.8 Spin transport
1.2.9 Spin injection and spin transport
1.2.10 Spin detection
1.2.11 Spin transfer
1.2.12 Spin coherence
1.2.13 Spin accumulation
1.2.14 Coherent spin transport through semiconductors and interfaces
1.2.15 Optical control of nuclear spins
1.2.16 Quantum dots—the artificial atoms in the solid state
1.3 Device principles
1.4 Different spintronics applications, device principles, and devices
1.4.1 Magnetic recording
1.4.2 Nonvolatile memories
1.4.3 MRAM (magnetoresistive random access)
Spin transistor
Quantum computer
1.5 Conclusion
References
Further reading
2 Introduction: carbon and carbon nanomaterials
2.1 Introduction
2.2 Carbon and carbon nanostructure materials in spintronics
2.3 Different forms of carbon and carbon nanostructure materials
2.3.1 Amorphous carbon, diamond-like carbon, and ultrananocrystalline diamond
2.3.2 Nanostructure carbon nanotube materials
2.3.3 Nanostructure graphene materials
2.3.4 Graphene oxide materials
2.3.5 Carbon nanoparticles and fullerene nanostructure materials
2.3.6 Other carbon–metal oxide/sulfide composites materials
2.4 Conclusion and perspectives of carbon and carbon nanostructure-based spintronics applications
References
Further reading
3 Magnetism and spintronics in amorphous/diamond-like carbon
3.1 Introduction
3.2 Magnetism of amorphous carbon and diamond-like carbon
3.3 Electrical and transport of amorphous carbon and diamond-like carbon
3.4 Magnetoresistance and spintronics of amorphous carbon and diamond-like carbon
3.4.1 Magnetoresistance of amorphous carbon films
3.5 Spin field effect transistor (FET)
3.6 Diamond-like carbon for magnetic storage disks
3.7 Conclusion and perspectives of amorphous carbon and diamond-like carbon in spintronics
References
Further reading
4 Magnetism and spintronics in carbon nanotubes
4.1 Introduction
4.2 Magnetism of carbon nanotubes
4.3 Spintronic devices
4.3.1 Magnetic tunnel junctions
4.3.2 Fabrication of magnetic tunnel junctions
4.3.3 Tunnel magnetoresistance in magnetic tunnel junctions
4.3.4 Application of magnetic tunnel junctions
4.4 Spin currents in magnetic tunnel junctions
4.4.1 Spin and charge transport
4.4.2 Spin polarization
4.5 Tunnel magnetoresistance in carbon nanotube–based spintronic devices
4.5.1 Spin-valve devices of carbon nanotubes
4.6 Conclusion and perspectives of carbon nanotubes in spintronics
References
Further reading
5 Magnetism and spintronics in graphene
5.1 Introduction
5.2 Making of graphene into magnetic materials
5.2.1 Ferromagnetism derived from hydrogenated zigzag-type pore edges graphene
5.2.2 Magnetism depending on pore edge termination by different foreign atoms
5.2.3 Doping and/or functionalization with transition metals
5.3 Spin generation and spin manipulation
5.3.1 Spin generation
5.3.2 Spin manipulation
5.4 Magnetism in graphene
5.5 Spin Hall effect and edge-derived spin phenomena
5.6 Spin injection, manipulation, and detection
5.7 Spin relaxation process
5.7.1 Hanle spin precession
5.7.2 Charged impurity scattering
5.7.3 Contact-induced spin relaxation
5.8 Spin relaxation in single-layer graphene and bilayer graphene
5.9 Electrical spin transport
5.9.1 Spin polarization
5.10 Spintronics magnetoresistance devices
5.11 Applications of graphene spintronics
5.11.1 Spin valve devices
5.11.2 Field-effect transistor
Ballistic transistors
5.11.3 Hall effect
5.11.4 Bipolar spintronics
5.12 Graphene-ferroelectric meta-devices
5.13 Conclusion and perspectives of graphene-based spintronics
References
Further reading
6 Magnetism and spintronics in graphene oxide
6.1 Introduction
6.2 Magnetization of graphene oxide
6.3 In-plane and out-of-plane magnetization/magnetic anisotropy
6.4 Magnetoresistance of graphene oxide
6.4.1 Mechanism of magnetic behaviors GO/reduced graphene oxide
6.4.2 Electrical transport mechanism of GO/reduced graphene oxide
6.5 Applications of graphene oxide spintronics
6.6 Conclusion and perspectives of graphene oxide–based spintronics
References
Further reading
7 Magnetism and spintronics in carbon nanoparticle/fullerene
7.1 Introduction
7.2 Carbon nanoparticle–based spintronics
7.3 Carbon nanosphere spintronics
7.4 Graphene-nano-dots spintronics
7.5 Fullerene-based spintronics
7.6 Conclusion and perspectives of carbon nanoparticles–based spintronics
References
8 Magnetism and spintronics in other carbon-based composite materials
8.1 Introduction
8.2 Carbon nanostructure-metal/nonmetal/metal-oxide composites in spintronics
8.3 Conclusion and perspectives of carbon-metal/nonmetal/composite-based spintronics
References
9 Challenges and emerging direction of carbon nanostructure materials in magnetism and spintronics
References
Index
Back Cover

Citation preview

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials Sekhar Chandra Ray Department of Physics, University of South Africa, Pretoria, South Africa

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. British Library Cataloguing-in-Publication Data A catalogue record for this book is available from the British Library Library of Congress Cataloging-in-Publication Data A catalog record for this book is available from the Library of Congress ISBN: 978-0-12-817680-1 For Information on all Elsevier publications visit our website at https://www.elsevier.com/books-and-journals

Publisher: Mathew Deans Acquisition Editor: Simon Holt Editorial Project Manager: Ana Claudia A. Garcia Production Project Manager: Selvaraj Raviraj Cover Designer: Harris, Greg Typeset by MPS Limited, Chennai, India

Dedication Dedicated to all the people for the betterment of their life

Contents About the author Preface Acknowledgments Introduction 1.

2.

3.

xi xiii xv xvii

The fundamental aspects of spintronics

1

1.1 Introduction

1

1.2 Fundamental aspects of spintronics

4

1.3 Device principles

12

1.4 Different spintronics applications, device principles, and devices

14

1.5 Conclusion

17

References

19

Further reading

21

Introduction: carbon and carbon nanomaterials

23

2.1 Introduction

23

2.2 Carbon and carbon nanostructure materials in spintronics

23

2.3 Different forms of carbon and carbon nanostructure materials

25

2.4 Conclusion and perspectives of carbon and carbon nanostructure-based spintronics applications

36

References

37

Further reading

45

Magnetism and spintronics in amorphous/diamond-like carbon

47

3.1 Introduction

47

3.2 Magnetism of amorphous carbon and diamond-like carbon

48 vii

viii

Contents

4.

5.

3.3 Electrical and transport of amorphous carbon and diamond-like carbon

52

3.4 Magnetoresistance and spintronics of amorphous carbon and diamond-like carbon

56

3.5 Spin field effect transistor (FET)

63

3.6 Diamond-like carbon for magnetic storage disks

67

3.7 Conclusion and perspectives of amorphous carbon and diamond-like carbon in spintronics

68

References

69

Further reading

73

Magnetism and spintronics in carbon nanotubes

75

4.1 Introduction

75

4.2 Magnetism of carbon nanotubes

77

4.3 Spintronic devices

83

4.4 Spin currents in magnetic tunnel junctions

87

4.5 Tunnel magnetoresistance in carbon nanotube based spintronic devices

93

4.6 Conclusion and perspectives of carbon nanotubes in spintronics

97

References

97

Further reading

102

Magnetism and spintronics in graphene

103

5.1 Introduction

103

5.2 Making of graphene into magnetic materials

105

5.3 Spin generation and spin manipulation

107

5.4 Magnetism in graphene

109

5.5 Spin Hall effect and edge-derived spin phenomena

113

5.6 Spin injection, manipulation, and detection

115

5.7 Spin relaxation process

117

Contents

6.

7.

ix

5.8 Spin relaxation in single-layer graphene and bilayer graphene

121

5.9 Electrical spin transport

123

5.10 Spintronics magnetoresistance devices

129

5.11 Applications of graphene spintronics

131

5.12 Graphene-ferroelectric meta-devices

138

5.13 Conclusion and perspectives of graphene-based spintronics

140

References

142

Further reading

150

Magnetism and spintronics in graphene oxide

151

6.1 Introduction

151

6.2 Magnetization of graphene oxide

153

6.3 In-plane and out-of-plane magnetization/magnetic anisotropy

165

6.4 Magnetoresistance of graphene oxide

165

6.5 Applications of graphene oxide spintronics

172

6.6 Conclusion and perspectives of graphene oxide based spintronics

176

References

176

Further reading

181

Magnetism and spintronics in carbon nanoparticle/fullerene

183

7.1 Introduction

183

7.2 Carbon nanoparticle based spintronics

184

7.3 Carbon nanosphere spintronics

186

7.4 Graphene-nano-dots spintronics

188

7.5 Fullerene-based spintronics

191

7.6 Conclusion and perspectives of carbon nanoparticles based spintronics

195

References

196

x

Contents

8.

9.

Magnetism and spintronics in other carbon-based composite materials

201

8.1 Introduction

201

8.2 Carbon nanostructure-metal/nonmetal/metal-oxide composites in spintronics

202

8.3 Conclusion and perspectives of carbon-metal/nonmetal/ composite-based spintronics

210

References

210

Challenges and emerging direction of carbon nanostructure materials in magnetism and spintronics 213 References

Index

214

217

About the author Sekhar Chandra Ray (PhD, University of the North Bengal, India) is a B-rated researcher in Physics from the National Research Foundation. He is currently a professor of physics at the University of South Africa in experimental condensed matter physics. Prof. Ray hails from India where he completed his doctoral studies which focused mainly on photovoltaic solar grade materials. He has worked as a research fellow and visiting scientist in Italy (INFM Fellow), Taiwan (NSC Fellow), Spain (ICMM, CSIC, Foreign Researcher Fellow, Ministry of Science and Technology, Spain), South Korea (Brain Pool Research Fellow, Govt. South Korea), and India (IACS, Visiting Scientist). His research group focuses on carbon nanostructure material in electronic structure/magnetic properties for the possible fabrication of spintronic devices application and bioimaging process. At present, Prof. Ray is working on different 2D-structure materials such as graphene, MoS2, stanene, silicene, and germanene. During his 22-year research career, he has published 125 peer-reviewed research articles, including seven in Nature Publishing Group (NPG) journal with more than 3800 citations in internationally recognized journals. Prof. Ray also acts as an editorial board member of Scientific Reports—NPG journal.

xi

Preface This book has been written to meet the basic requirement of researchers getting training in carbon and carbon nanomaterials for magnetism and spintronic applications. Carbon-based spintronics refers mainly to the spin injection and transport in carbon materials including carbon nanotubes, graphene, fullerene, and organic materials. In the last decade, extraordinary development has been achieved for carbon-based spintronics, and the spin transport has been studied in both local and nonlocal spin valve devices. A series of theoretical and experimental studies has been done to reveal the spin relaxation mechanisms and spin transport properties in carbon materials, mostly for graphene and carbon nanotubes. In this book, we provide a brief review on spin injection and transport in graphene, carbon nanotubes, fullerene, and organic thin films. The contents of the book are from different articles published in different journals from different research groups worldwide. I, the author of this book, am very much thankful to those publishers and the authors from whose publications we have collected all useful information that are presented in this book. This book consists of nine chapters namely, fundamental aspects of spintronics, carbon and carbon nanomaterials, magnetism and spintronics in amorphous/diamond-like carbon, magnetism and spintronics in carbon nanotubes, magnetism and spintronics in graphene, magnetism and spintronics in graphene oxide, magnetism and spintronics in carbon nanoparticles/fullerene, magnetism and spintronics in other carbon-based composites materials, and challenges and emergence direction of carbon nanostructure materials in magnetism and spintronics. We hope that this book will be useful for all researchers who are working on this research field. Any suggestions toward its further improvement will be thankfully acknowledged and incorporated in the next editions.

Sekhar Chandra Ray UNIVERSITY O F S OUTH AF RICA, JOHANNESBUR G, SOUTH AFRICA

xiii

Acknowledgments Professor Sekhar Chandra Ray wishes to extend his thanks to his friends, relatives, wife Mrs. Susmita Ray and son Master Shrishmoy Ray for their support in writing this book.

xv

Introduction The discovery and fabrication of new materials have opened the gate for new research fields in science and technology. The novel method of fabricating graphene, a purely 2D carbon lattice, and the discovery of the phenomenon of giant magnetoresistance in magnetic multilayers are not exceptions. The latter has brought about the creation of the new technological field of spintronics, which utilizes both spin and charge degrees of freedom of electrons. As for the former, many applications have been proposed; however, no practical devices have yet been developed in the field of spintronics. The aim of this book is to provide possible hints to overcome the difficulties in graphene applications in the field of spintronics by comparing the physical properties of graphene and magnetoresistive (MR) phenomena in spintronics. The book provides an overview and will be very useful for advanced undergraduate students and graduate students of physics, chemistry, and materials science and young researchers in nanotechnology and the field of spintronics. This book • covers the fundamental properties and various proposals for applications of graphene in spintronics, as well as an overview of charge and spin transport in graphene; • includes a description of physical phenomena in both graphene and spintronics; • provides illustrations and many references for easy understanding of the present status of graphene applications in spintronics; and • discusses the recent developments, future perspectives, and major challenges in these fields.

xvii

The fundamental aspects of spintronics

1

1.1 Introduction Information technology is one of the important issues in the 21st century. As the Moore law gradually loses its effect, conventional charge-based electronics will come to an end in the near future. Developing alternative high-speed and low-energy-consuming information technology is urgently needed. Many new methodologies have been proposed, such as molecular electronics, nano-electronics, spintronics, and quantum information techniques, among which spintronics is one of the most promising ones. Spintronics is a field of research exploiting the influence of the electron spin on electrical conduction. It is mainly known for the “giant magnetoresistance” (GMR) (Baibich et al., 1998; Binash et al., 1989) and the large increase of the hard disk capacity obtained with the read heads based on GMR. But the research on spintronics has also revealed many other interesting effects and is now developing along promising novel directions. Compared to other methodologies, spintronics is compatible with conventional electronics, thus many techniques used in conventional electronics can be directly extended to spintronics. “Spintronics,” known as spin electronics, involves the study of active control and manipulation of the intrinsic spin of the electrical charge of electron and its associated magnetic moment in solid-state system. The approach in the field of electronics is based on the up- or downspin of the carriers rather than on electrons or holes as in traditional semiconductor electronics. It is different from conventional electronics, which uses the electron’s charge degree of freedom for information processing; spintronics is devoted to incorporating the electron’s spin degree of freedom. In an ideal situation, there will purely be spin current and no charge current in the spintronics circuit, thus no heat will be created and wasted. There has been a great deal of recent interest in the concept of “spintronics” (Prinz, 1995, 1998). Spintronics is a multidisciplinary field whose central theme is the active manipulation of spin degree of freedom in solid-state system. Controlling and probing spin-polarized charge carrier (or manipulation of electron spin) in semiconductors and/or metals via electrical means, an attractive route toward the development of practical semiconductor/metal spintronic devices, which are expected to have a strong impact on future information processing and storage technologies. It is the use of a fundamental property of particles known as spin for information processing. It carries information in both the charge and spin of an electron, potentially offers devices with a great diversity of functionality in solid-state devices and other devices that exploit spin properties. In the case of the electron, the spin can in fact assume only the values 11/2 or 21/2: an eloquent invitation to use it to encode information, in analogy to bits “0” and “1” of the Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials. DOI: https://doi.org/10.1016/B978-0-12-817680-1.00001-9 © 2020 Elsevier Inc. All rights reserved.

1

2

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 1–1 Manipulation of electron movement (electron spin): spin-up (anticlockwise) and spin-down (clockwise).

binary code (Fig. 1
1). In ferromagnetic materials, the spin of the electrons can be modified from the outside, applying a magnetic field. When the magnetic field is removed, the new spin values are retained, that is, the encoded information remains firmly stored without need for power and without the risk of demagnetization. Spin transport and spin relaxation in semiconductors and metals are important solid-state physics issues that are included in the fundamental research along with new technology being implemented in the electronic storage technology. Spintronics helped in creating a prototype device that is used in the industry as a read head, and a memory-storage cell is the giantmagnetoresistive (GMR) sandwich structure, which consists of alternating ferromagnetic and nonmagnetic metal layers. Depending on the relative orientation of the magnetizations in the magnetic layers, the device resistance changes from small (parallel magnetizations) to large (antiparallel magnetizations). This change in resistance (also called magnetoresistance) is used to sense changes in magnetic fields. Electron spin can be identified as a magnetic field having one or two positions, known as “up” and “down”. This gives an extra two binary states to the conventional high and low-logic values, which are represented by simple currents. When the spin state is added to the mix, a bit can have four possible states, which can be called “up-high”, “up-low”, “down-high,” and “down-low.” These four states represent quantum bits or qubits.

▪▪▪

Why do we need spintronics? • Failure of Moore’s law • Low power consumption • Less electric current required

Chapter 1 • The fundamental aspects of spintronics

• • • • • •

Faster devices Larger storage capacity Smaller devices Less heat dissipation Spintronic memory is nonvolatile Spin manipulation is faster, hence greater reading and writing speed

3

▪▪▪

Metallic spintronics has already delivered functional devices (GMR read heads in largecapacity hard disk drives), and magnetic random access memory (RAM) (MRAM), of insulator spintronics [magnetic tunnel junctions (MTJs)]. The basic spin valve has evolved to a related thin-layered structure—MTJ—that displays giant tunneling magnetoresistance (TMR), a phenomenon where electrons tunnel through a thin insulator. This means the TMR signal is much larger than that from a GMR spin valve: indeed, it is almost 100 times larger. TMR is also the basis of magnetic RAM (MRAM), a nonvolatile memory that uses magnetic moments to retain data instead of electrical charges. However, the current basic physics research is mostly focused on semiconductor spintronics. Although creation of inhomogeneous spin distribution does not require energy penalty (in contrast to charge distributions of conventional electronics), spin is not conserved whereas charge is. Thus efforts in semiconductor spintronics research are focused on basic problems, such as coherent manipulation of electron spin at a given location, transporting spins between different locations within conventional semiconductor environment, all-electrical spin control via spin
orbit interactions, diluted magnetic semiconductors, and fixed or mobile spin qubits for quantum computing. Other possible applications of spintronics include high-speed magnetic filters, sensors, quantum transistors, and spin qubits for quantum computers (Steane, 1998; Loss et al., 1998; Burkard et al., 1999). Moreover, these “spintronic” devices might lead to quantum computer and quantum communication based on electronic solid-state devices, thus changing the perspective of information technology in the 21st century. More fundamental research will, however, be needed before a practical spintronic device can be demonstrated, as much remains to be understood about spin coherence, spin entanglement, spin dynamics, spin relaxation, spin transports, etc.—the different fundamental aspects of spintronics. Spintronics faces a number of challenges, including spin generation and injection, long distance spin transport, and manipulation and detection of spin orientation. In solving these issues, new concepts and spintronics materials were proposed one after another, such as half metals, spin-gapless semiconductors, and bipolar magnetic semiconductors. Topological insulators can also be viewed as a special class of spintronics materials, with their surface states used for pure spin generation and transportation. In designing these spintronics materials, first-principles calculations play a very important role. In this section, we attempt to give a brief discussion on the basic principles and theoretical design of these materials. Meanwhile, we also give some attention to antiferromagnetic (AFM) spintronics, which is mainly based on antiferromagnets and has aroused much interest in recent years.

4

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

1.2 Fundamental aspects of spintronics The fundamental aspects of spintronics are underlying the generation of carrier-spin polarization, spin coherence, spin entanglement, control spin and charge dynamics, spin injection, and spin-polarized transport in semiconducting/metallic electronic materials.

1.2.1 Spin polarization Spin polarization is the degree to which the spin, that is, the intrinsic angular momentum of elementary particles, is aligned with a given direction. This property may pertain to the spin, hence to the magnetic moment, of conduction electrons in ferromagnetic materials giving rise to spin-polarized currents. Spin polarization, not only of electrons, but also of holes, nuclei, and the excitations, can be defined as PX 5 XS/X; where the ratio of the difference XS 5 Xλ 2 X2λ and the sum X 5 Xλ 1 X2λ of the spin-resolved λ components for a particular quantity X. To avoid ambiguity as to what precisely is meant by spin polarization, both the choices of the spin-resolved components and the relevant physical quantity X need to be specified. Conventionally, λ is taken to be m or 1 (numerical value 11) for spin-up, and k or 2 (numerical value 21) for spin-down, with respect to the chosen axis of quantization (along the spin angular momentum, applied magnetic field, magnetization, or direction of light propagation). In ferromagnetic metals (FMs), it is customary to refer to m (k) as carriers with magnetic moment parallel (antiparallel) to the magnetization or, equivalently, as carriers with majority or minority spin (Tedrow et al., 1973). In semiconductors the terms majority and minority usually refer to relative populations of the carriers while m or 1 and k or 2 correspond to the quantum numbers mj with respect to the z axis taken along the direction of the light propagation or along the applied magnetic field (Meier and Zakharchenya, 1984; Jonker et al., 2003).

1.2.2 Spin relaxation Having established that one can generate spin-polarized earners, the most important issue is to determine how long these electrons remember their spin orientation. This is especially important for electronic applications, because if the spins relax too rapidly, the distances traversed by the spin-polarized current in a device will be too short to serve any practical purpose (Prinz, 1995). The spin
spin relaxation is the mechanism by which Mxy , the transverse component of the magnetization vector, exponentially decays toward its equilibrium value in nuclear magnetic resonance (NMR) and magnetic resonance imaging. It is characterized by the spin
spin relaxation time, known as T2, a time constant characterizing the signal decay. It is named in contrast to T1, the spin
lattice relaxation time. It is the time it takes for the magnetic resonance signal to irreversibly decay (37% i.e., 1/e) its initial value after its generation by tipping the longitudinal magnetization toward the magnetic transverse plane according to the following relation: Mxy ðt Þ 5 Mxy ð0Þe2t=T2 . T2 relaxation generally proceeds more rapidly than T1 recovery and different materials have different T2. When excited nuclear spins, lying partially in the transverse plane, interact with each other by the local magnetic field in-homogeneities on the micro- and nanoscales, their respective accumulated phases

Chapter 1 • The fundamental aspects of spintronics

5

deviate from expected values. While the slow- or nonvarying component of this deviation is reversible, some net signal will inevitably be lost due to short-lived interactions, such as collisions and diffusion, through heterogeneous space. T2 decay does not occur due to the tilting of the magnetization vector away from the transverse plane. Spin relaxation is very sensitive to the electronic band structure. Spins of conduction electrons decay because of the spin
orbit interaction and momentum scattering. At low temperatures (T # 20K), spin relaxation is caused by impurity scattering and is temperature independent. At higher temperatures, electrons lose spin coherence by colliding with phonons (phonons can induce a spin flip because in the presence of a spin
orbit coupling, electronic Bloch states are not spin Eigen states). Spin relaxation rate 1/T1 increases as temperature increases, with the growth becoming linear above the Debye temperature. This mechanism (Elliott, 1954; Yafet, 1963) is the most important spin relaxation mechanism in metals and semiconductors with inversion symmetry. Different spin relaxation processes have been found to be important in solids: • • • •

Elliot
Yafet D’yakonov
Perel’ Bir
Aronov
Pikus Hyperfine interaction

Nonequilibrium distribution of spins caused by interfaces or spin injections is brought back into equilibrium by these mechanisms, which can be obstacles for spintronics. Usually, suppressions of these effects are important issues in research and development though, sometimes, the relaxation would help fast device action.

1.2.3 Spin injection Spin injection in a nonmagnetic material is, in most cases, achieved by the creation of a nonequilibrium spin population (called spin accumulation) at the interface with a magnetic electrode. The rate of spin injection depends on the spin relaxation and dephasing mechanisms in the nonmagnetic material, which tends to restore the equilibrium in the accumulated spin population; the relatively long lifetimes of nonequilibrium electronic spins in semiconductors and metals, of about 1 ns, are essential for spintronic devices. The spin lifetimes can however increase to hundreds of nanoseconds in confined semiconductor heterostructures, which imply transport of coherent spin packets over hundreds of µm.

1.2.4 Ohmic injection In an FM the electrical conductivity of the majority-spin (spin-up) electrons differs substantially from minority spin (spin-down), resulting in a spin-polarized electric current. The most straightforward approach to spin injection is the formation of an ohmic contact between an FM and a semiconductor, anticipating a spin-polarized current in the semiconductor. However, typical metal
semiconductor ohmic contacts result from heavily doping the semiconductor surface, leading to spin
flip scattering and loss of the spin polarization.

6

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

Following up on earlier studies (Johnson et al., 1987; Van Son et al., 1987; Valet et al., 1993; HershÞeld et al., 1997) of diffusive spin transport, a recent work by Schmidt et al. (2000) has pointed out a fundamental problem regarding ohmic spin injection across ideal FM
nonferromagnet (NFM) interfaces. The effectiveness of the spin injection depends on the ratio of the (spin-dependent) conductivities of the FM and NFM electrodes, σF and σN, respectively. When σF # σN, as in the case of a typical metal, then efficient and substantial spin injection can occur, but when the NFM electrode is a semiconductor, σFcσN, and the spin-injection efficiency will be very low. Only for a ferromagnet where the conduction electrons are nearly 100% spin polarized can efficient spin injection be expected in the diffusive transport. A large number of materials apparently have such half-metal-ferromagnetic properties (De Grot et al., 1983; Kamper et al., 1988). Johnson et al. have proposed and pursued (Johnson 1998, 2001; Hammar et al., 1999, 2000a,b) an approach that may overcome this obstacle to spin injection by taking advantage of the splitting of the spin degeneracy of electrons confined in a semiconductor two-dimensional (2D) quantum well structure. The splitting is due to the spin
orbit effect that can arise from an asymmetry in the confining potential (Bychkov et al., 1994). The result can be an inducement of a nonequilibrium spin polarization if the 2D electron gas is carrying a current (Vorob’ev et al., 1979). However, as in the ohmic contact experiments, the small percentage change in device resistance, that is observed with changes in ferromagnet orientation, has led to suggestions of an alternative, local-Hall-effect explanation for the data and to other questions regarding this approach (Monzon et al., 2000; Wees, 2000; Hammar et al., 2000a,b).

1.2.5 Tunnel injection Alvarado and Renaud (Alvarado et al., 1992), using a scanning tunneling microscope (STM) with a ferromagnetic tip, showed that a vacuum tunneling process can effectively inject spins into a semiconductor. A recent extension of this has examined the effect of surface structure on spin-dependent STM tunneling (LaBella et al., 2001). The development of FM-insulator
FM-tunnel junctions with high magnetoresistance has also demonstrated that tunnel barriers can result in the conservation of the spin polarization during tunneling, suggesting that tunneling may be a much more effective means for achieving spin injection than diffusive transport. Theoretical work by Rashba (2000) has quantitatively developed the understanding of the potential effectiveness of tunnel injection. If the impedance of a barrier at an interface is sufficiently high, then the transport across that interface will be determined by the (spin-dependent) density of the electronic states of the two electrodes that are involved in the tunneling process. The current passing through the barrier is then sufficiently small enough for the electrodes to remain in equilibrium and the relative (spin-dependent) conductivities of the electrodes play no substantial role in defining or limiting the spin-dependent transport across the interface. Thus either a metal
insulator semiconductor tunnel diode or a metal
semiconductor Schottky barrier diode that uses an FM electrode can be expected to be an effective means for injecting spins into a semiconductor system.

Chapter 1 • The fundamental aspects of spintronics

7

1.2.6 Ballistic electron injection An alternative to tunnel injection is spin injection across ferromagnet
semiconductor interfaces in the ballistic regime, with the difference between the two spin conduction subbands of the FM and the conduction band of the semiconductor determining the spin-dependent interfacial ballistic electron transmission probability. It is generally assumed that the transverse momentum of an incident electron is conserved, and this determines the ballistic transmission and reflection probabilities of the interface (Kirczenow, 2001; Grundler, 2001). Also, once a spin-polarized electron enters the semiconductor electrode, the probability that it will be elastically scattered back into the ferromagnetic injector must be very small. If the device design also involves, for example, the spin-dependent capture of an injected carrier by another ferromagnetic electrode, then transport through the semiconductor region must be fully ballistic. However, if the objective is simply efficient spin injection, a three-dimensional ballistic point contact between a ferromagnet and a semiconductor should be effective. Recent experiments with point contacts formed between ferromagnetic and non-FMs have demonstrated the ballistic point-contact injection of high ( .40%) spin-polarized currents into the NFM (Upadhyay et al., 1998, 1999).

1.2.7 Hot electron injection Another spin technique involves the use of polarized “hot” electrons injection, having energies that are much greater than EF, by tunnel-injecting electrons into a ferromagnetic layer at energies cEF (Monsma et al., 1995; Jansen et al., 2001; Rippard et al., 2000). As the majority-spin and minority-spin electrons have much different inelastic mean free paths, hot electron passage through, for example, a 3-nm Co layer, is sufficient to result in a ballistic electron current that is more than 90% polarized (Rippard et al., 2000). This highly polarized hot electron current can then continue on to an underlying metal
semiconductor interface where a portion of the beam will enter the semiconductor, with the transmission probability being determined by energy and momentum constraints imposed by the band structure difference between the semiconductor and metal at the interface. If there is no substantial spin
flip scattering at the interface, the ballistic electron current entering the semiconductor will also be very highly polarized ( .90%), and the injection energy, relative to the bottom of the semiconductor conduction band, will be tunable through the tunnel injection bias. The disadvantage of hot electron injection is that the overall efficiency is low.

1.2.8 Spin transport Particular interest to the spin transport theory in semiconductor systems has been the question as to whether the quasiindependent electron model can adequately account for the experimental results, or whether many-body or correlated electron processes are important. The presence of spin-polarized carriers gives rise to both modified charge transport and

8

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 1–2 A schematic representation of the density of electronic states that are available to electrons in a normal metal and in a ferromagnetic metal whose majority-spin states are completely filled. E, The electron energy; EF, the Fermi level; N(E), density of states. Courtesy from Prinz, G.A., 1998. Magneto-electronics. Science 282, 1660
1663.

intrinsic spin transport, absent in the unpolarized case. Each of these aspects provides information about the degree of spin polarization, which can be utilized in spintronics. Spinpolarized transport will occur naturally in any material for which there is an imbalance of the spin populations at the Fermi level. This imbalance commonly occurs in FMs because the density of states available to spin-up and spin-down electrons is often nearly identical, but the states are shifted in energy with respect to each other (Fig. 1
2). This shift results in an unequal filling of the bands, which is the source of the net magnetic moment for the materials, but it can also cause the spin-up and spin-down carriers at the Fermi level to be unequal in number, character, and mobility. This inequality can produce a net spin polarization in a transport measurement, but the sign and magnitude of that polarization depends on the specific measurement being made. For example, an FM may be used as a source of spin-polarized carriers injected into a semiconductor, a superconductor, or a normal metal or can be used to tunnel through an insulating barrier. The nature of the specific spinpolarized carriers and the electronic energy states associated with each material must be identified in each case. The most dramatic effects are generally seen for the most highly polarized currents; therefore there are continuing efforts to find 100% spin-polarized conducting materials. These are materials that have only one occupied spin band at the Fermi level. Materials that are only partially polarized such as Fe, Co, Ni, and their alloys, which have a polarization P of 40%
50% (Soulen et al., 1998), are, however, adequate to develop technologically useful devices. Here the polarization P is defined in terms of the number of carriers n that have spin-up (nm) or spin-down (nk), as P 5 (nm
nk)/(nm 1 nk). Because of the spin polarization of an electron current, the effects seen in solid-state devices can be most readily visualized if one assumes that the current is 100% polarized (Fig. 1
2). In that case the only states that are available to the carriers are those for which the spins of the carriers are parallel to the spin direction of those states at the Fermi level. If the magnetization of the materials is reversed, the spin direction of those states also reverses.

Chapter 1 • The fundamental aspects of spintronics

9

Thus depending on the direction of magnetization of a material relative to the spin polarization of the current, a material can function as either a conductor or an insulator for electrons of a specific spin polarization. An analogy can be made with polarized light passing through an analyzer. However, in the optical case, crossing the polarization axis at 90 degrees prevents the transmission of the light, whereas for spin-polarized electrons, the magnetization must be rotated 180 degrees to stop electrical conduction. Flatté (2001) has extensively examined this issue in the diffusive transport regime and has concluded that an independent electron approach is quite capable of explaining measurements of spin lifetimes, particularly the room-temperature measurements. Sham et al. (2000) and Sham (1999) have been focusing on the very low temperature regime where collective electron processes may well be important in determining the spin relaxation rates and spin lifetimes, although experimental results in this regime are quite limited. On the device front, Flatté (2001) has considered the possibility of constructing unipolar electronic devices by using ferromagnetic semiconductor materials with variable magnetization directions. They have shown that such devices should behave very similarly to p
n diodes and bipolar transistors and suggest that they could be applicable for magnetic sensing, memory, and logic.

1.2.9 Spin injection and spin transport Spin-dependent electron transport, in bulk and nanostructures, is the presence of spinpolarized carriers that give rise to both modified charge transport and intrinsic spin transport and are absent in the unpolarized case. Each of these aspects provides information about the degree of spin polarization that can be utilized in spintronics. The transport of spinpolarized carriers across the semiconductor/metal interface where the metal is in the superconducting state. The study of semiconductor/superconductor hybrid structures has several important ramifications. In the context of spin unpolarized transport (Lambert et al., 1998; Beenakker, 1997), it has been demonstrated (De Franceschi et al., 1998) that this configuration can be used to examine the interfacial transparency that for a semiconductor/normal metal is typically limited by a native Schottky barrier. In the presence of spin-polarized carriers, semiconductor/superconductor structure can also serve to quantify the degree of spin polarization of a semiconductor and probe both potential and spin
flip interfacial scatˇ c´ et al., 1999). To understand such sensitivity to spin polarization and different tering (Zuti types of interfacial scattering, it is important to consider the process of Andreev reflection (Andreev, 1964), which governs the low bias transport. In this two-particle process, an incident electron of spin σ 5 m,k on a semiconductor/superconductor interface is reflected as a hole belonging to the opposite spin subband, back to the semiconductor region while a Cooper (1956) pair is transferred to the superconductor. The probability for Andreev reflection at low bias voltage is thus related to the square of the normal state transmission coefficient and can have stronger dependence on the junction transparency than the ordinary single-particle tunneling. For spin-polarized carriers with different populations in two spin subbands, only a fraction of the incident electrons from a majority subband will have a minority subband partner in order to be Andreev-reflected. In the superconducting state, for

10

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

an applied voltage smaller than the superconducting gap, single-particle tunneling is not allowed in the superconductor region and the modification of the Andreev reflection amplitude by spin polarization or junction transparency is manifested in transport measurements.

1.2.10 Spin detection The most obvious approach to the electrical detection of spin populations in semiconductors is to use the spin-dependent transport properties of semiconductor
ferromagnet interfaces. Experimental efforts, with this spin
valve detection scheme, have used ohmic contacts for the spin-collection electrode, but the same difficulties discussed earlier apply to spin collection, and it appears that for effective spin collection/detection, either a ballistic contact or a tunneling contact from the semiconductor to a ferromagnet will be required. If, however, for reasons of signal-to-noise, an efficient spin-dependent extraction of the injected spinpolarized current is required, the tunnel barrier has to be sufficiently thin so that (spin-dependent) tunneling transport into the ferromagnetic electrode is more probable than spin relaxation within the semiconductor (Rashba, 2000). An alternative spin detection technique is a potentiometric measurement, with a ferromagnetic electrode, of the chemical potential of the nonequilibrium spin populations (Hammar et al., 2000a,b; Johnson, 2001). With respect to the complete spin transistor device, an extensive analysis by Tang et al. (2000) has concluded that only for the case of ballistic transport throughout the device structure will the desired, electrical field
tunable spin precession be detectable as polarized electron transit through the semiconductor region. Moreover, they conclude, in accord with the initial suggestion of Datta et al. (1990), that a very narrow, single- or few-electron channel device structure will be required.

1.2.11 Spin transfer The spin-polarized current that flows from one relatively thick, and hence fixed, ferromagnetic layer, through a nonmagnetic layer, to another thin-film “free” nanomagnet can by spin-dependent scattering of the polarized current (Berger, 1996; Slonczewski, 1996, 1999) excite strong, uniform spin-wave processional modes in the nanomagnet (Tsoi et al., 1998; Myers et al., 1999; Rezende et al., 2000). In the absence of a strong external magnetic field, this spin-dependent scattering can also result in the reversal of the orientation of the magnetic moment of the free nanomagnet with the final orientation relative to the fixed layer being dependent on the direction of the current flow (Katine et al., 2000). This “spin-transfer” process opens up the possibility of new nanoscale devices for memory and other spin electronics applications (Weber et al., 2001). One application, in addition to direct current
addressable magnetic memory, might be the use of spin transfer to excite a uniform spin wave in a nanomagnet and then to use this nanomagnet as a precessing spin filter to inject a coherent spin pulse into a semiconductor structure.

Chapter 1 • The fundamental aspects of spintronics

11

1.2.12 Spin coherence Optical pulses are used to create a superposition of the basis spin states defined by an applied magnetic field and they follow the phase, amplitude, and location of the resulting electronic spin precession (coherence) in semiconductors, heterostructures, and quantum dots. The data identifies narrow ranges of doping concentrations where spin lifetimes in semiconductors are enhanced by orders of magnitude, culminating in the observation of spin lifetimes in semiconductors that exceed 100 ns. In heterostructures and quantum dots, nanosecond dynamics persist to room temperature, providing pathways toward practical coherent quantum magneto-electronics.

1.2.13 Spin accumulation When spin-polarized current is driven from a ferromagnetic film into a nonmagnetic film, faster than the spin polarization, it can diffuse away from the interface, and a nonequilibrium population of spin-polarized electrons builds up in a region of thickness Ls. This nonequilibrium magnetization is described as inequivalent chemical potentials for the upspin and downspin subbands of the normal metal. The chemical potential of the ferromagnet, however, is held in equilibrium by the intrinsic ferromagnetic
nonmagnetic metal interface; this is the same as an internal electric field, associated with the nonequilibrium spin accumulation that tries to drive electrons back across the interface and into the ferromagnet. Because spin and charge are both carried by the electron, a gradient of spin density results in an electric field, which can generate current flow or produce differences in voltage (Johnson et al., 1987). If the magnetic moments of the two ferromagnetic layers are parallel, spin accumulation in the base will create an electric field that pushes current into the collector, generating a positive current in the detector arm of the circuit. If, however, the magnetic moments are antiparallel, the spin-accumulation electric field at the base
collector interface has the opposite sign, current is pulled from the collector into the base, and a negative current is generated in the detector arm. The current flow through the detector can thus undergo bipolar modulation by modulating the direction of magnetization in the second layer. The device may be thought of as a nonvolatile computer memory element, storing information via the orientation of the second layer.

1.2.14 Coherent spin transport through semiconductors and interfaces Understanding the fundamental properties of spin transport in the solid state is essential for the development of semiconductor-based spintronics. In analogy to conventional devices, whose performance is characterized by carrier mobilities and lifetimes, spin mobilities and coherence times are figures of merit for spintronic devices. Recent theoretical work has shown that it is essential to consider the influence of electric fields induced by carrier motion to understand the motion of spin and that the room-temperature spin coherence times in bulk and quantum well structures appear to be dominated by precessional decoherence due to spin
orbit coupling (Flatté et al., 2000). These models describe how the low-field mobility

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

and diffusion of spin packets depend sensitively on the doping and reveal new opportunities to control spin interactions by engineering strain and crystal orientation (Ohno et al., 1999). The spatial selectivity and temporal resolution of optical techniques have been used to monitor the decoherence and dephasing of electron spin polarization during transport not only through bulk semiconductors but also across heterojunctions in engineered structures (Malajovich et al., 2000).

1.2.15 Optical control of nuclear spins Nuclear spins have been proposed as candidates for storing both classical and quantum information because of spin lifetimes that exceed those of electrons by at least several orders of magnitude and the degree of control provided by conventional NMR techniques. The experimental evidence shows that the ferromagnetic materials can be used to imprint nuclear spins in semiconductors (Kawakami et al., 2001), thereby offering an additional pathway for manipulating and storing information at the atomic scale.

1.2.16 Quantum dots—the artificial atoms in the solid state It has been proposed that the spin of an electron confined to quantum dots is a promising candidate for quantum bits and that array of quantum dots can be used in principle to implement a large-scale quantum computer (Loss et al., 1998; Burkard et al., 2000). Quantum operations in these proposals are provided by the coupling of electron spins in neighboring quantum dots by an exchange interaction between them. This interaction can be switched by applying controlled gate voltage pulses, thus allowing realization of fundamental quantum gates such as the exclusive OR. The readout of such a spin qubit can be performed efficiently as a spin-polarized electric current passing through the dot (Engel et al., 2001) or optically through integration in solid-state micro cavities (Imamoglu et al., 1999). Alternatively, qubit rotations can be implemented by local electrostatic shifting of the electron into a region with a different effective magnetic field, such as that which occurs at hetero-interfaces and in magnetic semiconductor structures.

1.3 Device principles The basic action in a spin-polarized device is shown in Fig. 1
3, where it is assumed that the electrons are traveling from an FM, through a normal metal, and into a second FM. When the magnetizations (or, equivalently, the magnetic moments) of the two FMs are in an aligned state, the resistance is low, whereas the resistance is high in the antialigned state. Actual devices are not generally fabricated in the orientation shown in Fig. 1
3, because they are made from thin films and the resistance perpendicular to the plane is too low. The common orientation, shown in Fig. 1
4, provides more useful resistance, but the physical picture of the spin-polarized transport is more complicated. The effect of the spin exclusion in antialigned films is still observed, but it results in high-interface scattering and

Chapter 1 • The fundamental aspects of spintronics

13

FIGURE 1–3 Schematic representations of spin-polarized transport from a ferromagnetic metal, through a normal metal, and into a second ferromagnetic metal for aligned and antialigned magnetic moments. [, Disallowed channel. Courtesy from Prinz, G.A., 1998. Magneto-electronics. Science 282, 1660
1663.

FIGURE 1–4 Schematic representations of transport that is parallel to the plane of a layered magnetic-metal sandwich structure for aligned (low resistance) and antialigned (high resistance) orientations. Courtesy from Prinz, G.A., 1998. Magneto-electronics. Science 282, 1660
1663.

“channeling” of the current into narrowed pathways (Fig. 1
4). When the films become aligned, both of these resistance-generating mechanisms are removed, and the device resistance decreases. This simple two-layer system is commonly referred to as a “spin valve” and is constructed so that the magnetic moment of one of the ferromagnetic layers is very difficult to reverse in an applied magnetic field, whereas the moment of the other layer is very easy to reverse. This easily reversed (or “soft”) layer then acts as the valve control and is sensitive to manipulation by an external field. The device can be used to measure or monitor those fields and can have numerous applications.

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

1.4 Different spintronics applications, device principles, and devices • Spintronic devices are used in the field of mass-storage devices. • It is used to compress massive amounts of data into a small area, as an instance, approximately one trillion bits per square inch (1.5 Gbit/mm2) or roughly, 1 TB data can be stored on a single-sided 3.5v diameter disk. • Spintronics is also used in the medical field to detect cancer. • Spintronics technology in general holds promises for digital electronics. It has been tested in mass-storage components, namely, hard drives.

1.4.1 Magnetic recording The first application to produce a substantially large economic impact was that for the read heads in magnetic disk recorders, which are components of every computer. The read head senses the magnetic bits that are stored on the media (disks or tapes). This information is stored as magnetized regions of the media, called magnetic domains, along tracks (Fig. 1
5). Magnetization is stored as a “0” in one direction and as a “1” in the other. Where two of these oppositely magnetized domains meet, there exists a domain wall, which is a microscopic region of 100
1000 Å (depending on the material used in the media). Although there is no magnetic field emanating from the interior of a magnetized domain itself, uncompensated magnetic poles in the vicinity of the domain walls generate magnetic fields that extend out of the media. It is these fields that are sensed by the GMR element. Where the heads of two domains meet, uncompensated positive poles generate a magnetic field directed out of the media, and where the tails of two domains meet, the walls contain uncompensated negative poles that generate a sink for magnetic lines of flux returning back into the media. The element is fabricated (Tang et al., 1994) so that the magnetic moment in the easily reversed layer lies parallel to the plane of the media in the absence of any applied fields. The magnetic moment in the fixed magnetic layer of the GMR element is oriented perpendicular to

FIGURE 1–5 A schematic representation of a GMR read head (green) that passes over recording media containing magnetized regions. The magnetization direction of the soft layer in the head responds to the fields that emanate from the media by rotating either up or down. The resulting change in the resistance is sensed by the current, I, passing through the GMR element. Courtesy from Prinz, G.A., 1998. Magneto-electronics. Science 282, 1660
1663.

Chapter 1 • The fundamental aspects of spintronics

15

the plane of the media. Thus when the head passes over a positive domain wall, the magnetic field pushes the easily reversed magnetic moment up, and when the head passes over a negative domain wall, the magnetic moment is pulled down. The measured resistance of the GMR element thus increases (for more antialigned layers) or decreases (for more aligned layers). The design goal for this element is to obtain a maximum rate of change in the resistance for a change in the sensed field. Typically, changes in resistance of 1% per oersted are reported.

1.4.2 Nonvolatile memories “Nonvolatile” refers to information storage that does not “evaporate” when power is removed from a system. Magnetic disks and tapes are the most widespread nonvolatile informationstorage media, because of their long storage lifetime, low cost, and lack of any wear-out mechanism. Computer core memory itself was nonvolatile before the introduction of semiconductor RAM in the early 1970s. The original core memory acquired its name because it was assembled from magnetic transformer cores that were fabricated out of insulating magnetic ferrite materials. These transformer cores were tiny toroidal rings that were threaded with fine copper wires. Current pulses through the wires could magnetize the cores either as right- or left-handed to store a 0 or a 1; each core was a bit. The information was read by current pulses that could test the core’s direction of magnetization through an inductively induced pulse in another wire. Although this memory was slow and expensive, and had low density by today’s standards, it was the industry standard during the 1950s and 1960s, and had the advantage that when power was removed, the stored information remained intact.

1.4.3 MRAM (magnetoresistive random access) Magnetoresistive random access memory (MRAM) is a nonvolatile and nondestructive readout memory, which is based on a magnetic anisotropy energy to retain information and the principle of magnetoresistance to retrieve information (Åkerman, 2005). Fig. 1
6 shows an illustration of MRAM architecture. Arrays of several MRAM cells form a memory device. A typical MRAM cell has a transistor and a magnetoresistive element, quite similar to a DRAM (dynamic RAM), which contains a transistor and a capacitor. While the charge stored in the capacitor of a DRAM defines its memory state, the resistance of the magnetoresistive element determines 1 and 0 states. A transistor for every MRAM cell is required as the absolute difference between the resistances, and hence, the voltages of two states are not high enough to function without a transistor. Moreover, the transistor also provides the current required for the write operation. A memory device follows at least three key requirements: (1) the proposed device should be able to store information. If the information is stored for long periods of time even without power, then it is called a nonvolatile memory device, (2) there should be mechanisms to read information from the device, and (3) there should be mechanisms to write information onto the device. To achieve these requirements in MRAM, researchers have designed and investigated various kinds of MRAM schemes in the past. In MRAM, these functions are performed as following: (1) The read operation is carried out by sensing the resistance

16

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 1–6 Schematic view of an array of MRAM cells in typical memory architecture. The orange box typically includes a transistor and a magnetic tunnel junction element. Courtesy from Bhatti, S., et al., 2017. Spintronics based random access memory: a review, Mater. Today 20 (9), 530
548.

difference between two states of a magnetoresistive device. (2) The storage of information relies on the magnetic retention properties, arising from the magnetic anisotropy of storage layer. (3) The write operation is performed by changing the orientation of storage layer magnetization, which can be achieved by inducing a magnetic field. The first storage element used in MRAM for storing information was based on spin valve structure, which mainly consisted of two ferromagnetic layers sandwiching a nonmagnetic conductive layer. The two ferromagnetic layers are called the free/soft layer and the hard/ pinned layer, respectively. An AFM layer is used in proximity of or in contact with the pinned layer in order to pin the magnetization direction of the layer, which should not be reversed during the operation of the memory device. The storage principle of an MRAM is based on the energy barrier (EB) required to switch the magnetization of a single-domain magnet from one direction to the other. The magnetization will be fixed in a particular direction if the EB for magnetization reversal is high enough to overcome the external stray fields and the thermally assisted reversal of magnetization. This storage principle is very similar to that used in magnetic recording, although the way the materials are designed and the information is written, are different. In an MRAM, the magnetization direction of the reference layer or pinning layers (PL) is fixed and only the magnetization direction of the free layer (FL) varies to store “0” and “1” states. Since the direction of the reference layer must never be changed, it is made of

Chapter 1 • The fundamental aspects of spintronics

17

materials that have a huge EB. The FL is designed with materials that have a magnetic anisotropy, just sufficiently high enough to store the magnetization for certain years (typically 10 years in the case of magnetic recording). The EB that helps store the information is typically proportional to KuV (where Ku is the magnetic anisotropy constant and V is the volume of the FL). This energy must be much larger (60 times, for storage time longer than 10 years) than the thermal energy kBT. In certain cases, the EB may be different from KuV and hence, the thermal stability factor is simply written as D(5EB/kBT). Although a high anisotropy is preferred for storing information, the anisotropy of these materials cannot be too high, as their magnetization direction needs to be oriented at will, to write 0 and 1 states. Based on the storage mechanism, the MTJs may be classified into two types: (1) in-plane MTJ, which has magnetization of ferromagnetic layers in the film plane and (2) perpendicular MTJ having the magnetization perpendicular to the film plane (Sbiaa et al., 2011).

Spin transistor It is a three-terminal bipolar device consisting of a normal metal sandwiched between two FM layers. Current is driven from the first ferromagnetic film (emitter) into the nonmagnetic metal (base) and back through the battery. A symmetric circuit arm connecting the second ferromagnetic film (collector) to the base contains a current detector.

Quantum computer A quantum computer is any device for computation that makes direct use of distinctively quantum-mechanical phenomena, such as superposition and entanglement, to perform operations on data. In a classical (or conventional) computer, information is stored as bits; in a quantum computer, it is stored as qubits (quantum bits). The basic principle of quantum computation is that the quantum properties can be used to represent and structure data and that quantum mechanisms can be devised and built to perform operations with this data.

1.5 Conclusion Spintronics materials incorporate a huge number of magnetic materials, including magnetic metals, topological insulators (TIs), and magnetic semiconductors. To overcome the challenges faced by spintronics, several new conceptual materials have emerged, such as half metals, spin-gapless semiconductors (SGSs), bipolar magnetic semiconductors (BMSs), and asymetric antiferromagnetic semiconductors (AAFMSs). The proposal of these materials is primarily based on first-principles calculations, and some of them have been verified by subsequent experiments, while others still left as theoretical models. Because of the abundance of family members and magneto-electronic properties, double perovskites, Heusler alloys, transition metal chalcogenides and pnictides, as well as graphene and graphitic nanostructures form fertile soils for the design of spintronics materials. Spintronics still has a long way

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

to go before its wide application in real life. Further development of spintronics is expected to focus on the following areas: 1. Proposing new concepts and spintronics materials. Up till now, most spintronics materials are featured by only one specific property or function. Nevertheless, it is possible to create a magnetic material with multifunctionality. This is not only interesting for fundamental research but also for fabricating multifunctional spintronic devices. 2. Seeking more candidate compounds for different classes of spintronics materials. For example, continuous efforts are needed to search for more SGSs, BMSs, and AAFMSs, since their family members are still much limited. In this respect, the Materials Genome Initiative serves as a good opportunity. Meanwhile, recently developed crystal-structure search methods, such as USPEX (Glass et al., 2006) and CALYPSO (Wang et al., 2010), provide us with very powerful tools. 3. Developing spintronics materials that can work at room temperature. To do this, efforts should be devoted to obtaining half metals with high Curie temperature and wide halfmetallic gap, TIs with sizable spin
orbit coupling gap, and magnetic semiconductors with room-temperature magnetic ordering and large spin polarization. 4. Designing low-dimensional spintronics materials to construct spintronic devices at nanoscale. This is particularly important for fabricating devices with high integration density and speed. 5. Understanding and engineering the interface between magnetic metals and semiconductors. In an ideal spintronic circuit, spin current will cross the interface without losing its polarization. Realizing such an ohmic contact for spins remains a challenge. Meanwhile, control of the magnetic interaction between two ferromagnetic components in the integrated circuit is also important. 6. Exploring AFM spintronics. Lots of works are still needed to improve the understanding and comprehension of AFM spintronics both from theoretical and experimental aspects. Also, developing novel device models and practical AFM materials is required. To continue the rapid pace of discoveries, considerable advances in our basic understanding of spin interactions in the solid state along with developments in materials science, lithography, miniaturization of optoelectronic elements, and device fabrication are necessary. The progress toward understanding and implementing the spin degree of freedom in metallic multilayers and, more recently, in semiconductors is gaining momentum as more researchers begin to address the relevant challenges from markedly different viewpoints. Spintronics read head sensors are already impacting a multibillion dollar industry and magnetic RAM using metallic elements will soon impact another multibillion dollar industry. With contributions from a diversity of countries and fields, including biology, chemistry, physics, electrical engineering, computer science, and mathematical information theory, the rapidly emerging field of spintronics promises to provide fundamentally new advances in both pure and applied sciences as well as have a substantial impact on future technology.

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2791. Katine, J.A., et al., 2000. Current-driven magnetization reversal and spin-wave excitations in Co/Cu/Co pillars. Phys. Rev. Lett. 84, 3149
3152. Kawakami, R.K., et al., 2001. Ferromagnetic imprinting of nuclear spins in semiconductors. Science 294, 131
134. Kirczenow, G., 2001. Ideal spin filters: a theoretical study of electron transmission through ordered and disordered interfaces between ferromagnetic metals and semiconductors. Phys. Rev. B 63, 054422:1
12. LaBella, V.P., et al., 2001. Spatially resolved spin-injection probability for gallium arsenide. Science 292, 1518
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941. Loss, D., et al., 1998. Quantum computation with quantum dots. Phys. Rev. A 57, 120
126. Malajovich, I., et al., 2000. Coherent transfer of spin through a semiconductor heterointerface. Phys. Rev. Lett. 84, 1015
1018. Meier, F., Zakharchenya, B.P., 1984. Optical Orientation. Elsevier, Northholland. Monsma, D.J., et al., 1995. Perpendicular hot electron spin-valve effect in a new magnetic field sensor: the spin-valve transistor. Phys. Rev. Lett. 74, 5260
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Schmidt, G., et al., 2000. Fundamental obstacle for electrical spin injection from a ferromagnetic metal into a diffusive semiconductor. Phys. Rev. B 62, R4790
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16325.

Further reading Meier, F., et al., 2003. Quantum computing with spin cluster qubits. Phys. Rev. Lett. 90, 047901:1
4.

2

Introduction: carbon and carbon nanomaterials 2.1 Introduction Traditional materials, including metals (Johnson et al., 1985), semiconductors (Baumberg et al., 1994; Lou et al., 2007) for spintronics, have been studied for many years and some of these efforts have already turned out to be industry products to serve modern industry. However, the carbon-based spintronics is just at its beginning. Theories have predicted many excellent spintronic properties in carbon-based materials, and this field has gained high attention. Electronics is based on the manipulation of electrons and other charge carriers, but in addition to charge, electrons possess a property known as spin. When spin is manipulated with magnetic and electric fields, the result is a spin-polarized current that carries more information than is possible with charge alone. Spin transport electronics, or spintronics, is a subject of active investigation in science to technology. Spintronics is one of the most promising next-generation information technologies, which uses the spins of electrons as information carriers and possesses potential advantages of speeding up data processing, high circuit integration density, and low energy consumption. As described in Chapter 1, The fundamental aspects of the spintronics, spintronics faces a number of challenges, including spin generation and injection, long-distance spin transport, manipulation and detection of spin orientation. In solving these issues, new concepts and spintronics materials were proposed one after another, such as half metals, spin gapless semiconductors, and bipolar magnetic semiconductors. Topological insulators can also be viewed as a special class of spintronics materials, with their surface states used for pure spin generation and transportation. In designing these spintronics materials, first-principles calculations play a very important role. This chapter attempts to give a brief review of the basic principles and theoretical design of these carbon and carbon nanostructure materials. In addition, synthesis process, functionalization process, doping process, and defects process of these materials are described briefly to use carbon and carbon nanostructures as magnetic materials and to use them in different magnetic applications.

2.2 Carbon and carbon nanostructure materials in spintronics Carbon-based nanostructure materials are believed to be better choices than the traditional metals and semiconductors in terms of their long spin relaxation distances. The spin current Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials. DOI: https://doi.org/10.1016/B978-0-12-817680-1.00002-0 © 2020 Elsevier Inc. All rights reserved.

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

can travel macroscopic distances in graphene and carbon nanotubes (CNTs) due to weak spin orbit and hyperfine couplings in carbon materials (Huertas-Hernando et al., 2006). Moreover, these carbon-based materials are also easy for the implementation of large-scale spintronic logic circuits. Carbon-based spintronics is an exciting topic with vast potential, most notably in fast memory devices, logic gates, transistors, and capacitors for computers, tablets, and handheld devices (Kuemmeth et al., 2010; Chen et al., 2013; Naber et al., 2007). This is a very attractive alternative to the existing technologies because the synthesis and processing of carbon-based nanostructures can be relatively inexpensive and easy. Furthermore, all hydrocarbon products can significantly reduce the need for critical materials used by the semiconductor industry (e.g., indium and gallium). Realizing this potential requires an understanding of the origins of magnetic coupling and improvement of this key property, enabling the design of new and significantly enhanced spintronic devices. A typical spintronic device is the so-called spin valve. The most simple spin valve device contains two ferromagnetic electrodes: using one ferromagnetic electrode for spin current injection and another for spin signal detection. For its practical use in industry, the spin signal should be large enough. A large spin signal can be received in weak spin relaxation system, where spin can travel a long distance without being scattered. Zigzag edges in sp2 carbon have demonstrated spin-ordered properties in such structures as graphene nanoribbons with bases of graphene (Dutta et al., 2010; Daniels et al., 2014) and coronene (de Aguiar, 2014); open nanotubes (Huang et al., 2010), nanowiggles (Costa Girão et al., 2011), nanorings (Potasz et al., 2010, 2011; Grujic´ et al., 2013), nanomeshes (Yang et al., 2011a,b), nanodots (Hod et al., 2008), and many more (Wang et al., 2008; Yazyev, 2010). While each of these individual structures has interesting properties, spin alignment between graphene layers gets coupled, allowing for more complicated device designs (Guinea, 2010). Simulations have demonstrated that external magnetic fields can change the spin ordering on zigzag edges (Szałowski, 2013), and even static electric field might have the potential to influence the spin properties of nanostructured carbon (Son et al., 2006; Sheng et al., 2014). Carbon chain links have been proposed as another method of transferring spin information between flakes (Zhou et al., 2011), and theoretical studies using metal contacts and side groups have been published (Deng et al., 2012; Zhang et al., 2014). However, the current literature on structural design often fails to recognize the existence of topological frustration, a frequent property of such devices that greatly decreases the likelihood of their experimental realization. Topological frustration occurs when all pz orbitals cannot form p bonds in sp2-hybridized carbon due to topological reasons, leading to open electronic shells (Longuethiggins, 1950). In classical chemistry, these structures cannot possess a complete Kekulé structure, as not every sp2 carbon is double-bonded exactly once, leading to free-radical formation. These are considered non-Kekuléan (or “concealed”) structures (Randic, 2003). Attempts at synthesis of non-Kekuléan hydrocarbon structures have resulted in unstable structures that decay at ambient conditions, destroying the desired properties at room temperature (Inoue et al., 2001) or which oxidize in a few days when exposed to ambient atmosphere (Tukada et al., 1992). Stable non-Kekulé structures exist but generally require the introduction of noncarbon atoms (Koide et al., 2010; Ueda et al., 2012).

Chapter 2 • Introduction: carbon and carbon nanomaterials

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While new methods for producing carbon nanostructures with spintronic properties and exacting edge designs are still being proposed (Gao et al., 2009) or experimentally realized (Campos et al., 2009), much of the current literature is generally unaware of the adverse effects of topological frustration in the final structures.

2.3 Different forms of carbon and carbon nanostructure materials Carbon-based materials possess versatile physical and chemical properties (Makarova et al., 2006). The outer 2s and 2p shells in a carbon atom can hybridize in three different ways to form sp, sp2, and sp3 orbital wave functions. For the carbon and carbon nanostructure materials consisting of sp2 bonding (crystalline graphitic), sp3 bonding (crystalline diamond), and a mixture of sp2 and sp3 bonding types (glassy carbon), a dominant factor in determining their physical properties is the sp2/sp3 ratio.

2.3.1 Amorphous carbon, diamond-like carbon, and ultrananocrystalline diamond Amorphous carbon (a-C) is a carbon material without long-range crystalline order. Shortrange order exists, but with deviations of the interatomic distances and/or interbonding angles with respect to the graphite lattice as well as to the diamond lattice. The term a-C is restricted to the description of carbon materials with localized π electrons as described by Anderson (1958). Deviations in the C C distances .5 (i.e., 6 Δx/X0 . 0.05, where X0 is the interatomic distance in the crystal lattice for the sp2 as well as for the sp3 configuration) occur in such materials, as well as deviations in the bond angles because of the presence of “dangling bonds.” a-C films could be doped with boron as well as nitrogen to produce pand n-type semiconductive materials for the use of different electronic and magnetic applications. Nitrogen doping of a-C films produces a wide bandgap n-type semiconducting material. Because of its excellent and unique properties, such as low or negative electron affinity, low work function, high thermal conductivity, and chemical inertness, nitrogendoped a-C films are being considered as potential materials for cold cathodes in microelectronic devices and flat-panel display applications. Mechanical properties of a-C films depend on methods and the amount of hydrogen atoms. It is necessary to clarify atomic structure and to understand the phenomena caused by the sliding friction and the mechanism. Diamond-like carbon (DLC) is an a-C material that contains a significant fraction of carbon bonded in sp3 hybridization (Robertson, 2002). DLC exists in several different forms of a-C materials that display some of the unique properties of diamond. DLC coatings can be amorphous, more or less flexible, hard, strong, and slick according to the composition and processing method. Film formation can be obtained by plasma-assisted chemical vapor deposition (CVD), ion beam deposition, sputter deposition, and RF plasma deposition, etc. The sp3-hybridized carbon provides diamond-like properties, such as high hardness, high wear resistance, low friction coefficient, and chemical inertness (Robertson, 2002), which

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makes them useful for industrial applications. For example, DLC is used as protective coating for mechanical tool heads, car engine parts, and hard disk drives (Hauert, 2004). Recently, there has been significant interest in modification of DLC for advanced applications, some of which include ultracapacitors, gas sensors, and dilute ferromagnetic semiconductors
Santana et al., 2009). (Candelaria et al., 2012; Markwitz et al., 2015; Paul et al., 2012; Colon Since DLC films are mechanically stable with tunable electrical conductivity (Grill, 1999), they act as an effective base for applications that require multifunctional advanced materials. A number of methods (Robertson, 2002; Murmu et al., 2014; Markwitz et al., 2014) have been used for the preparation, functionalization, and/or doping of DLC films. Ultrananocrystalline diamond (UNCD) is a unique form of carbon with grain sizes in the 3 5 nm region. This nanostructure has profound implications on electronic transport, as B10% of carbon is at the grain boundaries. Thus this material has significant π bonding that governs the majority of the electrical conductivity due to the lower energy gap of π πT transitions relative to σ σT transitions. The addition of nitrogen into the gas phase during deposition promotes n-type conductivity, due to the increase in the density of states associated with π bonding. This material is not doped in the conventional sense, and its applications lie in the electrode/metallic conductivity region rather than in the more moderately doped active device regime. The discovery of fine-grain (3 5 nm) nanostructured carbon films deposited from hydrogen poor plasmas has led to significant research in the properties and applications of this material, named UNCD (Gruen, 1999). By reducing the relative concentration of atomic hydrogen during the CVD of diamond, it is possible to increase the renucleation rate to the point where all columnar structure is lost. This results in films with grain sizes that are independent of film thicknesses, in stark contrast to conventional diamond deposition. Thus UNCD does not grow in the standard van der Drift regime (Van der Drift, 1967). The fine-grain size of this material results in a very high grain boundary concentration (B10%), and thus electronic transport phenomena are more complex than those based on classical infinite lattice-based band approximations. Disordered semiconductors also relax the constrictions of the infinite lattice on impurity elements, a particular issue for dense materials, such as diamond, as it allows impurities of low solubility to be taken up at grain boundaries. The density of UNCD is significantly lower than conventional diamond due to the high percentage of material at the grain boundaries, and this region may contain many impurities that are insoluble in diamond. The enhancement of the grain boundary volume and the simultaneous promotion of n-type conductivity when nitrogen was included in the growth chemistry led to the hypothesis that nitrogen can be incorporated in the grain boundary and produces n-type doping (Bhattacharyya et al., 2001). However, in conventional amorphous and disordered semiconductors, the relaxation of long-range order, in fact, reduces the possibilities of doping due to impurity atoms being able to take up their preferred chemical bonding. Thus doping is severely constrained in amorphous systems, as donors and acceptors must take substitutional sites for the most efficient doping with the least lattice perturbation. Substitutional n-type doping of diamond has been problematic principally due to the insolubility of donor atoms in the dense lattice. Phosphorus doping has been the most

Chapter 2 • Introduction: carbon and carbon nanomaterials

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successful till date, and this dopant has been used to demonstrate diamond p n junctions (Koizumi et al., 2001). However, the activation energy in diamond is rather high at 0.6 eV (Nesladek, 2005), leading to low room temperature conductivities. Despite these setbacks, phosphorus remains the n-type dopant of choice for active electronic devices. UNCD has applications where high conductivities are required at the cost of mobility, due to the disordered nature of the material. These applications include electrochemical electrodes, field emission, and heterostructures where the n-type UNCD acts as a source of electrons. UNCD can be made highly conductive by the addition of nitrogen into the gas phase during deposition and this conductivity is n-type. The conductivity is not due to doping, but due to the manipulation of the nanostructure of the material, leading to enhanced sp2 regions and midgap states. This leads to low-mobility hopping-type conduction processes and impurity band conduction, with very high carrier concentrations and low mobilities. UNCD can be highly useful where high carrier concentrations are required, such as in electrochemical electrodes, field emission, heterostructures, and high-temperature stable ohmic contacts.

2.3.2 Nanostructure carbon nanotube materials Carbon-based nanostructure materials, especially graphene and CNTs, are believed to be better choices than the traditional metals and semiconductors in terms of their long spin relaxation distances. The spin current can travel macroscopic distances in graphene and CNTs due to weak spin orbit and hyperfine couplings in carbon materials (HuertasHernando et al., 2006). Moreover, these carbon-based materials are also easy for implementation of large-scale spintronic logic circuits. Since 1999 spin current has been successfully injected into CNTs by using ferromagnetic electrodes (Tsukagoshi et al., 1999; Zhao et al., 2002; Kim et al., 2002; Krompiewski, 2005), and after 2006, spin transport experiments were performed for the separation of graphene (Hill et al., 2006; Tombros et al., 2007; Cho et al., 2007), which was first isolated from graphite by Novoselov et al. (2004). A typical spintronic device is the so-called spin valve. The most simple spin valve device contains two ferromagnetic electrodes: using one ferromagnetic electrode for spin current injection, and another for spin signal detection. For its practical use in industry the spin signal should be large enough. A large spin signal can be received in weak spin relaxation system, where spin can travel a long distance without being scattered. Carbon-based materials, especially graphene and CNTs, are believed to be better choices than the traditional metals and semiconductors in terms of their long spin relaxation distances. The spin current can travel macroscopic distances in graphene and CNTs due to weak spin orbit and hyperfine couplings in carbon materials. Moreover, these carbon-based materials are also easy for the implementation of large-scale spintronic logic circuits. CNTs are allotropes of carbon with a cylindrical nanostructure. These cylindrical carbon molecules have unusual properties, which are valuable for the nanotechnology, electronics, optics, and other fields of material science and technology. Owing to the material’s exceptional strength and stiffness, nanotubes have been constructed with a length-to-diameter ratio of up to 132,000,000:1 (Wang et al., 2009a,b), significantly larger than that for any other

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material. Since the discovery of CNTs by Iijima (1991), tremendous interest is shown for CNT-based spintronics. A single-walled CNT (SWCNT) can have a long charge coherence length and spin relaxation distance. CNTs can be synthesized by various methods such as arc discharge (Iijima, 1991) CVD (Cassell et al., 1999; Ren et al., 1998; Li et al., 1996), and laser ablation (Maser et al., 1998). There are two types of CNTs: SWCNTs and multiwalled CNTs (MWCNTs). SWCNTs can be seen as rolled-up cylinders of graphene strips. The electronic properties of SWCNTs can be either metallic or semiconducting, depending on their diameters and chiralities (Louie, 2001). MWCNTs can be seen as an array of coaxial SWCNTs. Different techniques have been developed to produce CNTs in sizable quantities, including arc discharge, laser ablation, CVD, and high-pressure carbon monoxide disproportionation (HiPCO). Among these, arc discharge, laser ablation, and CVD are batch-by-batch processes and HiPCO is a gas phase continuous process (Nikolaev, 2004). Most of these processes take place in a vacuum or with process gases. The CVD growth method is popular, as it yields high quantity and has a degree of control over diameter, length, and morphology. Using particulate catalysts, large quantities of nanotubes can be synthesized by these methods, but achieving the repeatability becomes a major problem with CVD growth (Schulz et al., 2009). The HiPCO process advances in catalysis and continuous growth are making CNTs more commercially viable (Takeuchi et al., 2014). The HiPCO process helps in producing high-purity SWCNTs in higher quantity. The HiPCO reactor operates at a high temperature of 900 C 1100 C and a high pressure of B30 50 bar (Bronikowski et al., 2001). It uses carbon monoxide as the carbon source and nickel/iron pentacarbonyl as the catalyst. This catalyst acts as the nucleation site for the nanotubes to grow (Nikolaev, 2004). Vertically aligned CNT arrays are grown by thermal CVD. A substrate (quartz, silicon, stainless steel, etc.) is coated with a catalytic metal (Fe, Co, and Ni) layer. Typically, that layer is iron and is deposited via sputtering to a thickness of 1 5 nm. A 10 50 nm under layer of alumina is also often put down on the substrate first. This imparts controllable wetting and good interfacial properties. When the substrate is heated to the growth temperature (B700 C), the continuous iron film breaks up into small islands. . . each island then nucleates a CNT. The sputtered thickness controls the island size, and this, in turn, determines the nanotube diameter. Thinner iron layers drive down the diameter of the islands, and they drive down the diameter of the nanotubes grown. The amount of time that the metal island can sit for at the growth temperature is limited, as they are mobile, and can merge into larger (but fewer) islands. Annealing at the growth temperature reduces the site density (number of CNT/ mm2) while increasing the catalyst diameter. However, as-prepared CNTs always have impurities, such as other forms of carbon like a-C; fullerene and noncarbonaceous impurities are metals that are used for catalysts (Itkis et al., 2003; Wang et al., 2014). These impurities need to be removed to make use of the CNTs in application (Eatemadi et al., 2014).

2.3.3 Nanostructure graphene materials Graphene is a strictly two-dimensional (2D) material with honeycomb carbon lattice. It has an unusual linear dispersion relation between energy and wave vector near the Dirac points

Chapter 2 • Introduction: carbon and carbon nanomaterials

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and its charge carriers behave as massless Dirac fermions with a high velocity of 1/300c (Castro Neto et al., 2009). Due to its unique band structure near the Dirac points, graphene has many unusual electrical properties, such as the anomalously quantized Hall effect, Klein paradox, and ballistic transport. The mobility µ of suspended graphene can be as high as 1,000,000 cm2/V/s, which is important for various graphene-based spintronic devices. The densities of charge carriers in graphene can be as high as 1013/cm22 and tunable by an electric field continuously between electrons and holes symmetrically (Novoselov et al., 2005; Zhang et al., 2005; Katsnelson et al., 2006; Du et al., 2008; Elias et al., 2011). Geim and Novoselov (2007) and Novoselov et al. (2004), for the first time, showed the repeatable synthesis of graphene through mechanical exfoliation in 2004, and this technique has been and is being widely used for the fabrication of graphene-based devices. Large varieties of methods have been developed in two possible directions: large-scale growth and large-scale exfoliation (Zhu et al., 2010; Wei et al., 2010; Allen et al., 2009). These methods for graphene synthesis can be divided into “bottom-up” and “top-down” approaches according to the classification of the synthesis method for nanomaterials. “Topdown” methods can be considered as variations of the “exfoliation” method, which produce graphene from graphite through breaking of these weak bonds (Quintana et al., 2012). Exfoliation method is basically a repeated peeling process, in which highly oriented pyrolytic graphite is peeled off into layers using Scotch tape and transferred onto a Si or SiO2 substrate, on which single- or few-layer graphene sheets can be found on examination by a microscope (Lu et al., 1999). The mechanical exfoliation technique for producing graphene sheets was found to be very easy for graphene device fabrication owing to the large size and good quality of the as-produced sheets. However, the limitation of the mechanical exfoliation method is that the output of this method is low, usually in “pieces,” and the size and the shape are not controllable. In “bottom-up” approach, CVD is used for the production of graphene films in the industry, in which a metal/semiconductor crystal substrate or thin film substrate (copper or nickel) is used as substrate and a hydrocarbon as the carbon source (methane, ethylene, etc.). By varying the experimental parameters (hydrocarbon, catalyst, gas flow, pressure, growth time, growth temperature, cooling rate, etc.), the thickness, size, and quality of graphene can be controlled (Wintterlin et al., 2009; Li et al., 2009). In CVD process graphene was first synthesized in 2006, in which camphor was evaporated at 180 C and then pyrolyzed in another chamber of the CVD furnace at 700 C 850 C using argon as the carrier gas. Upon natural cooling to room temperature, few-layer graphene sheets were observed on the Ni foils (Somani et al., 2006). The mechanism of the CVD growth of graphene lies in the diffusion of the carbon atoms into the substrate metals and the segregation of the substrate; the cooling rate during the CVD process is an important factor in the formation and quality of the as-produced graphene (Ruan et al., 2011; Forbeaux et al., 1998). At slow cooling rates, carbon atoms get sufficient time to diffuse into the bulk Ni and no segregation is found on the surface, while at a higher rate, carbon atoms segregate out of Ni but form a less crystalline, defective graphitic structure. At a moderate cooling rate, carbon atoms segregate and form graphene (Yu et al., 2008). In 2010 a large-area, transparent conductive sheet of 30 in. was demonstrated by single-layer graphene with good electrical and

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

optical performances. This large graphene film possessed an optical transmittance of 97% and a sheet resistance of 125 Ω/sq and exhibited the half-integer quantum Hall effect, indicating their high quality (Bae et al., 2010). The traditional CVD method of graphene requires temperatures as high as 1000 C (Li et al., 2009); therefore the CVD synthesis technique is modified in order to achieve high-speed, low-temperature deposition of graphene. Plasmaenhanced CVD (PECVD) extended the synthesis of graphene on any substrate without any special surface preparation operation or catalyst deposition. With the assistance of the plasma the temperature of the substrate could be lowered to 680 C (Wang et al., 2004a,b). It was predicted that a balance between deposition through surface diffusion of C-bearing growth species from the precursor gas and etching caused by atomic hydrogen could synthesize atomically thin graphene sheets. The CVD growth 2D-graphene films could be extended to three-dimensional structures. Three-dimensional interconnected graphene was prepared by using nickel foam as substrate and template, and graphene was deposited on the surface of the nickel foam under the condition of CVD growth of graphene (Chen et al., 2011). Highquality single- and few-layer graphene can be obtained by epitaxial growth through CVD of hydrocarbons on single-crystal metal substrates such as Ru(0 0 0 1) (Sutter et al., 2008; Zhang et al., 2009a,b,c) and Ir(1 1 1) (N’Diaye et al., 2006; Coraux et al., 2008). For the fabrication of wafer-based electronic device applications, graphene is grown on silicon carbide (SiC) substrate by the annealing process of single-crystal SiC under ultrahigh vacuum that leads to growth of ultrathin few-layer graphene on the Si-terminated surface, with the layer (Forbeaux et al., 1998; Levita et al., 2008). As a significant development in this technology, continuous films (millimeter scale) of graphene were synthesized on a Ni thin film coated SiC substrate, in which graphene formed on the upper nickel surface at a lower temperature (700 C 800 C) (Juang et al., 2009). Moreover, the basic material of graphene or graphene ribbon is benzene ring. Polycyclic aromatic hydrocarbons (PAHs) could be considered as nanosized graphenes with size between 1 and 10 nm because PAHs can be regarded as 2Dgraphene segments, which are composed of all sp2 carbons (Chio et al., 2011). Therefore graphene-like materials can be prepared from these graphitic precursors through chemical reactions (Quintana et al., 2012). Graphene is the hydrophobic conjugated carbon network that easily modified by various kinds of functional molecules and nanostructures through hydrophobic, π π stacking, and chemisorption interactions. The covalent modification of graphene sheets creates defects and disrupts the π-conjugated system of the graphene sheets. Supramolecular surface modification of graphene could be performed through π π interaction (Nduwimana et al., 2009; Zhang et al., 2010a,b,c), hydrophobic interaction (Zu et al., 2009), hydrogen bonding (Patil et al., 2009), and electrostatic interaction (Liang et al., 2009). Covalent functionalization of pristine graphene typically requires reactive species that can form covalent adducts with the sp2 carbon structures in graphene through cycloaddition (Park et al., 2013; Zhong et al., 2010; Sarkar et al., 2011, 2012; Xu et al., 2012; Georgakilas et al., 2010; Quintana et al., 2010, 2011), free-radical addition (Zhong et al., 2010), and substitution (Zu et al., 2009). Covalent modification could be done by using oxygen-containing groups (Liu et al., 2008, 2009a,b; Wang et al., 2009a,b; Xu et al., 2009; Zhang et al., 2009a,b,c; Ai et al., 2012; Stankovich et al., 2006a,b; Xu et al., 2008; Matsuo et al., 2004, 2005, 2007;

Chapter 2 • Introduction: carbon and carbon nanomaterials

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Burress et al., 2010; Gadipelli et al., 2011). However, each method has its own advantages and disadvantages. Among all these methods, CVD method is efficient for the production of graphene materials toward different applications. Graphene, produced in this method, was found to show better crystallinity than any other method. PECVD method has shown the versatility of synthesizing graphene on any substrate, thus expanding its field of applications including the spintronic devices. The bond length of C C in graphene is about 1.42 Å and has strong affinity in a particular layer, but a weak one between layers. The specific surface area of a single sheet of graphene is about 2630 m2/g (Stoller et al., 2008). Graphene has unique and outstanding optical properties ( . 97.7% transmittance) with a bandgap value of about 0 0.25 eV (Zhang et al., 2009a,b,c). Some other fascinating characteristics include high carrier mobility (B200,000 cm2/V s) (Geim and Novoselov, 2007; Geim, 2009) and high Young’s modulus (1.0 TPa). Graphene and its composite materials can be used as a semiconductor because of its extraordinary conducting properties. Graphene has been envisioned as the building block of all other important graphitic allotrope forms; fullerene-wrapped version of graphene, CNT-rolled version of graphene, and graphite-stacked version of graphene. Enoki and Kobayashi (2005) investigated the unique magnetic properties of nanographene, such as spin glass states, magnetic switching, and edge-state spin gas probing for the possible applications in electronic and magnetic devices. Chen et al. (2012) reported brief experimental studies about the effects of isotope effects on the thermal properties of graphene and found that the ratio of 12C and 13C play an important role on the thermal conductivity of graphene (Chen et al., 2012). All these exceptional properties make graphene a promising candidate for future electronic devices applications. Some of the potential properties are as follows: • • • • • • •

High Young’s modulus B1000 GPa. Effective moisture barrier. Electrical conductivity similar to copper. Density is four times lower than copper. Thermal conductivity is five times that of copper. Essentially an opened up CNT. High surface area of B2500 m2/g. Has a lower density than steel but can be up to 50 times stronger.

However, for a quick reference regarding the synthesis of graphene at different processes, grown/synthesized on different substrates, studies of their different properties using different measurements including their advantages/disadvantages are given in our earlier book (Ray, 2015).

2.3.4 Graphene oxide materials Graphite oxide has a similar layered structure to graphite, but oxygen-containing groups, which not only expand the interlayer distance but also make the atomic-thick layers hydrophilic, heavily decorate the plane of carbon atoms in graphite oxide. These oxidized layers

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could exfoliate in water under ultrasonication. If the exfoliated sheets contain only one or few layers of carbon atoms, such as graphene, these sheets are named graphene oxide (GO) (Novoselov et al., 2004). So, GO is a single-atomic-layered material made up of carbon, hydrogen, and oxygen molecules by the oxidation of graphite crystals, which are inexpensive and abundant. It is dispersible in water and easy to process. Most importantly, the GO can be (partly) reduced to graphene-like sheets by removing the oxygen-containing groups with the recovery of a conjugated structure. The reduced GO sheets are usually considered as one kind of chemically derived graphene and known as rGO. Some other names have also been given to rGO, such as functionalized graphene, chemically modified graphene, chemically converted graphene, or reduced graphene (Eda et al., 2010). GO has two important characteristics: (1) it can be produced using inexpensive graphite as raw material by cost-effective chemical methods with a high yield and (2) it is highly hydrophilic and can form stable aqueous colloids to facilitate the assembly of macroscopic structures by simple and cheap solution processes. Graphene sheet consists of only trigonally bonded sp2 carbon atoms and is perfectly flat (Lui et al., 2009), apart from microscopic ripples. The heavily decorated GO sheets consist partly of tetrahedrally bonded sp3 carbon atoms, which are displaced slightly above or below the graphene plane (Schniepp et al., 2006). Due to the structure deformation and the presence of covalently bonded functional groups, GO sheets are atomically rough (Paredes et al., 2009; Mkhoyan et al., 2009). Several researchers (Paredes et al., 2009; Kudin et al., 2007; Gomez-Navarro et al., 2010, 2007) have studied the surface of GO and observed highly defective regions, probably due to the presence of oxygen, and other areas are nearly intact. A report shows that the graphene-like honeycomb lattice in GO is preserved, albeit with disorder, that is, the carbon atoms attached to functional groups are slightly displaced but the overall size of the unit cell in GO remains similar to that of graphene (Pandey et al., 2008). Hence, GO can be described as a random distribution of oxidized areas with oxygen-containing functional groups, combined with nonoxidized regions where most of the carbon atoms preserve sp2 hybridization. So GO and rGO are hot topics in the research and development of graphene, especially about mass applications such as graphene. Reducing GO to produce reduced graphene oxide (rGO) is an extremely vital process as it has a large impact on the quality of the rGO produced, and therefore will determine how close rGO will come, in terms of structure, to pristine graphene. In large-scale operations where scientific engineers need to utilize large quantities of graphene for industrial applications, such as energy storage, rGO is the most obvious solution, due to the relative ease in creating sufficient quantities of graphene to desired quality levels. There are a number of ways reduction can be achieved, though they are all methods based on chemical, thermal, or electrochemical means. Some of these techniques are able to produce very high-quality rGO, similar to pristine graphene, but can be complex or time consuming to carry out. In the past, scientists have created rGO from GO by: • treating GO with hydrazine hydrate and maintaining the solution at 100 for 24 h; • exposing GO to hydrogen plasma for a few seconds;

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• exposing GO to another form of strong pulse light, such as those produced by xenon flashtubes; • heating GO in distilled water at varying degrees for different lengths of time; • combining GO with an expansion reduction agent, such as urea, and then heating the solution to cause the urea to release reducing gases, followed by cooling; • directly heating GO to very high levels in a furnace; and • linear sweep voltammetry. Reducing GO by using chemical reduction is a very scalable method, but unfortunately the rGO produced has often resulted in relatively poor yields in terms of surface area and electronic conductibility. Thermally reducing GO at temperatures of 1000 C or more creates rGO that has been shown to have a very high surface area, close to that of pristine graphene even. The heating process damages the structure of the graphene platelets as pressure builds up between them and carbon dioxide is released. This also causes a substantial reduction in the mass of the GO, creating imperfections and vacancies, and potentially also having an effect on the mechanical strength of the rGO produced. Electrochemical reduction of GO is a method that has been shown to produce very high-quality rGO, almost identical in terms of structure to pristine graphene. This process involves coating various substrates, such as indium tin oxide or glass, with a very thin layer of GO. Then, electrodes are placed at each end of the substrate, creating a circuit through the GO. In recent experiments the resulting electrochemically rGO showed a very high carbon to oxygen ratio and also electronic conductivity readings higher than that of silver (8500 S/m, compared to roughly 6300 S/m for silver). Other primary benefits of this technique are that there are no hazardous chemicals used, meaning no toxic waste to dispose of. Unfortunately, the scalability of these techniques has come into question due to the difficulty in depositing GO onto the electrodes in bulk form. In order for GO to be usable as an intermediary in the creation of monolayer or few-layer graphene sheets, it is important to develop an oxidization and reduction process that is able to separate individual carbon layers and then isolate them without modifying their structure. So far, while the chemical reduction of GO is currently seen as the most suitable method of mass production of graphene, it has been difficult for scientists to complete the task of producing graphene sheets of the same quality as mechanical exfoliation, for example, but on a much larger scale. Once this issue is overcome, we can expect to see graphene become much more widely used in commercial and industrial applications.

2.3.5 Carbon nanoparticles and fullerene nanostructure materials Carbon nanoparticle is considered as a new-generation green nanomaterial (Baker et al., 2010; Wang et al., 2010a,b; Li et al., 2010a,b; Cao et al., 2007, 2013; Krysmann et al., 2012; Yu et al., 2012; Guo et al., 2012; Zhu et al., 2012) and a promising alternative of fluorescent semiconductor nanocrystals (FCNs). FCN has been demonstrated as a potential optical detection probe (Zhu et al., 2012), bioimaging probe (Li et al., 2010a,b), light-emitting diode

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

material (Guo et al., 2012), efficient visible light active photocatalyst (Yu et al., 2012) and is used in other optoelectronic device applications (Gruber et al., 1997; Neugart et al., 2007; Batalov et al., 2009; Glinka et al., 1999; Zyubin et al., 2009; Yu et al., 2005; Sun et al., 2006; Zhou et al., 2007; Fu et al., 2007; Wee et al., 2007; Lim et al., 2009; Gao et al., 2007; Liu et al., 2007, 2009a,b; Zhao et al., 2008; Selvi et al., 2008; Bourlinos et al., 2008a,b; Mochalin and Gogotsi, 2009). These carbon nanoparticles are biocompatible and chemically inert (Sun et al., 2006; Zhou et al., 2007; Fu et al., 2007; Wee et al., 2007; Lim et al., 2009; Ushizawa et al., 2002; Cahalan et al., 2002; Huang et al., 2004; Kong et al., 2005a,b), which gives advantages over conventional cadmium-based quantum dots (Medintz et al., 2005; Kamat, 2008) and/or nanostructure materials. However, FCN is relatively less explored compared to other carbon-based materials such as fullerene (Diederich and Thilgen, 1996), CNT (Tasis et al., 2006), and graphene (Dreyer et al., 2010). In addition, the understanding of the origin of fluorescence in carbon nanoparticles is far from sufficient (Sun et al., 2006; Zhou et al., 2007; Gao et al., 2007; Zhao et al., 2008). Information on the microstructure and surface ligands remains unclear and details of the organic passivation is not sufficient to aid understanding of the surface states beneficial for light emission. FCN has been synthesized via physical methods such as high energy radiation based creation of point defect in diamond (Gruber et al., 1997; Neugart et al., 2007; Batalov et al., 2009; Yu et al., 2005) and laser ablation of graphite (Sun et al., 2006; Gao et al., 2007), or chemical methods such as oxidation of candle soot (Liu et al., 2007; Ray et al., 2009), carbonization of carbohydrate (Selvi et al., 2008; Peng and Travas-Sejdic, 2009; Zhang et al., 2010a,b,c; Yang et al., 2011a,b, 2012), thermal decomposition of small molecules (Selvi et al., 2008; Bourlinos et al., 2008a,b; Liu et al., 2009a,b), pyrolysis of polymers (Liu et al., 2009a,b), microwave-based pyrolysis (Li et al., 2011; Wang et al., 2011; Chandra et al., 2012), P2O5-based room-temperature dehydration of small molecule (Fang et al., 2012), electrochemical method (Li et al., 2010a,b; Bao et al., 2011), and chemical breakdown of carbon fiber (Peng et al., 2012)/graphene (Zhuo et al., 2012)/graphite (Zhao et al., 2008; Hens et al., 2012). A wide range of fluorescent carbon particles of different colors can be prepared by those approaches. However, there are four distinct limitations of currently available FCNs. First, most of the synthetic methods produce weakly fluorescent FCN with less than 1% quantum yield. Second, no methods are currently available for largescale synthesis of high-quality FCN. Although some methods report milligram-scale FCN with 5% 60% quantum yield, it requires sophisticated high energy radiation based synthesis followed by surface functionalization and size separation (Wang et al., 2010a,b). Third, although many methods report blue/green-emitting FCNs, only few methods report yellowand red-emitting FCNs. In addition, yellow/red-emitting FCNs are generally mixed with blue/green-emitting FCNs, and need to be isolated via specialized size separation methods (Wang et al., 2010a,b; Liu et al., 2007; Li et al., 2010a,b; Hens et al., 2012). Fourth, although functionalization of nanoprobe is essential in enhancing the labeling specificity, such strategies are little developed for FCN (Li et al. 2010a,b; Yu et al., 2005; Selvi et al., 2008; Ray et al., 2009; Bourlinos et al., 2008a,b; Fang et al., 2012; Peng et al., 2012). In addition, the synthetic methods are cumbersome and inefficient. Recent reports showed that surface passivation can lead to a significant increase in fluorescence quantum yield (4% 15%), however, the

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exact mechanism is not yet clear (Sun et al., 2006; Liu et al., 2009a,b). So, more efficient and large-scale synthesis of FCN and their isolation, purification, functionalization, and hence, bioimaging and applications are very challenging. Here we will describe some of the synthetic methods and functionalization approaches that we have developed in the context of current development of the field, and highlight some important approaches of others. Fullerene is an allotrope of carbon in the form zero-dimensional graphitic carbon (Yang et al., 2018). They are made of carbon atoms in hollow spherical, ellipsoid, and other shapes and sizes. Spherical fullerene is known as buckminsterfullerene or buckyball, resembling the balls used in association football. Structurally, fullerenes are identical to graphite and are composed of irregular stacked graphene sheets, and form hexagonal or pentagonal rings. Various forms of fullerenes have been found, and their sizes range from 30 to 3000 carbon atoms. Unless they are cylindrical, they must also contain pentagonal (or sometimes heptagonal) rings. Numerous fullerenes, including C70, C76, C80, C84, have been found. The most stable fullerene is C60. C60 fullerene molecule is composed of 12 pentagons and 20 hexagons of sp2-bonded carbon atoms. Cylindrical fullerenes are also called CNTs (bucky tubes). Fullerenes are similar in structure to graphite, which is composed of stacked graphene sheets of linked hexagonal rings. A fullerene molecule has unique chemical, physical, and physicochemical properties, which include the following: • The molecule can act as a semiconductor, conductor, and superconductor under specific conditions. • Fullerenes can display the photochromic effect, which is a change in light transmission based on intensity. • Ability to form compounds with many different sorts of materials including the ability to retain other substances inside the molecule and the ability to absorb free radicals. • Fullerenes are relatively safe and inert, and yet have properties that allow the substance to create active derivatives. This set of special characteristics differs depending on the type of materials with fullerenes or fulleroid fragments and offers a very broad scope for their application.

2.3.6 Other carbon metal oxide/sulfide composites materials Defect-induced room-temperature ferromagnetism (RTFM) in metal oxides has attracted significant attention because of their potential application of such oxides in spintronic devices (Dietl, 2010; Volnianska et al., 2010). A few reports have found that at the room temperature or below, RTFM is enhanced significantly when carbon and/or carbon nanostructure materials are functionalized/doped with those metallic oxide materials. Yang et al. (2018) observed that of carbon-doped ZnO nanowires (ZnO-C:NWs) via a mild C1 ion implantation method, the net spin polarization in the surface and bulk regions of ZnO-C:NWs is enhanced significantly. The C-doped ZnO exhibits intrinsic n-type ferromagnetic behavior with a Curie temperature that exceeds room temperature (Pan et al., 2007) and provides a novel way to dope

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

ZnO to form diluted magnetic semiconductors (Sato et al., 2010). The density functional theory calculation also reveals strong magnetic moments in C-doped ZnO (Dai et al., 2011) in which O atoms are substituted with, or replaced by, C atoms stabilizing the ferromagnetism (Peng and Ahuja, 2009). Calculations from first principles indicate that, in contrast, C atoms that substituted at Zn sites (CZn) in C-doped ZnO, which is thermodynamically more stable than that with C atoms that substituted at O sites (CO), form diatomic carbon complexes with a magnetic moment of 2 μB (Nayak et al., 2012). Hence, implanted C atoms in ZnO nanowires provide an ideal opportunity to examine the effect of C implantation on the bonding states and electronic structures of such NWs, based on an in-depth understanding of how the implanted C atoms modify magnetic behaviors in a ZnO host.

2.4 Conclusion and perspectives of carbon and carbon nanostructure-based spintronics applications In this chapter, we introduced the synthesis methods for graphene from two aspects according to the classic classification of synthesis method of nanomaterials, that is, “top-down” and “bottom-up” approaches. As for the “top-down” approaches the reduction of GO is the most popular method for the synthesis of graphene in the experimental study and is widely used in the graphene-based material fabrication. In the “bottom-up” approaches, CVD techniques for the synthesis of graphene were developed fastest and were considered as the most promising method for the large-scale production or large-sized graphene of high quality. Depending on the synthesis techniques, modification of graphene can increase the dispersibility, mechanical property, electronic property, and even biocompatibility, which would offer better performance for further applications. As a Dirac fermion system, graphene possesses unique electronic properties, such as a high integer quantum Hall effect, the Klein paradox, and an ambipolar electric field effect, along with ballistic conduction of charge carriers (Chen et al., 2010). Therefore graphenebased materials and devices have been widely used in the applications ranging from energy storage and conversion to electrochemical sensing, optoelectronics, and field-emission transistor devices. Considerable advances in this area have already been made. Nonetheless, there still remain intriguing issues, which need to be explored further. The investigation on the synthesis method is still one of the bottlenecks for the application of graphene-based devices and materials. Scalable synthesis will require a better understanding and optimization of the growth process, and new protocols or techniques should be developed for the controlled production of graphene in quality and quantity. The methods currently used for graphene synthesis have their own advantages and disadvantages, and which process is to be used should be guided by the end application of the material. The understanding of graphene at the molecular level would be beneficial for the tailing of graphene, such as opening the bandgap, tuning the conductivity, and improving the solubility and stability. Finally, systematic study of the assembly behavior, interaction, and reaction mechanism of graphene/GO is

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fundamental for the controllable design and fabrication of high-performance graphene-based devices and materials.

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Zhou, J., et al., 2011. Intrinsic ferromagnetism in two dimensional carbon structures: triangular graphene nanoflakes linked by carbon chains. Phys. Rev. B 84, 081402. Zhu, Y., et al., 2010. Graphene and graphene oxide: synthesis, properties, and applications. Adv. Mater. 22 (35), 3906 3924. Zhu, A., et al., 2012. Carbon-dot-based dual-emission nanohybrid produces a ratiometric fluorescent sensor for in vivo imaging of cellular copper ions. Angew. Chem. Int. Ed. 51, 7185 7189. Zhuo, S., et al., 2012. Upconversion and downconversion fluorescent graphene quantum dots: ultrasonic preparation and photocatalysis. ACS Nano 6, 1059 1064. Zu, S.-Z., et al., 2009. Aqueous dispersion of graphene sheets stabilized by pluronic copolymers: formation of supramolecular hydrogel. J. Phys. Chem. C 113 (31), 13651 13657. Zyubin, A.S., et al., 2009. Quantum chemical modeling of photoabsorption properties of two- and threenitrogen vacancy point defects in diamond. J. Phys. Chem. C. 113, 10432 10440.

Further reading Markwitz, A., 2014. Room temperature diamond-like carbon coatings produced by low energy ion implantation. Nucl. Instrum. Methods Phys. Res. B 331, 144 148. Pan, D., et al., 2010. Observation of pH-, solvent-, spin-, and excitation-dependent blue photoluminescence from carbon nanoparticles. Chem. Commun. 46, 3681 3683. Sun, X., Li, Y., 2004. Colloidal carbon spheres and their core/shell structures with noble metal nanoparticles. Angew. Chem. Int. Ed. 43, 597 601.

3

Magnetism and spintronics in amorphous/diamond-like carbon 3.1 Introduction There is considerable interest in semiconductor films containing magnetic ions due to the possibility of obtaining dilute ferromagnetic semiconductor (DMS) behavior. Dilute magnetic semiconductors can be used as a source of electronic spin-polarized carriers for spintronic applications (Ohno, 2010). Despite several decades of research, there are very few choices of materials suitable for DMS that could exhibit room-temperature ferromagnetism (Dietl, 2010; Silvestre et al., 2012; Matsumoto et al., 2001; Shinde et al., 2004). Diamond-like carbon (DLC) doped with ferromagnetic elements is a potential candidate for producing a DMS. Recently, there have been several studies aimed at realizing DMS through the addition of magnetic atoms into the DLC matrix (Paul et al., 2012; Colón et al., 2009; Siraj et al., 2011). Paul et al. reported ferromagnetism in nickel incorporated DLC. The origin of ferromagnetic order was however attributed to formation of Ni nanoparticles during the deposition process (Paul et al., 2012). Santana et al. carried out a similar investigation on the effects of chromium incorporation in DLC matrix. It was reported (Colón et al., 2009) that Cr-doped DLC exhibited low-temperature ferromagnetic order when Cr was present at low concentration and resulted in increased formation of antiferromagnetic carbide precipitates near the surface at higher concentrations. It has to be noted that in both the studies, DLC was synthesized and doped with transition-metal atoms using plasma-enhanced chemical vapor deposition (CVD). It would be interesting to find out if changing the doping mechanism of transition-metal atoms into DLC could produce a difference in the dopant
matrix interaction. Based on these mechanism, low-energy ion implantation is an alternative method that could be used to dope DLC films with magnetic atoms as it has proved to be an effective technique to control and modify the properties of thin films (Markwitz et al., 2015; Leveneur et al., 2011; Kennedy et al., 2013). Unlike other techniques, ion implantation is a localized process that offers precise control over material modification without any chemical contamination during the implantation process (Markwitz et al., 2015; Leveneur et al., 2011; Kennedy et al., 2013). By controlling the energy of the implantation, it is also possible to modify a selective region of the film, for example, only the surface region of a material unlike other doping techniques (Leveneur et al., 2011; Kennedy et al., 2013). Hohne et al. (2007) implanted diamond with iron and found that the magnetic properties were largely dominated by diamagnetic diamond. However, the implantation dose was limited to 10214
10215 at/cm2, which implies that the iron atoms were not placed close enough for Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials. DOI: https://doi.org/10.1016/B978-0-12-817680-1.00003-2 © 2020 Elsevier Inc. All rights reserved.

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

them to interact with each other to produce magnetic order. Increasing the magnetic ion concentration could lead to magnetic order. Cobalt may be a more preferred magnetic ion dopant as it is known to be incorporated in the diamond lattice (Lombardi, 2008) and semiconductor materials (Silvestre et al., 2012; Matsumoto et al., 2001; Shinde et al., 2004; Hyndman et al., 2009) and is believed to lead to a DMS (Silvestre et al., 2012; Matsumoto et al., 2001; Shinde et al., 2004; Lombardi, 2008, Hyndman et al., 2009). Hence, it is of great interest to study the effect of Co implantation on the magnetic and structural properties of DLC films, determine if this leads to magnetic order, and ascertain how the structural properties are affected by Co implantation. DLC films can be prepared through a number of methods (Robertson, 2002; Murmu, 2014; Markwitz et al., 2014a,b).

3.2 Magnetism of amorphous carbon and diamond-like carbon DLC is amorphous in nature. Magnetic ion implantation into DLC increase in disorder or an increase in the number of carbon rings in the DLC. These could occur in several pathways. One possibility is that the implanted ions, in this case Co, can lead to formation of ferromagnetic cobalt clusters that can cause magnetic ordering similar to the work of Paul et al. (2012). If the implantation fluence is high enough, interconnecting metallic Co networks with similar properties to that of bulk Co could be formed. Another possibility is that the implanted ions can form carbide clusters similar to that reported by Santana et al. from measurements on Cr-doped DLC (Colón et al., 2009). However, unlike chromium carbide, cobalt carbide is known to be ferromagnetic in nature (Harris et al., 2010; El-Gendy et al., 2014; Zhang et al., 2011). Yet another possibility is that the implanted Co atoms can be diluted into the DLC matrix. This can potentially lead to dilute ferromagnetic semiconductors (Paul et al., 2012; Colón et al., 2009; Siraj et al., 2011). The moment per Co atom (mCo) is plotted in Fig. 3
1A against the applied magnetic field for peak Co concentrations up to 25% after subtracting the diamagnetism from the Si substrate. There was no magnetic order observed for as-deposited DLC. It is apparent in Fig. 3
1 that mCo increases and saturates at B2 T, which is indicative of ferromagnetic order. The saturation mCo, mCo,Sat, is largest for 25% Co where it is 0.3 μB. This is significantly less than that found in Co metal (1.67 μB). It is also less than that found in Co nanoparticles where mCo,Sat is enhanced as the nanoparticle size is reduced and reaches 2.2 μB at 5K for a nanoparticle diameter of 0.8 nm (Osuna et al., 1996; Barea et al., 2005). However, the observed moment is comparable to that observed in cobalt-doped TiO2 (0.27
1.2 μB), which is believed to be a dilute ferromagnetic semiconductor (Silvestre et al., 2012; Matsumoto et al., 2001; Shinde et al., 2004; Hyndman et al., 2009). mCo,Sat plotted with respect to Co peak concentration can be seen in the inset to Fig. 3
1B. The saturated moment per cobalt atom was obtained by averaging mCo from 2 to 6 T. It is the same for Co concentrations up to 13% (0.22 μB), and it increases by a factor of 1.31 for 25% Co. This is not expected if the observed moment was attributed to small ferromagnetic nanoparticles. The dependence of the mCo,Sat on

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49

FIGURE 3–1 M
H loops @ cobalt atom (A) 5K at different concentration; (B) Co 25% implanted DLC at different temperature (5K, 150K, and 300K). M
T loops @ cobalt atom (C) 2.5 T for different Co concentrations (25% and 6.3%); (D) FC and ZFC of 25% Co concentration at 35 mT. DLC, Diamond-like carbon; FC, field cooling; ZFC, zerofield cooling. Reproduced with permission from Gupta et al., 2016. J. Phys. D: Appl. Phys. 49/5, 055002. Copyright IOP Publishing.

temperature, T, can be seen more clearly in Fig. 3
1C where mCo,Sat at 2.5 T is plotted against temperature for 25% Co and 6.3% Co. For both concentrations, there is a large decrease in mCo,Sat with temperature, but the temperature dependence is different where the decrease is more gradual for 25% Co. In both cases, mCo,Sat at is significantly smaller and has similar values at 300K. The measured temperature dependence of mCo,Sat plotted in Fig. 3
1C is not expected for bulk Co metal where mCo,Sat is known to follow the Bloch form where mCo,Sat (T) 5 mCo, n Sat(0)[1 2 D 3 T ], where D is the temperature prefactor, and n is 3/2 (Hendriksen et al., 1993; Crespo et al., 2004; Aquino et al., 2005). Different temperature dependences can occur in small nanoparticles where the number of available magnon modes is limited by the particle size. There is also a cutoff for the long-wavelength magnon modes, which leads to a gap in the magnon dispersion (Crespo et al., 2004). This can result in an increase in D and n (Hendriksen et al., 1993) and lead to n 5 2 (Aquino et al., 2005). However, this is not seen

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

in Fig. 3
3, and it has not been reported even for Co nanoparticles as small as 0.8 nm (Osuna et al., 1996; Barea et al., 2005). Fig. 3
1D shows a plot of mCo,Sat against temperature during zero-field cooling (ZFC) and field cooling (FC) for the Co sample with a peak Co concentration of 25% at an applied magnetic field of 35 mT. There is no evidence for a separation between the ZFC and FC curves for temperatures as low as 5K that has been reported for small Co (Osuna et al., 1996; Barea et al., 2005), where the separation followed by a peak in the ZFC curve at lower temperatures was attributed to superparamagnetic behavior. Superparamagnetism occurs when the thermal energy is greater than the magnetocrystalline anisotropy energy (Bean et al., 1959). The resultant blocking temperature, TB, is defined as the temperature at which there is a peak in the ZFC curve, and it can be written as (Bean et al., 1959): TB 5 KeffV/[75kB], where Keff is the effective magneto-crystalline energy density, V is the volume, and kB is Boltzmann’s constant. The data in Fig. 3
1D show that if any Co nanoparticles are present, they must have TB {5K. Using Keff 5 2.5 3 105 J/m3 for bulk hexagonal close-packed Co at low temperatures (Osuna et al., 1996), TB {5K implies that the radius, r, is r{1:2 nm. However, it is known that Keff calculated using the previous equation, for small Co nanoparticles, is also larger than the bulk value, and it can reach Keff 5 16 3 105 J/m3 for 1.5 nm diameter Co nanoparticles (Osuna et al., 1996). Thus there is no evidence that the ferromagnetic behavior is due to small Co nanoparticles based on the ZFC
FC data as well as the small mCo,Sat discussed earlier. It is also unlikely that it can be due to Co3C nanoparticles, which have a large Keff. For example, Keff 5 7.5 3 105 J/m3 for 8 nm Co3C nanoparticles (El-Gendy et al., 2014). Thus any Co3C nanoparticles if present must have r{0:83 nm (El-Gendy et al., 2014). Furthermore, mCo,Sat from Co3C [1.8 μB] (Carroll et al., 2012) is far greater than the measured mCo,Sat. Geng et al. (2004) studied the intrinsic magnetic properties of carbon-encapsulated magnetic Fe and Fe3C nanoparticles and found that the saturation magnetization (Ms), remanent magnetization (Mr), and coercive field (Hc) are measured as 58.8 emu/g, 5.0 emu/g, and 240 Oe, respectively. The stability of the encapsulated magnetic Fe and Fe3C nanoparticles was tested by heating the sample in air to 400 C, followed by a slow cooling procedure to room temperature over 12 h. They found that sample has no change in their weight or color, suggesting good thermal stability together with strong resistance to oxidation. This excellent property obviously arises from the completely closed carbon structures. The high thermal stability is also confirmed by the observation that no degradation occurs at room temperature for over 1 year, and no structural change happens in the heated sample using TEM investigations. Szczytko et al. (2007) also studied the carbon-encapsulated magnetic nanoparticles based on Fe, Mn, and Cr for spintronic applications. Szczytko et al. (2007) observed that the value of saturation magnetization of Fe-based nanoparticles is almost 60 times larger than the one of Mg or Cr and suggested that the Fe-based nanoparticles are more preferable for future spintronic devices. Colón et al. (2009) studied the magnetism of Cr-doped DLC. It was observed that below 12K, the system exhibits ordinary hysteresis loops, with a coercivity of order 0.8 mT (8 Oe), but at somewhat elevated temperatures (above 20K), the hysteresis loops are constricted.

Chapter 3 • Magnetism and spintronics in amorphous/diamond-like carbon

51

FIGURE 3–2 Hysteresis loops and virgin magnetization curves of Cr-DLC with 3% Cr at (A) 20K and (B) 10K. DLC, Diamond-like carbon. Reproduced with permission from Santana, J.A.C., et al., 2009. Magnetism of Cr-doped diamond like carbon. J. Appl. Phys. 105, 07A930. Copyright American Institute of Physics Publishing. (Santana et al., 2009).

Colón et al. (2009) showed the magnetization of Cr-DLC films with about 3% chromium as shown in Fig. 3
2A and B. Constricted loops frequently occur in inhomogeneous ferromagnets (Skomski, 2008) and probably reflect a cluster-size distribution ranging from very few interatomic distances to about 10 nm. Exchange interactions leading to Curie temperatures above 20K are common in magnetic oxides and not surprising in the present system, where the C 2p electrons strongly hybridize with the Cr 3d electrons. In fact, the strong overlap between 2p electron orbitals in elements such as B, C, and O means that 2p moments created by transition-metal ions and other impurities couple relatively rigidly to neighboring 2p atoms. Relatively extended orbitals of this type occur in some oxides (Andersen et al., 2000) and Co-doped semiconducting boron carbides (Carlson et al., 2007). Magnetism of Q-carbon: Q-carbon is a new phase of carbon is found to be over 40% harder than diamond. Nanosecond laser melting of amorphous carbon (a-C) and rapid quenching from the superundercooled state forms this phase. Closely packed atoms in molten metallic carbon are quenched into Q-carbon with 80%
85% sp3 and the rest with sp2. The number density of atoms in Q-carbon can vary from 40% to 60% higher than diamond cubic lattice, as the tetrahedral packing efficiency increases from 70% to 80%. Using this semiempirical approach, the corresponding increase in Q-carbon hardness is estimated to vary from 48% to 70% compared to diamond. It decreases only slightly less than 10% with temperature between 10 and 300K. From the M versus H and M versus T measurements, Curie temperature was estimated over 570K. Jagdish et al. (2018) showed that the undoped Q-carbon behaves as ferromagnetism with coercivity B150 Oe at 300K, which increases to B200 Oe at 10K. It is interesting to note saturation magnetization in M versus H curve (Fig. 3
3). Undoped Q-carbon exhibits n-type

52

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 3–3 M
H plots of Q-carbon at different temperatures (10K, 50K, 100K, 200K, and 300K) showing ferromagnetic characteristics. The upper inset shows MFM image of Q-carbon at 300K. The lower right inset shows the diamagnetic signal from as-deposited DLC, and the lower left inset shows finite coercivity (B100 Oe) at 100K. DLC, Diamond-like carbon; MFM, magnetic force microscope. By courtesy of Jagdish, N., et al., 2018. Progress in Q-carbon and related materials with extraordinary properties. Mater. Res. Lett. 6 (7), 353
364.

conductivity and extraordinary Hall effect. The origin of intrinsic ferromagnetism in Q-carbon is due to the electronic mixing of sp2 and sp3 orbitals during its formation. Some of the unpaired electrons remain in the Q-carbon structure, which render it to be ferromagnetic in nature. Further work is needed to understand the three-dimensional coupling mechanism of these spins to create ferromagnetism. The temperature-dependent magnetic anisotropy constant calculations in Q-carbon indicate a nonlinear dependence with an exponent value of 2.03, which is attributed to the presence of a mixture of sp2
sp3 hybridized states. The temperature-dependent saturation magnetization curve of Q-carbon follows the modified Bloch’s law with the Bloch exponent equal to 2.04, prefactor of 3.02 3 1026K22.04 and a Curie temperature of 570K (Bhaumik et al., 2018). The negative values of the ordinary Hall coefficient in the temperature range from 10K to 300K indicate that the electron band is more conducting than the hole band in Q-carbon. An exponent value of 1.95 6 0.05 that has been extracted from the Hall effect measurements in Q-carbon suggests that the electron scattering in the Hall resistivity in Q-carbon follows a nonclassical “side-jump mechanism.” This electronic scattering mechanism exists in Q-carbon due to the presence of nonconducting centers, which distort the wave function and create a local current density.

3.3 Electrical and transport of amorphous carbon and diamond-like carbon DLC films are very attractive for widespread microelectronic applications due to their optical, mechanical, and thermal properties (Robertson, 2002). In particular, they can be used to create semiinsulating passivation layers for high-voltage devices and terminations

Chapter 3 • Magnetism and spintronics in amorphous/diamond-like carbon

53

FIGURE 3–4 Schematic views and measurement configuration of the MIM and n-type MIS devices. MIM, Metal
insulator
metal; MIS, metal
insulator
semiconductor.

(Frischholz et al., 1993), or coatings for different applications (Basman et al., 2015). Most of the works concerning with DLC films concentrate on the characterization and comparison of the electrical properties, as DLC features are strongly sensitive on the used deposition method (Staryga et al., 2005). Reggiani et al. (2016) addressed the charge transport in DLC layers to identify the most relevant physical effects. In this study, Reggiani et al. (2016) have studied DLC films, deposited onto an Al-metal layer and n-type silicon substrate. An Al-metal was thermally evaporated on top to realize a metal
insulator
metal (MIM) device and a metal
insulator
semiconductor (MIS) diode. In addition, an H2
silicon interface treatment before the DLC growth was adopted on the MIS structure to improve the DLC/Si interface. Fig. 3
4 shows the schematic diagram of two measurement structures. The thicknesses of each deposited layer ranges from 190 to 240 nm and have optical bandgaps of 0.9
1.1 eV. Reggiani et al. (2016) studied the current
voltage characteristics of DLC films in metal
DLC
metal and metal
DLC
nSi structures. The current
voltage measurements were carried out in dark conditions by sweeping the top electrode bias to positive (V . 0) and negative (V , 0) values. Maximum current/voltage has been limited to safe operating regimes. For the different transport regimes the I
V characteristics were measured at B300K, B325K, B350K, and B375K and are shown in Fig. 3
5. Fig. 3
5A shows the I
V curves of the MIM and n-type MIS devices measured at room temperature (B300K). As expected, the n-type MIS diode shows a forward-bias behavior for V . 0 with current levels similar to the MIM device. In Fig. 3
5B the current at a fixed forward voltage of 1 V is reported as a function of the temperature for both devices showing very similar trends; the corresponding activation energies are close to 0.44
0.5 eV indicating that the energetic barriers for injection from metal and n-type Si should be very similar. As far as the boundary conditions are concerned, the metal Schottky barrier and the DLC affinity have been fixed, so as to reproduce the J
V experimental data of both MIM and MIS devices with the correct activation energies. The activation energy of MIM and MIS devices are 0.49 and 0.44 eV, respectively, which are extracted using the relation: J 5 exp (2EA/kBT), where kB is the Boltzmann constant (Fig. 3
5B). The mixture of sp3 and sp2 bonds gives rise to disordered molecules, which intrinsically lead to different activation energy. Fig. 3
5C and D shows the I
V

54

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 3–5 (A) J
V curves of the MIM and MIS devices as functions of applied voltage. As expected, the MIM curves are symmetric for positive and negative biases. Very similar current level is obtained for the MIS device in forward regime, while the reverse current is limited to lower values; (B) temperature dependence of the current density in forward-bias condition (V 5 1 V). The activation energy is extracted assuming JBexp(2EA/kBT), with kB the Boltzmann constant; (C) current-density characteristics of the MIM at different ambient temperatures; (D) current-density characteristics of the n-type MIS at different ambient temperatures. MIM, Metal
insulator
metal; MIS, metal
insulator
semiconductor. By courtesy of Reggiani, S., et al., 2016. TCAD-based investigation on transport properties of diamond-like carbon coatings for HV-ICs. In: IEDM16, pp. 926
929.

characteristics of MIM and MIS devices measured at different temperatures. It is worth observing that the deeply depleted silicon dominates the transport regime corresponding to the reverse biases in the MIS diode. For spin-electronics applications, Colón et al. (2009) have fabricated p-type Cr-DLC heterojunctions with silicon. Fig. 3
6A and B shows the I
V curves for heterojunctions made with 11 and 15 at.% Cr contents. At the low doping levels, heterojunction diodes can be made, but the capacitance is quite large and dominates the device properties, consistent with a-C films on ntype silicon (Valentinit et al., 2003; Konofaos et al., 2002). The best diode rectification is obtained for 11%
15% Cr doping as shown in Fig. 3
6A and B. With a Cr-doping concentration of 20% or more, the heterojunction diodes with n-type silicon show very large relative leakage currents in reverse bias and increasingly resemble a “bad” conventional resistor. The heterojunction diodes of n-type silicon and 11% and 15% Cr-doped DLC films as the p-type semiconductor show a strong negative magnetoresistance (MR) with the forward-bias

Chapter 3 • Magnetism and spintronics in amorphous/diamond-like carbon

55

FIGURE 3–6 The I
V curves from the 11 at.% Cr (A) and 15 at.% Cr (B) Cr-DLC film to n-type silicon heterojunction devices with changing applied magnetic field at room temperature. The change in forward current, as a function of the magnetic field, for Cr-DLC film to n-type silicon heterojunction devices at (C) 11.0 at.% Cr and (D) 15.0 at.% Cr has been plotted for a forward bias of 2 V. DLC, Diamond-like carbon. Reproduced with permission from Santana, J.A.C., et al., 2010. Appl. Phys. A 98, 811
819. Copyright Springer link.

current increasing with magnetic field, even at room temperature as shown in Fig. 3
6C and D. At 2 V forward bias the negative MR is as much as 50%
100% in an applied magnetic field as small as 300 Oe, as indicated in Fig. 3
6D. The negative MR saturates and shows little change at the higher applied magnetic fields, indicating that some magnetic ordering is the origin of this effect. What is the reason for the room-temperature MR of these heterojunction diodes? There is no evidence for ferromagnetic ordering in the nanoclusters, but antiferromagnetic order is well known to create uncompensated spins at the clusters’ surfaces. These cluster macrospins are likely to interact with each other, especially since the clusters are particularly concentrated at the DLC film surface, with the magnetic field and electric current. The coupling mechanisms for high concentrations would then be very different from the low-concentration mechanism. Narayan et al. (2018) studied the transport properties of p
n junction diamond grown by laser annealing (melting) of N-doped a-C. Fig. 3
7 shows the I
V plot of p
n junction diamond. The I
V plot shows the low turn-on voltage of B3 V, and a minimum of reverse-bias current shows superior diode characteristics. The temperature-dependent I
V plots also show the characteristic diode behavior as shown in Fig. 3
6B. With the increase in temperature, there is an increase in the forward-bias current. This occurs due to increased concentration of charge carriers with an increase in temperature. The ideality factor of the p
n junction can be calculated using the diode as shown in the following equation: logI 5

n VA logI0 2:303kB T

(3.1)

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 3–7 (A) I
V characteristics showing diamond p
n junction with the inset showing a schematic of the diode device (electrodes, B-doped diamond, and N-doped diamond); and (B) I
V characteristics at different temperatures (100K, 300K, and 350K) with the inset showing the ideality factor 5 2.077. By courtesy of Jagdish, N., et al., 2018. Progress in Q-carbon and related materials with extraordinary properties. Mater. Res. Lett. 6 (7), 353
364.

where, n, kB, T, VA, and I0 denote ideality factor, Boltzmann constant, temperature (of the p
n junction), applied voltage, and reverse saturation current, respectively. I0 and n are calculated to be B1.0 3 129 A and 2.077, respectively. The value of ideality factor is close to 2, which indicates that the fabricated diode is close to an ideal one.

3.4 Magnetoresistance and spintronics of amorphous carbon and diamond-like carbon Giant MR (GMR), tunneling MR (TMR), anisotropic MR (AMR), and colossal MR (CMR) have drawn much attention for their applications. They have potential applications in spintronic devices such as magnetic storage and magnetic sensors. GMR and TMR are mainly found in multilayer and granular films. In addition, spin valves and ballistic MR structures are also multilayered (Baibich et al., 1998; Berkowitz et al., 1992; Xiao et al., 1992; Dieny et al., 1991; García et al., 2001). Ferromagnetic layers or granules are indispensable parts in those sophistically designed structures. Large MR are also found in diluted magnetic semiconducting systems, such as (Ga, Mn)As and (Zn, Mn)O (Van Esch et al., 1997; Andrearczyk et al., 2005)

Chapter 3 • Magnetism and spintronics in amorphous/diamond-like carbon

57

AMR, which is lower than GMR, commonly exhibited in permalloy (a kind of ferromagnetic alloy compounded by Ni and Fe) (McGuire et al., 1975; Lee et al., 2000). These properties can be observe in a-C and/or DLC (with/without doped) materials and are discussed in the following sections.

3.4.1 Magnetoresistance of amorphous carbon films Semiconducting a-C has attracted a great attention due to its wide applications, such as field emission, pressure sensing, gas sensing, solar cells, electronics, and optical industry (Zhang et al., 2004; Tian et al., 2007; Caihua et al., 2011; Sagar et al., 2014a,b; Gopinadhan et al., 2015; Saleemi et al., 2016). Among these applications the MR is one of the most compelling properties due to its relation with spintronic applications (Sagar et al., 2017; Xin et al., 2009; Bai et al., 2010; Vianelli et al., 2015). The transport properties of doped a-C films have been studied intensively. Negative and positive MR (P-MR) was observed, and possible mechanisms were proposed (Xin et al., 2009; Wang et al., 2014; Saleemi et al., 2017a,b,c). However, for pure a-C films, the MR properties needed to investigate because it could help to reveal the influence of amorphous nature of such materials on their magneto-transport properties. Moreover, a clear understanding of transport mechanism of the disordered a-C thin films is still lacking. The a-C was reported to have different conduction mechanisms such as weak localization (WL) (Bayot et al., 1990), wave function shrinkage (Sagar et al., 2014a,b), spin blockade, electron
electron interaction (Wang et al., 2014), grain-boundary scattering (Chung et al., 2008), and variable range hopping (VRH) model (Wang et al., 2013). Therefore understanding of the electron transport properties of disordered a-C has great significance of study. MR is a sensitive local probe for studying the scattering process in disordered amorphous systems and highly depends on the temperature of conductivity (Iwase et al., 1999). Saleemi et al. (2017a,b,c) studied the MR of a-C films synthesized by the pulsed laser deposition (PLD) technique at different temperature, such as 450 C, 500 C, and 550 C, respectively. Saleemi et al. (2017a,b,c) also measured the MR at low temperature of 2K
40K under the magnetic field of 7 T, and a large MR of B46% was observed at temperature of 2K in the a-C synthesized at 500 C. The MR value of 19% and 32% was observed at temperature of 2K for the a-C synthesized at 450 C and 550 C, respectively. In this study, Saleemi et al. (2017a,b,c) showed that the MR decreases rapidly with the increase in measurement temperature under the magnetic field of 7 T. The MR vanishes after 40K indicating that the MR of a-C films depends on the growth temperature, because the growth temperature in PLD has a significant influence on the degree of disorder in a-C films, and the disorder plays important role on the MR of such a-C films (Wang et al., 2013). Fig. 3
8 shows that P-MR for all the previously mentioned three a-C films decreases rapidly with the increase in measurement temperature. MR decreased with the decrease in degree of disorder and changes with the change in deposition temperature. Several MR mechanisms may responsible for the large MR phenomenon. The ordinary MR (OMR) could not be the suitable mechanism for such a-C system because in OMR, MR 5 (μB)2, where m is the mobility of a-C films. If we suppose that MR of our a-C films was originated from OMR, then using the largest MR of B46%, the

58

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 3–8 MR
B curves at 2K
40K for the a-C synthesized at (A) 450 C; (B) 500 C; (C) 550 C. (D) For comparison study MR
B measured at 2K for all three a-C films. a-C, Amorphous carbon. Reproduced with permission from Saleemi, A.S., et al., 2017a. Structure dependent negative and positive magnetoresistance of amorphous carbon films. J. Appl. Phys. 121, 233903; Saleemi, A.S., et al., 2017b. Structure dependent negative magnetoresistance of amorphous carbon thin films. Diam. Rel. Mater. 72, 108
113; Saleemi, A.S., et al., 2017c. Large magnetoresistance of amorphous carbon films. Carbon 122, 122
127. Copyright Elsevier Publishing.

mobility, m, should be as large as B970 cm2/(V s). But the mobility of our a-C films measured at 2K was only 1.9 cm2/(V s), which is much smaller than the value we estimated to satisfy the OMR mechanism. The mobility was estimated by using the following equation: μ5

VH t ρIB

(3.2)

where VH is Hall voltage, t is sample thickness, r is resistivity, I is current, and B is the magnetic field. As the observed MR of a-C films was always positive, the WL effect and GMR/ TMR effect could not be the MR mechanisms for such a-C films. Mott-type, Efros
Shklovskii (ES)-type VRH and nearest-neighbor hopping (Inada et al., 2015) could be the possible mechanisms for such carbon system. Saleemi et al. (2017a,b,c) also studied the MR of annealed (B1000 C) a-C films synthesized by the CVD at 100 C. In this study, Saleemi et al. (2017a,b,c) used three different a-C films deposited at three different deposition time: 10, 20, and 30 min, respectively. The negative MR increases from 26% to 213% for the a-C films deposited for 10
30 min respectively as shown in Fig. 3
9A. Saleemi et al. (2017a,b,c) observed that the MR is negative up to B13%, when

Chapter 3 • Magnetism and spintronics in amorphous/diamond-like carbon

59

FIGURE 3–9 (A) MR% versus magnetic field at different temperature and (B) MR% versus temperature at different applied magnetic field for CVD deposited a-C film (deposited time 5 10 min). a-C, Amorphous carbon; CVD, chemical vapor deposition; MR, magnetoresistance. Reproduced with permission from Saleemi, A.S., et al., 2017a. Structure dependent negative and positive magnetoresistance of amorphous carbon films. J. Appl. Phys. 121, 233903; Saleemi, A.S., et al., 2017b. Structure dependent negative magnetoresistance of amorphous carbon thin films. Diam. Rel. Mater. 72, 108
113; Saleemi, A.S., et al., 2017c. Large magnetoresistance of amorphous carbon films. Carbon 122, 122
127. Copyright Elsevier Publishing.

measured at B2K under the magnetic field of 7 T as shown in Fig. 3
9A. Fig. 3
9B shows the MR versus temperature at different applied magnetic field for the a-C deposited for 30 min. The negative MR value decreases with the increase in temperatures, and in disordered systems it could have variety of mechanisms. The WL theory (WLT) is a possible mechanism for the negative MR that is due to the ordered graphitic like structures. The mechanism is explained by the WLT is for a lower temperature range of 2K
50K, whereas the grain-boundary scattering model for a higher temperature range of 50K
300K. The MR increases from 26% to 213% at 2K, respectively, for the a-C films deposited at 10
30 min as shown in Fig. 3
8A. The MR measured at different magnetic field (up to 7 T) within the temperature 2K
300K shown in Fig. 3
8B. The observed negative MR decreases with the increase in temperature

60

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

from 2K to 300K and increases with the increase in magnetic field from 1 to 7 T. MR of 213% is observed at 2K for 30 min synthesized a-C film, which has the most ordered structure and lowest resistivity. However, negative MR originated from WLT is usually a low-temperature (,100K) effect, because the inelastic scattering process due to electron
phonon or electron
electron interaction makes the phase incoherence between carrier waves, and this effect diminishes in high temperatures (Zhou et al., 2011; Hishiyama et al., 2001). In the case of P-MR observed in PLD, a-C techniques are used due to its disordered amorphous structure, and the Efros
Shklovskii-type VRH model explains the mechanism. Saleemi et al. (2016) also measured the angular-dependent magneto-transport properties of CVD deposited a-C thin films within the temperature range of 2K
300K under the magnetic field of 7 T. The angular MR (i.e., MR) was measured
by varying the angle according to the definition (Sagar et al., 2015): MR% 5 ðRθ 2 R0 Þ=R0 3 100, where Rθ and R0 are the resistance of the specimens at angle θ and 0 , respectively. Here θ is an angle between the applied magnetic field (B) and normal to the surface of specimen (n) as shown in the inset of Fig. 3
10D.

FIGURE 3–10 (A) Angular MR (%) versus position θ at B 5 7 T; (B) MR versus magnetic field at θ 5 0 from 2K to 300K; (C) angular MR (%) versus temperature at different angles at B 5 7 T of the a-C deposited in CVD for 5 min time. (D) Angular MR (%) versus temperature for at B 5 7 T at θ 5 90 of the a-C deposited in CVD for 5
25 min time. a-C, Amorphous carbon; MR, magnetoresistance. Reproduced with permission from Saleemi, A.S., et al., 2016. Angle dependent magneto transport in transfer-free amorphous carbon thin films. J. Phys. D: Appl. Phys. 49, 415005. Copyright IOP Publishing.

Chapter 3 • Magnetism and spintronics in amorphous/diamond-like carbon

61

The angular MR of undoped a-C thin films was studied from 0 to 180 at different temperatures from 2K to 300K as shown in Fig. 3
10A. Angular MR of B18% and B1.6% was observed at 2K and 300K, respectively, under the magnetic field of 7 T at 90 . Angular MR increased as the specimen rotated from the angle 0 to 90 , and its maximum value was observed at θ 5 90 , while angular MR decreased from 90 to 180 . Fig. 3
10B shows the field-dependent MR up to the magnetic field of 7 T of sample S1 for the temperature range 2K
300K. A maximum MR of 28.3% is observed at 2K temperature under the magnetic field of 7 T. The MR increases with the increase in magnetic field, and the behavior is the same for all of the temperatures. The value of MR was observed to decrease from 28.3% to 20.6% in conjunction with the increase in temperature from 2K to 300K, respectively. The magnitude of angular MR was found to be inversely proportional to the temperature for all of the angles as shown in Fig. 3
10C, and this behavior was the same for all of the specimens. The angular MR showed an exponential decay as temperature increased. The angular MR has no effect if either the specimen is rotated along the short axis or the long axis during the measurement, indicating that such a-C thin films do not have an anisotropic angular effect but remained unsaturated up to the magnetic field of 7 T. Tian et al. (2007) studied the MR of a-C and a-C:Fe films and found that the MR is higher in the case of a-C:Fe films. Tian et al. (2007) studied the field- and temperature-dependent MR of Fe-doped a-C (a-C:Fe) thin film synthesized on n-Si substrate. They found that the a-C:Fe and a-C films have very different conduction mechanisms, and their MR effect could not be explained by the known MR mechanisms. The MR of a-C:Fe and a-C films was measured at various temperatures by superconducting quantum interference device (SQUID) with both measuring current and magnetic field parallel to the film surface. Fig. 3
11A shows the MR measurement at room temperature. In low temperature (below 200K), of a-C: Fe and a-C films display a common feature of amorphous semiconductor that obeys the Mott VRH law (Xue et al., 2005). The MR of the a-C film at B 5 5 T is less than 2%. But at room temperature the Fe-doped film has a P-MR of 15% at B 5 1 T and 28% at B 5 5 T, respectively. The MR
T relation is unable to fit by linear or B2 exponential law. Fig. 3
11B shows the temperature dependence of magnetization in Fe-doped film. Magnetization decreased nearly linear from 10K to 300K without sudden drop. These results indicate that the small amount of Fe dopant fails to give rise to ferromagnetic phase in a-C film even in low temperature. Tian et al. (2007) noticed that the P-MR of a-C:Fe film occurs in the high temperature (HT) region, while the change of magnetoresistance between a-C film and a-C: Fe film is much bigger in HT region than in low temperature (LT) region. So the fact that the MR of a-C:Fe film is much higher than that of a-C film because Fe-doping changes the conduction mechanism significantly, and magnetic field dramatically increases the activation energy. This leads to increase in the resistance, leading to the P-MR. This P-MR effect is difficult to explain by the known MR models. a-C film and a-C:Fe films remain paramagnetic in the entire temperature region from low temperature to room temperature. So the mechanisms that were used in ferromagnetic films (like interface spin-dependent scattering), ferromagnetic semiconductors (like band spin-splitting), and ferromagnetic alloys (like easy axis in AMR alloys) are invalid.

62

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 3–11 (A) The MR dependence on magnetic field in a-C film and a-C:Fe film at room temperature. (B) The M
T relation (from 10K to 300K) of the a-C:Fe film. a-C, Amorphous carbon; MR, magnetoresistance. Reproduced with permission from Tian, P., et al., 2007. Enhanced room-temperature positive magnetoresistance of a-C: Fe film, Carbon 45, 1764
1768. Copyright Elsevier Publishing.

Besides, in a-C film and a-C:Fe film no metal
semiconductor transition was observed in the temperature region of MR. Moreover the peak temperature did not appears at the critical point between HT and LT region; therefore the MR effect could not be attributed to a phase transition like in some manganite perovskites (Salamon et al., 2001). Based on these results, Tian et al. (2007) proposed a mechanism to explain the conduction and MR in a-C:Fe film. This carbon-based MR material may be helpful for developing new kind of MR device. Jiang et al. (2012) observed a P-MR effect in Co-doped a-C films. The P-MR effect of the a-C:Co/Si junction, MR
H relations reach a peak value at V 5 5 V and T 5 65K, which is consistent with the result of Tian et al. (2007) obtained at room temperature. Under the reverse electric field the a-C:Co/Si junction has a colossal P-MR dependent on the bias voltage. The MR (26 V) reaches 100% at H 5 1 T and over 350% at H 5 5 T. When the voltage value deviates from 26 V, the MR significantly decreases, but it is still much higher than in the positive voltage. This phenomenon indicates that there are two very different PMR mechanisms dependent on the electric field direction. Although the d electrons of Co atoms mainly have exchange interactions with sp2 hybridized electrons in the disordered carbon system, the PMR of the a-C:Co/Si junction is unlikely to be induced only by such exchange interactions, especially under the reverse electric field (Tian et al., 2007). A possible explanation is that the effect of the applied magnetic field on the Co ions leads to the transition from

Chapter 3 • Magnetism and spintronics in amorphous/diamond-like carbon

63

sp2 to sp3 sites, increases the barrier width, and then decreases carriers passing through the potential barrier by tunneling. Akhgar et al. (2018) studied the MR of hydrogen-terminated diamond in the phase coherent backscattering (WL and antilocalization) at low temperatures. They found that the response to an external magnetic field is highly anisotropic, confirming the 2D nature of the carriers despite the short mean free path. By simultaneously applying perpendicular and parallel magnetic fields, they are able to probe the Zeeman interaction and micro-roughness of the quantum well at the diamond surface. Akhgar et al. (2018) have obtained a hole g-factor of  2:6 6 0.1 from MR curves at B2.5K and rms fluctuations in the width of the hole quantum well of about B3 nm over the phase coherence length of 33 nm. However, Akhgar et al. (2018) found that the magnetoconductivity (Δσ) of hydrogen-terminated diamond is reduced in the absence of an in-plane field as shown in Fig. 3
12A. The evolution of Δσ as a function of temperature shows a transition from WL to weak antilocalization (WAL) as the temperature is reduced, where WAL is characteristic of 2D systems in the presence of spin
orbit coupling. Akhgar et al. (2018) believed that a strong spin
orbit interaction occurs due to the strong electric field that accompanies the asymmetric 2D confinement potential at the diamond surface (Akhgar et al., 2016). Akhgar et al. (2018) also studied the magneto-conductance at different values of the constant in-plane field B! keeping the temperature constant as shown in Fig. 3
12B and C. At a temperature of 4K, where phase coherent backscattering dominates to give WL on account of the fact that the phase relaxation time is shorter than the spin
orbit scattering time, the magneto-conductance curves are largely unaffected by the application of the inplane magnetic field. In the case of magneto-conductance at 2.5K the signatures are dominated by WAL for B! 5 0 T and dependence on the in-plane field with the WAL suppressed as B! increases.

3.5 Spin field effect transistor (FET) Thin-film a-C has many attractive features for electronic use: a bandgap, which can be varied from B0.5 up to B4.0 eV, potential to produce n, p, and intrinsic material, and large area deposition at room temperature. a-Cs containing significant amounts of sp3 bonding are known as DLC. This is because of the presence of the sp3 bonds that confer the diamondlike properties of chemical inertness, mechanical hardness, and wider bandgap than one would find in an sp2-bonded material (Robertson, 1986). The most common form of DLC is hydrogenated a-C (a-C:H) and not possible to grow a pure 100% sp3-bonded a-C, which would correspond to “amorphous diamond” and be totally analogous to amorphous silicon. The highest sp3 content so far is B85%
90%. This material, because of its high percentage of tetrahedral bonding, is known as tetrahedrally bonded a-C (ta-C). Their high defect-state density and their associated low mobilities, however, have limited the use of a-C thin films in electronic devices. Their main application in the electronic field has been in their use as cold cathode-field emitters where their low threshold field has attracted much attention

64

Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 3–12 (A) Reduced magneto-conductivity, Δσ 5 σ(B) 2 σ(B0) (in units of G0 5 e2/πh), at different temperatures for B! 5 0 T. The solid lines fit to Rashba theory. Conductivity as a function of B with different applied in-plane magnetic fields (B) 2.5K and (C) 4K. Reproduced with permission from Akhgar, G., et al., 2018. G-factor and well width variations for the two-dimensional hole gas in surface conducting diamond, Appl. Phys. Lett. 112, 042102. Copyright American Institute of Physics Publishing Co.

(Milne et al., 2000). However, attempts have also been made to produce metal
semiconductor
metal structures (Egret et al., 1997), a-C/c-Si hetero-structures (Veerasamy et al., 1994) and thin-film transistors (Clough et al., 1996), with varying degrees of success. The manufacture of thin film transistors (TFTs) using DLC has attempted using both ta-C and ta-C:H as the active channel material. The ta-C has a sp2-rich surface layer present as a consequence of the subplantation deposition process. Such a layer exists both on the top of the film (because of the subplantation) and also at the substrate/film interface due to high incident ion energies. The thickness of both these layers is dependent upon ion energy used during deposition (Davis et al., 1995). The first series of experiments addressed the problem

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65

FIGURE 3–13 (A) Drain transfer characteristics for p-channel top-gate TFT; (B) comparison between a bottom-gate and a top-gate TFT. Reproduced with permission from Milne, W.I., 2003. Electronic devices from diamond-like carbon, Semicond. Sci. Technol. 18/3, S81
S85. Copyright IOP Publishing.

that this layer would lead to a pinning of the Fermi level, so a TFT structure that would allow easy removal of the sp2-rich layer was chosen. A top-gate coplanar structure was therefore adopted. At that time, there was also uncertainty as to whether the sp2-rich layer would instantly reform once removed, so a scheme was devised whereby the sp2-rich layer could be removed, and the gate insulator subsequently deposited all within the same pump down sequence. An reactive-ion etching (RIE) process to remove the sp2 layer using nitrous oxide as the etching gas was adopted and silane was then added immediately (Clough et al., 1996) to deposit the gate insulator. Subsequently, it was discovered that the sp2-rich once removed did not reform. Milne (2003) studied a typical undoped ta-C thin film transistor (TFT) operated in the pchannel accumulation mode confirming the p-type nature of the as-grown ta-C, but the on/ off ratio was only approximately two to three orders, and the hole mobility was estimated to be 1025
1026 cm2/V/s as shown in Fig. 3
13A. This indicates that extended-state transport is not the dominant conduction mechanism, and hopping in the band tails is more likely. In such a process the main factor in controlling the mobility is the degree of localization of tail states McKenzie (1996) who has predicted that the highly sp3-bonded material, a high

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

degree of localization of π-states at the band edges. In an effort to improve TFT performance, two alternative solutions were attempted (1) to investigate bottom-gate structures and (2) n-channel devices as these have been found to have significantly higher mobilities than p-channel devices made using a-Si:H. Bottom-gate TFTs were produced by depositing the ta-C films onto thermal oxide
covered c-Si wafers. The effect was to reduce the off current but gave no improvement in the on current as shown in Fig. 3
13B. The improvement in the off current is due to the fact that we are using thermal oxide as the gate insulator in this instance. n-Channel devices were then produced by adding nitrogen to the films during deposition as previously described (Veerasamy et al., 1993). Both top-gate and bottom-gate structures have been produced, and these operate in the n-channel enhancement mode as expected. However, the best of these only gave a maximum field-effect mobility value of B5 3 1027 cm2/V/s. Addition of hydrogen to ta-C films has previously been shown to reduce the electron spin resonance (ESR) spin density by approximately five times, and, therefore, in a final effort to increase the mobility, TFTs were produced using plasma beam source deposited ta-C:H as the active channel material in a bottom-gate structure. This had the effect of increasing the on/off ratio by a factor of 2. The best mobility obtained to date for any a-C-based TFTs is B1024
1025 cm2/V/s. This was measured on a bottom-gate p-channel TFT using a thermal oxide as the gate insulator and a ta-C film with 50/50 sp3/sp2 bonding. Much more work is needed in order to reduce the density of states (DOS) before a-Cs can be utilized in practical thin-film transistor structures. Maze et al. (2008) experimentally demonstrate nanoscale magnetic sensing, using coherent manipulation of an individual electronic spin qubit associated with a nitrogen-vacancy impurity in diamond at room temperature (Taylor et al., 2008). Maze et al. (2008) achieved the detection of 3 nT magnetic fields at kilohertz frequencies after 100 s of averaging in an ultrapure diamond. Maze et al. (2008) also demonstrate a sensitivity of 0.5 μT/Hz1/2 for a diamond nanocrystal with a diameter of 30 nm. Bhattacharyya (2004) studied the electrical conductivity of nitrogen-doped nanocrystalline diamond. Bhattacharyya (2004) observed that the nitrogenated nanocrystalline diamond increases the density of states at the Fermi level that helps to increase the conductivity. Low-temperature conductivity has been explained from a change over from Arrhenius behavior to Efros
Shklovskii
Pollak
Mott VRH conductivity. This approach also helps to improve the understanding the electronic structure and transport of conducting a-C by resolving some typical problems in the analysis of temperature-dependent conductivity. Bhattacharyya (2004) also claimed that the experimental results are in good agreement with theoretically predicated electronic structure of ultrananocrystalline diamond. Miyajima et al. (2009) studied the carbon-based bottom-gate thin-film transistors behaviors using pulsed laser deposited a-C and a-C nitride (a-CNx) films. Miyajima et al. (2009) observed that both series of devices have p-type conduction in the active channel at high electric fields, for which the conduction mechanism fitted to Poole
Frenkel-type behavior. The field-effect mobilities were 2.5 3 1023 cm2/V/S at high fields. The value of the mobility is almost one order higher than previous reports, since the conduction is at or close to the bottom of the extended states. Nitrogen inclusion does not appear to affect the performance of

Chapter 3 • Magnetism and spintronics in amorphous/diamond-like carbon

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the a-CNx at high fields since the current is postulated to be controlled by hole conduction states. It is noted that the mobility obtained in this study is that of the hole mobility through the edge of the valence band, where holes are trapped in trapping centers. The maximum reported density of trapping centers estimated from the Poole
Frenkel model analysis is 1021 cm23 (Miyajima et al., 2007). Therefore the hole mobility through valence band states with a less number of the trapping centers is thought to be much higher.

3.6 Diamond-like carbon for magnetic storage disks Magnetic storage is the most economic form of nonvolatile storage for many applications (Ferrari, 2004; Grochowski et al., 1994; Menon et al., 1999; Robertson, 2001, 2003; Goglia et al., 2001; Doerner et al., 2000; Wood, 2000). Flash is used when semiconductor storage is desired for close electrical integration with a microprocessor or random access memory. Optical storage in DVD is a demountable storage with higher storage densities, but it has a slower access time than magnetic storage. Its great advantage is that the storage density is increasing at a very rapid rate (Grochowski et al., 1994; Menon et al., 1999; Robertson, 2001, 2003; Goglia et al., 2001; Doerner et al., 2000). Recently, with the introduction of giant magnetoresistive heads, storage densities are increasing at 100% per year. This is much faster than the Moore’s law rate for silicon devices (50% per year). DLC films form a critical protective layer on magnetic hard disks and their reading heads. The ultimate limit to storage density is the superparamagnetic limit, where the thermal energy is able to overcome the coercive energy of the magnetic bit. Perpendicular recording should allow storage densities up to B1 Tbit/in.2 This requires the read head to approach closer to the magnetic layer and ever thinner layers of carbon 1
2 nm thick. Data are stored in a magnetic layer of Co
Cr
Pt alloy thin film (Doerner et al., 2000). A protective layer of DLC coating is applied over the Co layer, and one to two monolayers of a perfluoro-polyether such as ZDOL or Fomblin is used as a molecular lubricant. A read/write head flies above the rotating disk on an aerodynamic bearing. The ready write head consists of many layers of thin films and is also protected by a DLC film. The storage density is increased by reducing the area occupied by each bit of data. The areal density is the product of the tracks per inch and the bits per inch along a track. The ratio of tracks per inch and bits per inch is called the bit
aspect ratio. Initially, this ratio was approximately 20 (Goglia et al., 2001). The ultimate limit to storage density is the superparamagnetic limit, where the thermal energy kT can overcome the coercive energy of the magnetic bit (Ferrari, 2004; Grochowski et al., 1994; Menon et al., 1999; Robertson, 2001, 2003; Goglia et al., 2001; Doerner et al., 2000; Wood, 2000). In 1995 this limit was approximately 40 Gbits/in.2. The limit was increased to 200 Gbits/in.2 by reducing the bit
aspect ratio to 4 and by using materials of higher coercivity (Menon et al., 1999; Robertson, 2001, 2003; Goglia et al., 2001). Recently, Seagate and Fujitsu achieved a storage density of 100 Gbits/in.2 in laboratory demonstrations. One terabit per square inch may be possible, but using perpendicular recording, where the magnetization is normal to the film surface, in contrast to the standard longitudinal recording, where it is parallel to the surface

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

(Wood, 2000). A smaller bit size requires a smaller magnetic spacing, which is the vertical distance between the read head and the Co storage layer. The magnetic spacing is slightly greater than the fly height, which is the separation of head and disk. Reducing the fly height requires ever thinner carbon films. They are presently approximately 4 nm thick and need to reach 1
2 nm in the near future (Menon et al., 1999; Robertson, 2001). Indeed, in order to achieve the goal of B1 Tbit/in.2 the magnetic spacing must be reduced to 6.5 nm, which implies a B1 nm head and disk overcoat. This is only approximately seven-atomic-layer thick, and the performance of the carbon and the processes used to make it change dramatically when we approach 1 nm.

3.7 Conclusion and perspectives of amorphous carbon and diamond-like carbon in spintronics There have been many advances in the materials science and applications of a-C, DLC, and ultra-nanocrystalline diamond (UNCD)/diamonds. The development of previously introduced semiconducting materials has followed a well-established research and technological development progression. With the extreme and unique properties of DLC and UNCD/diamond, research has opened more doors rather than following an established development path. There are now many opportunities and challenges that will influence the broad range of technologies presented in this issue. It is necessary to improve doping efficiencies and reduce the effects of impurities and defects. Conversely, specific defect and impurity structures may enable their use for quantum devices. Identifying shallow dopants and new quantum defects and effects continues to be a challenge. In fact, research on charge transport and carrier dynamics has hardly been explored. For instance, the minority carrier lifetime, which is crucial in bipolar devices, has only been studied indirectly. High-quality heterostructures and metal interfaces have mostly been studied to enable specific devices, while indepth understanding is still missing. For instance, it was found that the Co-implanted DLC could be similar to dilute ferromagnetic semiconductors that are reported in other cobaltdoped semiconductors. It is noted that the bonding of Co to C in DLC is possible (Harris et al., 2010; El-Gendy et al., 2014; Zhang et al., 2011), and hence Co-doping of DLC behaves as the dilute ferromagnetic semiconductor that attributed to an Ruderman
Kittel
Kasuya
Yosida (RKKY) interaction and the Curie temperature depends on the carrier concentration (Dobrowolska et al., 2012). If this has occurred in Co-implanted DLC, then it is possible that there is an inhomogeneous Co distribution that leads to an inhomogeneous carrier concentration and hence a distribution in the Curie temperature that could result in mCo,Sat decreasing with increasing temperature. In this scenario the Co distribution would be different for low and high Co concentrations, and hence it would not be surprising if the temperature dependence of mCo,Sat was different. In the case of Cr-DLC and chromium carbide hydrogenated DLC alloys, a large coefficient of negative MR is observed in heterojunction devices with n-type silicon. The negative MR of the I
V curve, which is ascribed to uncompensated spins at the surface of the antiferromagnetic chromium carbide clusters, indicates that the

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material is suitable for spin-electronics applications. The surface hole-conducting channel continues to be a focus of basic research, as the fundamental properties of the layer are still not well characterized. These structures will enable a range of “unconventional” electronics that will bring together surface chemistry and electronics in a unique way.

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8192.

Magnetism and spintronics in carbon nanotubes

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4.1 Introduction Carbon nanotubes (CNT) have a curved, aromatic structure with electronic properties of metals to semiconductors of varying bandgaps as a function of chirality and display excellent electrical, thermal, and mechanical properties (Harris, 1999; Dresselhaus et al., 1996), large specific surface area, which are highly desirable for their enormous potential applications particularly for nano-electronics. Due to the covalently bonded structure, metallic CNTs are ballistic conductors more than 1 μm lengths and can carry current densities of 109 A/cm2 that can be designed to be the ultimate electronic circuits. Thus CNTs have been regarded as the most promising electronic material as silicon devices reach their fundamental scaling limitations. For the last three decades, since the discovery of CNTs by Iijima (1991), synthesizing useful quantities of analytically pure CNTs as well as characterizing and processing into forms suitable for specific applications are largely unsolved problems. To realize the potential advantages of CNTs in electronic circuits, besides controlling the materials characteristics, methods to assemble a large density of functional devices need to be developed. Due to low (small) spin
orbit coupling (SOC) interactions and long spin diffusion lengths ( . 100 nm) due to their ideal π-electron system, CNTs are considered very promising and lots of interests for CNT-based spintronic applications. There are two types of CNTs: single-walled CNTs (SWCNTs) and multiwalled CNTs (MWCNTs). CNTs containing nano-magnetic materials are able to form a perfect spin-transport medium, since electron transport in them is one-dimensional and ballistic with a long spin relaxation time and weak spin
orbital effects. Even pure CNTs, which are nonmagnetic materials, are characterized by a giant magnetoresistance (MR) (GMR) (Cottet et al., 2006; McIntosh et al., 2002). On the other hand, it is obvious that the modification of CNTs would lead to significant differences in their electronic structure and other different properties (Tyagi et al., 2005; Mykhailenko et al., 2007; Borowiak-Palen et al., 2005). Due to very large magnetic shape anisotropies, the encapsulation of magnetic phases in CNTs could provide a feasible approach to achieve magnetic order stabilization against thermal fluctuations in systems having extremely reduced dimensions. The ferromagnetic (FM) nanoclusters are expected to have much better magnetic properties than bulk metals due to their single-domain nature (Grobert et al., 1999) and are very useful for spintronic devices fabrication. Therefore it is desirable to produce CNTs not only with magnetic material inside of the tubes in a specific and controlled way but also further modification with specific species. Beyond the Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials. DOI: https://doi.org/10.1016/B978-0-12-817680-1.00004-4 © 2020 Elsevier Inc. All rights reserved.

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geometrical advantage of a quasi-one-dimensional CNT design, the carbon shells can provide an effective protection against oxidation. The growing of CNTs is, however, a catalyzed-determined process. The most used catalyst materials are Fe, Co, and/or Ni. All these metals show over a wide temperature range of FM properties. Therefore alternative techniques have to be developed in order to either remove completely the catalyst material from the nanotubes or to apply nonmagnetic catalysts. Various purification methods have been employed to remove magnetic impurities, such as chemical treatment, microwave heating, mechanical filtration, and heat treatment in a vacuum or oxidative environment (Rinzler et al., 1998; Strong et al., 2003; Harutyunyan et al., 2002; Zhou et al., 2001). However, a graphitic coating commonly found around FM catalyst particles shields the particles from acid dissolution. Attempts to remove this graphitic coating often result in damage or destruction of SWCNTs (Zhou et al., 2001). Although some groups applied magnetics filtration, the efficiency was low such that ferromagnetism still dominated the magnetic moment of the sample for fields of order a few Tesla (Thien-Nga et al., 2002; Wiltshire et al., 2005; Islam et al., 2005; Kim et al., 2005). To circumvent this problem, other researchers synthesized nanotubes using non-FM catalysts, such as Rh/Pd or Rh/Pt (Goze-Bac et al., 2002). Lipert et al. (2009) show two different ways to obtain CNTs having diamagnetic behaviors (nonmagnetic). In a CNT spin-valve device, spin-polarized electrons are generated at the FM/CNT interface, and the CNT provides a coherent path for the polarized electron, which can be detected at the other FM electrode. A versatile method for assembling CNTs on FM metal contacts from solution was already been developed. The quantum mechanical spin degree of freedom is now widely exploited to control current transport in electronic devices. Based on these regards, CNTs are considered to be a potential material for spin electronics device applications. CNTs exhibit a long electron mean free path (Bachtold et al., 2000; Yao et al., 2000) and weak SOC; thus the spin diffusion length is expected to be extremely large (Dresselhaus et al., 2001). Spin-dependent transport in CNTs was originally demonstrated in MWCNTs with FM contacts (Tsukagoshi et al., 1999). However, there are still several problems to overcome before their potential in this application can be realized. The directional placement of immobilized CNTs in aligned geometries on electrodes represents is one of the critical steps toward the creation of the spintronic devices. The MR effect (Alphenaar et al., 2001; Zhao et al., 2002) in a FM tunnel junction consists of two FM electrodes (source and sensor) separated by a non-FM layer (spacer), which is the key for spintronic devices. Because of the GMR at room temperature and the application for magnetic random access memories (MRAMs) (Ono et al., 1997, 1998), the MR effect is caused due to the spin polarization of the FM electrodes. It has been known since the work of Julliere (1975) that the resistance of single planar FM junctions decreases when the magnetizations of the electrodes rotate from parallel to antiparallel alignment. CNTs can be synthesized by various methods, such as arc discharge (Iijima, 1991), chemical vapor deposition (CVD) (Cassell et al., 1999; Ren et al., 1998; Li et al., 1996), and laser ablation (Maser et al., 1998). In this chapter, magnetization, MR, field-effect transistor (FET)
based transport properties of different catalysts (with/without) on the basis of SWCNT/MWCNTs are reviewed and discussed for spintronic applications.

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4.2 Magnetism of carbon nanotubes Ray et al. (2010) studied the magnetic properties of as-grown and postannealed MWCNTs with embedded iron nanoparticles at room temperature (B300K) and below room temperature (B5K) to investigate the hysteresis of the magnetic behavior. Anisotropy measurements were also performed by automated sample rotation; therefore a magnetic field (H) is applied both parallel and perpendicular to the tube axis. Fig. 4
1A and B and C and D plots the resulting magnetic moment M and magnetic field H measured at B305 and B5K, respectively, of the as-grown and postannealed MWCNTs. Fig. 4
1B shows that after the application of magnetic field parallel to the tube axis, there is an interface between FM and anti-FM materials in the postannealed MWCNTs network, so the coercivity is not symmetric. The FM behavior was apparently eliminated when the magnetic field is applied perpendicular to the tube axis. Ray et al. (2010) observed that MWCNTs exhibit significantly a higher coercivity B750 Oe than its bulk counterpart (Febulk  0.9 Oe), suggesting its potential use as lowdimensional, high-density magnetic recording media. The magnetic moment of the

FIGURE 4–1 Magnetic hysteresis (M
H) loop obtained at 300K when magnetic field is applied perpendicular and parallel direction of (A) as-grown and (B) postannealed MWCNTs. Magnetic hysteresis (M
H) loop obtained at 5K when magnetic field is applied perpendicular and parallel direction of (C) as-grown and (D) postannealed MWCNTs. MWCNTs, Multiwalled carbon nanotubes. Ray, S.C., et al., 2010. High coercivity magnetic multi-wall carbon nanotubes for low-dimensional high-density magnetic recording media. Diamond Relat. Mater. 19, 553
556. Copyright Elsevier Publishing. Reproduced with permission.

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postannealed MWCNTs is nearly 104 times lower than those of as-grown samples, and the magnetization, H, gradually decreases more than 2000 Oe, which is due to the diamagnetic contribution. Ray et al. (2010) strongly believe that the disappearance of FM behavior is attributed to the formation of Fe-carbides and Fe-oxides with a specific composition upon high-temperature annealing or the increase of kinetic bandwidth. The coercivity (Hc) of the as-grown sample significantly exceeds (B750 Oe) that of the bulk Fe counterpart (Febulk  0.9 Oe) (Bozorth, 1951) and Ni/Co nanowire arrays (Whitne et al., 1993; Ozhuharova et al., 2004), and it is comparable with the values obtained elsewhere (Kuo et al., 2003; Geng et al., 2006). These findings are encouraging for various technological applications and suggest a higher magnetic stability. In addition to the structure and alignment of the nanotubes, two other factors are responsible for the significant increase in Hc values: (1) First, Hc is well known to increase as the magnetization becomes higher and the feature size shrinks (Bertotti, 1998; Mills et al., 2006). The embedded Fe nanoparticles found at the bottom of the MWCNTs (Ray et al., 2010) are somewhat smaller than the typical domain of a size that can be magnetized by coherent rotation. It is noted that if the feature or particle size becomes comparable to the domain wall width (for Fe, it is B40 nm), the Hc begins to drop with feature size and magnetic hysteresis decreases as well, which seems quite possible for the postannealed MWCNTs. Moreover, increased surface-to-volume ratio for nanoparticles makes them susceptible to interaction with neighboring magnetic materials. (2) Second, large shape anisotropies of nanotubes can act on encapsulated Fe nanoparticles as shown in Fig. 4
1A. It is interesting to see, although the hysteresis loops look similar for parallel and perpendicular directions for the asgrown samples at the lowest recorded temperatures (B5K), they differ significantly for the postannealed MWCNTs establishing a high magnetic
anisotropic nature of iron particles in the MWCNTs, as shown in Fig. 4
1C and D. At low temperatures, Hc is found to increase for both MWCNTs; however, this dependence particularly for the postannealed samples can be found to be significantly different from previous reports. The coercivity is significantly enhanced at low temperature (B5K) and is nearly 2600 Oe for as-grown MWCNTs and B320 Oe for postannealed MWCNTs. The postannealed MWCNTs show high anisotropic behaviors at this low temperature (B5K) which can be seen in Fig. 4
1D. In fact, the shape of anisotropies helps to stabilize magnetic order against thermal fluctuations in low-dimensional systems. However, if the Fe particles used as the seeds for the nucleation of the nanotube are small enough, say, less than 20 nm (in the present case), the particles will be a single domain and will exhibit very high uniaxial anisotropy due to the stress of the nanotube and the shape. The possible mechanism is expected to be similar to the case studied by Geng et al. (2006) using Fe nanoparticles as catalysts for growing CNTs where the formation of the Fe3C was observed by the x-ray diffraction (XRD) measurements (Ray et al., 2010). These findings are encouraging for various technological applications and suggest higher magnetic stability. Geng et al. (2006) obtained the high coercivity and high loop squareness in Fe nanoparticles
filled CNTs array that could be an interest for magnetic recording as well as magneto-electronic sensors applications. Li et al. (2007) studied the magnetic characterization of Fe nanoparticles
encapsulated SWCNTs and compared with pristine SWCNTs at room temperature (B300K) and below

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room temperature (B5K). Li et al. (2007) found that the Fe nanoparticles
encapsulated SWCNTs have higher magnetization. To gain a better understanding of the magnetic behavior of Fe-filled SWCNTs, Li et al. (2007) performed zero-field-cooled (FC) (ZFC) and FC magnetization measurements. It was noticed that the Fe-filled SWCNTs striking sharp decrease in the FC curve at about 25K, corresponding to the existence of small magnetic particles in the superparamagnetic state. For ZFC curves, as the temperature increases, the magnetization shows an increase because the magnetic moment is thermally activated along the magnetic field direction. For Fe-filled SWCNTs a blocking temperature (Tb 5 transition temperature from FM to superparamagnetic state) peak can be observed in the ZFC curve at about 94K. By contrast, the Tb peak can be identified at near room temperature of 255K for pristine SWCNTs. This was similar to the magnetic behavior of iron-filled MWCNTs suggesting the presence of large magnetic particles in the case of our SWCNTs. Undoubtedly, the above results reveal an obvious difference between pristine and Fe-filled SWCNTs, and superparamagnetic properties of Fe-filled SWCNTs can be observed due to the small size distribution of Fe nanoparticles inside the SWCNTs. Zhang et al. (2001) studied the magnetic properties of Fe-nanoparticles trapped at the tips of the aligned CNTs in the temperature range of 5K
360K and found that the CNTs are highly FM in nature. They observed that the Fe particles behave ferromagnetically with Curie temperature much higher than 350K. Magnetization was measured in the magnetic field along the perpendicular and parallel to the film plane and found that the CNTs are highly anisotropy in nature. They obtained different magnetic parameters, namely, Mr/Ms ratio and coercive field Hc from M
H hysteresis loops and observed that these parameters decrease monotonically with the increase in temperature that can be attributed to the depinning of domain walls in the particles. In this study, HcB2.5 kOe was also obtained, which is very useful for the next-generation high-density recording media. Another advantage of the particles in the tips of the nanotubes is that the walls of the nanotubes act as a nonmagnetic separation, which is essential for the high magnetic recording media to eliminate the dipolar interaction between the neighboring particles. It is very important to note that the tube number density is one of the main factors determining the media recording density. Ray et al. (2017a,b) studied the magnetization of nitrogen-doped multiwall nitrogenated carbon nanotubes (MW-NCNTs) functionalized with chlorine as well as oxygen plasma atmosphere. It was observed that at room temperature, nonfunctionalized MW-NCNTs have diamagnetic behaviors, whereas chlorine- and oxygen-functionalized MW-NCNTs, respectively, hold paramagnetic and FM behaviors. The magnetization M
H hysteresis loop of MW-NCNTs obtained at 300K and 5K is shown in Fig. 4
2A. The behavior of spectral features unambiguously implies a pure diamagnetic behavior, although the nanotubes have contribution of strong magnetic Fe particles as a catalyst in the NCNTs. Lipert et al. (2009) observed the similar diamagnetic behavior of Fe-based MWCNTs, after the postannealing process at a very high temperature of B2500 C. Lipert et al. (2009) claimed that the FM behavior changed to diamagnetic due to complete evaporation of Fe-catalyst particles from the CNTs at this high-temperature annealing. They also observed the synthesis of

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FIGURE 4–2 M
H hysteresis loop of (A) MW-NCNTs, (B) chlorine-functionalized MW-NCNTs, and (C) oxygenfunctionalized MW-NCNTs (black solid ball and red hollow ball are at 300K and 5K, respectively). MFC and MZFC, M
T spectra of (D) MW-NCNTs, (E) chlorine-functionalized MW-NCNTs, and (F) oxygen-functionalized MW-NCNTs. Ray, S.C., et al., 2017a. Change of magnetic behaviour of nitrogenated carbon nanotubes on chlorination/ oxidation. Int. J. Nanotechnol. 14 (1
6), 356
366; Ray, S.C., et al., 2017b. Hall effect studies and magnetic behaviour in Fe-nanoparticle embedded multiwall CNTs. J. Nanosci. Nanotechnol. 17 (12), 9167
9171. By courtesy of Copyright Inder-Science Publication.

diamagnetic behavior for the CNTs using nonmagnetic Re as a catalyst. Ray et al. (2017a,b) expected that the diamagnetic behavior of MW-NCNTs may be due to the presence of nonmagnetic bonding that dominates the Fe catalyst in the MW-NCNTs structure. Furthermore, the oxygen-functionalized NCNTs show strong FM behavioral M
H hysteresis loop (opposite trend of pure NCNTs) unlike nonfunctionalized NCNTs, as shown in Fig. 4
2C. Del Bianco et al. (2009) observed the FM behavioral M
H loop for the core interface of oxygen-passivated Fe nanoparticles. Ray et al. (2017a,b) also expected that the FM behavior may occur due to oxygen passivation with the NCNTs on oxidation (oxygen plasma treatment). In case of chlorine-plasma-treated NCNTs, the M
H loops in Fig. 4
2B are not like either pure NCNTs or NCNTS:O, but interphase of those two (dia- and ferro-) that indicate the possible of paramagnetic behavior. Apart from that, they even expected that these magnetic behavioral changes occur due to the formation of different bonding with

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carbon/nitrogen/Fe catalyst on chlorine-/oxygen-plasma functionalization process. Ray et al. (2017a,b) also studied the thermal evolution of the magnetization of MW-NCNTs (:Cl/O) with temperature (T)-dependent magnetization (M) by the zero-field-cooling (MZFC) and field-cooling (MFC) procedures in an applied magnetic field of 1000 Oe in between 5K and 300K. Fig. 4
2D shows the M
T curve of N-CNTs, whereas N-CNT:Cl and N-CNTs:O are shown in Fig. 4
2E and F, respectively. Ray et al. (2017a,b) found that the MZFC curve gradually deviates from the MFC curve with the decrease of temperature at B255K (for MWNCNTs), B200K (for CNTs:Cl), and B300K (for CNTs:O), when the applied magnetic field is 1000 Oe. Upon further cooling the MZFC plot exhibits a cusp centered at B45K for the MWNCNTs and NCNTs:O, but not in NCNTs:Cl. This variable temperature magnetic data of MW-NCNTs and MW-NCNTs:O clearly indicate that the Fe/N-CNTs exhibit FM behavior below the room temperature, which is attributed to the uncompensated surface spin states or FM Fe clusters, although the M
H curve of nonfunctionalized MW-NCNTs shown in Fig. 4
2A is completely diamagnetic in behavior. It is believed that this FM performance of the Fe/N-CNTs:O comes from the FM Fe clusters with the formation of different bonding with carbon/nitrogen and the uncompensated surface spin states. In case of NCNTs:Cl the MZFC and MFC curves coincide up to B200K, as shown in Fig. 4
2D, when measured at an applied magnetic field of 1000 Oe, which get split below that temperature. A similar behavior has been observed by Del Bianco et al. (2009) to occur in oxygen-passivated Fe nanoparticles. In that case the anti-FM character of Fe2O3 was in the origin of the low-temperature irreversibility. This low-temperature FM phase magnetization is correlated to the fact that at lowest temperature and after MZFC process, the moments of magnetic particle Fe are not fully aligned with the applied field. Furthermore, no cusp is observed in the MZFC plot in N-CNT:Cl indicating non-FM nature. In case of N-CNT:O, it is found that the MZFC curve gradually deviated from the MFC curves with the decrease of temperature at about 300K as shown in Fig. 4
2F, when measured at an applied magnetic field of 1000 Oe. A similar behavior has been observed by Zhang et al. (2008) for CoO/CNTs core
shell nanostructures, when they measured at an applied magnetic field of 100 Oe between 2K and 300K. In our present case, upon further cooling, it is also observed that the MZFC plot exhibits a cusp centered at about  45K and the MFC data sequentially increases, indicating FM behavior at this temperature compared to NCNTs and NCNTs:Cl. It is believed that the FM behavior in NCNTs:O comes from the FM Fe clusters and uncompensated surface spin states owing to the formation of different bonding with carbon/nitrogen/Fe catalyst. Kapoor et al. (2018) studied the magnetization of 3d transition metals and oxides within CNTs by copyrolysis of metallocene and camphor and found that the Fe, Co, and Ni composites with CNTs have excellent magnetic behaviors and could be used for spintronic applications. They emphasized that the use of camphor enabled to obtain Ni- and Co-filled CNT with significantly improved Ms values (up to 12 emu/g), as compared to what one obtains for aerosol-assisted CVD (0.1
0.4 emu/g) reported earlier (Sun et al., 2013; Yang et al., 2016; Terrones et al., 2006). In the case of Ni@CNT, pyrolysis of only nickelocene in powder form is not known to yield well-formed CNT, though there exist reports of copyrolysis of nickelocene with ferrocene in aerosol-assisted CVD for the formation of Ni@CNT. Fig. 4
3A shows

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normalized Ms, limited to low-field scans, for a clear depiction of the coercivity, Hc, for each type of CNTs. The coercivities, Hc, for Fe@CNT and Co@CNT are approximately few 100 Oe and can be further tuned by variations in Tpyro, depending on the application-specific requirement. Here, the Ni@CNT exhibits very low Hc, almost mimicking a superparamagnetic type of behavior. However, the magnetization data is likely to contain contributions from both Ni and Fe in such cases (Yang et al., 2016). Copyrolysis of camphor with nickelocene, as shown in this figure, enables the formation of Ni@CNT with no such ambiguity in the magnetic data. Fig. 4
3B shows the magnetization as a function of temperature for Fe2O3@CNT, clearly exhibiting the Morin transition, which signifies the onset of weak ferromagnetism and piezomagnetism in this CNT
metal oxide hybrid. These CNTs could be synthesized by systematic variation of synthesis parameters and possible to obtain self-organized structures of filled CNT, with narrow diameter and length distribution and reduced residue particle density. The metal@CNT can be used as a template to form oxide@CNT that could be useful for numerous applications in spintronics, magneto-optics, and in the energy sector. Narayanan et al. (2009) studied the magnetic properties of MWCNT
SPION (superparamagnetic iron oxide nanoparticle) composites at low temperature, ZFC and FC experiments. The ZFC shows a blocking at B110K. A peculiar FM ordering is exhibited by the MWCNT
SPION composite above the room temperature due to the FM

FIGURE 4–3 (A) M
H loops at 300K limited to 6 1 kOe, depicting the coercivity (Hc) for Fe@CNTs, Co@CNTs, and Ni@CNTs. (B) M
T at 1 kOe, depicting a magnetic transition at B250K. Kapoor, A., et al., 2018. 3d Transition metals and oxides within carbon nanotubes by co-pyrolysis of metallocene & camphor: high filling efficiency and self-organized structures. Carbon, NY 132, 733
745. Copyright Elsevier Publishing. Reproduced with permission.

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interaction emanating from the clustering of superparamagnetic particles in the constrained volume of an MWCNT. This kind of MWCNT
SPION composite can be envisaged as a good agent for various biomedical applications. Titus et al. (2011) studied the magnetization behavior of nickel nanoparticle
deposited vertically aligned CNTs for the magnetic tunnel junction (MTJ) spintronic application. To probe the magnetic properties the field dependence of the magnetization was measured using a superconducting quantum interface device (SQUID) (Kenane et al., 2006). They observed that the vertically aligned CNTs exhibit FM behavior with a hysteresis loop at 2K and 300K, respectively. The saturation magnetization (Ms), coercivity (Hc), and remnant magnetization (Mr) for both temperatures were obtained from the M
H hysteresis loops. The saturation magnetization values and coercivity for nickel nanoparticle
deposited vertical carbon nanotubes (VCNTs) shows a decrease from 5.3 to 4.4 emu/g and 395 to 115 Oe, respectively, with the increase of temperature, indicating a characteristic FM behavior. To gain a better understanding of the magnetic behavior of nickel nanoparticle
deposited VCNTs, Titus et al. (2011) performed ZFC and FC magnetization measurements. For the ZFC measurement the nickel nanoparticle
deposited VCNTs sample is first cooled from 300K to 2K in zero magnetic field. On the other hand, for FC measurements the sample is cooled in the magnetic field (25 G) from 300K to 2K, and later the magnetization is measured in the warming cycle keeping the field on. The temperature dependence of ZFC and FC measurements under the applied magnetic field of 25 G for nickel nanoparticle
deposited VCNTs which exhibits the main features of FM behavior (Li et al., 2007). The blocking temperature TB (transition temperature from FM to superparamagnetic state) peak can be observed in ZFC curve at B44K. The low value of TB is directly in agreement with smaller size of nickel nanoparticles randomly deposited on the VCNTs (Fonseca et al., 2002; Linderoth et al., 1993).

4.3 Spintronic devices The quantum mechanical spin degree of freedom is exploited to control current transport in electronic devices. The readout of magnetic hard disks is based on the spin-valve effect, that is, the tunability of a conductance through the relative orientation of some FM polarizations (Prinz, 1998). Realizing of spin injection in nanostructures, for example, mesoscopic conductors or molecules, would allow to implement further functionalities. The realization of a “spin transistor” would allow electric field control of the spin-valve effect through an electrostatic gate (Datta et al., 1990; Schäpers et al., 2001). CNTs are particularly interesting, because they should exhibit a long spin lifetime and can be contacted with FM materials. This section presents the state of the art regarding the realization of spin transistor
like devices with CNTs. The most standard method to inject or detect spins in an insulating or conducting element M is to use the spin-valve geometry (Baibich, 1988; Binasch et al., 1989) in which M (CNTs) is connected to two FM leads L and R (Fig. 4
4, left). One has to measure the conductances GP and GAP of the spin valve for lead magnetizations in the parallel (P) and

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antiparallel (AP) configurations. This requires the use of two FMs with different coercive fields (HcL and HcR, respectively) for switching one magnetization with respect to the other with the help of an external magnetic field H (Fig. 4
4, right). The spin signal or MR is then defined as the relative difference MR 5 (GP 2 GAP)/GAP. Let us consider the situation in which the element placed between the two FM contacts is a tunneling barrier with a transmission probability independent of energy (Julliere, 1975). This case, usually referred to as Julliere’s model, describes the principle of magnetic memories and magnetic read From Fermi’s golden rule the transmission probability of the  heads. barrier for spins σE m; k is proportional to the electronic densities of states at the Fermi
energy (FE) Nl;σ 5 Nl 1 1 σɳl pl for spin s at the both contacts, with lAL; R and ɳ A 1 1; 2 1 the direction of the magnetization at contact l. Here, Nl is the spin-averaged density of states (DOS), and pl is the spin polarization at contact l. The conductance GP of the barrier in the parallel configuration is proportional to NLNR[(1 1 pL)(1 1 pR) 1 (1 2 pL)(1 2 pR)], whereas the conductance GAP in the antiparallel configuration is proportional to NLNR[(1 1 pL)(1 1 pR) 1 (1 2 pL)(1 2 pR)]. This leads to MR 5 2pL PR =ð1 2 2 pL PR Þ: If the spin polarizations pL and pR have the same sign the MR of the device is positive because the current flowing in the antiparallel configuration is lower due to the imbalance between NL,σ and NR,σ. The working principle of a spintronic device follows these steps: (1) information is stored into spins as an orientation (i.e., up or down), (2) spin information is carried by mobile electrons along a path or wire, and (3) the information is then read at a final point. The spin orientation of conduction electrons will exist for several nanoseconds making them useful in electronic circuit and chip design. The most basic method of creating a spin-polarized

FIGURE 4–4 Left: Spin-valve geometry. The CNT is connected to two ferromagnetic leads L and R with magnetic ~ is the applied magnetic field and Vsd is source
drain voltage. Right: Resistance curve r polarizations ~ P L and ~ P R. H (H) measured in the spin valve. Increasing [blue line] and then decreasing [red line] H. L and R have different coercive fields HcL and HcR, it is possible to selectively reverse the directions of ~ P L and ~ P R during this cycle. MR 5 (GP 2 GAP)/GAP and MR . 0. MR, Magnetoresistance. Cottet, A., et al., 2006. Nanospintronics with carbon nanotubes. Semicond. Sci. Technol. 21, S78
S95. Copyright IOP Publishing. Reproduced with permission.

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current is to transport current through a FM material and to transmit the electron spin carrying the information to the receiver point. Spin current is, therefore, an important tool to detect spin in spintronic devices. The important avenues for the development of spintronic devices are (1) fabrication of nanoscale nanostructures, including novel magnetic materials, thin films, hybrid structures, and functional materials; (2) research on spin effect (spin injection, and spin transport and detection); (3) demonstration of spintronic devices including GMR and tunnel MR (TMR) devices in MTJs; and (4) study of single electron tunneling (SET) in MTJs. Mohamed et al. (2008) describe an alternative method for realizing a CNT spin FET device by the direct synthesis of SWCNTs on substrates by alcohol catalytic CVD. They observed that the hysteretic MR at low temperatures due to spin-dependent transport and the maximum ratio in resistance variation of MR was found to be 1.8%.

4.3.1 Magnetic tunnel junctions A MTJ can be considered a spintronic device as it is composed of two FM materials, such as nickel, cobalt, or iron, separated by an ultrathin layer of insulator with a thickness of the order of nanometer (1029 m). It exhibits two resistances, low (RP) and high (RAP), depending on the relative direction of FM materials, parallel (P) or antiparallel (AP), respectively. The insulating layer is so thin that electrons can tunnel through the barrier if a bias voltage is applied between the two metal electrodes. In MTJs the tunneling current depends on the relative orientation of magnetizations of the two FM layers, which can be changed by an applied magnetic field. This phenomenon is called TMR. An important factor in TMR is the interaction between the electron spin (S) and angular momentum (L), that is, SOC. An example of SOC is splitting of hydrogen spectrum (Bratkovsky, 1998; Zhang et al., 1998; Dholabhai et al., 2008). The SOC deforms the electron shell as the direction of the magnetization rotates. This deformation also changes the amount of scattering undergone by the conduction electrons when traversing the lattice. There will be minimum resistance if the magnetizations are in parallel orientation, and it will go to maximum with opposite orientations (Fig. 4
4). Therefore such kind of junction can be easily switched between two states of electrical resistance—one with low and the other with very high resistance.

4.3.2 Fabrication of magnetic tunnel junctions The fabrication of MTJs with high TMR ratios is crucial in developing spintronic devices. With the advance of nanotechnology, there are various methods to deposit MTJs, such as molecular-beam epitaxy, magnetron sputtering, electron-beam evaporation and CVD, and so on. In detail, the MTJ’s main components are FM and insulator layers. The FM layers can be fabricated by sputter deposition (magnetron sputtering and ion-beam deposition). The fabrication issue is the magnetic alignment and thickness (deposition rates should be in the angstrom-per-second range). The best way of fabricating insulating layer is still under research. Some of the proven materials are Al2O3 tunnel barriers made by depositing a metallic aluminum layer in the range of 5
15 Å thickness. In addition, ion-beam oxidation,

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glow discharge, plasma, atomic-oxygen exposure, and ultraviolet-stimulated oxygen exposure are also alternate ways of insulator deposition. Since the first report on TMR by Julliere (1975), many studies have been performed to explore this property, especially on Al2O3 insulating layers. The necessity of controlling the magnetic properties of the magnetic layers introduces special requirements on the deposition process. The maintaining of inherent magnetic anisotropy is crucial in the deposition process. This can be set by applying magnetic field during deposition. The thickness and uniformity of the material, the coercivity, magneto-restriction, all are important in controlling the magnetic anisotropy.

4.3.3 Tunnel magnetoresistance in magnetic tunnel junctions Since its discovery, a large number of nano-devices such as single-electron transistor, light emitting diode, FET have been demonstrated. However, these devices are based on charge of the electron. TMR is a magneto-resistive effect that occurs in component consisting of two ferromagnets separated by a thin insulator (MTJ). The interest toward TMR is driven by the fact that MTJs with spin-dependent tunneling (SDT) are expected to provide technical promises that will allow the realization of nanoscale devices in more advanced spintronic applications. Moodera et al. (1995) fabricated the first reproducible TMR up to 24% at room temperature on CoFe/Al2O3/Co or NiFe junction. Today, reproducible TMR value up to 50% can be obtained with three-dimensional ferromagnets making them useful for industrial application (Parkin et al., 1999). TMR characteristics have already been measured in CNTs both experimentally and theoretically (Jensen et al., 2005). Tsukagoshi et al. (1999) demonstrated the MR in a single CNT contacted by FM metal electrodes. The spintronic devices exhibiting TMR using ferromagnet-contacted SWCNTs have been demonstrated by Jensen et al. (2005). Most of the reports on CNT-TMR system are mainly based on single CNT contacted to bulk FM material by an ex situ method (Xiong et al., 2004; Shimada et al., 1998; McEuen et al., 2002). The TMR effect is also known to be sensitive to the tunnel barrier/electrode interface. The barrier sensitivity may be more evident in a system with single CNT. De Teresa et al. (1999) studied FM MTJ with various barrier materials and found that even the sign of the TMR depends on the barrier materials. Yuasa et al. (2004) also investigated the effect of crystal anisotropy of the spin polarization on MTJ using single crystal iron electrodes of various crystal orientations. They found a clear crystal orientation dependence of the TMR; which might reflect the crystal anisotropy of the electronic states in the electrodes. TMR/GMR is known to originate from spin interaction between the magnetic and nonmagnetic particle at the interface and are related to the coercivity value (Bergenti et al., 2004).

4.3.4 Application of magnetic tunnel junctions With wider knowledge on how to manipulate spins (Chappert et al., 2007), we can build more state-of-the-art spintronic devices with extraordinary properties. Extended research into application possibilities of any spintronic effects is, therefore, crucial to realize more advanced spintronic devices. These devices made huge impact on computer technology by enabling higher storage of information in hard drives and faster reading of data in random access memories.

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The first successful application of MTJ was demonstrated in computer read head technology with Al2O3 barrier and MgO barrier MTJ. The magnetic recording density in hard disk drive increased (300
600 Gbit/in.2) considerably in these devices (Sanvito, 2007; Appelbaum et al., 2007; Tsymbal et al., 2007; Chantis et al., 2007). Another application of MTJ is to develop MRAM devices. It has been claimed that MRAM can exceed the speed of SRAM (static RAM), density of DRAM (dynamic RAM), and nonvolatility of flash memory. In addition, the nanodimension device has low power consumption and less heating. MRAM is an upgraded version of SRAM and DRAM where data is stored using spin instead of electrical charges. It overcomes one of the disadvantages of the conventional RAM, the loss of information by power failure. Leading companies, such as IBM, Motorola, and Honeywell, started the MRAM research in 1995, and they were supported by US Defense Advanced Research Projects Agency. The influence of spin transfer in MTJs can be observed by measuring resistive loops as a measure of external applied field and applied voltage. By sweeping the magnetic and electrical field, one can observe sharp drop in resistance which is attributed to the switching from parallel to antiparallel and vice versa. The drop of resistance is associated with the TMR. One of the factors that affect drop of resistance and TMR is DOS at the interface (Tsymbal et al., 2007; Chantis et al., 2007; Burton et al., 2007; Suzuura et al., 2000; Zhu et al., 2006). Spintronics aim to develop electronic devices whose resistance is controlled by the spin of the charge carriers that flow through them (Wolf et al., 2001; Zorpette, 2001; Zutic et al., 2004). This approach is illustrated by the operation of the most basic spintronic device, the spin valve (Julliere, 1975; Slonczewski, 1989; Moodera et al., 1995), which can be formed if two FM electrodes are separated by a thin tunneling barrier. In most cases, its resistance is greater when the two electrodes are magnetized in opposite directions than when they are magnetized in the same direction (Baibich, 1988; Binasch et al., 1989). The relative difference in resistance, the socalled magnetoresistance, is then positive. However, if the transport of carriers inside the device is spin- or energy dependent (Zutic et al., 2004), the opposite can occur and the MR is negative (George et al., 1994). The next step is to construct an analogous device to a FET by using this effect to control spin transport and MR with a voltage applied to a gate (Datta et al., 1990; Schapers et al., 2001). However, several spin relaxation mechanisms have been proposed theoretically for CNT (Semenov et al., 2007, 2010; Borysenko et al., 2008). Semenov et al. (2007, 2010) considered the hyperfine interaction with disordered nuclei spins I 5 1/2 of C13 isotopes (with the natural abundance of 1.10%) in semiconducting CNTs. The anticipated spin relaxation time is about 1 s at 4K, which is still much longer than the experimental observations for spin relaxation time (tens of nanoseconds). Borysenko et al. (2008) considered the anisotropy of the g tensor and flexural phonon modes in semiconducting CNTs. They found that the spin relaxation time can be tens of microseconds at room temperature.

4.4 Spin currents in magnetic tunnel junctions In the view of rapid progress in the fabrication of nanoscale MTJs, spin is a subject of great interest. Spin is a purely quantum mechanical quantity which provides an extra degree of

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freedom for the electron to interact with a magnetic field. In 1922 Stern and Gerlach demonstrated the most direct experimental evidence of the existence and of the quantized nature of the electron spin. The first experimental evidence of SDT was reported by Julliere (1975). Later, Berger (1978) proposed the idea that spin-polarized current act on local magnetization of ferromagnets and leads to GMR. The important property of spin is its weak interaction with the environment and with other spins, resulting in a long coherence or relaxation time, which is a very important parameter in the field of spin-transport and quantum computing. For the successful incorporation of spins into the currently existing electronics, one has to resolve issues such as efficient spin injection, spin transport, control and manipulation of spins, and finally detection of spin-polarized current. Spintronics without magnetism are an attractive pathway for designing semiconductor spintronic devices since SOC enables that the spin is generated and manipulated merely by electric field. By the application of electric field, the electrons move in the lattice generating a magnetic field which acts up on the spin. The spin
orbit interaction on mobile electrons was proved theoretically many decades ago. However, the practical harnessing of this concept is still at an early stage

4.4.1 Spin and charge transport Sahoo et al. (2005a,b) were the first to study the MR of MWCNT/SWCNT by contacting the two FM Pd0.3Ni0.7 strips with either MWCNT or SWCNT that allows to obtain devices with resistances as low as 5.6 kV at 300K. The yield of device resistances below 100 kV, at 300K, was around 50%. Fig. 4
5 shows typical MR curves for the SWCNT (Sahoo et al., 2005a,b). The MR observed is positive (MR 5 5.89%), for a gate voltage Vg 5 4.302 V as shown in Fig. 4
5A, whereas for the same device it is negative (MR 5 22.81%), for a gate voltage Vg 5 4.328 V as shown in Fig. 4
5B. The sensitivity S is of the order of 1%/T or less and can change sign for different Vg. From this figure, one can calculate the local field change ΔHloc required to obtain the observed hysteretic MR. For, Vg 5 23.1 V, one finds ΔHloc 5 22.9/0.2 5 214.5 T, which is negative and way beyond what can be obtained with microstrips. Furthermore, for Vg 5 23.3 V, one would need a positive ΔHloc, since both MR and S have the same negative sign. Such a sign change of the local magnetic field produced by two metallic ferromagnets for different gate voltages. Therefore stray field effects are not dominant in the MR signal for this type of F-MWCNT-F device. In addition, as one can see in Fig. 4
5, S is in general smaller for SWCNTs (Sahoo et al., 2005a,b; Nagabhirava et al., 2006; Man et al., 2006). Man et al. (2006) observed spin-induced MR in the SWCNTs contacted with high-transparency FM electrodes. The MR of SWCNT was measured for different values of the gate voltage, for an applied magnetic field increasing from 2700 to 700 mT (upsweep) and subsequently decreasing from 700 to 2700 mT (down sweep). The widths of the two PdNi contacts in this sample are 150 and 500 nm, resulting in magnetization reversal at magnetic field values of 250 and 125 mT, respectively. A hysteretic feature in the resistance is seen in the expected magnetic field range, corresponding to an antiparallel orientation of the magnetization in the PdNi electrodes. The

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FIGURE 4–5 MR curves for the SWCNT: (A) The MR observed is positive (MR 5 5.89%), for a gate voltage Vg 5 4.302 V. (B) The MR observed for the same device is negative (MR 5 22.81%), for a gate voltage Vg 5 4.328 V. (C) MR observed for an MWCNT connected to two PdNi leads, with different values of Vg. Depending on Vg, both signs of MR and sensitivity S are observed. The amplitude and the sign of S are not correlated with MR. Therefore the stray fields from the ferromagnetic electrodes cannot account for MR observed for this device. (D) Resistance of a Pd
SWCNT
PdNi device as a function of an external magnetic field for two values of Vg. Almost no hysteresis is observed. The maximum amplitude can be estimated (almost within the noise) as G/GB1%, more than an order of magnitude smaller than the observed signal with two PdNi electrodes. MR, Magnetoresistance; MWCNT, multiwalled carbon nanotubes; SWCNT, single-walled carbon nanotubes. Cottet, A., et al., 2006. Nanospintronics with carbon nanotubes. Semicond. Sci. Technol. 21, S78
S95. Copyright IOP Publishing. Reproduced with permission.

feature is superimposed on a smoother background, as expected. The absolute change in resistance has a magnitude of a few tenths of a kiloohm, which is much larger than the total resistance of the PdNi strips. This implies that the effect cannot be accounted for in terms of a change in the resistance of part of the PdNi contacts, which is much lower. The change in MR induced by the magnetization reversal is positive for most values of gate voltage, that is, the antiparallel orientation of the magnetization in the contacts results in an increase of the device resistance. One can conclude that stray field effects do not contribute substantially to MR observed in nanotubes, at least for the PdNi devices realized so far. Fig. 4
5C shows the MR of MWCNTs, whereas Fig. 4
5D shows the MR of SWCNTs at different applied gate voltages Vg. In the following, the results of the MWCNT device are discussed first. Fig. 4
5C shows single traces of the linear response resistance R as a function of the magnetic field H at 1.85K for two sweep directions and four different gate voltages Vg. For all cases the characteristic hysteretic behavior of a spin valve appears. On sweeping the magnetic field from

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2500 to 500 mT, the configuration becomes antiparallel between 0 and 100 mT, whereas it is always parallel for tHt . 100 mT. At Vg 5 23.1 V, for example, R increases from 49.7 to 51.5 kΩ when the sample switches from the parallel to the antiparallel configuration. This yields a normal positive TMR of 12.9%. In contrast, at Vg 5 23.3 V, R switches from 30.5 kΩ in the parallel configuration to a smaller resistance of 29.5 kΩ in the antiparallel configuration, yielding an anomalous negative TMR of 23.5%. Therefore the sign of the TMR changes with the gate voltage, demonstrating a gate-field-tunable MR. Fig. 4
5D shows the G/G measurement performed using a Pd
SWCNT
PdNi device, for two different values of gate voltages, one in the Coulomb valley and the other close to a resonance. The upper bound for G/G is 1.4% in amplitude, which is one order of magnitude lower than the maximum G/G.

4.4.2 Spin polarization As discussed above, spin polarization is an important factor in governing TMR along with the spin transport and spin injection. The spin polarization is a result of a subtle cancellation between two spin channels and is greatly influenced by the atomic, electronic and magnetic structures of the system. To build up on experimental findings, it is also essential to develop an accurate model of the spin polarization and transport of spin current through the FM/non-FM interface and finally into vacuum which is highly sensitive to the chemical and material details of the device. In this context, density functional theories (DFTs) (Arras et al., 2010) of MTJ system that can produce spin polarization effects in the FE are important. DFT is a widely used method for modeling charge/spin carrier transport semiconductors. There is plenty of literature on DFT-based calculations in studying SDT in MTJs (Caffrey et al., 2011; Stilling et al., 2007; Ke et al., 2008). The key components in the modeling are Schottky barrier (τb) and the applied voltage VA against current density. Ab initio Waldron et al. (2007) have demonstrated simulation of MTJs. Chung et al. (2009) report the effect of Schottky barrier profile on SDT in a ferromagnet
insulator
semiconductor system. Ray et al. (2010) studied the high coercivity magnetic MWCNTs for low-dimensional, high-density magnetic recording media. In this study, they used Fe-embedded MWCNTs with B80 mm in length and outer (inner) diameter of 20
50 nm (10
20 nm) and found the coercivity of 2600 and 732 Oe at 5K and 305K, respectively. These values are much higher than that of bulk iron (B0.9 Oe) and Fe/Co/Ni nanoparticles or nanowire arrays (B200
500 Oe) at the room temperature. This high coercivity and the structure of single-domain Fe nanoparticles isolated by anti-FM MWCNTs make it a promising candidate for low-dimensional, high-density magnetic recording media. Tsukagoshi et al. (1999) reported the injection of spin-polarized electrons from FM contacts into MWCNTs and observed the direct evidence for coherent transport of electron spins. A hysteretic MR in several nanotubes with a maximum resistance change of 9%, from which Tsukagoshi et al. (1999) estimated the spin-flip scattering length to be at least 130 nm, an encouraging result for the development of practical nanotube spin-electronic devices. In the MR measurements, Tsukagoshi et al. (1999) performed in 4.2K bath cryostat with the magnetic field (B) from a superconducting magnet directed in the plane of the substrate (BO). Two-terminal resistance is measured using an a.c. lock-in technique with an excitation

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voltage of 100 μV. Tsukagoshi et al. (1999) found that the lead resistance was negligible (B10 Ω) compared with the MWCNT resistance and the Co/MWCNT contact resistance. The field is swept slowly (,10 mT/min); at this sweep rate the MR of MWCNTs contacted by Au/ Pt shows no hysteresis in the applied field. Fig. 4
6 shows the two-terminal differential resistance of three different Co-contacted nanotubes as a function of magnetic field. Each device shows a large hysteretic MR peak. The field was swept first from 2100 to 100 mT (solid line) and then back to 2100 mT (dashed line). In each trace, a resistance peak appears as the magnetic field moves through 0 T. The width of the resistance peak is B50 mT, which is commensurate with the coercive field strength for a thin Co film (Rüdiger et al., 1999). Tsukagoshi et al. (1999) attribute the MR peak to spin-polarized injection (Prinz, 1998; Julliere, 1975; Moodera et al., 1995, 1998; Miyazaki et al., 1995; Meservey et al., 1994; Shang et al., 1998) between the FM contacts and the MWCNT. The magnetization direction of the left and right contacts is represented by the direction of the arrows at the top of the figure. When the magnetizations of the two contacts are parallel, the resistance is lower than when the magnetization of the two contacts is antiparallel. This explanation requires that the spinscattering length in the nanotube is of the order of the contact separation. In addition,

FIGURE 4–6 The solid (dashed) trace corresponds to the positive (negative) sweep direction. The differential resistance shows a large variation among devices—the device shown in (C) has a resistance an order of magnitude lower than the devices shown in (A) and (B). (The magnetic field BO is directed parallel to the substrate, and the temperature is 4.2K.) The percent difference ΔR/RAP between the tunnel resistance in the parallel and the antiparallel states is approximately 6% in (A), 9% in (B), and 2% in (C). Tsukagoshi, K., et al., 1999. Coherent transport of electron spin in a ferromagnetically contacted carbon nanotube. Nature 401, 572
574. Copyright Nature Publishing Group. Reproduced with permission.

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scattering at the contact interfaces must not completely randomize the spin. There is also a large hysteresis in the peak position (B 6 50 mT) between positive and negative sweep directions, indicating the probable influence of the contact magnetization. Similar hysteretic MR is observed in MTJs, where it has been attributed to spin-polarized electron tunneling (Prinz, 1998; Julliere, 1975; Moodera et al., 1995, 1998; Miyazaki et al., 1995; Meservey et al., 1994; Shang et al., 1998). MTJs consist of two FM contacts separated by a thin oxide layer. The conduction electrons within the FM contacts have a preferred spin direction, which is determined by the local magnetization. This causes the formation of majority and minority spin conduction bands with different densities of states at the FE. In the absence of spin scattering the resistance across the tunnel barrier is dependent on the relative alignment of the magnetization of the two contacts. In the antiparallel state the majority spin states are out of alignment, and the junction resistance is higher than in the parallel state in which the majority spin states are aligned. For the nanotube devices in Fig. 4
6 the contact magnetizations align parallel with the magnetic field at B 5 6 100 mT. As we sweep B through 0 T the magnetization polarity switches. The observed peak in the nanotube resistance suggests that the contact magnetizations switch separately and become misaligned as the field is swept. For a MTJ, misalignment occurs because different FM contact materials are used, with different coercivities—the magnetizations are misaligned when B lies between the coercive fields of the two contacts. But this does not explain the misalignment in the nanotube device, because the average coercivity of the two Co contacts should be very similar. In this case the misalignment may be caused by magnetization fluctuations that occur locally, on the scale of the nanotube diameter (30 nm). The average Co domain size, 50 nm (Rüdiger et al., 1999), is considered in the order of the width of the nanotube, so that the nanotube contacts only a small number of magnetic domains. The coercivity of each domain varies and depends on its geometry and the local energy conditions. Edge and surface effects are also important. It is reasonable, then, that there will be a range of B over which the magnetization at the two nanotube contacts will be misaligned. A resistance peak due to the misaligned state occurs, even though the average properties of the two contacts are similar. The small switches in the resistance, seen most clearly in Fig. 4
6, provide additional evidence for local magnetization fluctuations of individual domains. Tsukagoshi et al. (1999) include data in the figure taken from three different samples to give a clear indication of large sample-to-sample variation, typically observed in the magnetoresistance. In all CNTs, Co-contacted CNTs displayed hysteretic MR, with a peak height varying from 2% to 10%. Two CNTs showed a step-like resistance peak as seen in Fig. 4
6A, while the remaining samples showed a smoother peak as seen in Fig. 4
6B and C. It is likely that the sample-to-sample variations are due to inherent random variations in the surface potential over the small nanotube contact area. Previous experiments on nonmagnetically contacted nanotubes have observed large variations in the contact resistance (Langer et al., 1996; Bachtold et al., 1998). In addition, in the ferromagnetically contacted CNTs, it is impossible to control the particular domain structure in contact with the nanotube. Reproducibility might be improved by increasing the thickness and quality of the Co layers, which will increase the area of the magnetic domains. Unfortunately, our maximum film thickness (and hence the domain width) is limited to B100 nm by the

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electron-beam lithographic procedure. The spin-injection model for the nanotube MR requires a sufficiently small amount of spin scattering to occur both within the nanotube, and at the interfaces between the nanotube and the contacts. We can estimate the minimum spin-scattering length in the MWCNT using Julliere’s model for the MTJ. The difference between the tunnel resistance in the parallel (RP) and antiparallel (RAP) states (Julliere, 1975) is given by the following equation: ΔR RAP 2 RP 2P1 P2 5 5 RAP RAP 1 1 P1 P2

where P1 and P2 are the percentage of conduction electrons polarized in the majority spin band in the FM contacts 1 and 2. For Co the polarization has been determined (Meservey et al., 1994) to be 34% giving a maximum resistance change of 21%. In the present best case, ΔR/RAP reaches a maximum value of 9% (Fig. 4
6B), so that B14% of the spin-polarized electrons travel 250 nm through the nanotube without spin-flipping. The spin-scattering length, ls, can then be estimated by assuming that the spin polarization reduces as exp(2l/ls) within the nanotube. This gives ls 5 130 nm. Although fairly long, this is probably an underestimation. The spin-polarization near the ferromagnet/nanotube interface will depend on the interface quality and could be appreciably lower than 34%. Also, we do not take into account spin scattering at the ferromagnet/nanotube interface. Tsukagoshi et al. (1999) also studied the MR of a nanotube contacted with a double Co layer as a function of BO and observed that the technique improves the continuity of the Co film and reduces the contact resistance in comparison with the single Co-layer devices. However, the MR ratio, ΔR/RAP, and the coercive field are less than the single Co-layer devices, as shown in Fig. 4
6, which implies that the magnetization is averaged over many small magnetic domains in the double Co-layer devices and each has relatively weak coercivities. The MR is measured as a function of B perpendicular to the substrate (B\) for the CNTs. In this case the peaks are broader and shifted to higher fields when compared with the BO dependence. This is expected for a FM film with an in-plane easy axis of magnetization (Moodera et al., 1995). The temperaturedependent MR shows that the percentage difference between the resistance in the parallel and antiparallel configurations goes to zero as the temperature increases from 4.2K to 20K. ΔR/RAP decreases almost exponentially with the increase in the temperature. The exact mechanism for the temperature dependence is not yet known. Low atomic number of carbon, the spin
orbit scattering in the CNT (and hence its temperature dependence) should be negligible. However, the spin polarization at the interface will decrease with the increase in temperature if the nanotube/ferromagnet interface is of relatively poor quality.

4.5 Tunnel magnetoresistance in carbon nanotube
based spintronic devices CNTs are molecular tubes of carbon with outstanding properties (Min et al., 2006). They are among the stiffest and strongest materials known and have remarkable electronic behavior

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and many other unique properties. They are attractive for spintronic devices due to their nanoscale size, extremely large spin-flip scattering lengths, and because they can behave as one-dimensional ballistic quantum conductors (Glazman et al., 1988; Monsma et al., 2000; Kamper et al., 1987; Hofer et al., 2003). Experimental investigations on coherent spin transport through Co-contacted CNTs showed that spin could be coherently transported over 130 nm through the CNT (Tsukagoshi et al., 1999). Sahoo et al. (2005a,b) studied the gate dependence TMR of MWCNTs with PdNi electrodes, at T 5 1.85K). Fig. 4
7A shows the TMR observed to oscillate relatively regularly between 25% and 16% on a gate-voltage scale ΔVgTMR such that 0:4 V , ΔVgTMR , 0:75 V. The conductance of the MWCNTs has been studied at lower temperatures (T 5 300 mK), in order to resolve the single-electron states which could not be resolved at the temperature at which the MR was measured. A measurement of the differential conductance dV/dI as a function of source
drain Vsd and gate voltage Vg at T 5 300 mK is also studied for a relatively narrow Vg range. It displays the diamond-like pattern characteristic of single-electron tunneling in a quantum dot. The diamonds vary in size with single-electron addition energies ranging between 0.5 and 0.75 meV, in agreement with previous reports on MWCNT quantum dots with nonFM leads (Buitelaar et al., 2002), where coulomb blockade and energy level quantization were observed. The electron levels are nearly fourfold degenerate (including spin) and their evolution in magnetic field (Zeeman splitting) agrees with a g factor of 2. However, the TMR gate-voltage scale ΔVgTMR measured at T 5 1.85K is much larger than the scale Vge B25 mV for addition of single electrons: it corresponds to the addition of at least 16 electrons rather than 1. A gate-voltage scale that agrees with the TMR signal becomes visible if the linear conductance G at low temperatures is monitored over a wider gate-voltage range. The single-electron conductance peaks are strongly modulated in amplitude, leading to a regular beating pattern with the proper gate-voltage scale of ΔVg  0.4 V. Fig. 4
7B shows the conductance G and the TMR of a SWCNT device measured by Sahoo et al. (2005a,b). The quantum dot behavior is already observed at 1.85K, whereas this was only evident at 0.3K in the MWCNT device. This is consistent with the higher energy scales (both single-electron charging energy and level spacing) for SWCNTs as compared with MWCNTs. The typical single-electron addition energy amounts to B5 meV, whereas it was an order of magnitude smaller in the MWCNT device. The variation of the conductance G and the TMR are simultaneously shown in Fig. 4
6B. They attributed those behaviors to the quantum interference in CNT (Schäpers et al., 2001). First, the TMR changes sign on each conductance resonance. Furthermore, the line shape of the conductance resonances is symmetric, whereas that of the TMR dips is asymmetric. The jump in the G(Vg) data at Vg 5 4.325 V. The amplitude of the TMR ranges from 27% to 117%, which is a higher amplitude than for the MWCNTs. This might be due to the higher charging energy in SWCNTs (Barnas et al., 2000). Normal
SWCNTs
FM (N
SWCNTs
F) devices yield an order of magnitude lower signal proving that the current in the F
tube
F devices is indeed spin polarized. Gunnarsson et al. (2008) investigated the gate-voltage dependence of nonlocal spin signal. Nonlocal spin-valve devices of SWCNT were fabricated, and these devices showed clear quantum dot behavior at low temperatures. They found the nonlocal voltage oscillates around zero with a large amplitude about 1 μV. Similar phenomenon was also reported by Makarovski et al. (2007).

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FIGURE 4–7 (A) The TMR data of MWCNTs connected with two PdNi contacts were measured at T 5 1.85K. The MR oscillates with a period ΔVgTMR B0.4
0.75 V. (B) SWCNT device with a contact separation of L 5 500 nm measured at T 5 1.85K. Measurement (’) of the linear conductance G and the TMR around two resonances. The bars in reflect the error in deducing the TMR signal from R(B) curves. MWCNTs, Multiwalled carbon nanotubes; MR, magnetoresistance; SWCNT, single-walled carbon nanotubes; TMR, tunnel magnetoresistance. Sahoo, S., et al., 2005a. Electric field control of spin transport. Nat. Phys. 1, 99
102; Sahoo, S., et al., 2005b. Electrical spin injection in multiwall carbon nanotubes with transparent ferromagnetic contacts. Appl. Phys. Lett. 86, 112109. Copyright Nature Publishing Group. Reproduced with permission.

Mohamed et al. (2008) studied the MR of SWCNTs from FM electrodes. In Mohamed et al. (2008), the coercive force obtained from the hysteresis curve for both temperatures is B120 Oe. The current I versus voltage V characteristics for the devices was measured and found that resistance of the SWCNT devices is about 10
200 Ω at room temperature. Highresistance CNTs show good FET characteristics (Inami, 2008). These facts also suggest that the channels consist of metallic or semiconducting SWCNTs, but not other conductive carbon materials. In Mohamed et al. (2008), the resistance B186 Ω is almost constant against the magnetic field for a width channel of B500 μm at 300K. However, the resistance peak is B0 Oe at 4.5K, which varies slightly in the direction of the sweep. When the field was swept upward, a peak appeared at about 2110 Oe, whereas in the downward direction, a peak appeared at about 110 Oe. In the case of a device with a channel width of B250 μm, similar behavior to that in the device with the channel width of B500 μm was observed. To ensure

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that this MR effect is not governed by the MR of Co nanoparticles on the electrodes, a control experiment was conducted. Mohamed et al. (2008) measured the MR effect for one of the electrodes at the same temperatures, and no significant change in resistance was observed. This indicates that the hysteretic MR does not originate from the Co nanoparticles themselves. Although the Co FM component is independent of temperature, as mentioned above, MR effects were observed at a low temperature. For spin transport without spin scattering through a material connected to two FM electrodes, the resistance is high when the FM moments in the two electrodes are antiparallel and vice versa (Julliere, 1975). The spin-valve effects observed in MWCNTs and SWCNTs are well explained by this model (Sahoo et al., 2005a,b; Nagabhirava et al., 2006). SWCNTs seem to grow from Co nanoparticles, which form from Co thin film on Mo film while using alcohol catalytic chemical vapor deposition (ACCVD) (Inami et al., 2007). Therefore spin-dependent transport is expected to be governed by the magnetic properties of Co nanoparticles to which the SWCNTs are attached and to be strongly dependent on the size of the Co nanoparticles. Each SWCNT is attached to a Co nanoparticle with a different size. Thus it is expected that each SWCNT should show MR peaks at different field values where parallel and antiparallel magnetization configurations of the electrodes change. However, the magnetic property of an individual Co nanoparticle in contact with an SWCNT cannot be determined. In this study the average magnetic properties and MR effects of Co nanoparticles were observed. The sharp peaks of averaged MR at approximately 6 110 Oe correspond to the average coercive force measured by the SQUID. The origin of this unusual MR hysteresis still remains unclear, but the results of the MR effects are reproducible. A detailed analysis of the spin transport mechanism of these devices is to be carried out in the near future. The ratio of MR is defined by ΔR/R0, where R0 is the resistance in the saturation region. At 4.5K, ΔR/R0 is found to be about 0.7%
1.8%. This value agrees with the previous report on the MR of SWCNTs (Sagnes et al., 2003). To increase ΔR/R0, effective spin injection from the FM electrode to the SWCNTs, and spin-coherent transport in SWCNTs should be realized. For spin injection, improvement in the Co/SWCNT interface quality is necessary. For spin-coherent transport the growth of high-quality SWCNTs and the reduction of L between the electrodes are the most important factors. It is also interesting that the dependence of ΔR/ R0 on L can be used to clarify the spin diffusion length of SWCNTs.

4.5.1 Spin-valve devices of carbon nanotubes Tsukagoshi et al. (1999) fabricated the first two-terminal CNT spin valve device, and spindependent transport was demonstrated through MWCNTs and 9% MR at 4.2K was observed. The MWCNTs were synthesized from graphite rods by the arc discharge. Co electrodes 65nm thick deposited on the nanotube were used to inject spin current into MWCNTs. Later, spin-dependent transport properties in SWCNT or SWCNT network were also been studied in local or nonlocal geometry (Kim et al., 2002; Tombros et al., 2006; Yang et al., 2012; Jensen et al., 2005). Jensen et al. (2005) measured several kinds of two-terminal devices with different electrodes at low temperatures. They observed a large MR up to almost 100% and down to 2150% in the SWCNT devices with two Fe electrodes. However, a 10% MR was also

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observed in the device that only had one magnetic electrode. This phenomenon is quite confusing and need further exploring. Tombros et al. (2006) carried out the nonlocal spin transport measurements in SWCNTs. By separating the charge current from the spin current, they provide an ambitious prove of the spin accumulation induced MR in SWCNT.

4.6 Conclusion and perspectives of carbon nanotubes in spintronics Spin currents have been successfully injected to both SWCNTs and MWCNTs by using different FM electrodes, such as Co, Fe, PdNi, and lanthanum strontium manganite (LSMO). CNT-based spin valve devices show large MR up to 61% and a long spin diffusion length about 50 m at low temperatures, which make CNT a promising material for the spintronic devices. Several spin relaxation mechanisms were discussed, but systematic spin relaxation mechanisms need to be further investigated. The gate tunable properties were studied in local and nonlocal spin valves, indicating that the electric field can be used to control the spin transport in CNTs. By adding spin freedom to the traditional electronics, spintronic devices are expected to combine logic operations, data storage, and communications, making the devices faster and consume less electrical power than the conventional electronic devices. The carbon-based materials are outstanding candidates for this target and have already shown promising future from the recent experimental progress. Spin injection has been realized in graphene, CNT, fullerene, and organic semiconductors. Spin can transport a macroscopic distance in graphene and CNT. The spin valves fabricated from fullerene and organic semiconductors exhibit large MR at room temperature. It has been found in graphene that if a pinhole free and flatness SDT barrier could be obtained, the contact-induced spin relaxation can be reduced effectively. Complicated mechanisms were found in the spin precession in graphene and CNTs, and it needs to be systematically identified. Among these carbon materials, the two-dimensional graphene has drawn lots of attention and been developed quickly. Unlike CNTs, graphene can be tailored into particular shapes and the electronic and magnetic properties are sensitive to the edge structure of graphene. Theory has predicted that the zigzag-edged GNRs are very useful in the future spintronics and many novel spintronic devices based on them have been proposed. However, till now, there are a less number of experimental studies for these proposed devices, which should be the next focus in this field.

References Alphenaar, B.W., et al., 2001. Spin transport in nanotubes. J. Appl. Phys. 89, 6863
6867. Appelbaum, I., et al., 2007. Electronic measurement and control of spin transport in silicon. Nature 447, 295
298. Arras, R., et al., 2010. Interface states in the full-oxide Fe3O4-MgO-Fe3O4 magnetic tunnel junction. IEEE. Trans. Magn. 46, 1730
1732.

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420.

Magnetism and spintronics in graphene

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5.1 Introduction Electronics and spintronics are two fields of technology, which are very strongly coupled. This is due to the fact that both electronic and spintronic devices use the same elementary particles, electrons, for their operation. However, each field uses a different fundamental property of the particle; in electronics, it is the charge, and in spintronics, the angular momentum, also better known as spin. For the latter, placing the particle in a magnetic field results in a coupling of its magnetic moment (generated by its spin) to the magnetic field. If now the spin is measured, we obtain only two possible states, the spin-up state and the spin-down state. This last spintronic property can be used to perform Boolean logic operations in a similar way as is already done nowadays in computer chips. For this type of logical operations, two states are needed, a 0 and a 1 state. This can be easily found in a spintronic device just by assigning the spin-down state to the 0 and the spin-up state to the 1 state. However, to be able to create a computer chip containing only spintronic devices, it is necessary to build fundamental spintronic devices in which the spin state can be manipulated. A variety of magnetic materials have been discovered and devices utilizing them to industrial systems in spintronic applications. In all cases, rare magnetic elements are indispensable to provide polarized electron spins, which yield magnetism. It has been known for more than last two decades that the carbon-based spx orbital systems can lead to spin polarization based on edge-localized electrons (Nakada et al., 1996; Fujita et al., 1996; Kusakabe et al., 2003; Okada et al., 2001; Lee et al., 2004, 2005; Veiga et al., 2008). In particular, a specified atomic structure, the so-called zigzag type, at graphene edges has attracted much attention for spintronics applications. Graphene has proved to be an attractive material for spintronics (Tombros et al., 2007, 2008; Wang et al., 2008a,b,c,d; Józsa et al., 2008, 2009a,b; Popinciuc et al., 2009; Han et al., 2010, 2011, 2012; Maassen et al., 2011, 2012; Guimarães et al., 2012; Vera-Marun et al., 2012; Dlubak et al., 2012; Liu et al., 2013; Hill et al., 2006; Ohishi et al., 2007; Shiraishi et al., 2009; Nishioka et al., 2007; Cho et al., 2007; Goto et al., 2008; Han et al., 2009a,b,c; Pi et al., 2010). Graphene has a low spin
orbit (SO) interaction (SOI), which in principle should translate into a long spin lifetime. Together with the high charge carrier mobility (Bolotin et al., 2008), it implies a long distance over which the spin information can be transported. Other aspects that make graphene a unique system for spintronics include its tunable carrier concentration. The lack of surface depletion region enables modification by surface interaction with metal or chemical doping Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials. DOI: https://doi.org/10.1016/B978-0-12-817680-1.00005-6 © 2020 Elsevier Inc. All rights reserved.

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(Pi et al., 2010; Huard et al., 2008; Pi et al., 2009; Schedin et al., 2007; Farmer et al., 2009) and prediction of novel spin-dependent behavior such as fully spin-polarized magnetic ordering in nanoribbons (Son et al., 2006a,b). The graphene nanoribbon (GNR) is a onedimensional strip line of graphene with two edges on both sides along the longitudinal direction. It is assumed that the perfect edge of atomic structures without any defects allows the electron spins localized at the zigzag edges (Nakada et al., 1996; Niimi et al., 2006) to be stabilized toward the polarization [i.e., (anti)ferromagnetism] as a result of the exchange interaction between the two edges, which yields a maximum spin ordering in these orbitals, in a GNR (Nakada et al., 1996; Fujita et al., 1996; Kusakabe et al., 2003; Okada et al., 2001; Lee et al., 2004, 2005; Veiga et al., 2008) graphene nanomeshes (GNMs) with hexagonal nanopore arrays (Shima et al., 1993; Yang, 2011), and graphene nanoflakes (Rosser et al., 2007). It is similar to the case of Hund’s rule for atoms. Moreover, the spin configuration is highly sensitive to the type and number of foreign atoms [e.g., hydrogen (H) and oxygen (O)] that terminate the edge carbon dangling bonds (Kusakabe et al., 2003; Soriano et al., 2011). Again, according to the Lieb’s theorem, the presence of low-concentration defects in ensembles of carbon atoms (e.g., graphene flakes) results in the appearance of a net magnetism. It predicts the emergence of ferromagnetism by an increase in the difference between the numbers of removed A and B sites (ΔAB) of the graphene bipartite lattice at the zigzag edges (Rosser et al., 2007; Yang, 2011; Soriano et al., 2011). The magnitude of ferromagnetism increases with increasing values of ΔAB. These two models provide a very interesting spin configuration and predict the emergence of the spin polarization in allcarbon materials. However, the spintronics is a highly promising issue as a key technology for the next generation of devices (Moodera et al., 1995; Yuasa et al., 2004; Hayakawa et al., 2006; Munekata et al., 1989; Ohno et al., 1992; Hai et al., 2009; Tada et al., 2012, 2013; Hashimoto et al., 2014). Some kinds of structures have been developed for spintronics (Moodera et al., 1995; Yuasa et al., 2004; Hayakawa et al., 2006; Munekata et al., 1989; Ohno et al., 1992; Hai et al., 2009; Tada et al., 2012, 2013; Hashimoto et al., 2014), such as giant magnetoresistance (MR) (GMR) (Moodera et al., 1995), tunneling MR (TMR) (Yuasa et al., 2004; Hayakawa et al., 2006), and spin valve structures. In particular, the TMR structure has realized high efficiency of the TMR ratio, ΔR/R0, defined as (RAP 2 RP)/RP, where AP and P refer to antiparallel and parallel orientations of the spin configurations (magnetizations) of the two electrodes. Therefore polarized spins in carbon materials as mentioned before become highly important and desirable. It has been predicted that the spin-based phenomena will be realizable using graphene edges. The spin-filtering (rectifying) effect predicted that GNRs with antiferromagnetic spin alignment on two edges can manipulate only electron spins with the same moment by applying in-plane electric fields (Son et al., 2006a,b). Realization of (quantum) spin Hall effect (SHE) was also predicted by resolving the double degeneracy of edge spin bands (e.g., by introducing SOI and controlling two spins with opposite moments existing in two different bands by applying electric fields (Kane et al., 2005; Kane, 2007; Balakrishnan et al., 2013). Indeed, the observation of large spin diffusion current in high-quality graphenes fabricated on hexagonal boron nitride (BN) (h-BN) substrate (Abanin et al., 2011) and also SHE in

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hydrogenated graphenes by introducing SOI (Baibich et al., 1988) have been experimentally confirmed. They are opening the door to novel all-carbon spintronics without using rare metals. The large spin coherence of graphene must lead to high-efficiency spintronic devices. In this chapter, we will discuss about the magnetism and spin-based phenomena of graphene materials and their probable spintronics applications: • Covers the fundamental properties and various proposals for applications of graphene in spintronics, as well as an overview of charge and spin transport in graphene • Includes a description of physical phenomena in both graphene and spintronics • Provides illustrations and many references for easy understanding of the present status of graphene applications in spintronics

5.2 Making of graphene into magnetic materials Among the many applications predicted for graphene, its use as a source of magnetism is the most unexpected one and an attractive challenge for its active role in spintronic devices (Han, 2014). Generally, magnetism is associated to a large degree of electron localization and strong SOI. Both premises are absent in graphene, a strongly diamagnetic material. The simplest method to induce magnetism in graphene is to create an imbalance in the number of carbon atoms in each of the two sublattices. Turning graphene magnetic is a promising challenge to make it an active material for spintronics. Predictions state that graphene structures with specific shapes can spontaneously develop magnetism driven by Coulomb repulsion of π-electrons, but its experimental verification is demanding. The observation and manipulation of individual magnetic moments in graphene open-shell nanostructures could be done in different process. The presence of single electron spins localized around certain zigzag sites of the carbon backbone via the Kondo effect. Nearby spins coupled into a singlet ground state (GS) and quantify their exchange interaction via singlettriplet inelastic electron excitations. Theoretical simulations picture also shows how electron correlations result in spin-polarized radical states with the experimentally observed spatial distributions. Extra hydrogen atoms bound to radical sites quench their magnetic moment and switch the spin of the nanostructure in half-integer amounts. The zigzag atomic structure of graphene edges also leads to spontaneous spin polarization (Shimizu et al., 2012; Castro Neto et al., 2009a,b). Not many works, however, have reported on the experimental observation of magnetism and spin-based phenomena derived from graphene zigzag edges. One reason is that edge-related phenomena are easily destroyed by disorder (damage, defects) and by contamination introduced during the fabrication process (e.g., by lithographic methods). We have, therefore, developed two original nonlithographic fabrication methods of graphene edges: (1) GNRs derived by unzipping carbon nanotubes combined with air blowing and three-step annealing and (2) GNMs fabricated by etching graphene using nanoporous alumina template as masks (Shimizu et al., 2012; Castro Neto et al., 2009a,b).

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In this section, observation of magnetism and spin-based phenomena are reviewed on the basis of the latter method.

5.2.1 Ferromagnetism derived from hydrogenated zigzag-type pore edges graphene Bulk graphenes without any pores have no such features even after H2 annealing, implying no contribution of parasitic factors (e.g., defects, impurities) of bulk graphenes. The presence of less damage/impurities confirms the most of bulk graphene regions, because mechanically exfoliated bulk graphenes show extremely low sp2/sp3 content ({0.1). These results strongly suggest that the observed ferromagnetism is associated with polarized spins localized at the H-terminated zigzag pore edges. It is surprising that the ferromagnetism observed at 2K appears even at room temperature (RT) with a larger magnitude of the hysteresis loops. The correlation between the interpore spacing (corresponding to the width of the GNR) and the magnetization is that the magnitude of the hysteresis loop decreases with increasing the interpore spacing. In particular, it is revealed that the residual magnetization is inversely proportional to the interpore spacing. This result is qualitatively consistent with theories for the GNR model according to which the edge spin stability and ordering of a zigzag-edged GNR are determined by the exchange interaction between the two edges, leading to vanishing of the ferromagnetic (FM) edge spin ordering with increase of the interpore spacing (Fujita et al., 1996; Lee et al., 2005). Such a correlation cannot be understood by the ferromagnetism originating from the defects located only at pore edges or in the bulk graphene between pores. Again, the ferromagnetism amplitude should be mostly independent of the interpore spacing, because defects locate only at the pore edges. In the latter case, ferromagnetism amplitude should increase with an increase of the interpore spacing because the density of the defects increases. Consequently, it is conclude that the observed ferromagnetism is not due to parasitic origins (e.g., defects, impurities) but can be truly attributed to H-terminated zigzag pore edges.

5.2.2 Magnetism depending on pore edge termination by different foreign atoms Oxygen: It has been reported (Veiga et al., 2008) that the formation of a spin-paired C
O chemical bond drastically reduces the local atomic magnetic moment of carbon at the zigzag edge of GNRs and suppresses the emergence of ferromagnetism. O-terminated GNMs exhibit a diamagnetism-like weak hysteresis loop. This is consistent with the work of Veiga et al. (2008), which reported that the formation of a spin-paired C
O chemical bond drastically reduces the local atomic magnetic moment of carbon at the zigzag edge of GNRs and suppresses the emergence of ferromagnetism. Moreover, the diamagnetism of graphene has mostly disappeared due to the formation of the nanomesh, because such an array drastically reduces the bulk graphene area available for the presence of loop currents to produce diamagnetism at the currently applied magnetic field range.

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5.2.3 Doping and/or functionalization with transition metals Pristine graphene is nonmagnetic. GNR presents spin polarization on the two edges. Based on spin polarization control on the edges by applying an electric field (Son et al., 2006a,b) or a magnetic field (Yazyev et al., 2008), the nanoribbons can be used for spintronics applications. To induce magnetism in a nonmagnetic material like the graphene, a method of doping with transition metals (elements located in d-block of the periodic table with partially filled d sub-shell) is used (Cramer et al., 2009). These are adsorbed as interstitial impurities or are embedded in vacancies (Lee et al., 2013). The electric and magnetic properties of the nanoribbon are affected by the direction of spin from the edges (FM or AFM).

5.3 Spin generation and spin manipulation 5.3.1 Spin generation Among the different key parts of spintronics, spin generation is the most important one and has different methods of spin generation in graphene. They are mainly (1) making magnetic graphene by the electric field, doping, and defect engineering; (2) tunneling injection through FM materials; and (3) thermal spin polarization by heat. The simplest method to generate spin-polarized current in graphene is making graphene itself magnetic. Once the graphene is made FM, the injected normal charge current via a bias voltage will be spin-polarized. Among the others, FM orderings are various defects on graphene structures, such as vacancies, topological defects, and hydrogen chemisorption defects (Wang et al., 2008a,b,c,d). One of the exciting results is the experimental observation of ferromagnetism in the point defects of graphene at RT using bulk, solution-processable, functionalized graphene materials (Kim et al., 2008). It promises real applications of graphene-based spintronic devices. However, the number of vacancy defects in solutionprocessable graphene oxide can hardly be controlled precisely. Creating radiation damage with high-energy particles (protons) is another way to produce several types of point defects in graphite (Esquinazi et al., 2003). The number of defects due to the radiation damage process can be well controlled compared to the solution-process method. During the synthesis of graphene, reactive particles may produce chemisorption defects, such as hydrogen, resulting in an sp3 state. It has been theoretically predicted that hydrogen chemisorption defects, either partially hydrogenated (Yazyev et al., 2007) or semihydrogenated (Zhou et al., 2009a,b), show strong spin polarization. RT ferromagnetism in partially hydrogenated graphene has been observed experimentally recently (Xie et al., 2011). Direct external magnetic source is a magnetic field that can make injected normal current spin polarization or change the magnetic direction of spin-polarized current. Actually, this method is widely used in the modern hard disk drives and memory by utilizing the GMR effect, which makes the large-scale and industrial applications of spintronics. The spintronic device is a spin valve, in which a spin-polarized current injected from one FM electrode goes through the graphene before being detected by the other electrode.

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This measurement is rapidly followed by several other spin-transport measurements, including two-terminal “local” and four-terminal “nonlocal” measurements (Tombros et al., 2007; Popinciuc et al., 2009; Goto et al., 2008). The Seebeck effect describes that electric currents can be induced by a temperature difference. Following Seebeck effect, spin-up and spin-down currents flowing in opposite directions can be generated by a temperature gradient and can be converted in to a spin voltage via an inverse SHE (Uchida et al., 2010; Bauer et al., 2010). These findings indicate that spintronic components may be operated with thermally induced spin currents in the absence of external bias voltage and represent a milestone for spin caloritronics, a new exciting research field taking advantage of both spintronics and thermos electronics (Bauer et al., 2010, 2012). Recently, Zeng et al. (2011a,b) investigated graphene-based spin caloritronics and found the spin Seebeck effect in magnetized Zigzag graphene nano ribbons (ZGNRs) magnetic Zigzag graphene nano ribbons (M-ZGNRs), providing a new method to create spin currents in the graphene system.

5.3.2 Spin manipulation To design graphene-based spintronic devices such as spin field effect transistor (FETs), the effective control of spin currents needs to be resolved after spin injection into graphene. For example, the negligible SO coupling (SOC) in graphene is a strong limitation for designing the Datta-type spin FETs (Datta et al., 1990). Besides, the absence of a bandgap and the difficulty in p-type or n-type doping make graphene-based spin diodes difficult to realize. Recently, Zeng et al. (2011a,b) designed a ZGNR-based spin diode by utilizing a spindependent transmission selection rule, which originates from the wave function symmetry of ZGNRs. This design can eliminate the requirement of SOC for spin control. Moreover, the proposed spin diode allows a bottom-up design of full range of spintronic devices, including spin current or voltage amplifiers, spin logic gates, and complex calculators. A ZGNR at its GS has edge states with FM ordering at one edge and AFM coupling between two edges (Son et al., 2006a,b). Applying an external magnetic field can force the edge states of ZGNRs to switch to FM coupling (Kim et al., 2008). Moreover, spin stiffness and spin correlation length at one-dimensional edge are limited. As a result, environmental perturbations, such as high temperature, may destroy spin ordering of the strongly localized edge states (Yazyev et al., 2007). Therefore three magnetic states of ZGNRs can be exploited for designing spintronic devices. These are AFM GSs, FM metastable states, and nonmagnetic states. The band structure of GS-ZGNRs shows a direct bandgap and energy splitting at the X point. FM ZGNR is metallic, with two spin subbands crossing each other at the Fermi level. NM-ZGNRs are semimetallic and with degenerate spin subbands. Moreover, N-ZGNRs (N is the number of zigzag chains) show a transmission selection rule that is related to the wave function symmetry when N is even. This transmission selection rule is unaffected by the magnetic state of ZGNRs (Kim et al., 2008). The flexible control over spin current makes it possible to use the ZGNR spin diode to build multifunctional spintronic devices, such as spin transistors. The ZGNR-based spin diode can also be exploited as a building block for spin logic gates that perform logic operations with spin as the operation variable.

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5.4 Magnetism in graphene The long-standing interest in magnetic behavior of pure carbon-based systems has been reported at RT magnetic ordering in highly oriented pyrolytic graphite (HOPG) (Esquinazi et al., 2002, 2003), nanographites (Enoki et al., 2009), nanodiamonds (Talapatra et al., 2005), and disordered carbon films (Rode et al., 2004; Ohldag et al., 2007). Although in these studies magnetization signals M were small (typically, less than B0.1 emu/g, that is, less than 0.1% of the magnetization of iron), a consensus is emerging that, despite the absence of d or f electrons, magnetism in carbon systems may exist under a variety of experimental conditions. Theoretically it was found that atomic scale defects in graphene-based materials, for example, adatoms and vacancies, can carry a magnetic moment μ of about one Bohr magneton, μB (Krasheninnikov et al., 2009; Yazyev, 2008; Lopez-Sancho et al., 2009; Faccio et al., 2008). Also, extended defects such as edges can give rise to M (Harigaya et al., 2002; Fujita et al., 1996; Kobayashi et al., 2006). The possibility of long-range magnetic ordering has been predicted for randomly distributed point defects and grain boundaries (Yazyev, 2008; Faccio et al., 2008), and bilayer graphene (BLG) was suggested to exhibit spontaneous many-body ferromagnetism (Castro et al., 2008). All this leaves little doubt that magnetism in graphenebased systems can in principle exist, although the whole subject remains highly controversial, especially as concerns (1) the role of environment and magnetic contamination (Sato et al., 2003) and (2) the mechanism that could lead to the strong interaction required for ferromagnetism at RT. Graphene has inevitably led to the question of possible ferromagnetism in this novel material too, especially due to the fact that it presents the basic structural element for all other graphitic forms (Geim et al., 2007). The first experiments reported RT ferromagnetism in bulk samples obtained by conversion of nanodiamond and arc evaporation of graphite (Ramakrishna Matte et al., 2009) and in graphene oxide (Wang et al., 2009). In both studies, magnetic signals were again small (saturation magnetization MSB0:1
1 emu/g) and have left open the same questions that haunt the previous reports of RT ferromagnetism in carbon materials. Sepioni et al. (2010) studied the magnetism of exfoliated graphene. The magnetization M as a function of H and T shows low-temperature paramagnetism behaviors as shown in Fig. 5
1. Below 20K, the magnetization response in parallel H becomes positive (Fig. 5
1A). As T is lowered further, a typical paramagnetic behavior emerges, with low-field susceptibility χ 5 M/H following the Curie law χ~1/T (Fig. 5
1B). In perpendicular H, magnetization was dominated by diamagnetism. Nevertheless, after subtracting the linear background, ΔM (H, T) curves showed exactly the same paramagnetic contribution as in parallel H (Fig. 5
1B); that is, the paramagnetism is isotropic. To characterize the magnetic species contributing to the observed behavior, we plot M as a function of the reduced field H/T. However, the origin of the detected moments, unlike RT ferromagnetism, intrinsic paramagnetism with J 5 1/2 would agree with the existing theories because vacancies, adatoms, and edges can carry localized moments (Krasheninnikov et al., 2009; Yazyev, 2008; LopezSancho et al., 2009; Faccio et al., 2008; Harigaya et al., 2002; Fujita et al., 1996; Kobayashi et al., 2006). Magnetic moments in graphene can be associated not only with point defects

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FIGURE 5–1 (A) Magnetic moment M as a function of parallel H at different T. (B) M (T) in parallel H. Inset: Excess moment ΔM after subtracting the diamagnetic background. Reprinted with permission from Sepioni, M., et al., 2010. Limits on intrinsic magnetism in graphene. Phys. Rev. Lett. 105, 207205. Copyright American Physical Society.

but also with extended ones such as zigzag edges (Harigaya et al., 2002; Fujita et al., 1996; Kobayashi et al., 2006). In this case, the magnetic moment would depend on the total length of zigzag segments and, in principle, can be arbitrarily large. At first glance, this mechanism seems to lack an explanation for the value of MS being much smaller than the available broken bonds could generate. Ray et al. (2014) investigated the magnetic behavior of partially/fully hydrogenated (hydrogen plasma treated at different temperature) vertically aligned few-layers graphene (FLG) synthesized by microwave plasma
enhanced chemical vapor deposition In this report, the FLGs are hydrogenated at different substrate temperatures to alter the degree of hydrogenation. The unique morphology of the structure gives rise to a unique geometry in which graphane/graphone (one side “H” bonded graphene) is supported by graphene layers in the bulk, which is very different from other widely studied structures such as onedimensional nanoribbons. In this study the FM interactions seems to be predominant and the presence of antiferromagnetic interaction was presence. Free spins available via the conversion of sp2 to sp3-hybridized structures and the possibility of unpaired electrons from defects induced upon hydrogenation are thought to be likely mechanisms for the observed FM orders. Ray et al. (2014) measured the magnetic properties of the FLG and FLG:H within the range of 22 kOe , H , 2 kOe at temperatures of 300K and 40K, respectively. The measured magnetic

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FIGURE 5–2 (Left) Magnetic hysteresis loops obtained for FLG and FLG:H at 300K and 40K, respectively. (Right): Magnetic force microscopy images of pristine (A
C) and hydrogenated FLG [(D
F) @ 50 C and (G
I) @200 C]. FLG, Few-layers graphene. Ray et al., 2014. Graphene supported graphone/graphane bilayer nanostructure material for spintronics. Sci. Rep. 4, 3862. By courtesy of NPG Publication.

hysteresis loops are shown in Fig. 5
2(left), with the FLG:H@50 C showing the most expressed FM behavior with maximum field hysteretic features and highest saturation moments (Ms  13.94 3 1024 emu/g), while other FLG:H show more confined hysteretic features and lower saturation moment. As compared to pristine FLG, the magnetic moment values of FLG:H@100 C and FLG:H@200 C (Ms  6.10 3 1024 emu/g) are slightly higher due to hydrogen incorporation in the FLG (Ms  3.47 3 1024 emu/g) resulting in the formation of sp3-hybridized carbon structure through mono and possible di-hydrogen termination. These FLGs are free from any catalyst remnants and detectable foreign magnetic impurities, the observed magnetism in the FLG can be attributed to (1) defects and vacancies created during synthesis and (2) creation of sp3-hybridized structures (Matte et al., 2009; Yang et al., 2011a,b; Rout et al., 2011). The ID/IG ratio trend of pristine and hydrogenated FLG obtained from Raman spectra shows that the FLG:H@50 C have the highest defect ratio along with the highest content of hydrogen as measured from X-ray absorption near edge structure spectroscopy. The FLG:H@50 C was expected and indeed do show the highest magnetization signals due to the relatively lower temperatures of 50 C; the hydrogenation may possibly occur on only the top-most surface layer of the FLG, thereby favoring the higher observed magnetic moment (Bunch et al., 2008). Similar to the results reported by others, we observed maxima in the magnetization at lower temperatures especially for FLG:H@50 C (Ning et al., 2013). Based on different hydrogen attachment on graphene, Yazyev et al. (2007) predicted that the orthodimers and para-dimers are nonmagnetic, while single hydrogen attachment (monomer) to be magnetic (Xie et al., 2011). This may explain why FLG:H@50 C is more magnetic than

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FLG:H@100 C (200 C). The magnetism observed in FLG:H is attributed to an intrinsic mechanism beyond reasonable doubt; it is important to discuss the role of hydrogen in enhancing the magnetism in nanostructured carbon and how it can be promoted during the synthesis itself. It has been shown theoretically that hydrogenation is an efficient way in which to introduce and enhance magnetism in graphene sheets. The addition of hydrogen leads to the rupturing of the delocalized π bonding network of graphene, leaving the 2pz electrons in the unhydrogenated carbon atoms unpaired, and thereby extending the p
p interactions resulting in the long-range FM coupling with a putative higher Curie temperature and a more homogeneous magnetism (Zhou et al., 2009a,b, 2012). Similar to the functionalization strategies of other nanomaterials, the synthesis of hydrogenated graphene can be done via a wet chemistry route or by plasma-based processes. The wet chemistry approach includes solution-based Birch reduction of graphite oxide to yield graphane or by liquid phase hydrogenation/exfoliation of graphite (Eng et al., 2013; Pumera et al., 2013). The plasma functionalization route involves hydrogenation of sp2 carbon materials such as carbon nanotubes (CNTs), graphene, or graphene oxide in a hydrogen gas/plasma environment (Pumera et al., 2013). Arc-discharge of graphite in a hydrogen rich environment has also been shown as an effective method for the synthesis of graphane (Subrahmanyam et al., 2011). However, theoretical calculations have suggested that the formation of graphene via hydrogenation of graphene will not yield large graphitic domains, since uncorrelated H frustrated domains are expected to be formed during the early stages of hydrogenation reaction (Pumera et al., 2013). This will invariably lead to the shrinkage of the graphene sheet leading to extensive sheet corrugations, thus making the direct deposition of graphane more desirable (Pumera et al., 2013). Zhou et al. (2012) have proposed a physical method to fabricate a semihydrogenated graphene sheet. Their idea revolves around the use of graphane as a substrate to support the BN sheet, after which the BN sheet is fluorinated. As the binding of the F with N is highly unstable, the FBN configuration can be easily achieved. Due to the presence of unpaired electrons, the N atoms are quite reactive in nature and upon the application of pressure will pick up the H atoms from graphane. When the applied pressure is removed, the resultant structure is semihydrogenated in nature (Zhou et al., 2012). In our present case, we are depositing FLG from the gas phase in plasma; the direct deposition of hydrogenated graphene via plasma deposition similar to work reported by Wang et al. can be a feasible route for enhancing the magnetic properties of FLG during the synthesis itself (Wang et al., 2010). The process reported by Wang et al. involves the use of remote discharged 13.5 MHz radiofrequency plasma inside an ultrahigh-vacuum source (Wang et al., 2010). The precracking of the gaseous precursors to generate the reactive free radicals in gas phase allows for lower substrate temperatures and also limits the damage caused by energetic plasma ions during the growth of film. The growth process was carried out using a premixed 5% CH4 in H2, resulting in an excess of atomic hydrogen in the gas phase and the inevitable hydrogenation of graphene films with formation of graphane (Wang et al., 2010). In literature the role of hydrogen during the noncatalytic growth of FLG has been linked to the etching of amorphous carbon films, which may occur during the initial nucleation stages (Burgess et al., 2011; Yuan et al., 2009). Thus for the formation of magnetic graphene structures, careful tuning of plasma parameters such as gas conditions, plasma power, temperature, ion energy, and bias in the microwave plasma will be required. The FM order arises from the

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free spins available via the conversion of sp2 to sp3-hybridized structures and/or from the unpaired spin electron from the defects induced upon hydrogenation (Wang et al., 2009). Both these factors may in principle be responsible for producing fundamental magnetic species. The FM ordering of the spins is energetically preferable for the AA distribution in the graphene plane. Therefore it can be stated that the FM exchange of spins of the localized states in graphane is possible only among the H-vacancy defects located on the exchange neighboring carbon atoms (Lee et al., 2005; Berashevich et al., 2009). Defects in our BLG/trilayer graphene (TLG) break the translational symmetry of the lattice and create localized states at the Fermi energy to produce an effective self-doping, where charge is transferred from defects to the bulk. In the presence of local electron
electron interactions, these localized states become spin-polarized, leading to the formation of pseudo-local moments (Lehtinen et al., 2004). Most of the theoretical (Rout et al., 2011; Harigaya, 2001; Esquinazi et al., 2003) and experimental (Harigaya, 2001; Esquinazi et al., 2003) works find that the net spin is stable within a large conjugation system in unit structures of graphene at RT and their stability is due to the huge p-conjugation in these molecules. If indeed the long-range orderly magnetic coupling of these spins may arise via either intramolecular interaction in individual graphene sheets or intermolecular interaction between neighboring graphene sheets then stable ferromagnetism could arise (Rout et al., 2011). We agree that the RT ferromagnetism is an intrinsic property of graphene-based materials, and for direct and conclusive evidence, we have performed further magnetic force microscopy (MFM) analysis. Low moment magnetic probes with Co/Cr coating were used to detect magnetic domains in the pristine and hydrogenated FLG. Fig. 5
2 (right) shows topographic (height), amplitude, and phase signals imaged simultaneously for MFM to assess correlation of surface features, identify and eliminate possible artifacts, and to assess effects of magnetization. The magnetized Co/Cr-coated probe interacts with magnetic field gradients generated by magnetic domains within the prepared sample resulting in changes in the phase and amplitude of the oscillating cantilever. Therefore from the amplitude and phase images, the existence of magnetic domains in the samples should be evident. The MFM phase and amplitude images show very good correlation in magnetic domain positions. For all samples the magnetic domains appear as dark and bright-localized regions in phase and amplitude images, respectively. Images clearly show that the domains in the FLG:H are more localized than in the case of FLG. A simple scaling of MFM phase data suggests that pristine FLG has the weakest magnetization, whereas the FLG@50 C has the strongest magnetization effect, which is consistent with the M
H magnetization results, described before.

5.5 Spin Hall effect and edge-derived spin phenomena Theoretically, it has been known for more than two decades that the carbon-based spx orbital systems can lead to spin polarization based on edge-localized electrons (Nakada et al., 1996; Fujita et al., 1996; Kusakabe et al., 2003; Okada et al., 2001; Lee et al., 2004, 2005; Veiga et al., 2008). A specified atomic structure, the so-called zigzag type, at graphene edges has attracted much attention. One-dimensional GNRs are made of two edges on the both sides of GNR along the longitudinal direction. It assumes the perfect edge atomic structures

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without any defects and allows the electron spins localized at the zigzag edges (Nakada et al., 1996, Niimi et al., 2006) to be stabilized toward polarization [i.e., (anti)ferromagnetism] as a result of the exchange interaction between the two edges, which yields a maximum spin ordering in these orbitals, in a GNR (Nakada et al., 1996; Fujita et al., 1996; Kusakabe et al., 2003; Okada et al., 2001; Lee et al., 2004, 2005; Veiga et al., 2008), GNMs with hexagonal nanopore arrays (Shima et al., 1993; Yang et al., 2011a,b), and graphene nanoflakes (Rosser et al., 2007). Moreover, the spin configuration is highly sensitive to the type and number of foreign atoms [e.g., hydrogen (H) and oxygen (O)] that terminate the edge carbon dangling bonds (Kusakabe et al., 2003; Soriano et al., 2011). It is assumes according to Lieb’s theorem that the presence of low-concentration defects in ensembles of carbon atoms (e.g., graphene flakes), resulting in the appearance of a net magnetism. It predicts the emergence of ferromagnetism by an increase in the difference between the number of removed A and B sites (ΔAB) of the graphene bipartite lattice at the zigzag edges (Rosser et al., 2007, Yang et al., 2011a,b; Soriano et al., 2011). The magnitude of ferromagnetism increases with increasing values of ΔAB. Different kinds of structures have been developed for spintronics (Moodera et al., 1995; Yuasa et al., 2004; Hayakawa et al., 2006; Munekata et al., 1989; Ohno et al., 1992; Hai et al., 2009; Tada et al., 2012, 2013; Hashimoto et al., 2014), such as GMR (Moodera et al., 1995), TMR (Yuasa et al., 2004; Hayakawa et al., 2006), and spin valve structures. In particular the TMR structure has realized high efficiency of the TMR ratio, ΔR/R0, defined as (RAP 2 RP)/ RP, where AP and P refer to antiparallel and parallel orientations of the spin configurations (magnetizations) of the two electrodes. Even a TMR ratio as high as over 1000% has been obtained by using the CoFeB/MgO/CoFeB junction (Munekata et al., 1989). Moreover, a variety of materials have been developed for spintronic devices, for example, FM metals [e.g., cobalt (Co), iron (Fe), chromium (Cr), manganese (Mn)] (Moodera et al., 1995; Yuasa et al., 2004; Hayakawa et al., 2006; Munekata et al., 1989), FM semiconductors [(In, Mn), As], and so on (Ohno et al., 1992; Hai et al., 2009; Tada et al., 2012, 2013; Hashimoto et al., 2014). In contrast, in any case, rare metals are required for realizing spintronics. This is a critical problem with the current limitation of rare material resources. Therefore polarized spins in carbon materials as mentioned before become highly important and desirable. Some theoretical works have predicted spin-based phenomena realizable using graphene edges. For instance the spin-filtering (rectifying) effect predicted that GNRs with antiferromagnetic spin alignment on two edges can manipulate only electron spins with the same moment by applying in-plane electric fields (Son et al., 2006a,b). Realization of (quantum) SHE was also predicted by resolving the double degeneracy of edge spin bands (e.g., by introducing SOI) and controlling two spins with opposite moments existing in two different bands by applying electric fields (Kane et al., 2005; Kane, 2007; Balakrishnan et al., 2013). Indeed, the observation of large spin diffusion current in high-quality graphenes fabricated on h-BN substrate (Abanin et al., 2011) and also SHE in hydrogenated graphenes by introducing SOI (Baibich et al., 1988) have been experimentally confirmed recently. They are opening the door to novel all-carbon spintronics without using rare metals. The large spin coherence of graphene must lead to high-efficiency spintronic devices.

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5.6 Spin injection, manipulation, and detection One of the key challenges in graphene spintronics is to inject spin current efficiently into graphene. Due to the conductance mismatch problem between graphene and FM metal electrodes, the spin polarization of the injected charges becomes very small. In order to overcome this problem a spin-dependent tunneling barrier was proposed at the spin-injection interface. The spin-injection efficiency can be estimated from the MR measurements of lateral spin-valves (LSVs). The MR measurements can be performed using either a local configuration where the spin injection and detection paths are the same or a nonlocal configuration where injection and detection paths are different. The nonlocal configuration usually yields a higher signal-to-noise ratio (SNR), because the spin accumulation is detected as a spindependent voltage difference with respect to the FM reference electrode without the involvement of charge current, excluding the large nonspin-related background signals. The spin-valve signal generated depends on the spin-injection efficiency, which is strongly limited by the conductivity mismatch between FM metals and graphene (Rashba, 2000; Fert et al., 2001). The 2D-graphene could be a promising candidate due to its unusual spindependent physical properties. Graphene, a single or few layers of carbon atoms, exhibits very low SOC, ultrahigh mobility, and the gate tunable conductivities (Geim et al., 2007; Das Sarma et al., 2011; Castro Neto et al., 2009a,b). Han et al. (2012) using transparent contacts (Co/SLG) performed spin injection and tunneling contacts (Co/MgO/SLG). With tunneling contacts, the nonlocal MR was increased by a factor of B1000, and the spin-injection/detection efficiency was greatly enhanced from B1% (transparent contacts) to B30%. The typical nonlocal MR curves studied by Han et al. (2012) are shown in Fig. 5
3. The arrows in Fig. 5
3A indicate the magnetic directions of the four Co electrodes. Han et al. (2012) investigated the dependence of the spin transport and spin precession as a function of distance using a single-layer graphene (SLG) sheet contacted by several Co electrodes at various spacings [(A) L 5 1 μm, (B) L 5 2 μm, and (C) L 5 3 μm)] and measured at 300K with Vg 5 0 V. The nonlocal MR decreases from 100 to 2 mΩ as the spacing increases from 1 to 3 μm. The spin-injection efficiency is calculated to be B1%, based on the following equation: ΔRNL 5

1 P 2 λG 2L=λG e σG W

(5.1)

where P is the spin-injection/detection efficiency, and σG, W, and λG are the conductivity, width, and spin diffusion length of the SLG, respectively. This low spin-injection efficiency of 1% is expected due to the conductance mismatch between Co and SLG (Schmidt et al., 2000). Fig. 5
3D shows the dependence of the nonlocal MR on the spacing between two center electrodes. The nonlocal spin signal decreases as a function of spacing and the data is fitted using the preceding equation, and the best fitted results, P 5 0.013, was obtained at λs 5 1.6 μm that provides reasonable representation of the data. Han et al. (2012) also studied the Ti-seeded MgO barrier for the fabrication of SLG spin valves with tunneling contacts. It was observed that the current
voltage (I
V) curves between electrodes were highly nonlinear and differential contact resistance showed a very

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FIGURE 5–3 The SLG spin valve device with transparent contacts Co/SLG. Nonlocal MR scans for the three different configurations (A) L 5 1 μm, (B) L 5 2 μm, and (C) L 5 3 μm, respectively. The arrows show the magnetization of the four Co electrodes, and the black curves are taken while H is increasing (decreasing). (D) The dependence of nonlocal MR on the spacing between the central injector and detector electrodes. The red/black curve is a fitted curve of all three data points and the black curve is a fit without the L 5 3 μm data. SLG, Single-layer graphene. Han, W., et al., 2012. Spin transport and relaxation in graphene. J. Magn. Magn. Mater. 324, 369
381. Copyright Elsevier Publishing. Reproduced with permission.

sharp peak at Idc 5 0 μA, which is an indication of tunneling between the Co electrodes and SLG. A 130 Ω nonlocal MR was observed for a SLG spin valve measured at 300K with Vg 5 0 V. The spacing between injector and detector (L)B2.1 μm, and the width of SLG (W)B2.2 μm, the spin-injection efficiency, P, was calculated to be 26%
30% using the preceding equation with experimental values of σG 5 0.35 mS and typical experimental values of λG 5 2.5
3.0 μm. This high spin-injection/detection efficiency highlighted the high quality of the MgO tunnel barrier, which alleviated the conductance mismatch between Co and SLG in spin valves having transparent contacts (Rashba, 2000; Fert et al., 2001). This compares favorably with the tunneling spin polarization of 35%
42% measured by spin-dependent tunneling from Co into a superconductor across polycrystalline Al2O3 barriers (Meservey et al., 1994; Kant et al., 2004; Kaiser et al., 2005). In this measurements the nonlocal and local MR loops measured at 4K and the signal were B100 and B200 Ω, respectively, which is precisely the behavior expected theoretically (Fert et al., 1996; Jedema et al., 2003). The nonlocal MR loop has much better SNR and is more sensitive to detect the spin signal compared to the local MR. Han et al. (2010) achieved tunneling spin injection from Co into SLG using TiO2-seeded MgO barriers. Liu et al. (2013) reported the electrical injection and detection of

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spin accumulation in TLG/MgO/Permalloy LSV structure. Józsa et al. (2009a,b) controlled the efficiency of spin injection into graphene at RT using nonlocal spin valve measurements in cobalt/aluminum-oxide/graphene structures with an injection efficiency as high as 18%, where electrical contact is achieved through relatively transparent regions in the oxide. This value is further enhanced to 31% by applying a dc current bias on the injector electrodes, which causes carrier drift away from the contact. A reverse bias reduces the ac spin valve signal to zero or negative values. Józsa et al. (2009a,b) introduce a model that quantitatively predicts the behaviors of the spin accumulation in the graphene under such circumstances, showing a good agreement with our measurements. Liu et al. (2013) clearly observed in nonlocal spin valve signal in the LSV, indicating that spin coherence extends underneath all FM contacts. Liu et al. (2013) also show that the low-resistivity graphene/MgO/Py junctions enable efficient spin injection and detection in LSV with high applied current density, which leads to large spin accumulation of 120 μV at RT. A spin diffusion length of 1.5 μm was obtained for the injector
detector separation dependence of spin valve signal measurements carried out at RT, while at T 5 10K, the diffusion length increases to 2.3 μm. Ohishi et al. (2007) demonstrate spin injection into a graphene thin film with high reliability by using nonlocal MR measurements, in which the electric current path is completely separated from the spin current path. Using those nonlocal measurements, an obvious MR effect was observed at RT. Ohishi et al. (2007) described the MR effect is due to reversal of FM electrodes.

5.7 Spin relaxation process Graphene is a promising material for spintronics because it is predicted to have long spin relaxation times and long spin diffusion length due to the low intrinsic SO and hyperfine couplings. The measured spin lifetimes in SLG (50
200 ps) were orders of magnitude shorter than expected from the intrinsic SOC. Spin qubits (Trauzettel et al., 2007) and many other spintronic devices based on graphene could become available due to the fact that in intrinsic graphene spins are expected to relax very slowly (Kane et al., 2005; Huertas-Hernando et al., 2006; Honki Min et al., 2006). The reason behind this is the low hyperfine interaction of the spins with the carbon nuclei (only 1% of the nuclei are 13C and have spin) and the weak SOI due to the low atomic number. Spin relaxation in graphene is investigate using electrical graphene spin valve devices in the nonlocal geometry. As the SO and hyperfine coupling are weak in graphene, spin traveling distance is expected to be macroscopic. Usually, the contact between electrodes and graphene could induce a significant relaxation of spin current at the interface. FM electrodes with in-plane magnetizations inject spins parallel to the graphene layer. They are subject to Hanle spin precession under a magnetic field B applied perpendicular to the graphene layer. By introducing a pinhole-free and flatness spin-dependent tunnel barrier, a long spin diffusion length that can exceed 100 μm has been observed in the epitaxial graphene on SiC. In spite of the contact-induced spin precession, some other spin precession mechanisms were also studied. When spin transport in graphene, there are mainly two

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kinds of spin relaxation mechanisms in graphene: Elliott
Yafet (E-Y) mechanism (Elliott, 1954; Yafet, 1963) and Dyakonov
Perel (D-P) mechanism (Dyakonov et al., 1972; Dymnikov et al., 1976). In case of the E-Y mechanism, spin relaxation occurs with a momentum scattering event. It can originate from the charge impurities scattering, phonon scattering, short-range scattering, edge scattering, and so on. Among these scattering factors the charge impurity scattering is mainly considered in a relatively dirty graphene sample at low temperature, and in this case, spin life time τ s is proportional to momentum scattering time, τ p . While in case of the D-P mechanism, spin relaxation occurs between momentum scattering events. Spin relaxation could come from ripples in graphene. Spin lifetime τ s and momentum scattering time τ p have an inverse relationship as τ s ~1=τ p : Tombros et al. (2008) studied anisotropic spin relaxation in graphene using the fields more than 1.5 T force the magnetization direction of the FM contacts to align to the field, allowing injection of spins perpendicular to the graphene plane. Tombros et al. (2008) compare the spin signals at B 5 0 and B 5 2 T and found a 20% decrease in spin relaxation time for spins perpendicular to the graphene layer compared to spins parallel to the layer. Tombros et al. (2008) analyzed the results in terms of the different strengths of the SO effective fields in the in-plane and out-of-plane directions and discuss the role of the E-Y and D-P mechanisms for spin relaxation. Han et al. (2011) investigated spin relaxation in SLG and BLG and observed that they are strongly contrasting behaviors for SLG and BLG. In SLG, the spin lifetime (τ s) varies linearly with the momentum scattering time (τ p) as carrier concentration is varied, indicating the dominance of E-Y spin relaxation at low temperatures. In BLG, τ s and τ p exhibit an inverse dependence, which indicates the dominance of D-P spin relaxation at low temperatures. The different behaviors is due to enhanced screening and/or reduced surface sensitivity of BLG, which greatly reduces the impurity-induced E-Y spin relaxation.

5.7.1 Hanle spin precession The Hanle effect provides an independent measure of the spin diffusion length and also yields the values of the spin lifetime and diffusion constant. This is achieved by applying an out-of-plane magnetic field (H\) that induces spin precession at a Larmor frequency of ωL 5 gμBH\/h ¯ ωL, where g is the g-factor, μB is the Bohr magneton, and h ¯ is the reduced Planck’s constant. Fig. 5
4A
C shows Hanle spin precession curves which are obtained by measuring the nonlocal resistance as a function of H\ for different configurations. The top branches (red/gray curves) are for the parallel magnetization state of the central electrodes, and the bottom branches (black curves) are for the antiparallel magnetization state. The characteristic reduction in the spin signal with increasing magnitude of H\ is a result of spin precession induced by the out-of-plane field, which reduces the spin polarization reaching the detector electrode. For L 5 3 μm a nearly complete Hanle curve is obtained (Fig. 5
4C). For the smaller spacings, the transit time is reduced so that the Hanle peak is broadened and cannot be fully measured within the range of our electromagnet (Fig. 5
4A and B).

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FIGURE 5–4 (A
C) RNL as a function of the out-of-plane magnetic field. (D) Conductivity at different Au deposition. (E) τ S as a function of td: nh 5 2.9 3 1012 cm22 (blue circles), ne 5 2.9 3 1012 cm22 (red squares), and Dirac point (black triangles). For SLG tunneling contacts (F) L 5 4.5 μm, (G) L 5 5.5 μm. (H) Pinhole contacts for L 5 2.0 μm. Red (black) circles are parallel (antiparallel). SLG, Single-layer graphene. Han, W., et al., 2012. Spin transport and relaxation in graphene. J. Magn. Magn. Mat. 324, 369
381. Copyright Elsevier Publishing. Reproduced with permission.

Quantitatively, the Hanle curve depends on spin precession, spin diffusion, and spin relaxation and is given by RNL ~ 6

ðN 0

    1 L2 t pffiffiffiffiffiffiffiffiffiffiffi exp 2 dt; cosðωL t Þexp 2 τS 4Dt 4πDt

where the 1 (2) sign is for the parallel (antiparallel) magnetization state, L is the distance from injector to detector and ϖL 5 gμB HL =h ¯ is the Larmor frequency, in which g is the g-factor, μB is the Bohr magneton, and ¯h is the reduced Planck’s constant (Jedema et al., 2002). Using the preceding above equation, Han et al. (2012) fitted the parallel and antiparallel Hanle curves for SLG and the fitting parameters obtained were D 5 2.5 3 1022 m2/s and pffiffiffiffiffiffiffiffiffi τ s 5 84 ps, which corresponded to a spin diffusion length of λG 5 Dτ S 5 1:5 μm. This Hanle lifetime represents a lower bound of spin lifetime due to contact-induced spin relaxation.

5.7.2 Charged impurity scattering Han et al. (2012) systematically introduce additional sources of charged impurity scattering and monitor their effect on spin lifetime. Gold impurities were selected for this purpose

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because they were shown to behave as charged impurity scatterers with 1/r Coulomb potential when deposited at low temperature without clustering (18K) (McCreary et al., 2010) and were not expected to generate other effects such as resonant scattering, wave function hybridization, or chemical bonding. Fig. 5
4D shows the effect of Au doping (at 18K) on the conductivity of SLG as a function of gate voltage. The mobility decreased and the Dirac point shifted to more negative gate voltages as the doping increased. The shifting of the Dirac point indicates that the Au atoms donate electrons to the graphene and become positively charged scattering centers. Fig. 5
4E shows the best fit values of τ s and D, achieved by fitting Hanle data for device E, as a function of Au coverage at the Dirac point (black squares), for an electron concentration of 2.9 3 1012 cm22 (red circles) and for a hole concentration of 2.9 3 1012 cm22 (blue triangles). In all three cases, τ s did not decrease with increasing Au coverage while the corresponding values of D decreased as a function of Au coverage (Fig. 5
4E inset). These results clearly show that for spin lifetimes on the order of 100 ps, charged impurity scattering is not the primary source of spin relaxation in graphene, even though it is very effective at generating momentum scattering.

5.7.3 Contact-induced spin relaxation Measurement of Hanle spin precession in SLG spin valves having transparent, pinhole, and tunneling contacts shows that contact-induced effects are very important to the spin relaxation in SLG. For SLG spin valves with tunneling contacts, the spin lifetimes measured at the Dirac point were 771 and 448 ps with 4.5 and 5.5 μm spacing, respectively, at 300K, as indicated in Fig. 5
4F and G. These are much longer than the spin lifetime of 134 ps measured for pinhole contacts (Fig. 5
4H) and 84 ps for transparent contacts (Fig. 5
4C at 300K, which are consistent with the values reported in previous studies (50
200 ps) (Tombros et al., 2007; Han et al., 2009a,b,c; Józsa et al., 2009a,b; Popinciuc et al., 2009). Due to the pffiffiffiffiffiffiffiffiffi increased spin lifetimes, the spin diffusion lengths from the Hanle fits (λG 5 Dτ S ) are significantly larger for tunneling contacts (2.5
4.0 μm) than for transparent and pinhole contacts (1.2
1.4 μm). The longer spin lifetimes and spin diffusion lengths with tunneling contacts indicate that the effect of the contact-induced relaxation is substantial for transparent and pinhole contacts. The FM contacts can theoretically introduce spin relaxation through a number of mechanisms. One mechanism is through the inhomogeneous magnetic fringe fields. Roughness of the Co film could produce inhomogeneous local magnetic fields that vary with the morphology, which will generate spin relaxation through inhomogeneous spin precession about the spatially varying local fields that proposed by Dash et al. (2011). A second mechanism is interfacial spin scattering, which is possible because of the direct contact between the FM and graphene for the transparent and pinhole contacts. A third mechanism for contact-induced spin relaxation is related to the Hanle measurement itself. With metallic Co in contact with graphene the spins diffuse from the graphene to the Co with a characteristic escape time, τ esc. Due to the conductance mismatch between Co and graphene, the escape time can become comparable to or less than the actual spin lifetime

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(denoted as the spin-flip time, τ sf). In this case the Hanle lifetime (τ s) in the preceding equation is determined by both the spin-flip time and the escape time, with a simple relationship 21 21 of τ 21 s 5 τ sf 1 τ esc when the spin diffusion length is much larger than the sample size (Zaffalon et al., 2005). Because the Hanle lifetime is determined mostly by the smaller of τ sf and τ esc, it can only provide a lower bound of the spin lifetime. More accurate measurement of the true spin lifetime in graphene is made possible by the insertion of good tunnel barriers 21 to reduce the out-diffusion of spins. This greatly increases τ esc to yield τ 21 s  τ sf when τ esccτ sf, so that the measured Hanle lifetime τ s more accurately measures the actual spinflip time in graphene τ sf. We note that it is possible to model this type of contact-induced spin relaxation numerically (Popinciuc et al., 2009), but an analytic expression for the Hanle curves including escape time effects is currently unavailable. Based on the Hanle data in Fig. 5
4C, Fig. 5
4F
H with much longer lifetimes for tunneling contacts compared to pinhole and transparent contacts, it is likely that for spin valves with lifetimes in the 50
200 ps range, the dominant spin relaxation is generated by the contacts. Thus future studies of spin relaxation will require the use of tunneling contacts to suppress the contact-induced spin relaxation.

5.8 Spin relaxation in single-layer graphene and bilayer graphene Han et al. (2012) studied the spin relaxation in SLG spin valves using the tunnel barrier devices to suppress the contact-induced spin relaxation at RT (300K) and below RT (4K). No correlation between τ s and D is observed at 300K, but at 4K, it shows both quantities increasing with carrier concentration. The correlation of τ s and D implies a linear relation between τ s and the momentum scattering time, τ p (Józsa et al., 2009a,b; Maassen et al., 2011; Fabian, 2007). This indicates that at low temperatures the spin scattering is dominated by momentum scattering through the E-Y mechanism (i.e., finite probability of a spin-flip during a momentum scattering event) (Huertas-Hernando et al., 2009; Elliott, 1954; Meier and Zachachrenya, 1984). The temperature dependences of τ s and D at different carrier concentrations are shown in Fig. 5
5A and B. As the temperature decreases from 300 to 4K, τ s shows a modest increase at higher carrier densities (e.g., from B0.5 to B1 ns for Vg 2 VCNP 5 160 V) and little variation for lower carrier densities. The temperature dependence of D shows a similar behavior as τ s. To analyze the relationship between the spin scattering and momentum scattering, Han et al. (2012) plotted τ s versus D for temperatures lower than, higher than, or equal to 100K temperatures and found that the τ s scales are linear with D at lower temperature, which indicates that an E-Y spin relaxation mechanism is dominant at lower temperatures (#100K), whereas for higher temperatures, τ s and D do not follow the linear relationship, which suggests that multiple sources of spin scattering were present. BLG differs from SLG not just in thickness but also in band structure (linear for SLG, massless fermions versus hyperbolic for BLG, massive fermions) and intrinsic SOC (Abanin et al., 2007; Guinea, 2010).

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FIGURE 5–5 Spin relaxation in SLG: (A) Temperature dependence of spin lifetime and (B) spin diffusion coefficient at different gate voltages relative to the charge neutrality point. Spin relaxation in BLG: (C) Temperature dependence of spin lifetime and (D) spin diffusion coefficient at different gate voltages. Han, W., et al., 2012. Spin transport and relaxation in graphene. J. Magn. Magn. Mat. 324, 369
381. Copyright Elsevier Publishing. Reproduced with permission.

Han et al. (2012) studied the different gate voltage dependence of τ s and D for BLG spin valve. It was observed that longer spin lifetimes are observed in BLG (up to 6.2 ns) than in SLG (up to 1.0 ns). Theoretically, the intrinsic SOC in BLG is an order of magnitude larger than in SLG, which is predicted to result in shorter spin lifetimes for BLG (Guinea, 2010). The opposite experimental trend verifies that the spin relaxation in graphene is of extrinsic origin in the SLG. It is also observed that the spin lifetimes at 20K varies differently from RT results, where τ s varies from 250 to 350 ps as a function of gate voltage and has no correlation with D. Measurement at 4K, gate voltage dependence D exhibits lower values near the charge neutrality point and increasing values at higher carrier densities. The opposite behaviors of τ s and D suggest the importance of D-P spin relaxation (i.e., spin relaxation via precession in internal SO fields) where τ s scales inversely with τ p (Meier and Zachachrenya, 1984; D’yakonov et al., 1972). Fig. 5
5C and D shows the temperature dependences of τ s and D, respectively, for BLG device. At low temperatures, τ s is enhanced while D is reduced, which is different from SLG as shown in Fig. 5
5A and B where both D and τ s increase as temperature decreases for most gate voltages. The opposite trends of the temperature dependences of τ s and D suggest the strong contributions of spin relaxation mechanisms of the D-P type, which is also suggested by Yang et al. (2011a,b). Possible sources of extrinsic E-Y spin relaxation include long-range (Coulomb) impurity scattering and short-range impurity scattering (Castro Neto et al., 2009a,b), while an extrinsic D-P spin relaxation could arise

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from curvature of the graphene film (e.g., random SOCs) (Huertas-Hernando et al., 2006, 2009; Dugaev et al., 2011). The transition from E-Y-dominated SLG to the D-P-dominated BLG could be due to a strong reduction of the E-Y contribution because of enhanced screening of the impurity potential in thicker graphene (Guinea, 2007) and the smaller surface-tovolume ratio. Drögeler et al. (2014) studied the spin lifetimes versus charge mobility for various types of graphene (SLG, BLG, TLG) and substrates (h-BN and SiO2) and found that a significant increase of the spin lifetime (exceeding 3 ns) and length (exceeding 10 μm at 300K) as compared with graphene on SiO2. The graphene layer fully encapsulated by using h-BN layers showed that the ratio of spin lifetimes for spin directions perpendicular and parallel to the graphene sheet could be tuned by the use of top and bottom gate electrodes. These results are in agreement with an electrically induced Rashba-type SOC and open up new possibilities for electric control of spin transport in graphene. A careful procedure was considered to eliminate the effects of the nonencapsulated outer regions, where values of more than 12 μm were obtained for the spin relaxation length in the central encapsulated region, thus establishing a new record for spin relaxation length at RT (Guimarães et al., 2014).

5.9 Electrical spin transport The introduction of magnetism into graphene or the surface states of TIs is particularly interesting toward the quantum anomalous Hall effect (QAHE) (Han, 2014; Pesin et al., 2012; Tse et al., 2011; Zhang et al., 2012; Wang et al., 2013, 2015; Liu et al., 2008; Yu et al., 2010; Culcer et al., 2011; Chang et al., 2013; Kou et al., 2014; Checkelsky et al., 2014; Wei et al., 2013; Alegria et al., 2014; Lang et al., 2014; Jiang et al., 2015). The enhanced spin orbit couplings in hydrogen-doped graphene, silicene, germanane, tin are potential candidates for quantum SHE (QSHE) (Han, 2014; Kane et al., 2005; Liu et al., 2011; Xu et al., 2013; Qian et al., 2014). The spin orbit torque at the TIs/ferromagnet interface has been demonstrated to be significantly larger than conventional heavy metals (Mellnik et al., 2014; Fan et al., 2014). Inducing magnetism in graphene and the surface states of thermal insulator (TIs) holds the potentials toward the realization of QAHE. Very intriguingly, the QAHE has been observed on magnetic TIs recently by several groups (Chang et al., 2013, 2015; Kou et al., 2014; Checkelsky et al., 2014) at ultralow temperatures. Generally speaking, there are two routes to induce magnetism in a nonmagnetic material. The first one is doping another element, such as Mn or Cr, or creating defects. The other one is exchange coupling via proximity effect in adjacent with FM materials, with spin-polarized d or f orbitals. Graphene spintronics was envisioned early on as a promising direction for work and innovation, owing to the combination of the unique electronic band structure of so-called massless Dirac fermions (DFs), weakly sensitive to backscattering and traveling at very high speed over very wide distances at RT. In addition, the weak SOC in sp2 carbon also suggested that electron spin should be carried nearly unaffected over unprecedented distances, making feasible practical applications of lateral spintronics (Roche et al., 2014; Han, 2014; Dery et al., 2012). Pioneering works initially echoed such high expectations (Tombros et al., 2007), while

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FIGURE 5–6 (A) Schematics of the typical four-terminal spin valve geometry for measuring spin lifetime in graphene. (B) Hanle experimental curve from which spin lifetimes are extracted using a semiclassical spin diffusion equation. Reprinted with permission from Drögeler, M., et al., 2014. Nanosecond spin lifetimes in single- and fewlayer graphene-hBN heterostructures at room temperature. Nano Lett. 14, 6050
6055. Copyright 2014 American Chemical Society.

more recent experimental results have confirmed the potential of graphene for transporting spin signal at RT over tens of micrometers, which is more than enough for technological realization (Dlubak et al., 2012; Seneor et al., 2012; Guimarães et al., 2014; Drögeler et al., 2014). Electrical spin transport was studied in nonlocal spin valve field-effect devices by Drögeler et al. (2014) as shown in Fig. 5
6A, where the graphene was attached to, or supported by, a h-BN (Fu et al., 2014). Drögeler et al. (2014) observed that the resistance of the contacts should be higher than the spin resistance of the medium in which the spins are injected to prevent the backflow of the spins into the contact. This backflow would be accompanied by spin relaxation in the metallic ferromagnet (which has a very short spin relaxation time), thus reducing the spin-injection efficiency. It was shown that this had to be balanced with the dwell time of the spins in the semiconductor channel, giving only a narrow window for the contact resistance in devices (Fert et al., 2007). The effective spin relaxation time extracted from Hanle spin precession measurements (see the typical Hanle curve in Fig. 5
6B) and the values of contact resistances (Guimarães et al., 2014; Drögeler et al., 2014; Fert et al., 2007; Kamalakar et al., 2015; Volmer et al., 2013, 2014, 2015; Idzuchi et al., 2015). This identifies the one of the factors that could be responsible for the observed short spin relaxation times. Another bottleneck for spin transport is inhomogeneous oxide barriers which may result in conducting pinholes within the otherwise insulating spin injection and detection barriers. In graphene/MgO/Co spin-valve devices, it was found that the pinholes cause inhomogeneous current flow through the MgO barrier. The actual spin signal results in an additional charge accumulation signal which was measured as a

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magnetic field-dependent background signal in the nonlocal voltage (Volmer et al., 2015). This background signal, which is often observed in nonlocal spin transport studies, is thus a hallmark for the quality of the oxide barriers. Han et al. (2010) observed a nonlocal MR (ΔRNL) of 130 is at RT, which was the largest value in any material. Investigating (ΔRNL) versus SLG conductivity from the transparent to the tunneling contact regimes demonstrates the contrasting behaviors predicted by the driftdiffusion theory of spin transport. Han et al. (2012) also studied introducing Au as charged impurity scattering and monitor their effect on spin lifetime of SLG and found that the mobility decreased and the Dirac point shifted to more negative gate voltages as the Au doping increased. The best fit values of tS and DS achieved by fitting Hanle data for the device as a function of Au coverage at the Dirac point for an electron concentration of 2.9 3 1012 cm22 and for a hole concentration of 2.9 3 1012 cm22. Han et al. (2012) found that the measurement of Hanle spin precession in SLG spin valves having transparent, pinhole, and tunneling contacts has very important effects on induced contact to the spin relaxation time. For SLG spin valves with tunneling contacts the spin lifetimes measured at the Dirac point were 771 and 448 ps with 4.5 and 5.5 μm spacing, respectively, at 300K. These are much longer than the spin lifetime of 134 ps measured for pinhole contacts and 84 ps for transparent contacts at 300K. Due to the increased spin lifetimes, the spin diffusion lengths pffiffiffiffiffiffiffiffiffi from the Hanle fits λG 5 Dτ S are significantly larger for tunneling contacts (2.5
3.0 μm) than for transparent and pinhole contacts (1.2
1.4 μm). The longer spin lifetimes and spin diffusion lengths with tunneling contacts indicate that the effect of the contact-induced relaxation is substantial for transparent and pinhole contacts. Maassen et al. (2012) developed an easy, upscalable process to prepare LSV devices on epitaxially grown monolayer graphene on SiC(0001) and perform nonlocal spin transport measurements. Maassen et al. (2012) have observed that the longest spin relaxation times τ S in monolayer graphene, while the spin diffusion coefficient DS is strongly reduced compared to typical results on exfoliated graphene. Rein et al. (2015) studied the MR and charge transport of CVD grown nitrogen-doped graphene. The nitrogen doping instead leads to a sixfold increase in the charge carrier concentration up to 4 3 1013 cm22 at RT, indicating highly effective doping. The magnetotransport exhibits a conspicuous sign change from positive Lorentz MR in undoped to large negative MR that we can attribute to the doping-induced disorder. At low magnetic fields, Rein et al. (2015) use quantum transport signals to quantify the transport properties. For the study of transport properties, values for the charge carrier mobility μ and the charge carrier density n are derived from Hall measurements and are given by n 5 1=RH e and μ 5 1=eρs n, where the Hall coefficient RH is the slope of the Hall resistance as a function of the applied magnetic field (B), ρs is the sheet resistance, and e is the fundamental charge of an electron. Rein et al. (2015) obtained the charge carrier mobility for undoped graphene is of (1014 6 4) cm2/V s and a charge carrier density of 6.42 6 0.03 3 1012 cm22 at 279K. At 2.5K the mobility amounts to 1122 6 32 cm2/V s and the density to 6.22 6 0.03 3 1012 cm22. These nearly temperatureindependent values are typical for good-quality transferred graphene in line with what has previously been reported (Tan et al., 2007; Song et al., 2012; Li et al., 2009; Wang et al., 2012). In case of the nitrogen-doped graphene, at 279K, the mobility of 23 6 4 cm2/V s is

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strongly reduced, whereas the charge carrier density of 39.25 6 0.15 3 1012 cm22 is strongly increased by more than a factor six compared to the undoped graphene. For the lowtemperature measurement at 2.5K the resulting values are 11 6 3 cm2/V s and 22.00 6 0.15 3 1012 cm22, respectively. On doping the modification of the graphene lattice with additional nitrogen atoms increases the charge carrier density significantly as each nitrogen atom provides an additional electron. Furthermore, these defects form additional scattering centers and decrease the mobility. To quantify the transport properties, Rein et al. (2015) studied the MR of nitrogen-doped and undoped graphene using the relation: MRðBÞ 5

ρsheet ðBÞ 2 ρsheet ð0Þ ; ρsheet ð0Þ

where ρsheet ðBÞ is the sheet resistance at a given magnetic field B. For the undoped graphene a positive MR (Fig. 5
7A and B) was found with a small peak for the low-temperature measurement at fields smaller than 0.5 T. The positive high-field part of the MR at both 2.3 and 279K is well described by the Lorentz MR developed by Kohler (1938): MRBðωc τÞα with ωc the cyclotron frequency, τ the scattering time, and α 5 2, in line with previous measurements [36] on pristine graphene on SiO2. The small positive MR peak at low fields (,0.5 T) is visible only at low temperatures that indicates the phase coherent transport phenomenon of weak localization, which is prominent only at low temperatures. In case of nitrogen-doped graphene exhibits a very different MR signal with the opposite sign as shown in Fig. 5
7C and D, the negative MR is detected at all observed fields and at temperatures of 279 and 2.6K. Both the slope and magnitude of the MR increase at lower temperatures. At 279K the sheet resistance at zero field is ρsheet(0 T) 5 7 kΩ, and we find a 4% decrease (MR) at 8 T. At 2.6K the sheet resistance amounts to ρsheet(0 T) 5 36 kΩ with a much larger decrease of 38% (MR) at 8 T. This observation of negative MR and a strong temperature dependence shows, strikingly, that the nitrogen doping completely changes the nature of the transport. The decrease at small fields can be attributed to weak localization as seen in the low-temperature measurement of the undoped graphene. Tombros et al. (2007) provided the first unambiguous spin-dependent transport measurements in graphene. The spin valve signals and precession measurements revealed a spin relaxation length (λG) of 1.5
2 μm. The spin signals are found to be weakly dependent on temperature or the charge carrier density (which is determined by electrical gating). Later on, spin signal has been measured with electrode distance up to 10 μm (actual λG 5 3.9 μm) at RT, and it has been found that few layer graphene exhibits a longer spin lifetime than SLG (Maassen et al., 2011; Nishioka et al., 2007) due to the screening effect of outer layers which reduce the influence of external scatters. Further improvements could be possible with suspended graphene which mobility exceeds 100,000 cm2/V/s (Guimarães et al., 2012; Bolotin et al., 2008; Du et al., 2010). Up to now, spin relaxation obtained in such devices can reach around 5 μm (Guimarães et al., 2012), but further improvement is foreseen when the influence of the nonsuspended contact part of the graphene sheet is reduced. Although highquality graphene layers can be made by different types of methods and the effect of substrate

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FIGURE 5–7 Magnetoresistance for doped and undoped graphene (magnetic fields applied perpendicular to the sample) for different temperatures. The upper line shows the undoped case for 2.3K (A) and 279K (B). At low temperature, the weak localization peak is clearly visible at zero field. Classical Lorentz magnetoresistance was fitted to the data (white lines). Graphs (C) and (D) show the same measurements for the doped case where the sign and the shape of the curves change compared to the pristine graphene. Reprinted with permission from Rein, M., et al., 2015. Magnetoresistance and charge transport in graphene governed by nitrogen dopants. ACS Nano 2, 1360
1366. Copyright 2015 American Chemical Society.

can be reduced by making graphene suspended between electrodes, little progress has been made in the development of suitable contacts with high spin-injection efficiency and low contact resistance. This obstacle must be overcome first before the full potential of graphene can be utilized for spintronics applications. An alternative way is to make graphene itself magnetic, for example, through molecular doping and proximity effects (Garnica et al., 2013). Graphene edge magnetization is particularly interesting since it turns graphene into a half-metal. However, experimental observations of this theoretically predicted edge magnetism is still lacking (Son et al., 2006a,b). For graphene, hydrogen doping and vacancy defects have been used to induce magnetism. Paramagnetic moments were first observed in graphene with vacancy defects by SQUID measurement and later in graphene with hydrogen doping or vacancy defects probed by pure spin current (Nair et al., 2012; McCreary et al., 2012). In the later one, McCreary et al. used hydrogen atomic source to dope the graphene and performed in situ spin transport measurement. As shown in Fig. 5
8(top), the dip at zero magnetic field for the nonlocal spin transport measurement indicates that the diffusive spins interact with the local hydrogen-induced magnetic moments via exchanging coupling. Up to date, however, longrange FM order in doped graphene is still missing. On the other hand, proximity-induced FM graphene has been observed in the heterostructures of graphene and FM insulator, yttrium iron garnet (YIG) (Wang et al., 2015).

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FIGURE 5–8 (Top) The detection of paramagnetic moments in hydrogen-doped graphene via pure spin. (Bottom) The anomalous Hall resistance measurements on magnetic graphene at various temperatures. Reprinted with permission from McCreary, K.M., et al., 2012. Magnetic moment formation in graphene detected by scattering of pure spin currents. Phys. Rev. Lett. 109, 186604; Wang, Z., et al., 2015. Proximity-induced ferromagnetism in graphene revealed by the anomalous Hall effect. Phys. Rev. Lett. 114, 016603. Copyright American Physical Society.

In their study the graphene device was transferred onto YIG thin films grown via pulsed laser deposition, and a top gate was used to tune the carrier density in graphene. As shown in Fig. 5
8(bottom), anomalous Hall effect has been observed up to 250K in graphene on YIG. The spin orbit coupling in graphene could be enhanced by doping (Castro Neto et al., 2009a,b) of different atoms. A Hall bar device on hydrogen-doped (functionalized) graphene exhibited large colossal SHE (Balakrishnan et al., 2013). Balakrishnan et al. (2013) measured the current flows from IS to ID, and the voltage is measured on the other two bars in a nonlocal geometry. The measured giant nonlocal resistance indicates the largely enhanced spin orbit coupling in graphene. However, there are also reports stating that this nonlocal voltage could be associated with some unknown mechanisms that are not related to spin (Kaverzin et al., 2015). Pi et al. (2010) studied the effects of surface chemical doping on spin transport in graphene by performing nonlocal measurements in ultrahigh vacuum while depositing gold adsorbates. Pi et al. (2010) manipulation of the gate-dependent nonlocal spin signal as a function of “Au”-coverage and observed that that Au-charged impurity scattering is not the dominant mechanism for spin relaxation in graphene, despite its importance for momentum

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FIGURE 5–9 In-plane and out-of-plane magnetization curves of SLG/YIG measured at 100K. The inset shows the in-plane magnetization curves of SLG/YIG at 100 and 300K in the low field region. SLG, Single-layer graphene; YIG, yttrium iron garnet. Reprinted with permission from Sakai, S., et al., 2018. Dirac cone spin polarization of graphene by magnetic insulator proximity effect probed with outermost surface spin spectroscopy. Adv. Funct. Mater. 28, 1800462. Copyright Wiley online library.

scattering. They have found unexpected enhancements of the spin lifetime illustrate the complex nature of spin relaxation in graphene.

5.9.1 Spin polarization Sakai et al. (2018) studied the spin polarization of graphene by magnetic insulator proximity effect probed with outermost surface spin spectroscopy. In the magnetization measurements, the SLG/YIG Y3Fe5O12 exhibits an in-plane easy axis of magnetization due to the YIG thin film (MYIG) as shown in Fig. 5
9. Under the in-plane magnetization conditions, the coercive field (HC) and the remanent magnetization (MR) normalized to the saturation magnetization (MS) are 4 Oe and  0.35 at 100K and 0.3 Oe and  0.2 at 300K, respectively (inset of Fig. 5
9). The saturation fields in the out-of-plane magnetization curves are 2400 (100K) and 2000 Oe (300K), which are interpreted as the shape anisotropy of the YIG thin film, 4πMYIG.

5.10 Spintronics magnetoresistance devices MR is the change in electrical resistance under an external magnetic field at the atomic level, which is of great interest both fundamentally and technologically. Graphene and other twodimensional layered materials provide an unprecedented opportunity to explore MR at its nascent stage of structural formation. Investigation of MR, the changes in electric resistance under an external magnetic field (B), is of interest both from fundamental and technological viewpoints. MR sensors are widely used in day-to-day applications (Ripka et al., 2010), where the MR value is an important figure of merit. Several attempts have been made to understand the MR in graphene. A large value of MR in weak magnetic fields at RT (300K) is of special attention in producing MR sensors. MR in SLG- and multilayer graphene in the temperature range 1.9
400K and in magnetic fields up to 14 T was extensively studied

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(Gonalez et al., 2007; Friedman et al., 2010; Gopinadhan et al., 2013, 2015; Liao et al., 2012). The MR values from 20% to 40% in field 0.5 T at 300K depending on the sample cross-size were measured on HOPG with different thickness (Gonalez et al., 2007). The MR value for multilayer epitaxial graphene grown on a SiC substrate was less than 5% (Friedman et al., 2010). The MR dependence on temperature and magnetic field was measured on a SLG in Gopinadhan et al. (2013). It was found that MR does not reach saturation and equals to 275% at 300K in the 9 T field. At 0.5 T the MR value does not exceed 10%. If the current is perpendicular to the graphene layer plane, the MR value is less than 1% in this field (Liao et al., 2012). Recently, an extremely large MR (880% at 9 T and 400K) was measured in few-layer graphene/BN heterostructures (Gopinadhan et al., 2015). However, at 300K and a magnetic field of 0.5 T, MR value was about 25%. A linear MR of 80%
250% was reported in epitaxial multilayer graphene at 2K in a normal magnetic field of 12 T (Friedman et al., 2010). The explanation of the temperature-independent linear MR based on the quantum theory (Abrikosov et al., 2000) fails to account for the observed MR at RT as no quantum effects are expected at RT. A linear and quadratic MR of 60%, at 300K, and a magnetic field of 14 T are reported in chemical vapor deposition grown few-layer graphene with the current perpendicular to the film plane (Liao et al., 2012). In magnetic storage applications the data are retrieved from a magnetic hard disk with a MR read sensor that is extremely sensitive to low (stray) magnetic fields (Parkin et al., 2004; Novoselov et al., 2005). Prompted by the huge demand for MR sensors with a high sensitivity, low energy consumption, low cost, and ready availability, researchers are investigating various new materials. Graphene, a stack of singlelayer carbon atoms arranged in a hexagonal periodic lattice with weak van der Waals interlayer interaction, can be a wonder electronic material that can exhibit large MR values. Fundamentally, atomically thin structures provide the simplest system to understand the origin of MR, thanks to the discovery of semimetallic graphene and other low-dimensional conducting structures (Novoselov et al., 2005; Castro Neto et al., 2009a,b; Zhang et al., 2005). In addition, the structure suffers from defects, grain boundaries, and impurities, which limit the mobility of the carriers reducing the MR. By increasing the mobility with graphene peeled from kish graphite (Morozov et al., 2008), a larger RT MR can be achieved. However, it is found that highly ordered pyrolytic graphite samples with sizes of 100 mm exhibit a very small MR (González et al., 2007), and the smaller MR observed in the past in few-layer graphene samples is explained by the size effect. Recently, a sizeable MR is reported at 300K from SLG (Gopinadhan et al., 2013) owing to enhanced scattering from charged impurities. A negative MR is also predicted (Kim et al., 2008) and observed (Bai et al., 2010) in GNRs, but the MR value is limited (B100%). A finite MR has been reported recently on graphene/ BN vertical heterostructures (Kamalakar et al., 2014; Chen et al., 2013). Now, there are various scenarios, such as local Hall bar, van der Pauw, and nonlocal geometry, where the multilayer graphene can show large MR values. In a local geometry, small-sized (few microns) few-layer graphene samples can exhibit an extremely large MR of 2000% at 400K in graphene/BN heterostructures. Experimental evidence shows that fewlayer graphene can be an electronic material for MR sensors at both low and high magnetic fields, and at practical device operating temperatures of up to 400K, with a very small

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temperature coefficient of resistance. The MR is electric-field tunable, thus providing additional functionalities to the sensor. Furthermore, we show that Ettingshausen
Nernst effect based nonlocal MR sensors can exhibit MR 490,000% at RT. Gopinadhan et al. (2015) studied extremely large MR in few-layer graphene/BN heterostructures. In this study, Gopinadhan et al. (2015) reported an extremely large local MR of B2000% at 400K and a nonlocal MR of .490000% in an applied magnetic field of 9 T at 300K in few-layer graphene/BN heterostructures. It was predicted that the local MR is aroused from large differential transport parameters, such as the carrier mobility, across various layers of few-layer graphene upon a normal magnetic field, whereas the nonlocal MR is due to the magnetic field
induced Ettingshausen
Nernst effect. Nonlocal MR suggests the possibility of a graphene-based gate tunable thermal switch. Gopinadhan et al. (2015) also claimed that the graphene heterostructures may be promising for magnetic field
sensing applications. Matveev et al. (2018) studied the MR of graphene at RT and found that the maximum positive MR was 100% in 0.5 T magnetic field normal to the graphene plane. The measured MR was positive with quadratic magnetic-field (B) dependence in the fields up to B0.07 T and quasilinear dependence in the fields up to 0.5 T. It was also found that the larger was the area of monolayer graphene in the grown graphene films, the higher was the MR value.

5.11 Applications of graphene spintronics 5.11.1 Spin valve devices There are mainly two kinds of geometries for graphene spin valve devices: local geometry and nonlocal geometry. The local geometry has two FM electrodes; the spin current is injected from one electrode, then transported through graphene, and detected by another electrode. In this geometry the spin current is mixed with the charge current. The spin signal is detected as the difference in resistance between the parallel state and antiparallel state of two magnetic electrodes. Generally, it can be expressed as MR 5 (ΔR/Rp) 3 100%. ΔR 5 RA 2 Rp. Rp and RA denote the resistance of parallel state and antiparallel state, respectively. Usually, the two magnetic electrodes should be different in geometric aspect ratios to obtain different coercive fields; thus the two electrodes can be aligned to antiparallel or parallel states by sweeping the magnetic field. Wang et al. (2008a,b,c,d) reported the MR properties of quasi-two-dimensional mesoscopic graphene (MG) spin valve devices consisting of MG flakes contacted by FM electrodes. For the devices, ultrathin magnesium oxide (MgO) tunnel barrier is inserted at the FM/MG interface; the spin valve effect has been observed, with MR magnitudes up to 12% at 7K and signals persisting up to temperatures as high as 60K. In contrast, the spin valve effect has not been seen in devices without MgO, suggesting the importance of spin-dependent interfacial resistance for spin injection into MG. In addition an investigation of the voltage bias dependence and gate voltage dependence of MR has been performed. Fig. 5
10A
C shows the magnetic field dependences of resistance for three MG spin valve devices A, B,

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FIGURE 5–10 (A
C) Two-terminal resistances measured during the magnetic field scan for samples A, B, and C, respectively. The thicknesses of MG in samples A, B, and C are 40, 30, and 10 nm, respectively. Black (red) curve corresponds to upward (downward) sweeping of the magnetic field. (D)
(F) Temperature dependences of magnetoresistance for samples A, B, and C, respectively. Reprinted with permission from Wang, W.H., et al., 2008a. Magneto-transport properties of mesoscopic graphite spin valves Phys. Rev. B 77, 020402(R); Wang, W.H., et al., 2008b. Growth of atomically smooth MgO films on graphene by molecular beam epitaxy. Appl. Phys. Lett. 93, 183107; Wang, X., et al., 2008c. Atomic layer deposition of metal oxides on pristine and functionalized graphene. J. Am. Chem. Soc. 130, 8152
8153; Wang, Y., et al., 2008d. Room-temperature ferromagnetism of graphene. Nano Lett., 9(1), 220
224. Copyright American Physical Society.

and C. The thicknesses of MG in samples A, B, and C are 40, 30, and 10 nm, respectively, where the thicknesses are determined by AFM after all magnetotransport measurements have been performed. These three samples have electrode gaps of 200 nm and have RT resistances of 45, 17, and 11 kΩ, respectively. As the magnetic field is ramped up or down, we observe two transitions in the device resistance corresponding to the two distinct coercivities of the FM electrodes. This MR is attributed to the spin valve effect in which the resistance depends on the relative magnetization alignment of the two FM electrodes due to spin-polarized transport across the MG. For samples A and C the MR is negative—the device resistance is lower for the antiparallel magnetization state of the FM electrodes. However, we observe positive MR for sample B, where device resistance is higher when FM electrodes are in an antiparallel magnetization state. Considering that the bulk Fermi-level spin

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polarizations for Co and Fe (35% and 40%, respectively) have the same sign (Meservey et al., 1994), one would expect from a Julliere-type analysis (Julliere, 1975) that the MR should always be positive. However, in the case of a double-barrier structure with spin-preserving conduction channel, both positive and negative MR are possible due to quantum interferences in the channel. In addition, it is possible that oxidized Co or Fe at the FM/MgO interface can change the sign of MR by changing the Fermi-level spin polarization (Belashchenko et al., 2004). While the exact cause of the different MR signs is not yet clear, this behavior is consistent with the related CNT spin valves where the MR can be positive (Tsukagoshi et al., 1999), negative (Thamankar et al., 2006), or controlled by gate voltage (Sahoo et al., 2005; Man et al., 2006; Nagabhirava et al., 2006). We note that there are small resistance jumps in the MR scans, which do not repeat on subsequent scans. Surface antiferromagnetic oxides are likely to form at the edges of the FM electrodes where there is no gold cap. This could generate domain wall pinning sites to produce a gradual switching (as seen in sample B). Surface antiferromagnetic oxides may also generate asymmetric switching (as seen in sample A) through the exchange bias effect. Randomness in homogeneities in the MgO barrier could generate a spatially varying tunneling rate, which would make the MR signals more sensitive to particular FM regions and could produce sample-to-sample variations in device resistance. To improve the MR characteristics and reduce sample-to-sample variations, optimization of the MgO barrier and the FM capping procedures are needed. The temperature dependence of MR for samples A, B, and C are shown in Fig. 5
10D
F. Sample A, which exhibits the largest value of MR among our sample set, exhibits MR up to at least 50K, but unfortunately the temperature dependence was not completed due to sample failure. The MR of sample A is 9% at 7K and gradually decreases at a roughly uniform rate as temperature increases. The MR of sample B is 3.6% at 7K and decreases in a nonlinear manner as temperature rises and disappears at 65K. The MR for sample C is 2.7% at 1.7K and decreases almost linearly as temperature increases, eventually disappearing at 40K. It is arguable that the observed spin valve signals are not due to spin-polarized transport through MG but to other effects such as anisotropic MR (A-MR). This type of MR, however, should show in devices even with a single FM electrode contact. Hill et al. (2006) studied the graphene spin valve device using soft magnetic Ni
Fe electrodes into graphene and inject polarized spins changes 10% in their resistance when the electrode switches from the parallel to the antiparallel state within the applied field is swept between 450 and 450 G at the RT. This coupled with the fact that a field-effect electrode can modulate the conductivity of these graphene films makes them exciting potential candidates for spin electronic devices. The negative ΔR is mainly due to the A-MR of the electrodes. Maassen et al. (2011) studied the spin-transport properties of FLG. Fig. 5
11A shows a typical nonlocal spin-valve measurement (Tombros et al., 2007) on FLG. Sending a generates a spin accumulation at the electrode. The spins diffuse on both sides of the electrode along the flake and generate a voltage drop Vnl between electrodes, defining the nonlocal resistance Rnl 5 Vnl/I. Switching the magnetization of one of the inner electrodes using an in-plane magnetic field results in a sign change of Rnl (see Fig. 5
11A). When the outer contacts are located within the spin-relaxation length, additional switches can be observed

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FIGURE 5–11 (A) Nonlocal spin-valve signal of a seven-layer graphene having electrode distance L 5 8 μm. The red and black arrows show sweep directions of the magnetic field. (B) Hanle precession measurements of a five-layer graphene at Vg 5 V0 for L 5 2.8 μm (black, largest amplitude), L 5 5.4 μm (blue, smallest amplitude), and hole-doped state at Vg 5 V0 2 60 V for L 5 5.4 μm (red, intermediate amplitude). The precession is measured for the P and AP configuration of the inner contacts. The curve for L 5 2.8 μm shows a switch from the P to the AP state at 2140 mT. AP, Antiparallel; P, parallel. Reprinted with permission from Maassen, T., et al., 2011. Comparison between charge and spin transport in few-layer graphene. Phys. Rev. B 83, 115410. Copyright American Physical Society.

(Tombros et al., 2007). The spin-valve measurement in Fig. 5
11A is taken on a seven-layer graphene with an inner contact distance of L 5 8 μm. In Fig. 5
11B, three Hanle measurements on a five-layer graphene sample are presented. Each curve consists of the nonlocal signal acquired for the parallel (P) and antiparallel (AP) orientation of the inner contacts. The black and the blue dots represent the measurements for L 5 2.8 μm and L 5 5.4 μm, respectively, at the gate voltage V0. The red curve is measured on the longer distance at Vg 5 V0 2 60 V, where electron charges are induced by the gate. The amplitude for the measurement with increased L is smaller due to additional spin relaxation, as the spins have to travel a longer distance resulting in a longer time interval for spin relaxation. In addition to the change in the amplitude, a shift in the B-field values for the crossing points of the parallel and the antiparallel precession curve is visible for the two curves measured at Vg 5 V0. The crossing points represent the B-field value where the spins have, on average, precessed for 90 degrees, resulting in both configurations in a signal of Rnl  0. An increased distance L, corresponding to an increased travel time for the spins, therefore decreases the B field which results in 90 degrees precession (Popinciuc et al., 2009; Fabian, 2007). The measurements show that the spin signal can be enhanced by inducing more charge carriers (see enhanced spin signal comparing the measurement at Vg 5 V0 2 60 V and Vg 5 V0) that was also observed by Józsa et al. (2009a,b) in SLG.

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5.11.2 Field-effect transistor Józsa et al. (2008) studied the drift of electron spins under an applied dc electric field in SLG spin valves in a field-effect transport at RT. In this system the graphene layer is contacted by a series of Co electrodes used for the injection or detection of spins and the application of the dc electric fields. Different dimensions of the contacts and their spacing were used by Józsa et al. (2008) for the spin valves in field-effect transport. Józsa et al. (2008) have fabricate using the graphene flakes typically between 0.2 and 4 μm wide and up to 50 μm long. This results are shown in Fig. 5
11A; where the variable width of the Co electrodes yields a difference in their coercive fields (20
50 mT), allowing for separate switching in the external magnetic field. Contacts with similar widths, however, can sometimes interchange in their switching order depending on, for example, the domain wall nucleation, on the sweeping direction of the field (positive to negative values and vice versa). Such an effect is observed in the measurement shown in Fig. 5
12A. Each discrete resistance level seen in Fig. 5
12(top) corresponds to a combination of magnetically parallel or antiparallel electrodes. The step heights carry information about the importance of that specific electrode’s contribution to the spin injection or detection process. Unlike in semiconductors, in graphene, due to its band structure, electrostatic gating allows for switching from hole to electron conduction while keeping the carrier mobility, diffusion constant, Fermi velocity, electric conductivity, and other parameters approximately unchanged. The gating effect is reflected in the graphene resistivity as plotted in Fig. 5
12(bottom). The position of the Dirac neutrality point separating hole and electron conduction regimes, as identified by the minimum in conductance (Novoselov et al., 2005; Tan et al., 2007) shows a small hysteresis (19 to 111 V) for the two sweep directions. The measurement also reveals a carrier mobility of μ 5 0.25 m2/V s.

Ballistic transistors The availability of high-mobility graphene up to RT makes ballistic transport in nanodevices achievable. The p
n
p transistors in the ballistic regime give access to Klein tunneling physics and allow the realization of devices exploiting the optics-like behavior of DFs as in the Veselago lens or the Fabry
Pérot cavity. Wilmart et al. (2014) propose a Klein tunneling transistor based on the geometrical optics of DFs. They considered a prismatic active region delimited by a triangular gate, where total internal reflection may occur, which leads to the tunable suppression of transistor transmission. Liang et al. (2007) also studied the performance projections for ballistic GNR FET and found that it behaves as high-mobility digital switches, with the potential to outperform the silicon MOSFET.

5.11.3 Hall effect The introduction of magnetism into graphene or the surface states of TIs is particularly interesting toward the QAHE (Han, 2014; Pesin et al., 2012; Tse et al., 2011; Zhang et al., 2012; Wang et al., 2013, 2015; Liu et al., 2008; Yu et al., 2010; Culcer et al., 2011; Chang et al., 2013; Kou et al., 2014; Checkelsky et al., 2014; Wei et al., 2013; Alegria et al., 2014; Lang et al., 2014; Jiang et al., 2015). The enhanced spin orbit coupling in hydrogen-doped graphene,

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FIGURE 5–12 (Top) Measurement of nonlocal spin valve where arrows represent the magnetic orientation of different electrodes. (Bottom) 4-point measurement of the graphene resistivity between two electrodes contacts versus the gate voltage. Reprinted with permission from Józsa, C., et al., 2008. Electronic spin drift in graphene field-effect transistors. Phys. Rev. Lett. 100, 236603. Copyright American Physical Society.

silicene, germanane, tin, are potential candidates for QSHE (Han, 2014; Kane et al., 2005; Liu et al., 2011; Xu et al., 2013; Qian et al., 2014). The spin orbit torque at the TIs/ferromagnet interface has been demonstrated to be significantly larger than conventional heavy metals (Mellnik et al., 2014; Fan et al., 2014). Inducing magnetism in graphene, the surface states of TIs holds the potentials toward the realization of QAHE. Very intriguingly, the QAHE has been observed on magnetic TIs recently by several groups (Chang et al., 2013, 2015; Kou et al., 2014; Checkelsky et al., 2014) at ultralow temperatures. Generally speaking, there are two routes to induce magnetism in a nonmagnetic material. The first one is doping another element, such as Mn or Cr, or creating defects. The other one is exchange coupling via proximity effect in adjacent with FM materials, with spin-polarized d or f orbitals. Sakai et al. (2018) measured the hall resistance to study the effects of the proximity contact in SLG/YIG Y3Fe5O12 on the carrier transport property of SLG. Fig. 5
13 shows an almost linear magnetic field dependence of the Hall resistance RHE that is attributed to the ordinary Hall effect by the Lorentz force. The type, concentration, and mobility of carriers in SLG are determined to be p-type,  1 3 1012 cm22, and  900 cm2/V/s, respectively, from the slope of the curves. The carrier mobility is relatively low compared to the reported values for the devices of CVD-grown SLG with SiO2 dielectrics (1000
7000 cm2/V/s) (Hwang et al., 2010; Chan et al., 2012). In light of the nonchemical nature of adhesion of SLG on the YIG surface and the effective dielectric screening in YIG with a high dielectric constant (  19.5) (Chen et al., 2008; Fallahazad et al., 2010), it is reasonably expected that higher carrier mobility would be achieved in graphene/YIG devices by the reduction of the corrugation of graphene by using the YIG film with atomically flat surface (Tang et al., 2016) and by further improvement in the graphene transfer process which enables to prevent the crack formation. Moreover, a small but nonnegligible nonlinear resistance is identified in the RHE
B curve at

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FIGURE 5–13 Hall resistance RHE as a function of the magnetic field measured at 20, 100, and 300K. Nonlinear components RNL in the Hall resistance are plotted in the inset. Reprinted with permission from Sakai, S., et al., 2018. Dirac cone spin polarization of graphene by magnetic insulator proximity effect probed with outermost surface spin spectroscopy. Adv. Funct. Mater. 28, 1800462. Copyright Wiley online library.

each temperature (20, 100, and 300K). The nonlinear resistance component (RNL) obtained by subtracting the linear component from the RHE
B curve is plotted in the inset of Fig. 5
12. RNL shows a sigmoidal curve whose magnitude changes depending on the temperature. It has been reported that the graphene devices composed of a SLG/YIG heterostructure exhibit a nonlinear Hall resistance attributed to AHE (Wang et al., 2015).

5.11.4 Bipolar spintronics To use graphene for bipolar spin transport, Han et al. (2012) studied the nonlocal electron and hole spin transport under different dc bias current: Idc 5 1300 μA (squares), 0 μA (circles), and 2300 μA (triangles) in the SLG device (VD 5 234 V and L 5 1 μm) at 300K as shown in Fig. 5
14C. The AC modulation for lock-in detection was 30 μA. For positive bias, which means that the current is going from Co to SLG (illustrated in Fig. 5
14A, the gatedependence of ΔRNL follows the zero bias data. On the other hand, when the bias is negative and the carriers are holes (triangles, Vg , VD), a strong reduction of ΔRNL was observed. In this case the holes in the SLG are driven toward the Co electrode E2 and become spin polarized due to spin-dependent reflection from the FM interface (i.e., spin extraction; Dery et al., 2007), as shown in Fig. 5
14B. A very interesting aspect is that the reduction of ΔRNL was observed for spin extraction of holes but not for the spin extraction of electrons. For further studies, this reduction of the ΔRNL was measured using nonlocal MR as a function of dc bias current at fixed gate voltages: Vg 5 0 V (electrons, solid squares) and for Vg 5 270 V (holes, open squares) at 300K, as shown in Fig. 5
14D. For electrons, there is only a slight variation in ΔRNL as a function of Idc. For holes at positive bias the behavior of ΔRNL is similar to the electron case. For holes at negative bias, however, there is a significantly stronger variation of ΔRNL as a function of dc current bias, with decreasing ΔRNL at larger negative biases. The ΔRNL is measured at 300K as a function of both gate voltage and dc current bias. The two main trends, namely, the roughly constant ΔRNL versus Idc for

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FIGURE 5–14 Bipolar spin transport in SLG spin valves with transparent contacts at 300K. (A and B) Schematic drawing of spin injection under positive bias and spin extraction under negative dc bias current. (C) Nonlocal MR as a function of gate voltage; at Idc 5 0 A, Idc 5 300 A, and Idc 5 2300 A. (D) Nonlocal MR as a function of dc bias current; at Vg 5 0 V (electrons, solid squares) and 270 V (holes, open squares). (E) Nonlocal MR as a function of dc bias current at the hole (the dashed line shows the zero value). MR, Magnetoresistance; SLG, single-layer graphene. Han, W., et al., 2012. Spin transport and relaxation in graphene. J. Magn. Magn. Mat. 324, 369
381. Copyright Elsevier Publishing. Reproduced with permission.

electrons and the reduced ΔRNL for hole spin extraction, can be clearly seen. These behaviors have been observed in all devices. In one particular SLG device (DeviceN, transparent contacts, VD 5 232 V, L 5 1 μm), there was a very strong change of ΔRNL as a function of dc current bias for hole doping (Vg 5 250 V). By increasing the negative dc bias current to 2660 μA, a sign reversal of ΔRNL at B450 μA was observed that can be seen in Fig. 5
14E. These observations is observed in many different systems (Lou et al., 2007; Valenzuela et al., 2005; Dash et al., 2009), but the interesting aspect of SLG is that this effect is seen for holes but not for electrons, despite having symmetric electron and hole bands. Therefore SLG provides a unique system to investigate this phenomenon.

5.12 Graphene-ferroelectric meta-devices Memory metamaterials are artificial media that sustain transformed electromagnetic properties without persistent external stimuli. Earlier memory metamaterials were realized with phase-change materials, such as vanadium dioxide or chalcogenide glasses, which exhibit memory behavior with respect to electrically/optically induced thermal stimuli. But these materials require a thermally isolated environment for longer retention or strong optical pump for phase-change.

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Kim et al. (2015) demonstrate electrically programmable nonvolatile memory metadevices realized by the hybridization of graphene, a ferroelectric meta-atom/meta-molecule (MM). These meta-devices are superior in nonvolatility and logic-gate operation at RT, which provides a new pathways for emerging optoelectronic applications. For a memory metadevice having a single electrical input, amplitude, phase, and even the polarization multistates were clearly distinguishable with a retention time of over 10 years at RT. Furthermore, logic-gate functionalities were demonstrated with reconfigurable logic-gate meta-devices having two electrical inputs, with each connected to separate ferroelectric layers that act as the multilevel controller for the doping level of the sandwiched graphene layer. In these graphene-ferroelectric nonvolatile memory meta-device (GF-NMM) fabrications, an array of hexagonal MAs, SLG, a ferroelectric polymer [poly(vinylidene fluoride-cotrifluoro-ethylene) or P(VDF-TrFE)], and a terahertz transparent electrode (TTE) composed of periodical subwavelength-scale metallic strips are placed sequentially on a polyimide substrate (Kim et al., 2015). An array of the hexagonal metallic pattern exhibiting polarization-independent inductance-capacitance resonance was chosen as the MA structure to intensify light
matter interaction. A large-sized SLG synthesized by chemical vapor deposition (Kim et al., 2009) on a Cu foil is transferred onto the array of MAs on polyimide using a ferroelectric polymer as a mechanically supporting film (Ni et al., 2012). As the dipoles in the ferroelectric polymer (PVDF-TrFE) consist of weakly electronegative hydrogen atoms and strongly electronegative fluorine atoms, the application of an external gate voltage (VG) over the coercive voltage (VC) aligns the dipoles in the ferroelectric. These aligned dipoles, as a result, induce polarization (P) at the surface of the ferroelectric and exert an electrostatic force consistently on the charge carriers in the graphene layer. The TTE was carefully designed to apply a uniform electric field to the ferroelectric and transmit broadband terahertz waves vertically incident on the TTE without much loss (Lee et al., 2012). Pulsed external gate voltage [VG,pulse (V)] lasting for 1 s was applied between the TTE and the graphene/MAs in the measurement. Fig. 5
15A shows the GF-NMM fabricated terahertz time domain spectroscopy. On applying VG,pulse (1200 V) P, corresponding to positive remanent polarization (1PR), depletes the same polar charges out of graphene; THz transmission spectra through the GF-NMM showed a resonance dip at 1.1 THz. With the subsequent application of VG,pulse (2200 V), P changes to negative remanent polarization (2PR); the resonance frequency then shifted to 0.8 THz, and the bandwidth was observed to slightly broaden. Fig. 5
15B shows the measured transmission amplitude (TA) through the GFNMM and the Fermi level (EF) of graphene in the GF-NMM as a function of VG,pulse. Hysteretic variation of the spectral features is attributed to the change in graphene doping level resulting from the reversal of ferroelectric. It shows gradual transmission change near positive and negative VC. The coupled chiral MMs is employed in the fabrication of the chiral GF-NMM. Fig. 5
15C shows the azimuthal polarization rotation angle (θ) for two distinct VG, pulse values, where θ is extracted from the phase difference between the two circular polarizations. The Δθ attains the maximum value of 8 degrees at 1.1 THz. To trace the hysteretic behavior in the polarization states more clearly, θ measurement was carried out at 1.1 THz and is plotted in Fig. 5
15D. Because of ferroelectricity, θ changes gradually near positive

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FIGURE 5–15 (A) Measured (open circles) and simulated (solid lines) THz transmission spectra for VG,pulse lasting for 1 s [Red: VG,pulse(1200 V), Blue: VG,pulse(2200 V), and Yellow: VG,pulse(2120 V)]. (B) Hysteresis transmission amplitude (TA). VC1 and VC2 are the positive and negative coercive voltages, respectively. Arrow refers to the VG,pulse sweep direction. Logic states denoted as 00, 01, 10, and 11 correspond to the multilevel transmission amplitudes for the retention time measurement. (C) Polarization rotation angle (θ) through chiral memory metadevice after the application of an external pulsed gating voltage (VG,pulse) lasting for 1 s, and the difference (Δθ) between θ1200 V and θ2200 V. (D) Polarization rotation angle (θ) in the VG,pulse within a range of 1200 and 2200 V at 1.1 THz. Arrows refer to the VG,pulse sweep direction. CDZM, Conjugated double Z meta-molecules. Reprinted with permission from Kim, W. Y. et al. (2016). Graphene
ferroelectric metadevices for nonvolatile memory and reconfigurable logic-gate operations. Nat. Commun. 7, 10429. Copyright Nature Publishing Groups.

and negative VC. Multilevel polarization states can also be stored for over 105 s without much degradation. Ni et al. (2012) also studied the nonvolatile ferroelectric polymer gating using monolayer graphene.

5.13 Conclusion and perspectives of graphene-based spintronics Spin injections into graphene have been realized experimentally. Local spin valve devices show about 10% MR change at low temperatures; nonlocal spin valve devices show RT gate tunable spin transport properties. With the help of high-quality tunnel barrier, spin current was observed to diffuse up to 100 μm in epitaxial graphene on SiC. Several spin relaxation mechanism in graphene have been studied. The spin relaxation in single layer and double layer was found to be dominated by different mechanisms, and some experiments showed complicated spin relaxation mechanisms. In order to enhance the spin polarization and spin

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life time, further studies on the spin relaxation mechanisms are needed. The one-dimensional Z-GNRs show interesting magnetic properties. Some novel graphene spintronics devices based on Z-GNR have been proposed in theory, but more efforts should be made for experimental studies. One of the challenges for the fabrication technology for producing spintronic prototype devices based on large-area graphene (mainly CVD-grown polycrystalline graphene). This requires the study of device variability on the same integration chip, the clarification of the impact of grain boundaries on spin lifetimes, and the achievement of long spin lifetimes. Although the electronic quality of the CVD form is compromised by the polycrystalline morphology that displays a variety of grain boundaries, scaffolds for chemical contamination, reasonable spin lifetimes of about 200 ps, and spin diffusion lengths around 2 μm with channel lengths up to 4 μm have already been reported at RT. The long-distance spin transport capability of CVD graphene with channels extending up to 16 μm on SiO2/Si substrates has been discussed. Nonlocal spin transport and Hanle precession measurements in such long channels give rise to a spin lifetime of 1.2 ns with RT spin diffusion length of 6 μm (Roche et al., 2015). Such long spin transport capability of CVD graphene opens up possibilities for lateral spintronic memory and logic technologies. Experimental and theoretical efforts will be further maximized and a focus will be made on engineering efficient device fabrication techniques embedding spin manipulation functionality such as the SHE and spin torque, to produce practical (RT) applications of graphene spintronic devices. Spin current generation by means of the SHE in heavy metals for spin torque applications has gained great attention in recent years and magnetic memory prototypes have been demonstrated using this technology. Spin Hall angles larger than 10% have been recently reported in graphene after the addition of adatoms, metallic particles, or by contacting graphene with WS2, suggesting that the SHE (in modified graphene) can also be an efficient way of generating spin currents. To implement strong enough SOC heterostructures of graphene with other 2D materials having large SOC is another target, since they offer larger capability to manipulate spin by electric field. For spin injection and detection in graphene the conventional metal-oxide barriers with FM contacts are often used. However, the growth of oxide barriers on graphene is challenging and suffers from problems such as pinholes, interface roughness, and defects related to oxygen vacancies and doping in graphene under the contacts. The 2D nature of insulating h-BN makes it appealing for spin-tunnel barriers because of the absence of surface states, which can minimize the charge or spin traps at the interface. Graphene has a strong potential for vertical geometry spintronic devices since it can efficiently filter one of two spin channels or it can significantly improve interfacial spin polarization giving rise to high tunnel MR values. An important next step is to explore the behavior of noncollinear spintronic phenomena such as spin transfer torque, important for development of STT-MRAM and spin torque nano-oscillators [38]. The chemical functionalization of graphene by heavy adatoms is also a research direction of interest, given the predictions of the formation of a QSHE. In conclusion the type of graphene used in spintronic devices (mechanically exfoliated, epitaxially grown, or CVD grown) is important to achieve long spin diffusion lengths, but the preservation of its quality must be enforced by using encapsulation with h-BN layers or special treatments to avoid contact contamination, impedance mismatch, or too-invasive

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Further reading Ingla-Aynes, J., et al., 2015. 24-μm spin relaxation length in boron nitride encapsulated bilayer graphene. Phys. Rev. B 92, 201410. Lee, B., et al., 2008. Conformal Al2O3 dielectric layer deposited by atomic layer deposition for graphene-based nanoelectronics. Appl. Phys. Lett. 92, 203102. Zou, K., et al., 2010. Deposition of high-quality HfO2 on graphene and the effect of remote oxide phonon scattering. Phys. Rev. Lett. 105, 126601.

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6.1 Introduction The current state of research indicates that solution-based approach, involving chemical reduction of graphite oxide (GO), yielding reduced graphene oxide (rGO) or chemically modified graphene, is simple and has the advantages of being scalable, rapid, and cost effective. For different practical applications, researchers modified the structure and chemical composition of graphene two-dimensional (2D) material by introducing different oxygen-containing functional groups using various physical and chemical routes (Stankovich et al., 2007). Depending on the amount of oxygen-containing functional groups and the ratio of sp2/sp3 hybridized carbon atoms, the newly formed products are known as graphene oxide (GO) or reduced GO (rGO). The decoration of the graphene structure with the functional groups helps to introduce a band gap in this zero band gap material. This change also brings some other significant alternation in its properties such as giving rise to its mechanical strength (Dikin et al., 2007), molecular-level chemical sensing (Robinson et al., 2008), and solubility in a variety of solvents (Stankovich et al., 2006). However, the processing of graphene for the functionalization purpose brings imperfections and disorderness in the lattice structure, which deteriorates its excellent electrical and optical properties. The amount of degradation in optoelectronic properties increases with an increasing number of processing steps. Various growth techniques and investigation of different properties of rGO have been reported, but detailed studies correlating the structure with magnetic and electrical properties are lacking. Therefore an alternative method that can comprehend graphene-based thin film without the need for a multistep functionalization process would be very useful for high-throughput GO/rGO preparation to investigate the structural, optical, electronic, and magnetic properties. Another interesting challenge is to understand the magnetic property of graphene-based materials even though they are nominally known as nonmagnetic materials. Exploring the intrinsic magnetism of graphene has been a long standing interest, and the magnetic graphene breakthrough could lead to superfast, superefficient electronic devices based on spintronics, in which the magnetic properties of a material as well as its electrical charge are manipulated. Many researchers are inspired to study the magnetoresistance (MR) property in the graphene-based materials. The applications of graphene, GO, and rGO in diverse fields; a full understanding of the charge carrier transport mechanism and the MR effect is yet to be achieved; and more research studies could widen the potential applications of these Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials. DOI: https://doi.org/10.1016/B978-0-12-817680-1.00006-8 © 2020 Elsevier Inc. All rights reserved.

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carbon-based materials in the field of high-sensitive magneto-sensors. The mechanism of magnetism in graphene and related materials, even in the absence of d and f electrons, has gained large interest in recent years (Boukhvalov et al., 2011; Yazyev et al., 2008; Tombros et al., 2007; Kou et al., 2011; Li et al., 2012). Carbon-based materials are considered to be very promising for spintronic applications owing to their weak spin-orbit coupling and their potential to have a long spin lifetime (Feng et al., 2013). Symmetry breaking at the edges of the sheet, defects/vacancies, and the substitution of atoms and hydrogen chemisorption are widely accepted scenarios to elucidate the origin of magnetism in graphene and related materials (Yazyev et al., 2007, 2008; Banhart et al., 2011; Peres et al., 2005; Singh et al., 2009; Akhukov et al., 2012; Kobayashi et al., 2006; Wimmer et al., 2010; Chen et al., 2013; Zhou et al., 2009). A large numbers of reports have suggested that oxygen-containing [carbonyl (C 5 O), carboxyl (
COOH), epoxy (C
O
C), etc.] and/or hydroxyl (
OH) groups are the origin of magnetism in graphene and related materials (Boukhvalov, 2010, 2013; Ghaderi et al., 2010; Santos et al., 2012; Wang et al., 2011a,b; Tang et al., 2014, 2015; Ray et al., 2014). Boukhvalov et al. (2011) suggested that the presence of hydroxyl clusters favors magnetism in graphene and proposed that the most stable magnetic configuration in graphene sheets involves the high-spin hydroxyl groups that are formed on the top of wrinkles or ripples (Boukhvalov et al., 2011). Santos et al. (2012) used density functional theory (DFT) to calculate the local spin moments of the carboxyl and hydroxyl groups that are adsorbed on the surface of graphene are 1 and 0.56 μB, respectively. Wang et al. (2011a,b) also used DFT calculations to reveal that the hydroxyl group is mostly responsible for ferromagnetism in GO. The presence of two hydroxyl groups bound to nonneighboring carbon atoms that are separated by one carbon atom favors the magnetic moment in GO. However, Bagani et al. (2014a,b) presented opposing arguments for various magnetism between GO and rGO, the density of wrinkles in the GO sheet decreased upon chemical reduction at high temperature (600 C) owing to the removal of many epoxy groups, increasing the number of zigzag edges/ edge states, causing rGO to have greater magnetism than GO. The increase in magnetic moment is due to the increase in the number of zigzag edges/edge states after annealing of GO, which are stable with the same spin to minimize the Coulomb-repulsion energy. The role of oxygen-containing and hydroxyl groups in inducing magnetization in GO and rGO sheets remains a matter of controversy, and no spatially resolved experimental measurement to compare chemical states (or oxygen-containing and hydroxyl groups) between wrinkle and flat regions before and after chemical reduction have been conducted. Specifically, no measurement has provided any clear evidence concerning whether the high-spin hydroxyl clusters (or oxygen-containing groups) are truly responsible for the high magnetization on the top of wrinkles on GO sheets, or whether the number of oxygen-containing and hydroxyl groups at the wrinkle and/or flat regions can be reduced by the reduction process, therefore, either to enhance or to reduce the magnetic moment in GO or rGO. In several studies, researchers have observed the magnetic properties in GO at low temperatures (Zhou et al., 2011a,b; Chen et al., 2012a,b), and in very few studies, the room temperature MR effect has been observed (Qin et al., 2010; Muchharla et al., 2014). The lowtemperature negative and positive MR effects have been explained by weak localization and

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weak antilocalization effects, respectively (Zhou et al., 2011a,b). On the other the hand, hightemperature MR effects are attributed to the vacancy and disorder-induced magnetic moments (Muchharla et al., 2014), the orbital parity mismatching between the σ and πT bands, and the enhanced exchange splitting due to the dangling bonds (Qin et al., 2010). In most of the MR studies on rGO samples, researchers have reported negative MR with very small intensities, which is unfavorable for practical applications (Muchharla et al., 2014; Wang et al., 2011a,b). Only negative MR at low temperature has been observed in chemically reduced rGO with the application of a magnetic field up to 4 T (Wang et al., 2011a,b). Although another group reported a negative MR in rGO films at room temperature, the obtained MR value is very low (8.6% under 0.5 T) (Peng et al., 2016). Vianelli et al. (2015) have observed both positive and negative MR (up to 60% at 2K under 7 T magnetic field) in rGO films. However, the crossover in the MR has only been observed below the temperature of liquid He. This requirement prohibits the use of such magnetic rGO in practical devices that operate at room temperature.

6.2 Magnetization of graphene oxide The long spin diffusion length makes graphene very attractive for novel spintronic devices, and thus has triggered a quest for integrating the charge and spin degrees of freedom. Ideal graphene is intrinsically nonmagnetic, due to a delocalized π bonding network. Therefore synthesis of ferromagnetic graphene or its derivatives with high magnetization is urgent due to both fundamental and technological importance. Despite these major developments, graphene’s magnetization in most cases is rather low, as defect-induced magnetization relies primarily on the creation of magnetic moments at edge sites of the graphene sheet, where ad-atom clustering and vacancy reconstruction, as well as thermal stability and structural integrity issues, impede a continuous increase of the underlying spin density (Nair et al., 2012). On the other hand the rich defect structure of GO, partly retained after reduction to rGO, offered an appealing alternative for boosting the density of magnetic moments in graphene-related materials. A significant increase of the magnetization in the range of 0.5
1.66 emu/g has been accordingly reported for N-doped rGO and GO (Liu et al., 2013a,b). A further increase of the magnetization for GO-based materials to  2.4 emu/g has been realized based on the formation of sp3-type defects on the basal plane of the graphene sheet, circumventing limitations of edge-type defect magnetism (Tang et al., 2015). Specifically, high magnetization has been reported for rGO by the introduction of hydroxylinduced magnetic moments in GO whose density can be tailored by annealing at high temperatures. In that case, as well as in the case of ultrasonic exfoliated graphene laminates (Sepioni et al., 2010) and pristine GO (Tang et al., 2014), the presence of high spin states (S 5 2 and 5/2) was derived from the analysis of the isothermal magnetization data, indicative of significant deviations from the spin-half paramagnetism induced by point defects in graphene materials (Nair et al., 2012). The ensuing large magnetic moments were earlier predicted by theoretical calculations pointing to the stabilization of seven OH-group clusters

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

on the basal graphene plane with a high spin S 5 5/2 ground state (Boukhvalov et al., 2011). However, the origin of magnetic moments and their relation to oxygen-containing and hydroxyl groups in GO and rGO sheets remained elusive. Recent work on photothermally reduced GO indicated that shows magnetic behavior upon reduction of GO to rGO is mainly associated with C 2p(σ )-derived states involving edge defects/vacancies rather than C 2p (π ) states related to the oxygen functional groups (Bagani et al., 2014a,b; Wang et al., 2015). The inevitable scatter on the magnetic response of GO reflected the diversity of oxidation conditions and the concomitant variation of its defective structure. Experimental evidence for the presence of high spin magnetic clusters in GO-based magnetic materials by spin-sensitive techniques such as electron spin resonance (ESR) would be valuable for the elucidation and tailoring of their magnetic properties (Singamaneni et al., 2014). Diamantopoulou et al. (2017) studied the evolution of magnetism for graphene oxide (GO) and chemically rGO by means of static magnetization and ESR spectroscopy. They have found that the strong paramagnetism with a saturation magnetization of B1.2 emu/g and weak antiferromagnetic interactions were identified in pristine GO. Apart from spin-half defect centers, ESR spectroscopy results shows the excitation of high spin states, consistently with the high spin (S 5 2) magnetic moments derived from the magnetization analysis, corroborating the formation of spatially “isolated” magnetic clusters in GO. A marked reduction of GO’s magnetization (B0.17 emu/g) along with an appreciable rise of diamagnetism (22.4 3 106 emu/g Oe) was detected after chemical reduction by sodium borohydride, reflecting the drastic removal of paramagnetic defects and the concomitant growth of sp2 domains in rGO. Sarkar et al. (2017) studied the magnetic properties of graphite oxide and rGO. Fig. 6
1A and C shows the MZFC and MFC dc magnetization as a function of temperature under 1 kOe magnetic field for the graphite oxide and rGO, respectively. A diamagnetic behavior (negative magnetization) has been found over the temperature range of 300K
50K. However, a clear indication of the presence of a paramagnetic signal (positive magnetization) has been observed below B50K where the magnetization increases steadily with decreasing temperature. No branching between the MZFC versus T and MFC versus T curves has been observed. The observed temperature dependence of magnetization obeys the Curie law indicating a true paramagnetic behavior of the sample down to the lowest measured temperature of 1.6K. The magnetization curves as a function of magnetic field at 1.6K and 5K for the graphite oxide and rGO are shown in Fig. 6
1B and D. A significant diamagnetic contribution is quite evident from the M(H) curves. Fig. 6
1D shows the M(H) curves after correcting for the diamagnetic contribution as obtained from the M(H) study at room temperature. The s-shaped nature of the M(H) curves in Fig. 6
1D, with a tendency toward its saturation at 1.6K, is evident. Besides, no hysteresis has been observed down to 1.6K. The observed M(H) behavior is, therefore, consistent with the paramagnetic behavior found in Fig. 6
1A. However, it is interesting to note here that a moderate magnetic field of 50 kOe is nearly adequate to saturate the paramagnetic moments of the graphite oxide. In the case of rGO a significantly enhanced signal is observed. For GO, at the lowest temperature of 1.6K the observed maximum magnetization is B0.075 emu/g, whereas it is 0.17 emu/g for the rGO.

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FIGURE 6–1 (A) ZFC (solid circles) and FC (open triangles): magnetization as a function of temperature under 1 kOe magnetic field for graphite oxide. (B) The dc magnetization at 1.6K and 5K for graphite oxide after correcting for the diamagnetic contribution. The inset enlarges the low-field region of the M(H) curves. (C) ZFC (solid circles) and FC (open triangles): magnetization as a function of temperature under 1 kOe magnetic field for reduced graphene oxide. (D) The dc magnetization at 1.6K and 5K for reduced graphene oxide after correcting the diamagnetic contribution. The inset enlarges the low-field region of the M(H) curves. FC, field cooling; ZFC, zero field cooling. By courtesy of Sarkar, S.K., et al., 2017. Investigation of graphite oxide and reduced graphene oxide: magnetic properties revisited. Asian J. Mater. Chem. 1 (3
4), 66
74.

Sarkar et al. (2017) also studied the magnetization as a function of reduced field (H/T) for both GO and rGO as shown in Fig. 6
2. The observed magnetization is fitted using the Brillouin function: M 5 NgJμB

     2J 1 1 ð2J 1 1Þx 1 x ctnh 2 ctnh 2J 2J 2J 2J

where x 5 gJμB H/kBT is the ratio of the Zeeman energy of the magnetic moment in the effective field Hc 5 H 1 Hm, where Hm is the molecular field proportional to magnetic moment M, kB is the Boltzmann constant, g is the Lande g factor, J is the angular momentum number, and N is the number of present magnetic moments. Assuming g 5 2, the best fit with the Brillouin function is obtained with J 5 3/2 for both GO and rGO. The values of N are derived to be 1.91 3 1018 and 5.80 3 1018/g for the GO and rGO, respectively. These lower values of N (induced spin centers) are typical of a spin system with a weak paramagnetic

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 6–2 Magnetization as a function of H/T at 1.6K for (A) graphite oxide and (B) reduced graphene oxide samples. The solid lines are fits with the Brillouin function. By courtesy of Sarkar, S.K., et al., 2017. Investigation of graphite oxide and reduced graphene oxide: magnetic properties revisited. Asian J. Mater. Chem. 1 (3
4), 66
74.

behavior, while representative materials with a strong paramagnetic behavior would have N values in the range of 1021
1022. The derived value of g(52) indicates that spin only moments (i.e., without any orbital moments) contribute to the magnetism of these GO/rGO. As mentioned above, Wang et al. (2015) studied the magnetization of graphite oxide, photothermal moderately rGO (M-rGO) and heavily rGO (H-rGO) using the synchrotronbased X-ray microscopic and spectroscopic techniques, including scanning transmission X-ray microscopy (STXM), X-ray absorption near-edge structure (XANES) spectroscopy, valence band photoemission spectroscopy (VB-PES), and X-ray magnetic circular dichroism (XMCD). Element-specific XMCD provides evidence of ferromagnetic behavior in graphite oxide. The results of C K-edge STXM
XANES provide clear evidence that the higher number of C 2p(σ )-derived defect/vacancies states, rather than of the C 2p(π ) states, is bound with oxygen-containing and/or hydroxyl groups on the graphite oxide surface. This feature is related to the change of magnetic behavior of ferromagnetic graphite oxide to that of paramagnetic M-rGO oxide and H-rGO. Fig. 6
3 shows the normalized magnetization
hysteresis (M
H) curves of GO, M-rGO, and H-rGO at room temperature after the diamagnetic contribution from the Si substrate is subtracted (Wang et al., 2015). The inset plots the M
H curves of GO, M-rGO, and H-rGO and the Si substrate before the diamagnetic Si contribution was subtracted, revealing that the ferromagnetic coercivity and saturated magnetic field of GO were B150 and 3000 Oe,

Chapter 6 • Magnetism and spintronics in graphene oxide

157

1.0 GO M-rGO H-rGO

0.0 M (10–4emu/cm2)

M (H)/M(1T)

0.5

–0.5

Si GO M-rGO H-rGO

2 1 0 –1 –2

–1.0 –1.0

–1

–0.5

0.0

0

0.5

1

1.0

H (T) FIGURE 6–3 Room-temperature [M(H)/M(1T)]
H curves of GO, M-rGO, and H-rGO after subtraction of diamagnetic background that arises from silicon substrate. Inset M
H curves (without background subtraction) of GO, M-rGO, and H-rGO. H-rGO, Heavily reduced graphene oxide; M
H, magnetization
hysteresis; M-rGO, moderately reduced graphene oxide. Wang, Y. F., et al., 2015. Visualizing chemical states and defects induced magnetism of graphene oxide by spatially resolved-X-ray microscopy and spectroscopy. Sci. Rep. 5, 15439. By courtesy of NPG publication.

respectively. The ferromagnetic behavior of GO gradually weakens as the photothermal (PT)reduction proceeds, exhibiting a paramagnetic behavior for M-rGO and to an even greater extent for H-rGO, although the GO is typically considered to be as being spin-half paramagnetic (Liu et al., 2013a,b). The variation of M
H curves in Fig. 6
3, as described in the literature (Boukhvalov, 2010, 2013; Ghaderi et al., 2010; Santos et al., 2012; Wang et al., 2011a,b; Tang et al., 2015; Ray et al., 2014), if the ferromagnetism of the GO is dominated by the oxygen-containing and/or hydroxyl groups in the GO sheets, especially in the wrinkle regions, then the PT-reduction of GO to M-rGO and H-rGO, must have removed a rising proportion of oxygen-containing and/or hydroxyl groups from GO sheets, so the M-rGO and HrGO sheets, with fewer oxygen-containing and/or hydroxyl groups, are paramagnetic. If this argument is true, the magnetism in GO is primarily caused by the C that is π-bonded with oxygen-containing and/or hydroxyl groups, changing the ferromagnetic behavior of GO into the paramagnetic behavior of M-rGO and H-rGO, as the proportion of oxygen-containing and/or hydroxyl groups varies with the degree of PT-reduction. To understand better the origin of ferromagnetic behavior in GO and its gradually giving way to a paramagnetic behavior upon PT-reduction to form M-rGO and H-rGO, the STXM
XANES, VB-PES, and XMCD are used by Wang et al. (2015). Fig. 6
4 presents optical density (OD) images (panel I), C K-edge STXM stack mappings (panel II), and decomposed STXM mappings (panels III
VI) of the surfaces of randomly selected single sheets of GO, M-rGO, and H-rGO. The bright areas in the OD images represent thick regions, dim areas represent thin regions, and gray areas represent the regions of intermediate thickness, as observed in GO [panel I(a)], M-rGO [panel I(b)], and H-rGO [panel I(c)], respectively. Based on the OD, the selected regions of GO, M-rGO, and H-rGO

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 6–4 OD images and corresponding stack mapping from STXM images of GO, M-rGO, and H-rGO are shown in panels I and II. Panels III
VI present stack mappings from C K-edge STXM images of GO, M-rGO, and H-rGO, which are decomposed into blue, yellow, red, and green regions that are associated with the different thicknesses of samples. Spectra of all samples typically present background (blue), flat (yellow), medium (red), and wrinkle (green) regions. H-rGO, Heavily reduced graphene oxide; M-rGO, moderately reduced graphene oxide; OD, optical density; STXM, scanning transmission X-ray microscopy. Ray, et al., 2015. Sci. Rep. 5, 15439. By courtesy of NPG publication.

are typically attributed to wrinkle, medium, and flat regions of the GO, M-rGO, and H-rGO sheets. As presented in panels I(a)
I(c), the brightest region of H-rGO has a higher average OD (1.29) than does GO (0.93) or M-rGO (0.64), suggesting that the thickest regions were preferably formed in the H-rGO sheets, even though they were the most heavily reduced. The various colors shown in the C K-edge STXM stack mapping in panels II(a)
II(c) of Fig. 6
4 correspond to the randomly varying thickness of GO, M-rGO, and H-rGO. The decomposed STXM stack mappings (panels III
VI) are shown in blue (background), yellow (flat), red (medium), and green (wrinkle), which relate directly to the regions of the samples with various thicknesses (Singh et al., 2014; Zhou et al., 2011a,b). The maps were divided into four regions by principle component analysis (PCA) for cluster analysis, based on spectroscopic differences. The PCA spectrum of each region is the average from all image pixels in that region. The background is shown in blue; the OD or absorbance of the background is nearly zero, corresponding to the near-null intensity of the C K-edge STXM spectrum. A more intense average spectrum generally indicates a thicker sample, with the thickness increasing

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159

from flat through medium to wrinkle regions. All chemical species in these regions can affect thickness, and the thick regions are typically attributed to wrinkle regions of GO sheets. As shown in panels IV
VI in Fig. 6
4, the flat, medium, and wrinkle regions are present at random locations on the surface of GO, M-rGO, and H-rGO. GO cannot be formed with a perfectly flat geometry because the wrinkle geometry of GO sheets is generally more stable than their flat geometry (Schniepp et al., 2006; Shen et al., 2014), so the formation of wrinkle regions of GO sheets is simply observed in both GO and rGOs. More details concerning STXM
XANES measurement can be found elsewhere (Wang et al., 2015). If the magnetism in GO is mainly determined by the presence of oxygen-containing and/or hydroxyl groups, the intensity of the corresponding features in the C K-edge STXM
XANES spectra should be significantly affected by the transformation from ferromagnetic GO to paramagnetic M-rGO and H-rGO. But there is no such difference observed in the STXM
XANES spectra (Wang et al., 2015). The intensities of features associated with oxygen-containing and hydroxyl groups in the wrinkle regions of H-rGO surfaces are close to those of GO after heavy PT-reduction of GO, whereas those in flat regions are higher than those of GO, suggesting that the presence of the C 2p(π ) states that are bound with oxygen-containing and hydroxyl groups may not be the main cause of the ferromagnetic behavior in GO. However, the intensities of features associated with the C 2p(σ )-derived states and the features decrease with the PT-reduction of GO to M-rGO and then to H-rGO, implying the correlation between the numbers of C 2p(σ )-derived states and the transformation of ferromagnetic GO into paramagnetic M-rGO and then H-rGO. Lee et al. (2015) engineered magnetism in GO with various chemical groups, namely, epoxy, ketone, hydroxyl, and C
O groups on GO surface. Destroying the epoxy group with heat treatment or chemical treatment diminishes the magnetism in the GO. Liu et al. (2013a, b) studied the magnetic properties of N-doped/undoped graphene-oxide (GO) for potential applications in spintronic devices. The magnetic behaviors, as shown in Fig. 6
5A, are the dependence of susceptibility χ 5 M/H on temperature T and it fits well with the Curie law χ 5 C/T. Inset is the corresponding 1/χ
T curve, which demonstrates a linear, purely Curielike paramagnetic behavior, a solid evidence of the existence of localized magnetic moment. As shown in Fig. 6
5B, no significant ferromagnetic signal is observed even at 2K. Liu et al. (2013a,b) fitted the M
H curve using the Brillouin function. M 5 Ms

     2J 1 1 2J 1 1 1 x coth x 2 coth 2J 2J 2J 2J

(6.1)

where x 5 ðgJμB H=kB TÞ, MS 5 NgJμB , kB is the Boltzmann constant, N is the number of present magnetic moments, J the angular momentum number, and g is the Landau factor which is assumed to be 2. As shown by the fitting curve, the Brillouin function provides good fits for J 5 S 5 1/2. The intrinsic spin-half paramagnetism agrees well with the existing theories for the contributions of point defects such as vacancies, ad-atoms, and edges (Nair et al., 2012; Krasheninnikov et al., 2009; Yazyev, 2008). Also by fitting the curve, Ms can be obtained which is 0.11 emu/g,

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 6–5 (A) Typical χ
T curve measured from 2K to 300K under the applied field H 5 3 kOe. Inset 1/χ
T curve. (B) M
H curve measured at 2K. (C) Typical M
H curve measured at 2K. Inset magnetization curve. (D) M
T curve of NGO measured from 2K to 300K under the applied field H 5 500 Oe. (E) M
H curves of NGO measured at 80K and 110K. Inset magnetization curves. (F) M
H curve of NGO at 2K. Inset ferromagnetic mass magnetization. M
H, magnetization
hysteresis; NGO, nitrogenated graphene oxide. Liu et al., 2013. Realization of ferromagnetic graphene oxide with high magnetization by doping graphene oxide with nitrogen. Sci. Rep. 3, 2566. Copyright NPG Publishing. Reproduced with permission.

similar to the value reported (Sepioni et al., 2010). Liu et al. (2013a,b) also carried out the magnetic measurements for nitrogenated graphene oxide (NGO), where no positive magnetic signal is observed, and only purely diamagnetism can be observed in both GO and NGO at 300K. Fig. 6
5C shows the typical mass magnetization (M
H) curve of NGO measured at 2K. One can find that the coercive field (Hc) and remnant magnetization (Mr) are 160 Oe and 0.039 emu/g (inset of Fig. 6
5C), a solid evidence for ferromagnetism. Subsequently, Liu et al. (2013a,b) performed the M
T measurement of NGO and found a clear Tc at c.100.2K (Fig. 6
5D), implying that the Tc is above liquid N2 temperature of 77K. It also can be confirmed by the two M
H curves measured at 80K and 110K (Fig. 6
5E). Liu et al. (2013a,b) found that the M
T curve of NGO shows typically paramagnetic behavior at low temperature below 8K (Fig. 6
5E). Namely, the magnetism at 2K is composed of two parts: paramagnetism and ferromagnetism, which can be expressed as Mtotal 5 Mpara 1 Mferro. Considering the fact that ferromagnetic mass magnetization can saturate at a high applied field, one have to set paramagnetic M0 of the graphen oxide as 1.39 emu/g and J as 1.11 to fit the curve by Brillouin function (Fig. 6
5F). As reported, J can be 0.5, 1, or 2.5, etc., which generally should be the integral multiple of 0.5 and corresponds to magnetically coupled unpaired electrons (Nair et al., 2012; Sepioni et al., 2010; Ney et al., 2011). However, the magnetic

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structure of our sample may be uneven, and different J may exist in different sheets in NGO. Theoretically, the N atom can induce a net magnetic moments of 1.8, 0.95, or 0.67 μB for specific structures (Li et al., 2009; Dai et al., 2010). It also may be the reason that J 5 1.11 in NGO. By subtracting the paramagnetic signal from the observed data, one can obtain the remaining ferromagnetic moment. The magnetization is approximately saturated at 6 kOe (inset of Fig. 6
5F), and the saturation magnetization of ferromagnetic signal is ca. 0.27 emu/g at 2K. Sarma et al. (2017) studied the magnetization of pristine GO and nitrogen functionalized GO. In this study, the M
H curves for the pristine GO and GO:Nx were evaluated from 210 kOe , H , 10 kOe at room temperatures (300K). Fig. 6
6 shows the measured M
H loops for the GO:N1.30 displaying the strongest ferromagnetic behavior with a saturation magnetization (Ms) value of 5.3 3 1023 emu/g and a coercivity (Hc) of 10 Oe. On the other hand the and GO:N0.65 shows the Ms of 4.9 3 1024 emu/g and an Hc value of 19 Oe, higher than that of pristine GO having Ms 5 6.8 3 1025 emu/g with coercivity (Hc) of 38 Oe. The origins of magnetization in N-doped GO/graphene are still controversial wherein ferromagnetism in N-doped GO has been ascribed to pyrrolic groups which can provide a net magnetic moment of 0.95 μB/N atom (Miao et al., 2016). In contrast, Ito et al. (2015) have observed that the presence of pyrrolic groups lead to a reduction in the magnetization values. In this case, we have found that the GO:N1.30 have higher graphitic in nature and is responsible for the higher magnetization. Besides the N content, GO:N1.30 become more

FIGURE 6–6 Magnetic hysteresis loops obtained for GO, GO:N0.65 and GO:N1.30 at 300K. Sarma, S., et al., 2017. Electronic and magnetic properties of nitrogen functionalized graphene oxide. Diamond Relat. Mater. 79, 1
6. Copyright Elsevier Publishing. Reproduced with permission.

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disordered, that is, the enhancement of sp2-content in the GO structure and hence magnetization is enhanced. The magnetization in sp2-carbon structures also originate from defects as well. Later, Ghosh et al. (2018) studied the magnetization of “GO-Nx” and “GO-Nx:Fe” composite measuring the M
H loops within the magnetic field range 210,000 Oe , H , 1 10,000 Oe at room temperatures (300K) as shown in Fig. 6
7. Ghosh et al. (2018) found that the magnetization of “GO-Nx” does not only depend on the content of carbon and/or nitrogen but also strongly depends on different N-precursor. Fig. 6
7 shows that the highest magnetization was obtained in “GO-Nx” synthesized with C6H12N4-precursor, with the saturation magnetization (Ms) value of 0.0057 emu/g and a coercivity (Hc) of 53 Oe; whereas the lowest magnetization was found (Ms 5 0.0031 emu/g and Hc 5 31 Oe) in “GO-Nx” synthesized with CH4N2O-precursor. This magnetic behavior also depends on the formation of different N-functional groups, such as pyridine-N, pyrrolic-N, graphitic-N, and N-oxide, different carbon
carbon bond and carbon
oxygen functional groups, such as
OH,
COOH, C
O/ C5O that are attached with the “GO-Nx” matrix during the formation of “GO-Nx” (Sun et al., 2014). However, the origin of this magnetization in N-doped GO is still controversial. There are some reports on ferromagnetism in N-doped GO that has been ascribed due to pyrrolic group which can provide a net magnetic moment of 0.95 μB/N atom (Miao et al., 2016). In contrast, Ito et al. (2015) have observed that the presence of pyrrolic group leads to a reduction in the net magnetization value of nitrogenated GO. In our present case, we have found that “GO-Nx” synthesized with C6H12N4-precursor have higher graphitic content (sp2) and exhibits higher magnetization. This enhancement of sp2-content makes this “GO-Nx” structure more disordered and hence magnetization is enhanced significantly. The magnetization in sp2-carbon structures also originates from the defects as well. A similar phenomenon was also observed by Sarma et al. (2017), where XPS and microstructural Raman study shows that the sp2-content as well as ID/IG ratios (microstructural sp2/sp3 ratio and/or their defects)

FIGURE 6–7 Magnetization M
H loops of NH4OH-based “GO-Nx”, C6H12N4-based “GO-Nx”, C2H3N-based “GO-Nx,” and CH4N2O-based “GO-Nx”. M
H, magnetization
hysteresis. Ghosh, et al., 2018. Tuning of magnetic behaviour in nitrogenated graphene oxide functionalized with iron oxide. Diamond Relat. Mater. 89, 35
42. Copyright Elsevier Publishing. Reproduced with permission.

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increased with increase of nitrogen concentration in the GO-Nx structure, as a result the magnetization enhances significantly. It is clearly seen that the highest magnetization was observed in C6H12N4-precursor based “GO-Nx”, which is due to the higher content of graphitic carbon as well as different graphitic nitrogen bonds present in the “GO-Nx” structure. Majumder et al. (2018) reported about the enhancement of magnetization and blocking temperature of the Fe3O4-rGO nanocomposite, which was formed by embedding Fe3O4 nanoparticles onto rGO sheets. The magnetic properties of rGO and Fe3O4-rGO nanocomposite are described by the magnetization versus temperature (M
T) at 100 Oe magnetic field, shown in Fig. 6
8A and the magnetization versus applied magnetic field or M
H plots at temperatures 10K, 100K, and 300K shown in Fig. 6
8B
D. The M
T plots of rGO and Fe3O4-rGO nanocomposite (Fig. 6
8A) suggest that the magnetic moment of Fe3O4-rGO nanocomposite is greater than rGO. This implies that GO encapsulation on bare Fe3O4 NPs enhances its magnetic moment. The blocking temperature (the maximum of the

FIGURE 6–8 (A) M
T plot at 100 Oe magnetic field of both samples. Magnetization versus applied magnetic field (M
H) plots of both samples at different temperatures (B) 10K, (C) 100K, and (D) 300K. M
H, Magnetization
hysteresis; M
T, magnetization versus temperature. Majumder, S., et al., 2018. Magnetization enhancement of Fe3O4 by attaching onto graphene oxide: an interfacial effect. J. Phys. Chem. C 122, 21356
21365. Copyright ACS Publications. Reproduced with permission.

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

magnetization vs temperature curve) of Fe3O4-rGO nanocomposite is greater than that of GO. Thus encapsulation of Fe3O4-rGO nanocomposite results in an enhancement of the blocking temperature. Moreover, Majumder et al. (2018) have observed that for all temperatures (10K, 100K, and 300K), saturation magnetization (Ms) of Fe3O4-rGO nanocomposite is greater than that of GO. Magnetite (Fe3O4) is a member of the spinel family of minerals; crystallographically, magnetite takes a cubic inverse spinel form and can be represented as (Fe[1])A[Fe[2], Fe[3]]BO4, where Fe[1] having four nearest neighbor O22 ions occupies tetrahedral interstices and Fe[2] and Fe[3] having six nearest neighbor O22 ions fill the octahedral interstices. The oxygen ions form a close-packed cubic lattice with the iron ions located at interstices between the oxygen ions. A sites are populated by Fe31 ions only and B sites by both Fe31 and Fe21 ions, with twice as many B sites occupied than the A sites. Thus Fe[1] is in Fe31 (d5) with the spin moment S 5 5/2 at tetrahedral interstices, and Fe[2] and Fe[3] at the octahedral interstices are in Fe31 (d5) with spin moment S 5 5/2 and Fe31 (d6) with spin moment S 5 2. It is wellknown that the strongest Fe
Fe super exchange antiferromagnetic interactions are between Fe[1]
Fe[2] sites and Fe[1]
Fe[3] sites, whereas for Fe[2]
Fe[3] sites, because of unequal number of electrons at the nearest neighboring sites, d6
d5 forms a double-exchange bond giving rise to ferromagnetic interactions among them. So, these d electrons hop around ferromagnetically, enhancing the electrical conductivity leading to metallicity beyond the Verwey temperature, Tv (Imada et al., 1998). Below the Verwey temperature, Tv, the material is insulating where the extra d electron at the d6 site stops moving and a charge ordered d6
d5 insulating state emerges. So, the formula unit of Fe3O4 gives a net spin of S 5 2, and its magnetic moment is 4 μB. It is well-known that GO mainly consists of functional groups such as epoxy (C
O
C), carbonyl (C 5 O), hydroxyl (OH), or carboxyl (COOH) on the basal plane and at the edges of the graphene lattice. Here, we can argue out that there are several ways by which Fe3O4 can be bonded with GO, where only the low spin state Fe21(d6, S 5 2) will bond to form a high spin state Fe31(d5, S 5 5/2). In all these cases, sp in value at the Fe31 ion increases from S 5 2 to S 5 5/2. Majumder et al. (2018) have estimated that the increase in the magnetic moment/magnetization per formula unit is (5/2
2) 3 100/2  25%. However, only Fe21 (d6) residing on the surface of the particles can form singlet bonds; if this takes place, this increase will be less than 25%. In Fe3O4 the lattice spacings are a 5 b 5 c 5 6 Å and the unit cell volume is Ba3. The total number of Fe3O4 units in an NP of radius r is NV 5 4πr3/(3a3), and the total number of Fe3O4 units on the surface will be NS 5 4πr2a/(a3); hence, if we assume r  50 Å, then NS/NV  3a/r  3 3 6/50  1/3. Therefore the increase in the magnetic moment/magnetization will be  25% 3 1/3  8%, which is of the same order of magnitude as obtained experimentally (B4.3%). Finally, the magnetic structure both above and below Tv is the same here, that is, for d5, S 5 5/2 indicated as (n), and for d6, S 5 2 indicated as (m); then the magnetic ordering at the lowest temperature is (n)A [l,k]B O4, which determines the net magnetization. These observed phenomena of enhancement of blocking temperature and magnetization can be attributed to the increase in the magnetic domain size because of the added interface of rGO and Fe3O4. Teo Peik-See et al. (2014) also studied the magnetization of rGO/iron oxide

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165

nanocomposite materials for environmental remediation. In this study, the bare Fe3O4 nanoparticles and the rGO/Fe3O4 nanocomposites showed typical S-like curve M
H loops with no coercivity, inferring that they exhibit superparamagnetism, while their magnetization behaviors were removed in the absence of the applied magnetic field. The saturation magnetization (Ms) of the rGO/Fe3O4 nanocomposites increased from 1.63 to 30.30 emu/g with an increase in the content of Fe3O4 nanoparticles in the rGO sheets. The saturated magnetization value observed for the pristine Fe3O4 nanoparticles was 58.70 emu/g, which was higher than that of the magnetic rGO/Fe3O4 nanocomposites. This can be attributed to the presence of graphene in the nanocomposites (Ren et al., 2011).

6.3 In-plane and out-of-plane magnetization/magnetic anisotropy Ning et al. (2017) studied the magnetic anisotropy of graphene oxide membranes. In this study, they have found that the improvement of the perpendicular magnetic anisotropy (PMA) of CoFeB thin films by applying a coating of GO membranes. Ning et al. (2017) observe that the PMA of the CoFeB/MgAl
O stacks is strongly enhanced by the coating of GO membranes and even reaches 0.6 mJ/m2 at room temperature after an annealing process. The critical thickness of the membrane-coated CoFeB for switching the magnetization from the out-of-plane to the in-plane axis exceeds 1.6 nm. The contribution of the GO membranes to the magnetic anisotropy energy is due to changes in the hybridization of 3d orbitals, varying the location of the C atomic layer with Co changes the contribution of the Co
C stacks to PMA. Thus the large PMA achieved with GO membranes can be attributed to the orbital hybridization of the C and O atoms with the Co orbitals. These results provide a comprehensive understanding of the PMA and point toward opportunities to achieve multifunctional graphene-composite spintronic devices. Sepioni et al. (2010) studied the magnetic anisotropic behavior of highly oriented pyrolytic graphite (HOPG) and found that the ferromagnetisms are temperature independent between 2K and 300K temperature as shown in Fig. 6
9, therefore denoting a Curie temperature above 300K. While the signal is clearly visible and measurable in the orientation perpendicular to the graphite ĉ-axis, in the parallel orientation, it becomes over shadowed by the diamagnetic background (inset of Fig. 6
9). This is due to the large anisotropy of the diamagnetic susceptibility in graphite, which causes a diamagnetism  30 times stronger in the direction perpendicular to the graphitic planes.

6.4 Magnetoresistance of graphene oxide Haque et al. (2018a,b) investigate the magneto-resistance and magneto-transport properties of rGO at room temperature (Bhaumik et al., 2017; Haque et al., 2018a,b). The rGO films show a nonsaturating negative and positive MR at low and high fields, respectively. In this

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 6–9 ΔM(H) curves after the subtraction of the diamagnetic background for HOPG-ZYA in the orientation perpendicular to the graphite ĉ-axis at 2K, 10K and 300K. In the inset M(H) in direction parallel to the ĉ-axis is plotted at 300K in the low-field region. HOPG-ZYA, highly oriented pyrolytic graphite -high grade. Sepioni, M., et al., 2010. Limits on intrinsic magnetism in graphene. Phys. Rev. Lett. 105, 207205. Copyright APS Publications. Reproduced with permission.

study the maximum value of the MR exceeds 160% under 1.2 T magnetic field at room temperature. Room temperature MR measurements were performed by a four-probe measurement system where an external magnetic field was applied perpendicular to the current and main axis of the rGO device. The current was injected using two adjacent electrodes, and the rest two probes were used to measure the voltage. Arbitrary adjacent pairs can be chosen with equal performance as the device is fourfold symmetric. A schematic of this measurement setup is shown in the Fig. 6
10F and G. In the measurement setup, resistance is defined as follows: R5

V1 2 V2 I

(6.2)

Fig. 6
10A
D shows the change in resistance over a range of magnetic field from 21.5 to 1.5 T for different input currents. The MR was calculated using the following equation: MR% 5

RB 2 RB0 x 100 RB0

(6.3)

It shows that the rGO thin-film devices have a negative MR effect when the applied magnetic field is low, and after a certain range (B0.5 T), they exhibit a positive MR effect. That are revealed with reported negative MR at the low-applied magnetic field and positive MR at the high magnetic field in graphene (Ciuk et al., 2012). All the data are symmetric and proves that the sources of the MR are consistent in both polarities of the applied magnetic field. At low temperature the source of the negative MR can be attributed to the weak localization effect. The exchange coupling between the spin of the localized electrons, and the

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167

FIGURE 6–10 R
T at (A) I 5 0.01 μA. (B) I 5 0.1 μA. (C) I 5 1 μA. (D) I 5 10 μA. (E) Positive and negative MR at different currents. (F) Schematics of the four-probe measurement system. (G) Connection of the van der Pauw technique. MR, Magnetoresistance. By courtesy Haque, A., et al., 2018a. Large magnetoresistance and electrical transport properties in reduced graphene oxide thin film. IEEE Trans. Magnets 54, 1000209.

surrounding magnetic moment is responsible for an increase in the bare binding energy of the localized charge carriers. Under the applied magnetic field the probability of the alignment of the magnetic polaron increases; hence, the bare binding energy decreases. Therefore the rate of hopping and percolation of the charge carriers increases, which causes a decrease in the resistance in the rGO film and/or may be the diffuse scattering at crystallite boundaries (Fujita et al., 1968). The observed large positive MR at a higher magnetic field could be the outcome of carrier mobility fluctuations: the nonuniform spatial distribution of carrier mobility at the higher magnetic field. Changing the amount of applied current into the rGO device affects the magnitude of the MR in the whole applied magnetic field range. The variation in the positive and negative MR with sending current is shown in Fig. 6
10E. At the maximum applied magnetic field (1.2 T) the calculated positive MR was as high as 160% for a sending current of 0.01 μA. The maximum negative MR was as high as B18% for the same current but at a low magnetic field of around 0.45 T. On the other hand, when the sending current was high (10 μA), the MR effect almost decreased to zero. With increasing sending current, both positive and negative MR effects continuously weakened. Hence, it can be inferred that the MR effect is dependent on the current density that was sent to measure the percent of MR. However, while varying the applied current, the resistance does not change at zero magnetic field. Because of only under perpendicular applied magnetic field, the movement of the free charge carriers is hindered at different extent when the density of

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the free charge carrier is varied. Peng et al. (2016) investigate the magneto-transport behaviors of high sulfur concentration rGO (HS-rGO) and low sulfur concentration rGO (LS-rGO) nanosheets. In order to investigate the magneto-transport behavior of HS-rGO and LS-rGO nanosheets, standard four-terminal nanosheet devices in a commercial physical property measurement system were used. In this case the graphene system is actively pursued in spintronics for its nontrivial sp electron magnetism and its potential for the flexible surface chemical tuning of magneto-electronic functionality. The MR of graphene can be effectively tuned under high magnetic fields at cryogenic temperatures, but it remains a challenge to achieve sensitive magnetoelectric response under ambient conditions. Peng et al. (2016) report the use of surface modulation to realize superparamagnetism in rGO with sensitive magnetic field response. The superparamagnetic rGO was obtained by a mild oxidation process to partially remove the thiol groups covalently bound to the carbon framework, which brings about large low-field negative MR at room temperature (28.6%, 500 Oe, 300K). This strategy provides a new approach for optimizing the intrinsic magnetoelectric properties of 2D materials. Fig. 6
11A and B shows the typical temperature dependent resistivity of the HS-rGO and LS-rGO nanosheets with three atomic layers under a series of external perpendicular

FIGURE 6–11 Magneto-transport properties of HS-rGO and LS-rGO. Temperature-dependent resistance under various magnetic fields at the temperature range from 270K to 300K for (A) HS-rGO and (B) LS-rGO, respectively. Magnetic-field dependence of the MR for (C) HS-rGO and (D) LS-rGO at 300K. HS-rGO, High sulfur concentration rGO; LS-rGO, low sulfur concentration rGO; MR, magnetoresistance. Peng, J., et al., 2016. Superparamagnetic reduced graphene oxide with large magnetoresistance: a surface modulation strategy. Angew. Chem. Int. Ed. 55, 3176
3180. Copyright Wiley online Publications. Reproduced with permission.

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magnetic fields ranging from zero to 1 Tesla. Different magneto-transport behaviors are observed: LS-rGO has a large negative magnetic field coefficient of resistivity, while HS-rGO has a negligible value. The field-dependent MR for HS-rGO and LS-rGO are plotted in Fig. 6
11C and D, respectively. Here, the field-dependent MR is defined as MR(H)% 5 [ρ(H) 2 ρ(0)]/ρ(0), where ρ(H) and ρ(0) are the resistivity under the magnetic field H and zero field, respectively. The MR value of HS-rGO was rather small at room temperature; however, LS-rGO exhibited large negative MR at the same temperature. For LS-rGO the external magnetic field greatly suppressed electrical resistivity under a very low magnetic field: For instance, the negative MR values were as large as 28.6% under 500 Oe and 214.9% under 1T at 300K, as depicted in Fig. 6
3D. Moreover, the MR curve of LS-rGO exhibits two distinct regions: the low-field region (within 2000 Oe) where the MR shows a rapid increase with the magnetic field, and the high-field region (above 2000 Oe), where the MR displays a much slower increase with the field. Furthermore, the coercivity field and MR value can be modulated by the oxygen-treatment time and finally both of them reached a saturated value. The rGO nanosheets with superparamagnetic properties (LS-rGO) show a large low-field negative MR effect at room temperature (28.6%, 500 Oe, 300K), representing a new graphene-based magnetoelectric system. This strategy provides a new approach for optimizing the intrinsic magnetoelectric properties of 2D materials.

6.4.1 Mechanism of magnetic behaviors GO/reduced graphene oxide In this section, it is discuss the possible mechanism for the observed magnetic behavior of the studied GO/rGO. A defect free graphene plane, that is, a single layer of graphite should be diamagnetic since graphite itself is diamagnetic. Wang et al. (2009) have reported room temperature ferromagnetism in graphene, while Matte et al. (2009) have found the presence of both ferromagnetic and antiferromagnetic features in graphene. But, Sepioni et al. (2010) have found that the graphene at any temperature down to 2K have no ferromagnetism, but a strong diamagnetism above 50K (similar to graphite) and a weak paramagnetism below 50K. Judging from these reports, it appears that the graphene, prepared by various methods, can behave differently, and these observations could be better represented by the nature of defects in these graphene. A real graphene has several defects such as topological defects (pentagons, heptagons in their combination), atomic defects (vacancies such as missing C-atoms and ad-atoms), and extended defects (zigzag edges, cracks, etc.). Further, it may be mentioned that a real graphene is also not perfectly planar; rather corrugations, ripples, wrinkles, etc., occur on it. Literature provides extensive theoretical and experimental studies on magnetic behavior of graphene based on both defects and zigzag states. Yazyev et al. (2007) have shown that a magnetic moment of about 1 μB can develop due to the defect-induced extended states for one vacancy defect or one hydrogen chemisorption defect. Yazyev (2008) has also shown that single-atom defects can induce ferromagnetism in the disordered graphene and proton-irradiated graphite. Many theoretical studies revealed that the defects could induce a magnetic moment in a small defect region of graphene. If they are coupled by Ruderman
Kittel
Kasuya
Yosida (RKYY) or any other exchange

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

interactions ferromagnetism can appear. If they are well separated so that they are not coupled, a weak or strong paramagnetism can appear. Though the mechanism of strong exchange interaction required for the magnetic moments to appear at room temperature is not clearly understood, the magnetic behavior should depend on the concentration of the defects in graphene. It appears that the concentrations of the defects in the graphene of references (Yang et al., 2009; Matte et al., 2009) are high, while it is small for the graphene studied (Sepioni et al., 2010). Applying Lieb theorem (1989) the epoxide groups cannot induce local magnetic moments, but the O atom of the hydroxyl group bonded to only one C atom of either sublattice can induce local magnetic moments. Hence, such moments, developed due to the hydroxyl groups, are to be considered along with the moments developed due to vacancy defects. If these moments are well separated so that they do not interact, paramagnetism is expected to appear, which we have observed for the GO and annealed rGO. The defects in the GO and rGO very much depend on the starting graphite material used during the preparation. Therefore the magnetic properties, originated from the defects, vary from GO/rGO to GO/rGO and are depending on synthesis process. Numerous reports have suggested that oxygen containing [functional groups, for example, carbonyl (C5O), carboxyl (
COOH), and epoxy (C
O
C)] and/or hydroxyl (
OH) groups are responsible for magnetism in graphene and related materials (Boukhvalov, 2010, 2013; Santos et al., 2012; Wang et al., 2011a,b; Tang et al., 2014, 2015; Ray et al., 2014). Boukhvalov (2010, 2013) suggested that the hydroxyl clusters favors magnetism in graphene, and the most stable magnetic configuration in graphene sheets involve the high spin hydroxyl groups that are formed on top of wrinkles or ripples. Santos et al. (2012) applied DFT to calculate the local spin moments of the carboxyl and hydroxyl groups as 1 and 0.56 μB, respectively, that are adsorbed on the surface of grapheme. Similar DFT calculation by Wang et al. (2011a,b) further revealed that the hydroxyl group is mostly responsible for ferromagnetism in graphite oxide, and they further proposed that the presence of two hydroxyl groups bound to nonneighboring carbon atoms separated by one carbon atom favors the magnetic moment in graphene oxide. Various functionalized groups (
OH,
O
,
COOH, C5O, etc.) enter into the graphene skeleton, breaking the π bond of graphene structure (Mkhoyan et al., 2009) during the oxidation process, but the exact decoration of the functionalized groups on the graphene skeleton remain uncertain. However, nuclear magnetic resonance (NMR) study of graphene oxide has shown that the carbonyl groups are located in the periphery of the graphene oxide sheet (Lerf et al., 1998; Cai et al., 2008) Therefore it may be inferred that only hydroxyl (
OH) and epoxy (
O
) groups are abundant in the interior region of the graphene sheets. After the chemical reduction process, the magnetic moment of graphene oxide is expected to decrease due to the removal of such groups. On annealing rGO the concentration of epoxy groups reduces further and the magnetic moment should decrease, but we have observed an increase in moment after annealing rGO at 600 C. This shows that apart from removing the functional groups from graphite oxide, annealing does something more. The π electrons are energetically degenerate at the zigzag edges and have highly localized edge states. These

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edge states are populated with the same spin to minimize the Coulomb-repulsion energy, leading to large moments at the zigzag edge boundary (Kumazaki et al., 2008; Bhowmick et al., 2008; Lee et al., 2005; Banerjee et al., 2005, 2006; Panigrahi et al., 2011; Bagani et al., 2014a,b). Upon chemical reduction followed by annealing at 600 C the density of wrinkles in the graphene oxide sheet decreased owing to the removal of many epoxy groups, increasing the number of zigzag edges/edge states, causing annealed rGO to have greater magnetism than graphite oxide. Element-specific high-spatial-resolution chemical analysis is a desirable tool for directly examining the role of oxygen containing and hydroxyl groups in particular regions and to elucidate the difference between chemical states in specific (wrinkle or flat) regions on the surfaces of graphite oxide and rGO.

6.4.2 Electrical transport mechanism of GO/reduced graphene oxide Hall measurement is an effective technique to investigate the electrical transport properties of graphene oxide materials, which has been conducted in different research studies (Xu et al., 2013; Rein et al., 2015; Pei et al., 2012). Hall measurements semiconductor properties, such as the charge carrier density n and Hall mobility μ, were calculated using following equations (Rein et al., 2015): n5

1 and RH e

μ5

1 ; eρs n

where the Hall coefficient RH is the slope of the Hall resistance as a function of the applied magnetic field (B), ρs is the sheet resistance, and e is the fundamental charge of an electron. The positive value of RH, that is, the positive slope type, determined from the Hall measurements in all the rGO samples is an indication of the hole-dependent charge carrier transportation. The majority of the charge carriers in the rGO thin films are p-type. The calculated values of the room temperature charge carrier concentration are in-between 1.84 3 1012 and 1.26 3 1013 cm22, which indicates a high level of doping in the rGO films (Friedman et al., 2010). To study the conduction mechanism the temperature-dependent resistance measurements were carried out for the different type of rGO produced using different number of LASER shots with a wide range of temperature B5K , T , 350K as shown in Fig. 6
12. With decreasing temperature the trend of nonlinear increment in the resistance is consistent with many other semiconducting 2D systems (Seung et al., 2011; Kravchenko et al., 2004). The essence of the resistance dependence on the temperature can be split into two distinct regions. At high-temperature regime the resistance of the rGO thin films increases slowly with decreasing temperature. On the other hand, at the low-temperature regime, the increment of the resistance was faster with the same rate of temperature drop. At high-temperature region the resistance versus temperature curve correlates best with the Arrhenius-like temperature dependence, whereas below the temperature of 210K, the thermal activation energy is not sufficient to energize significant amount of charge carriers to the conduction level. In this case the temperature-dependent resistance

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 6–12 Resistance as a function of temperature for the rGO produced using different number of LASER shots, namely, 300, 2000, 5000, 10,000, and 20,000. By courtesy Haque, A., et al., 2018b. Temperature dependent electrical transport properties of high career mobility reduced graphene oxide thin film devices. IEEE Trans. Semicond. Manuf. 31, 535
544.

behavior changes from Arrhenius-like temperature dependence to the 2D variable range hopping system.

6.5 Applications of graphene oxide spintronics Several devices have been fabricated using GO as a starting material for at least one of the components. One such device is a graphene-based field-effect transistor (Su et al., 2010; Wang et al., 2010), field-effect transistors (FETs) that employ rGO have been used as chemical sensors (Lu et al., 2011; Chen et al., 2012a,b; He et al., 2012) and biosensors. Su et al. (2010) studied the bottom-gate-operated transistors fabricated by evaporating Au electrodes directly on top of the reduced GO sheets, which were previously deposited on SiO2/Si substrates. Fig. 6
13A demonstrates the typical output characteristics (drain current Id vs drain voltage Vd) for the device prepared from a single-layer GO sheet after 1000 C alcohol reduction. Inset shows the photograph of the device, where the graphene edge is indicated by dotted lines. Only GO sheets with regular shapes and suitable sizes were used for transistor fabrications because it is convenient to extract their field-effect carrier mobility. Fig. 6
13B shows the transfer curves (Id vs gate voltage Vg) for the device. No neutrality point (Dirac point) is observed for the as-prepared device within Vg scanning range, likely because the reduced GO sheet is still heavily p-doped.

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FIGURE 6–13 (A) Output characteristics Id versus Vd and (B) transfer curves for a single-layer GO sheet after 1000 C alcohol reduction with or without DMF treatment. (C) Statistical mobility data for the devices made from monolayered GO sheets reduced at various temperatures. DMF, dimethylformamide. Su, C.-Y., et al., 2010. Highly efficient restoration of graphitic structure in graphene oxide using alcohol vapors. ACS Nano 4 (9), 5285
5292. Copyright ACS Publications. Reproduced with permission.

The Dirac point of the device can be shifted back to 0
30 V once the device is soaked with dimethylformamide (DMF) vapors for 8 h, where the DMF dopes the device with electrons and left-shifts the Dirac point. In addition, it is reasonable that the Id is consequently reduced, as shown in Fig. 6
13A, due to the decrease in the hole carrier concentration.

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The field-effect mobility of holes was extracted based on the slope ΔId/ΔVg fitted to the linear regime of the transfer curves using the equation μ 5 (L/WCoxVd) 3 (ΔId/ΔVg), where L and W are the channel length and width and Cox are the gate capacitance. The field-effect mobility of the alcohol-reduced GO sheets increases to B210 cm2/V s in ambient condition, which is at least two orders of magnitude higher than the reported B0.1
1 cm2/V s after hydrazine chemical reduction. Fig. 6
13C compiles the mobility data for the devices made from mono-layered GO sheets reduced at various temperatures, where the number in parentheses indicates the number of devices prepared. The statistical results clearly demonstrate that the mobility increases with a reduction in temperature. Moreover, the device mobility for the GO sheet reduced from H2 gas at 1000 C is much lower than those from alcohol reduction at the same temperature. Lu et al. (2011) studied the transport characteristics of rGO devices that strongly influenced by the gas adsorption. Fig. 6
14A and B shows the transport characteristics of an rGO device before any sensing tests and after the analyte (NO2/NH3) exposure (immediately measured at the end of a sensing cycle). The Ids
Vg curves (Ids is the drain
source current and Vg is the gate voltage) were obtained with the device stored in ambient conditions. Before gas sensing the rGO device demonstrated a threshold voltage (Vth) at about 120 V

FIGURE 6–14 Transport characteristics (Ids
Vg) of an rGO device (A) before any sensing and at the end of an NH3 sensing cycle; (B) after an NO2 sensing cycle (Vds 5 0.1 V). (C) Ids
Vds curves before sensing, after NH3, and after NO2 sensing. Lu, G., et al., 2011. Toward practical gas sensing with highly reduced graphene oxide: a new signal processing method to circumvent run-to-run and device-to-device variations. ACS Nano 5 (2), 1154
1164. Copyright ACS Publications. Reproduced with permission.

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175

and p-dominated conductance within the Vg scan range ( 6 40 V) and for both upward and downward Vg scans. This p-dominated semiconducting behavior agrees with the results acquired for room-environment exposed graphene prepared by micromechanical cleavage of graphite (Romero et al., 2008) and chemical (Gilje et al., 2007), and thermal (Lu et al., 2009; Dan et al., 2009) reduction of graphene oxide. The p-type behavior could be attributed to the polarization of adsorbed molecules (e.g., water) and/or defects introduced on the graphene sheets during the preparation or reduction process (Kim et al., 2008). The observed hysteresis for Ids
Vg curves is typical of back-gated carbon nanotube (CNT) and graphene FETs and is generally ascribed to the polarization of adsorbed molecules (such as water) in the applied electric field (Dan et al., 2009; Kim et al., 2003; Lohmann et al., 2009). Fig. 6
14A also includes the Ids
Vg curves obtained immediately after an NH3 sensing cycle (10 minutes air/ 15 minutes 1% NH3/and 25 minutes air); the curves are almost symmetric (v-shaped) with Vth at about 0 V, indicating the ambipolar characteristic of the device after the NH3 exposure. The conductance of the device at Vg 5 0 slightly decreased, which is consistent with the incomplete recovery at the end of the sensing cycle, as shown in Fig. 6
14B. NH3 could serve as an n-type dopant, leading to a lowered hole-density and the shift of the threshold voltage toward the negative regime (Lohmann et al., 2009). In contrast, Fig. 6
14B shows that NO2 exposure not only rendered the rGO device completely p-type within the Vg scan range (evidenced by a significant shift of Vth toward a much more positive Vg) but also significantly increased its conductance. NO2 has an unpaired electron and is a strong oxidizer with electron-withdrawing power (Leenaerts et al., 2008). The electron transfer from rGO to NO2 could increase free hole-density in rGO, thereby enhancing its conductance. In addition, because of their polarity, NH3 and NO2 adsorbed on rGO could contribute to the hysteresis behavior of Ids
Vg curves. The shift of transport curves could be used to estimate the total charge transfer (ΔQ) using ΔQ 5 C ΔVth, (Hecht et al., 2006; Peng et al., 2009), where C is the SiO2 gate capacitance and ΔVth is the shift of the threshold voltage Vth for the rGO FET device. Assuming negligible change in the capacitance C, the more dramatic change in Vth after NO2 sensing suggests that more charge transfer occurs between NO2 and rGO than that between NH3 and rGO even with the NH3 concentration (1%) much higher than NO2 (100 ppm). Our observation agreed with the first-principles study (Leenaerts et al., 2008) that estimated higher charge transfer between NO2 and graphene than that between NH3 and graphene. Fig. 6
14C compares the Ids
Vds (Vds is the drain
source bias) curves before and after sensing. The conductance of the rGO device slightly decreased after the NH3 sensing, while it was still more than threefold higher than that in air after NO2 sensing. These curves are symmetric and mostly linear, which could imply possible ohmic contact between rGO sheets and metal electrodes. The Vth shift after the NO2/NH3 exposure Fig. 6
14A and B would have been unlikely if a Schottky contact had dominated the rGO FET. Single-walled CNTs (SWCNTs) can be regarded as seamless cylinders by rolling up graphene sheets. It was found that Au can make good ohmic contact with p-type SWCNTs with a contact resistance of about 10
50 kΩ (Yaish et al., 2004). Contact resistance of similar magnitude may be expected for graphene and Au since CNT and graphene are closely related. Furthermore,

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compared with the cylinder-plane contact for SWCNTs bottom-contacted with the electrode, the planar contact area between an rGO platelet and the Au electrode is relatively larger, which could lower contact resistance. The conductance of the air-dried rGO used in our study is B1700 S/m (Park et al., 2009), which corresponds with an estimated resistance of B590 kΩ for a 1 nm thick square rGO platelet with 1 μm sides. The junctions formed between rGO platelets could also add to the resistance of rGO platelet network. The resistance of rGO devices is most likely dominated by rGO platelets; therefore we speculate that the contact has limited role in the sensing process and that the charge transfer between the analyte (NO2/NH3) and the rGO primarily contributes to the sensing response.

6.6 Conclusion and perspectives of graphene oxide
based spintronics As in many other fields, the research of GO/rGO applications has seen dramatic progress, and is expanding rapidly. The advances made in this area so far are exciting and encouraging and the challenges, however, are also very huge and must be overcome. One of such challenges is thorough and profound understanding of graphene-oxide spintronics. Such knowledge certainly facilitates development of more efficient GO-based nano-platform for spintronic and other applications. These goals can only be reached by the joint efforts from materials sciences and nanotechnology. Development of suitable chemical synthesis and functionalization approaches for precise control over size, size distribution, morphology, structural defects, and oxygen-containing groups of GO is urgently needed, as this is closely correlated to the performance of the GO-based nanomaterials for those applications, and the safety issues as well. For GO-based FET, tuning electronic property of GO by controllable modification and reduction of originally prepared GO, and development of techniques to integrate GO into practical devices having high sensitivity, selectivity with acceptable reproducibility, reliability and low cost remain a big challenge. GO will bring disruptive solutions to the current industrial challenges related to energy generation and storage applications, first in nano-enhanced products and then in radically new nano-enabled products. GO-based systems for energy production (photovoltaics, fuel cells), energy storage (supercapacitors, batteries), and hydrogen storage will be developed via relevant proof of concept demonstrators that will progress toward the targeted technology readiness levels required for industrial uptake.

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Further reading Qin, S., et al., 2014. Strong ferromagnetism of reduced graphene oxide. Carbon 78, 559
565.

Magnetism and spintronics in carbon nanoparticle/fullerene

7

7.1 Introduction Magnetic properties of carbon-based materials attract considerable attention due to both scientific and practical reasons (Makarova et al., 2006). The fact that such materials, being metal-free or containing only negligible amounts of metallic impurities, demonstrate a ferromagnetic (FM) behavior even above the room temperature (RT) is intriguing by itself. On the other hand, a board utilization area, stretching from spintronics and light metals in techniques to applications, is expectable for such unusual carbon-based magnetism (Makarova et al., 2006). The most known bulk carbon modifications are graphite and diamond, containing negligible intrinsic disorder, and exhibit only diamagnetism (Stamenov et al., 2005). Therefore the unconventional magnetic properties of carbon-based materials are attributable to defects or disorders introduced into an ordered most matrixes. In particular, this has been demonstrated directly by intentional incorporation of defects by proton irradiation of graphite, leading to pronounced FM properties (Makarova et al., 2006; Barzola-Quiquia et al., 2008). Experimental evidence of the intrinsic magnetism in defect-rich carbon materials (Makarova et al., 2006; Stamenov et al., 2005; Barzola-Quiquia et al., 2008) is accompanied by the theoretical predictions of the FM behavior in such structures as negatively curved graphite surface (Park et al., 2003), a mixture of carbon atoms with alternation of sp2
sp3 bonds (Ovchinnikov et al., 1988), the graphene zigzag edges, and disordered graphite with random atomic defects (Yazyev et al., 2008a,b). Similarly, carbon nanoparticles (CNPs), fullerene, carbon nanosphere, carbon nano-dots, etc. are also representing case of the defect-rich carbon media due to numerous surface defects and high surface-to-volume ratio and behave FM-like magnetic materials, which could be very useful for magnetic device applications mainly for the spintronics. For the spintronics, electron spin states are an attractive realization of a quantum bit (qubit) as they can undergo a transition between the spin-up and spin-down quantum states (DiVincenzo, 1995). The most commonly used technique for manipulating electron spin is electron spin resonance (ESR) (Poole et al., 1971). ESR is the physical process, whereby electron spins are polarized in an external magnetic field B0 and rotated by an oscillating magnetic field B1 (perpendicularly to B0, of frequency f ), which is resonant with the spin precession frequency in an external magnetic field f 5 gμBB0/h (μB is the Bohr magneton and g the electron spin g-factor, h is the Planck’s constant). ESR is result of the coherence of the precession of electrons over the spin
lattice and spin
spin relaxation times, T1 and T2, Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials. DOI: https://doi.org/10.1016/B978-0-12-817680-1.00007-X © 2020 Elsevier Inc. All rights reserved.

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respectively (Bloch, 1946). T1 is characterized by a number of spin
lattice relaxation processes that depend on the spin
orbit coupling to connect the spin of an electron with the lattice vibrational spectrum of the solid. The second important relaxation time, T2, is set by the probability of spin
spin relaxation. T2 is concerned with the local magnetic field contribution by one magnetic atom on others and represents the phase coherence of a set of spins. The dominant relaxation time is the shorter of T1 and T2. In magnetically homogenous itinerant systems (e.g., metals), the condition T1 5 T2 is often met and represents the longest period of time that in-phase precession electron spins, and magnetization can propagate as a uniform mode (Jánossy, 1980). Electron spin states, therefore, need to be robust against decoherence. The feasibility of applications involving classical or quantum information processing is hence critically dependent on T1 and T2 relaxation times (Ardavan et al., 2007; Awschalom et al., 2007; Zwanenburg et al., 2013; Wolf et al., 2001; Yang et al., 2013; Dzhioev et al., 2002; Yakushiji et al., 2005; Hai et al., 2010). The prerequisite for T1 and T2 relaxation times is B100 ns, as this is the state-of-the-art lower bound for signal processing times in quantum electronic devices (Colless et al., 2013; Morton et al., 2006). Advances in the fields of inorganic (Loss et al., 1998; Awschalom et al., 2013) and molecular (Warner et al., 2013; Aromi et al., 2012) quantum dots (QDs) have made the electron spin system promising for practical application in spintronic and quantum information processing.

7.2 Carbon nanoparticle
based spintronics Lähderanta et al. (2012) studied the magnetic properties of different CNPs that are shown in Fig. 7
1; where the origin of magnetic magnetism in CNPs is due to intrinsic near-surface defects. For investigations the CNPs were put into a plastic ampoule filled with helium. Magnetization M (B) was measured with a superconducting quantum interface device (SQUID) magnetometer in fields B up to 5 T by increasing and decreasing the field. The dependence of M (T) was measured in a constant magnetic field between 1 mT and 5 T, after cooling the CNPs from 300K down to 3K in zero field (zero field
cooled magnetization, MZFC) or in the applied field (field-cooled magnetization, MFC). Thermoremanent magnetization (TRM) was investigated after cooling the CNPs from 300K to 3K and reducing the field to zero. The magnetization data are presented after subtraction of the diamagnetic contribution as shown in Fig. 7
1. Temperature dependence of the magnetization in all the investigated CNPs in weak fields of 1 2 50 mT is characterized by a strong irreversibility or deviation of MZFC (T) from MFC (T) and by TRM decreasing with T. Magnetic irreversibility is decreased with increasing the field and vanishes above BB1 T. As can be seen in the lower figure of left panel of Fig. 7
1, M (B) exhibits a saturation already above BB2 T at high temperatures, whereas at low TB3K 2 5K, a deviation from such behavior is observed. Hysteresis is observed already at the RT. Hysteresis loops are similar as shown in the right panel of Fig. 7
1, while strong magnetism is observed as shown in the left panel of Fig. 7
1. Only a small deviation of TRM(T) from the difference of MFC(T) 2 MZFC(T) (left panel of Fig. 7
1)

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185

FIGURE 7–1 The plots of MZFC and MFC versus T and the dependence of M and B of carbon nanoparticles and the hysteresis curve and plots of BC versus Tn of carbon nanoparticles. By courtesy Lähderanta, et al., 2012. IOP Conf. Ser. Mater. Sci. Eng. 38, 012010.

indicates an insignificant nonsphericity of magnetic particles. Saturation of the magnetization at high temperatures (lower figure of left panel of Fig. 7
1) characterizes the system of large magnetic particles with sizes BRav, provided that Tb exceeds the RT. On the other hand a substantial deviation from saturation at low TB3K 2 5K is evident from the lower figure of left panel of Fig. 7
1 as well. Therefore they exhibit a paramagnetic or superparamagnetic response and make only a small contribution to the saturation magnetization of the material, Ms(0), which characterizes the blocked regime of the system of magnetic nanoparticles at B . BK and T{Tb . It can be shown that contribution of such small particles can be excluded by fitting the plots of M(T) at B $ 2 T with the Curie
Weiss law, M(T)  Ms(0) 1 CB/(T 2 θ), yielding the values of Ms(0)  5.6, 4.9, 4.7, and 7.2 (in units of 1024 A m2/g). Hysteresis loop in the right panel of Fig. 7
1 is typical of a system of blocked magnetic nanoparticles, whereas in the right panel, it is broaden considerably and has a more rectangular shape, indicating a FM response from Co ions and/or clusters. It was observed a substantial decrease of the coercive field, BC(T), with increasing T. In a system of magnetic nanoparticles, in the blocked regime, the temperature dependence of the coercivity is given by the expression BC(T) 5 BC(0) [1 2 (T/Tb) n], where n and BC(0) depend on the magnetization reversal mode and on the applicability of the single domain (SD) regime. For SD particles,

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

n 5 1/2 and 2/3 for coherent rotation and curling, respectively; otherwise the value of n 5 1 is connected to domain wall effects. The analysis of BC(T) with the expression previously mentioned is done by minimizing the plots of the standard deviation versus n, evaluated with a linear fit of the plots of BC(T) versus Tn by variation of n. We have found that the minimum of these plots yields n  0.80 and 0.81, giving the best linearization of BC versus Tn displayed in the left panel of Fig. 7
1. The value of n lying between 2/3 and 1 suggests a crossover of a SD curling mode of the magnetization reversal and a multidomain regime, which is expectable due to the relatively large values of particle radii and the broad size distribution evident in the bottom panel of Fig. 7
1. On the other hand the regime of curling is supported also by the relation of BC ð0Þ{BK, whereas for coherent rotation, these values should be comparable.

7.3 Carbon nanosphere spintronics The carbon nanospheres are easily synthesized (Choucair et al., 2012) and readily processable, yielding a homogenous material that is structurally highly noncrystalline. These carbon nanospheres can be reliably employed for spintronic applications with minimal processing and without the need for fabricating a well-defined crystal structure to achieve long T1 and T2. In a conducting carbon nanosphere qubit system the rich chemistry of carbon can allow for a myriad of noncovalent and covalent interactions to connect the nanospheres to conducting electronic device surfaces (Kahlert et al., 2014, 2016). The carbon nanospheres are of a size that can be isolated on a surface from the “top-down” using micromanipulator probe tips (Harneit, 2002). The time window for processing electron spin information (spintronics) in solid-state quantum electronic devices is determined by the spin
lattice and spin
spin relaxation times of electrons. Minimizing the effects of spin
orbit coupling and the local magnetic contributions of neighboring atoms on spin
lattice and spin
spin relaxation times at RT remains substantial challenges to practical spintronics. In this case the conduction-electron spin
lattice and spin
spin relaxation times of 175 ns at 300K in 37 6 7 nm carbon spheres, which is remarkably long for any conducting solid-state material of comparable size. Following the observation of spin polarization by ESR, it was able to control the quantum state of the electron spin by applying short bursts of an oscillating magnetic field and observe coherent oscillations of the spin state. These results demonstrate the feasibility of operating electron spins in conducting carbon nanospheres as quantum bits at RT. Náfrádi et al. (2016) manipulate long lifetime spins using carbon nanospheres at RT using ESR measurements. They have obtained that the continuous-wave ESR linewidth (peak-to-peak) is ΔH 5 0.056 mT (Fig. 7
2A), and the g-value is 2.00225 (see Fig. 7
2). This is a remarkably narrow conduction-electron spin ESR line testifying the long spin relaxation times. The observed spectra had, to a high precision, homogeneously broadened Lorentzian line shapes, and the deviation from the Lorentzian line shape in the entire spectra was ,5%, which reveals the itinerant nature of the spins. The observed linewidth determined by

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187

FIGURE 7–2 Characterization of the spin system of the carbon spheres by ESR. (A) Temperature dependence of T1 and T2 at nZ 5 9.5 GHz. (B) Temperature dependence of the spin susceptibility, wspin, measured by ESR at nZ1/4 9.4 GHz with an overlaying Curie
Weiss line, characteristic to small paramagnetic particles. (C) The ESR linewidth is plotted as a function of the Zeeman energy, EZ1/4hnZ measured by a multifrequency ESR at 300K. The linear fit (straight solid line) using Eq. (7.1), with T1 5 T2 5 175 ns gives δ 5 1 meV. (D) The temperature-independent g-factor shift Dg relative to the free electron g-value, in good agreement with a material exhibiting very weak spin
orbit coupling. Error bars represent the confidence interval of least square fits to the spectra. ESR, Electron spin resonance. Reproduced with permission from Náfrádi, B., et al., 2016. Room temperature manipulation of long lifetime spins in metallic-like carbon nanospheres. Nat. Commun. 7, 12232. Copyright NPG publications.

continuous-wave ESR is identical within the experimental error with the T2-derived Lorentzian width. Note that the size distribution of the carbon nanospheres has a negligible effect on the linewidth at 9.4 GHz because of the motional narrowing of conduction electrons (Náfrádi et al., 2014). The hyperfine ESR lines of 13C were also absent due to motional narrowing of conduction electrons (Náfrádi et al., 2014). The g-factor is characteristic to conduction electrons of carbon, and it does not originate from metallic inclusions or from localized paramagnetic “dangling” bonds of carbon (commonly with g 5 2.00282) (Shames et al., 2002). In the ESR experiment, detection and spin rotation occur at the same time, and Náfrádi et al. (2014) extended their experiments to probe the spin relaxation dynamics of T1 and T2 independently using pulsed ESR (Fig. 7
2A). At 9.5 GHz frequency and 300 K, with good approximation, Náfrádi et al. (2014) found that the intrinsic T1 5 T2 5 175 ns. Pulsed ESR, therefore, simultaneously validated our continuous-wave ESR results and verified that the line shape obtained by continuous-wave ESR was indeed homogenous as expected for

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

itinerant electrons. The temperature-dependent properties of the ESR spectra (Fig. 7
2) support that conduction electrons are confined within the carbon nanospheres. The g-factor was temperature independent (Fig. 7
2D), which is in good agreement with general observations in metals with weak spin
orbit coupling and in graphitic nanoparticles (Andersson et al., 1998). The resulting spin susceptibility is temperature dependent, following Curie
Weiss dependence, as one may expect for nanoparticles of metals (Fig. 7
2B) (Halperin, 1986; Kubo, 1962). Multifrequency ESR in the 4
420 GHz frequency and 2K
300K temperature range also confirmed that conduction electrons are confined within the CNPs (Fig. 7
2) (Náfrádi et al., 2008a,b). The ESR linewidth revealed a linear increase with increasing magnetic field at 300K (Fig. 7
2C). Note that in the case of bulk metals, ΔH is solely determined by spin
orbit coupling making it independent of the magnetic field (Halperin, 1986; Elliott, 1954). However, the behavior observed when the carbon nanospheres experience a variation in external magnetic field is characteristic to conduction electrons enclosed in nanoparticles where T1 and T2 are determined by both the spin
orbit interaction and electron confinement (Halperin, 1986). This broadening of ΔH for itinerant electrons confined on small particles is as follows (Kawabata, 1970): ΔH 5 EZ =δγ e T2 , where EZ 5 hν Z is the Zeeman energy, δ is the average electronic energy level spacing, and γ e is the electron gyromagnetic ratio (Kawabata, 1970). As the temperature was decreased, T2 reached 300 ns at 4K, while T1 reached 450 ns (Fig. 7
4A). There is a deviation from the T1 5 T2 dependence below B100K. During the delineation of T1 and T2 below B100K, T1 and T2 nevertheless continue to increase at different rates. The electron spin dynamics of T1 is directly related to phonon dampening in disordered graphitic sheets (Náfrádi et al., 2014; Andersson et al., 1998) and the existence of Wallis-type (Wallis, 1959) local phonon modes. Luo et al. (2016) explore the influence of the six lyophilized vesicular nanospheres selfassembly on the spin state of the Fe(II) coordination complexes using temperaturedependent magnetic measurements in the temperature range of 400K
200K, and the results were compared with the corresponding bulk materials. The results are interesting that the spin-transition temperatures of these vesicular nanospheres show a negative correlation with the alkyl chain length, that is, the longer the alkyl tails, the lower the transition temperatures, which is in contrast to bulk materials.

7.4 Graphene-nano-dots spintronics Graphene quantum dots (GQDs) with a tiny size of only several nanometers present extraordinary properties due to quantum confinement (Ponomarenko et al., 2008) and edge effects (Ritter et al., 2009), which attract a great deal of interest lately in the fields of magnetic applications because of their spin-polarized edge states and potential applications in the spintronic devices (Wang et al., 2009; Li et al., 2010; Enoki et al., 2008; Wang et al., 2008; S¸ ahin et al., 2010), has aroused continual and tremendous interest. Due to the high edge-tosurface area ratio and the possible substantial spin-polarized edge states at the zigzag

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segments. The spin-polarized edge states on zigzag edges of single narrow graphene nanoribbon have shown RT magnetic ordering, raising hopes of spintronic devices operating under ambient conditions (Magda et al., 2014). Theoretical researches predicted that the geometric shape plays an important role in the magnetic properties of GQD (Fernández-Rossier et al., 2007; Ezawa, 2007; Espinosa-Ortega et al., 2013; Heiskanen et al., 2008). In addition, both defects and reconstruction at the edge were reported to switch off the spin polarization of the edge states (Jiang et al., 2008; Huang et al., 2008; Kunstmann et al., 2011). Despite this, the spin-polarized edge states have also been predicted to be robust against shape irregularity and edge roughness (Heiskanen et al., 2008; Akola et al., 2008; Wimmer et al., 2010; Espinosa-Ortega et al., 2011; Bhowmick et al., 2008; Jaworowski et al., 2013). Clearly, the theoretical researches on the edge states magnetism of GQDs with naturally complicated boundary are highly controversial. Therefore to unveil the edge states magnetism in GQDs, experimental investigation on the magnetism of intrinsic GQDs is of great significance and highly anticipated (Espinosa-Ortega et al., 2013; Yazyev, 2010). Sun et al. (2017) reported the intrinsic magnetism of GQDs (c.2.04 nm) with mostly pristine edges and nearly perfect basal plane. The GQDs show the purely Curie-like paramagnetism with the local moment of 1.2 μB at 2K. They proposed that the majority of GQDs are nonmagnetic, which may attribute to defects and/or reconstruction at the edge arising from the high-temperature annealing. The magnetism and hence magnetic moment of GQDs may mainly originate from the residual zigzag edge states passivated by hydroxyl groups. The ratio of nonmagnetic GQDs is approximately 6/7, with most of the magnetic edge states suppressed by edge defects and/or edge reconstruction arising from the high-temperature annealing. Fig. 7
3A shows the temperature dependence of the mass magnetization M from 2K to 300K of GQDs. It is found that there is no magnetic ordering signal detected at any temperature. Despite the absence of magnetic ordering, GQDs exhibit noticeable paramagnetism at low temperature and diamagnetism at high temperature. After subtracting the linear diamagnetism the 1/χ
T curve was plotted, which fits well with the Curie law χ 5 C/T (Fig. 7
3B). To characterize the species contributing to the localized magnetic moments, measurements of the mass magnetization M versus magnetic field H at 300K and 2K were performed. As noticed from the M
H curve shown in Fig. 7
3C the GQDs show linear diamagnetism at 300K. Since the diamagnetic magnetization varies linearly with magnetic field and is independent of temperature, the net mass magnetization ΔM at 2K (Fig. 7
3D) was obtained by subtracting the diamagnetic magnetization at 300K. The corresponding ΔM versus H curve of GQDs at 2K with error bars given by repeating the measurements. The ΔM
H curve of GQDs at 2K can be described by the standard Brillouin function: M 5 MS

     2J 1 1 2J 1 1 1 x Coth x 2 Coth 2J 2J 2J 2J

(7.1)

where x 5 gJμB H=kB T, the saturated magnetization MS 5 NgJμB , g is the land factor, kB is the Boltzmann constant, J is the angular momentum number, and N the number of present

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 7–3 Magnetic properties of GQDs. (A) M
T and (B) 1/(χ
T) under the applied field H 5 1 kOe. (C) M
H measured at 300K. (D) ΔM
H measured at 2K. The coefficient of determination for the fitting is 0.9997 indicating the goodness of fitting. GQDs, Graphene quantum dots. Reproduced with permission from Sun, Y., et al., 2017. Magnetism of graphene quantum dots. Quantum Mater. 5, 1
7. Copyright NPG publications.

spins. The ΔM value of GQDs at low magnetic field is small and prone to be affected by thermal fluctuations, while the ΔM value at high magnetic field is more accurate. Then our Brillouin function fitting is done on the ΔM
H curve at high magnetic field. Assuming g 5 2, the Brillouin function fitted J and Ms of GQDs are 0.6 and 0.589 emu/g, respectively. One can find that the J value is close to the quantum number 1/2. It is mentioned that after the magnetic measurements exclude the magnetic contribution of 3d impurities. The 3d impurity elements of the GQDs are negligible with Fe , 65 ppm, Cr , 35 ppm, Mn, Ni, and Co lower than the detection limit, and the detected magnetism of GQDs are intrinsic. QDs in graphene are quasizero-dimensional regions to which charge carriers can be confined (Güçlü et al., 2014; Recher et al., 2010; Tarucha et al., 1996; Kouwenhoven et al., 2001). With electrostatic gate electrodes that can be used to adjust the confinement potential and with electric contacts for transport, QDs allow in principle for full control over the individual electron. QDs provide a controllable playground to investigate the behavior of spins as well as sources of decoherence and methods to increase coherence times (Awschalom et al., 2013; Hanson et al., 2007; Petta et al., 2005). The main sources of spin decoherence are interactions with lattice excitations and the hyperfine interaction with present nuclear spins. Droth et al. (2016) studied and reviewed the effective spin
phonon coupling in detail and present a generic power law for the spin relaxation time T1 as a function of the magnetic field. They discuss the spin relaxation in detail. The Heisenberg exchange interaction is

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191

paramount for coherent spin qubit operation and addressed in the context of magnetism in graphene nanoflakes. Nuclear spins in the host and surrounding material can be considered by several means, and the influence of 13C nuclei has been studied in detail. Impressive advances in general spintronics and the fabrication of graphene devices are likely to spark significant advances in spintronics with QDs in the near future.

7.5 Fullerene-based spintronics Fullerenes (C60), the first discovered fullerene, were chosen considering its several properties that make it ideal for organic spintronic devices. First, C60 molecules can be sublimated in ultrahigh vacuum and can, therefore, be cleanly sandwiched between FM metallic thin films and integrated in vertical devices. Second, C60 molecules are very robust and can sustain the top metallic electrode without being damaged, unlike other organic materials. Third, hyperfine interaction is supposed to be very weak in fullerenes due to the absence of polarized hydrogen nuclei and because of the small natural abundance of the 13C nuclear spin (,2%) (Wang et al., 2010). Finally, C60 lowest unoccupied molecular orbital (LUMO) is quite well matched with the Fermi energy of common FM metals, such as cobalt or permalloy (Fe80Ni20, Py), making possible a relatively easy current injection from magnetic electrodes while keeping a moderate energy injection barrier. Fullerenes have susceptibilities of the magnitude expected for the atomic diamagnetism of carbon. The large five- and six-membered rings do not contribute substantially, because clockwise currents in one five-membered ring, for example, induce anticlockwise currents in the neighboring six-membered ring. In C60 the induced currents in the five- and six-membered rings compensate entirely (Haddon, 1994). The susceptibility increases slowly as the size of the fullerene molecule increases, approaching that of graphite, for C5000 (Haddon, 1994). Rhombohedral C60, prepared in a narrow range of temperatures and pressures, has been found to exhibit small FMlike magnetization of around 0.05 Am2/kg at 10K (Makarova et al., 2001). Similar moments have been found in other samples (Makarova, 2004; Makarova et al., 2001), and in partially graphitized glassy carbon (Wang et al., 2002). Gobbi et al. (2011) studied the RT spin transport in C60-based spin valves (SVs). Gobbi et al. (2011) fabricated Co/AlOx/C60/Py hybrids structure deposited on Si/SiO2 substrates for the study of vertical SV systems. The electronic transport properties, RT current
voltage (I
V) curves of the vertical SVs of different thickness C60 systems are shown in Fig. 7
4. The I
V traces become progressively more asymmetric as the C60 thickness is increased. Above a thickness of 20 nm, the low-bias resistance (measured at 10 mV) increases typically by 40% while lowering the temperature to 200K; the value remains almost constant below that temperature. Gobbi et al. (2011) showed that below this temperature, the behavior is compatible with the conductivity dominated by quantum-mechanical tunneling from molecule to molecule. Gobbi et al. (2011) highlight that each tunneling process is inherently spin conserving; so interpretation of the electronic transport between molecules, together with the very small intramolecular spin relaxation mechanisms, suggests that coherent spin transport over relatively long distances should be observed in C60-based SVs.

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 7–4 (A and B) Room temperature current
voltage (I
V) traces for 8 and 28 nm of C60. (C and D) T
R measured at 10 mV with 8 and 28 nm C60 thicknesses. (E) Room temperature low-bias resistance versus thickness of C60 layer. (F) Tunnel barrier thicknesses obtained by fitting the room temperature I
V traces with the Simmons equation as a function of the nominal C60 thickness. Reproduced with permission from Gobbi, M., et al., 2011. Room-temperature spin transport in C60-based spin valves. Adv. Mater. 2011, 23, 1609
1613. Copyright Wiley Publications.

In RT magnetoresistance (MR) study, Gobbi et al. (2011) used the C60 having thickness 5 and 28 nm, respectively, which are represented in Fig. 7
5A and B. In this thickness range, MR values as high as 10% are measured at RT and do not increase significantly at 80K. The rounded shape of the MR and the low coercive fields suggest that the antiparallel state is not well stabilized, most likely due to the small magnetic shape anisotropy created by our crossbar junction configuration and a possible magnetic coupling between electrodes at such C60 thickness. These steps are most probably related to an increase in magnetic pinning sites as a consequence of increasing surface roughness with thickness. The pinning sites alter the intrinsic coercive field of the magnetic layers, giving rise to pseudo-stochastic behavior in the MR (Majumdar et al., 2006; Chan et al., 2010). The MR values become even larger (  8.5%) at lower temperatures (80K) for the C60 having thickness 28 nm. These results compare very positively with the data available in the literature in two different aspects: on the one hand, RT MR is usually negligible, electrodes of which are highly polarized magnetic oxides (such as manganites). On the other hand, RT

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193

FIGURE 7–5 (A and B) Room temperature magnetoresistance in samples with 5 and 28 nm thickness of C60, respectively. MR values of 9% and 5.5% are measured. The two traces are typical for thin (,10 nm) and thick ( . 15 nm) C60 layers. (C) Bias voltage dependence of the MR for the sample with 18 nm of C60 thickness measured at 80K. MR, Magnetoresistance. Reproduced with permission from Gobbi, M., et al., 2011. Room-temperature spin transport in C60-based spin valves. Adv. Mater. 23, 1609
1613. Copyright Wiley publications.

MR values for thicknesses at least one order of magnitude are higher than that composed of 3d-FM metals. It is expected that the significant MR values for C60-based hybrid SVs are up to an approximate thickness of 100 nm. As relevant as the RT MR values for thick C60 is the voltage dependence of that MR (%) (see Fig. 7
5C) for a typical curve measured from the sample with 18 nm of C60. At 80K the MR at low bias reaches 13%, and at a high bias of 21 V the MR is still 0.9%. The slow decay of the MR with applied bias is important since the overall output current increases simultaneously with the bias, and relatively large current values are needed for possible applications of spin devices, such as long-distance information transport (Hueso et al., 2007). The obtained spin coherent transport in fullerenes at RT was observed from the large values ( . 5%) of RT MR in relatively thick ( . 25 nm) fullerenebased SVs. Zare-Kolsarakia et al. (2004) studied that a new situation may arise if the Co clusters are embedded in an insulating matrix built from large molecules, for example, C60 fullerenes (so-called buckyballs). It is not clear at all if the electrons will tunnel through such large molecules. Furthermore, it is well known that the charge transfer occurring between C60 molecules and metal surface (Hunt et al., 1995; Hoogenboom et al., 1998) leads to doped fullerenes, which may show metallic conductivity in the case of heavy doping. The magnetotransport properties of a very similar system, namely, of a carbon nanotube contacted with

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

FIGURE 7–6 ΔR/R as a function of magnetic field μ0H for fullerene Co volume fraction (vCo) 5 0.29 at 4K. Reproduced with permission from Zare-Kolsarakia, H., et al., 2004. Spin-dependent transport in films composed of Co clusters and C60 fullerenes. Eur. Phys. J. B 40, 103
109. Copyright Springer Link.

FM Co electrodes, have been studied (Yang et al., 2003; Tsukagoshi et al., 1999). The observation of a MR has been interpreted as evidence for coherent spin transport of the spinpolarized electrons, injected at the Co contacts, through the nanotube having a length of about 250 nm. In the MR study the resistance R was measured in a sweeping magnetic field at temperatures T # 60K. The resistance change ΔR showed the well-known hysteretic behavior typical for the tunneling resistance of a granular FM system below its superparamagnetic blocking temperature. Fig. 7
6 gives the relative resistance change ΔR/R for Co volume fraction (v
Co) 5 0.29, having a TMR 5 17.5% at T 5 4K. The TMR is defined as TMR 5 RHC 2RHS =RHC with Hc and Hs being the coercive and saturation field, respectively. This TMR seems to be representative for Co clusters embedded in matrices. The most interesting and result with respect to the MR is the observation that the TMR in Co/C60 mixtures strongly depends on the Co-cluster volume fraction vCo. Zhou et al. (2017) studied the vertical SVs fabricated with C60 interlayer between the graphene/cobalt FM electrodes as shown in Fig. 7
7A. The inset of Fig. 7
7A shows the RT current
voltage (I
V) curves of the C60-based device, which displays nonlinear and symmetric characteristics, as well as a parabolic voltage dependence of differential conductance (dI/dV). Furthermore, the junction resistance of device decays monotonically with increasing temperature. These results indicate that the C60 act as a tunnel barrier in the SVs (JöbssonÅkerman et al., 2000). Fig. 7
7B shows the MR traces of the SV devices of C60. A negative MR is also found, which originates from negative spin polarization of the graphenepassivated FM electrodes. The MR ratio decreases slowly with increasing temperature for both kinds of SV and the value is 20.63% at 5K for the C60-based SV that may be due to the poorer quality of the organic
FM interface formed during the top-electrode fabrication. The similar temperature dependence of the MR as that of the device resistance further evidences the spin-dependent tunneling through the interlayer of C60. The thickness of 10
15 nm of

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195

FIGURE 7–7 (A) Temperature dependence of the RA product for the SV of fullerene C60 and inset shows the RT current
voltage curves; (B) MR traces measured at different temperatures for the SV structure with an organic interlayer of C60. MR, Magnetoresistance; RT, room temperature; SV, spin valve. Reproduced with permission from Zhou, G., et al., 2017. Graphene-passivated cobalt as a spin-polarized electrode: growth and application to organic spintronics. J. Phys. D: Appl. Phys. 50, 095001. Copyright IOP Publications.

organic interlayers, the multistep tunneling of spin-polarized electrons via the intermediate states (defect states) inside the highest occupied molecular orbital
lowest unoccupied molecular orbital (HOMO
LUMO) gap, will serve as the main channel of the device current (Tran et al., 2012). Spin-polarized (minority spin) electrons, injected from the graphenepassivated FM electrodes, will undergo a few of the tunneling steps from molecule to molecule until they reach the second FM electrode.

7.6 Conclusion and perspectives of carbon nanoparticles
based spintronics We have observed both the large MR values and the small decrease in MR with applied bias are related to the robust intrinsic properties of fullerenes/CNPs for spin transport. Therefore it is expected that many more spin fullerenes/CNPs-based hybrid devices will follow from ultrathick ( . 100 nm) SVs to the development of more advanced organic spintronic devices, such as the organic spin transistor. The MR in mixtures of Co clusters and C60 fullerenes with a low Co-cluster volume fraction (vCo 5 0:23) has been found to be (1) strongly

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Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials

enhanced compared to that in the system Co/Kr(Xe) and (2) very similar to that observed, for example, for Co clusters coated with CO molecules. These two experimental facts indicate that it is the interaction of the C atoms with the Co cluster surface, which is responsible for the enhancement of the TMR. It is expect that the Co clusters embedded in an amorphous carbon matrix may be another interesting system showing an enhanced TMR. Since amorphous carbon does not contain C atoms with double and triple bonds, as they exist in C60 and CO, respectively, one also would get information on the possible importance of C double (triple) bonds for the enhancement of the TMR. Experiments are in progress in order to examine the TMR in this system. The additional observation, namely, the strong reduction of the TMR with increasing Co-cluster volume fraction in Co-cluster/C60 mixtures is not yet quite understood. The previously explanation of this reduction as caused by spin
flip processes occurring in the electron-doped tunneling barrier is rather speculative. The abovementioned experiments with Co clusters in amorphous carbon may help to clear up this point since the charge transfer between Co clusters and C atoms in amorphous carbon will not lead to an electron-doped tunneling barrier. We believe that with the increasing quality of single crystals, significant progress toward a better understanding of the mixed state properties lies immediately ahead, and that magnetization measurements and their comparison with other techniques will play a key role in this development. These nanoparticles show FM or superparamagnetic behavior at high/low temperatures, which is demanded for nanospintronic applications.

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8

Magnetism and spintronics in other carbon-based composite materials 8.1 Introduction Defect-induced room-temperature ferromagnetism (RTFM) in metal oxides, such as ZnO, TiO2, In2O3, HfO2, and Cu2O, has attracted significant attention because of their potential application in spintronic devices (Dietl, 2010; Volnianska et al., 2010; Hong et al., 2006; Venkatesan et al., 2004; Kim et al., 2015). The RTFM of these metal oxides arise by their defects. These defects may arise from their lattice defects—either “O” vacancies and/or interstitial “metal.” These defects facilitated the ferromagnetic (FM) coupling in these materials. For example, in the case of pure ZnO thin film or nanoparticles and/or nanowires, magnetic behavior is shown due to “O” vacancies and/or Zn interstitials (Phan et al., 2013; Ong et al., 2012; Wang et al., 2010). Organic molecule
capped ZnO nanoparticles (Garcia et al., 2007), nonmagnetic ion-doped ZnO thin films/nanoparticles (Pan et al., 2007), and partially oxidized Zn nanowires (Li et al., 2010) also exhibit RTFM behavior. Even the absorption of some organic molecules by ZnO nanoparticles can induce FM behavior (Guglieri et al., 2012). Ferromagnetism was also observed in carbon adatoms in carbon nanotubes (Lehtinen et al., 2004) and carbon-substitutional doping in boron nitride nanotubes (Wu et al., 2005). Some of these studies have speculated that intrinsic carbon defects could be responsible for the observed magnetic properties. Again, dilute magnetic semiconductor (DMS), such as ZnO and GaN, have attracted wide interest, and there has been a major effort to produce DMSs with Curie temperatures (TC) at or above room temperature (RT). If non-transition metal (TM) dopants are (such as carbon) incorporated in ZnO, they induce magnetism (Wang et al., 2018; Pan et al., 2007). The density functional theory (DFT) calculation shows that the origin of magnetism does not depend on Zn 3d electrons but the unpaired/dangling 2p states of O atoms in the immediate vicinity of VZn. That induces the spin polarization at the top of the valence band (Peng et al., 2009), and the vacancy-induced magnetism preferentially resides on the surface of ZnO (Wang et al., 2008). Wong et al. (2013) observed the spincoupling phenomena in crystalline CoFeB/graphite interfaces, which can be used for carbonbased spintronics applications. Monodispersed iron oxide nanoparticle
reduced graphene oxide (rGO) composites also exhibit FM behavior at RT with small coercivity (Guoxin et al., 2014). Therefore it is worthy to review and discuss the magnetic properties of different carbonbased composite materials and use it for probable spintronics applications. Because the spintronics is an emerging technology that exploits the intrinsic spin degree of freedom of the electron, we have reviewed and discussed different carbon-based materials in the preceding Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials. DOI: https://doi.org/10.1016/B978-0-12-817680-1.00008-1 © 2020 Elsevier Inc. All rights reserved.

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chapters for probable spintronics applications. However, different carbon-based composite materials, as discussed above, also have potentiality for these applications. In the following sections, we discuss the magnetic properties of different carbon- and carbon nanostructuremetal/nonmetal/metal-oxide composite materials for probable spintronics applications.

8.2 Carbon nanostructure-metal/nonmetal/metal-oxide composites in spintronics With regard to spintronics, hybrid structures that combine carbon-based materials and FM materials are basic building blocks. For instance, spin transport in graphene monolayers, bilayers, and multilayers along with other carbon-based materials, such as carbon nanotubes and diamond-like carbon, have already been reviewed and discussed. However, a lack of reliable ways to fabricate clean and structurally ordered FM/carbon-based interfaces remains a major challenge in this field (Chen et al., 2008; Bolotin et al., 2008). Karpan et al. (2007) have suggested using epitaxial sandwich structures containing Co and/or Ni electrodes separated by multilayer graphene/graphite. However, a key issue to be tackled is the practical realization of an epitaxial FM electrode on top of graphene/graphite. Due to the large difference in surface energy of 3d FM metals and carbon-based materials, FM metals deposited onto graphene/graphite usually exhibit a 3D growth mode (Bäumer et al., 1995; Poon et al., 2006; Wong et al., 2011). But it was observed that the results in poor epitaxy, unless strongly out of equilibrium conditions, are used. Based on these situations, the spintronics community has begun to appreciate the technological importance of amorphous FM (a-FM) alloys (Hasegawa, 1983) for achieving novel spin-dependent phenomena relying on latticematched interfaces. For instance, magnetic tunnel junctions, incorporating the ternary alloy CoFeB in conjunction with an MgO barrier exhibit giant tunneling magnetoresistance values at RT (Djayaprawira et al., 2005; Lee et al., 2007). The primary mechanism governing this effect is based on the band structure symmetry and the requisite coherent (001)-textured body-centered cubic crystal structure involved at the CoFeB/MgO interface (Butler et al., 2001). In addition, a-FM alloys feature many desirable properties (Egami, 1984), such as magnetic softness, due to their amorphous/nanocrystalline nature and tunability of their electronic, magnetic, and structural properties by varying elemental compositions, which makes them unquestionably important materials for spintronic applications. Dong et al. (2007) studied the magnetic properties of mesostructured γ-Fe2O3/carbon composites with γ-Fe2O3 nanoparticles embedded in the wall of ordered mesoporous carbon materials and found that the coercivity is almost zero, and the remanence indicates the superparamagnetic behavior of the composites. In another study, Guoxin et al. (2014) found different saturation magnetization (Ms) of different monodispersed iron oxide nanoparticlerGO composites formed by self-assembly in aqueous phase and are completely dependent on the size and density of FeO
NPs
rGO. The higher densities of iron oxide nanoparticles (FeO
NP) in the composites lead naturally to bigger value of specific Ms. In this case the

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iron oxide nanoparticles were synthesized using FeCl3 as the iron source and then selfassembled on rGO sheets through electrostatic attraction. Fig. 8
1 shows the different values of saturation magnetization. These values are 38.3, 19.5, and 7.7 emu/g, respectively, for different composites prepared with the initial volume ratios (FeONP:rGO) at 4:1, 1:1, and 1:4, respectively. The magnetization for all composites is almost saturated in a low magnetic field (1500 Oe), indicating that the magnetocrystalline anisotropy energy is small (Kusakari et al., 2007). It shows that the iron oxide nanoparticles would change from ferromagnetism to superparamagnetism when the size of particle was smaller than 25 nm (Lee et al., 2009). The FM behavior of the FeONP
rGO composites correlates with the various interactions (exchange interactions, dipolar interparticle interactions, and interfacial cross
grain
boundary interactions) between iron oxide nanoparticles and rGO (Tromsdorf et al., 2007). Recently, Ghosh et al. (2018) studied the magnetic behaviors for the different nitrogenated graphene oxide (GO-Nx) and “GO-Nx” functionalized with different iron oxides (GO-Nx/Fe-oxide) at RT (B300K). Measured magnetic
hysteresis (M
H) loops within the magnetic field range 210,000 Oe , H ,1 10,000 Oe at RT (300K) of “GO-Nx” and GO-Nx/ Fe-oxide are shown in Fig. 8
2. Ghosh et al. (2018) found that the magnetization of “GO-Nx” not only depends on the content of carbon and/or nitrogen but also strongly on different N-precursors. Fig. 8
2A shows that the highest magnetization was obtained in “GO-Nx” synthesized with C6H12N4precursor, with the Ms value of 0.0057 emu/g and a coercivity (Hc) of 53 Oe, whereas the lowest magnetization was found (Ms 5 0.0031 emu/g and Hc 5 31 Oe) in “GO-Nx” synthesized with CH4N2O-precursor. This magnetic behavior also depends on the formation of different N-functional groups, such as pyridine-N, pyrrolic-N, graphitic-N, and N-oxide, different carbon
carbon bonds and carbon
oxygen functional groups, such as
OH,

FIGURE 8–1 Magnetization curves of FeONP
rGO composites measured at 300K with different densities of iron oxide nanoparticles. Inset is the corresponding zoom of the magnetization curves. rGO, Reduced-graphene oxide. Guoxin H., et al., Monodispersed iron oxide nanoparticle-reduced graphene oxide composites formed by selfassembly in aqueous phase, Fullerenes Nanotubes Carbon Nanostruct. 23, 2014, 283
289.

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FIGURE 8–2 Magnetization of M
H loops of (A) NH4OH-based “GO-Nx”, C6H12N4-based “GO-Nx”, C2H3N-based “GO-Nx,” and CH4N2O-based “GO-Nx”. (B) Magnetization of M
H loops of Fe2O3-based “GO-Nx:Fe” composites, FeOOH-based “GO-Nx:Fe” composites, and Fe3O4-based “GO-Nx:Fe” composites. Ghosh, B., et al., 2018. Tuning of magnetic behaviour in nitrogenated graphene oxide functionalized with iron oxide. Diam. Relat. Mater. 89, 35
42. Copyright Elsevier publications. Reproduced with permission.


COOH, C
O/C 5 O, that are attached with the “GO-Nx” matrix during the formation of “GO-Nx” (Sun et al., 2014). However, the origin of this magnetization in N-doped GO is still controversial. There are some reports on ferromagnetism in N-doped GO that has been ascribed due to pyrrolic group, which can provide a net magnetic moment of 0.95 µB/N atom (Miao et al., 2016). In contrast, Ito et al. (2015) have observed that the presence of pyrrolic group leads to a reduction in the net magnetization value of nitrogenated GO. Ghosh et al. (2018) found that “GO-Nx” synthesized with C6H12N4-precursor have higher graphitic content (sp2) and exhibits higher magnetization. This enhancement of sp2 makes this “GONx” structure more disordered, and magnetization is enhanced significantly. The magnetization in sp2-carbon structures also originates from the defects as well. A similar phenomenon, also observed by Sarma et al. (2017) from XPS and microstructural Raman study, concludes that the sp2 as well as ID/IG ratios (microstructural sp2/sp3 ratio and/or their defects) increased with increase of nitrogen concentration in the GO-Nx structure; as a result, the magnetization was enhanced significantly. To tune this magnetization of “GO-Nx”, Ghosh et al. (2018) have further studied the “GO-Nx:Fe” composites. For this study, Ghosh et al. (2018) used the highest magnetized “GO-Nx” (synthesized with C6H12N4-precursor) and

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functionalized with three different iron oxides, namely, Fe2O3, FeOOH, and Fe3O4, for the preparation of “GO-Nx:Fe” composites. The M
H loops of three “GO-Nx:Fe” composites along with pure “GO-Nx” (prepared from C6H12N4-precursor) are shown in Fig. 8
2B. It is found that the magnetization is enhanced more than an order in all of these three “GO-Nx: Fe” composites. The “GO-Nx:Fe” composites prepared with Fe2O3 show the highest magnetic moment with MsB0.043 emu/g, which is nearly eight times higher than pure “GO-Nx” (prepared with C6H12N4-precursor). In the case of ZnO, when “C” is doped or irradiated, their magnetic properties change drastically. Pan et al. (2007) studied the RTFM in carbon-doped ZnO films and found that the ferromagnetism behavior with Curie temperatures is higher than 400K. The magnetic moment based on the content of carbide in the films was (1:5
3:0) µB per carbon atom, which are in agreement with the theoretical prediction. The magnetism is due to the Zn
C system in the ZnO environment that was formed by the substitution of oxygen atom and formation of Zn
C bonds in the carbon-doped ZnO films (Ramqvist et al., 1969). Recently, Wang et al. (2018) studied in detail about the origin of magnetic properties in carbon (C)-implanted ZnO nanowires (ZnO-NWs). Fig. 8
3 shows the RT M
H curves of both “C”-implanted ZnO-NWs (ZnO
C:NW) and pure ZnO-NW, where the magnetic field was applied parallel to the direction of growth of the NW. Fig. 8
3 reveals that the

FIGURE 8–3 Room temperature M
H curves of ZnO
C:NW and ZnO-NW. Upper and lower insets present magnetic field applied parallel to growth direction (c-axis) and magnified M
H loops of ZnO
C:NW and ZnO-NW, respectively, at 300K. M
H, Magnetic
hysteresis. Wang, Y.F., et al., 2018. Origin of magnetic properties in carbon implanted ZnO nanowires. Sci. Rep. 8, 7758:1
13. Copyright Nature Springer publications. Reproduced with permission.

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magnetization of ZnO
C:NW is significantly greater than that of ZnO-NW. The lower inset in Fig. 8
3 clearly reveals the enhanced magnetization in ZnO
C:NW, for which Ms and coercivity of ZnO
C:NW (ZnO-NW) are B2.7 emu/cm3 (0.3 emu/cm3) and B200 Oe (100 Oe), respectively. The Ms for ZnO
C:NW is about 10 times that of ZnO-NW, and this effect is attributable to the implanted C atoms in the ZnO-NW matrix. All these findings suggest that the difference between the effects of implanted C-atoms in the surface and bulk regions are responsible for the magnetic behaviors in the ZnO
C:NW relative to those of ZnO-NW. Wang et al. (2018) also use the X-ray magnetic circular dichroism (XMCD) measurement for the study of magnetization of ZnO
C:NW and ZnO-NW. Fig. 8
4A and B presents the O K-edge and Zn L3,2-edge XANES spectra of ZnO
C:NW and ZnO-NW. The lower panels display the corresponding XMCD spectra obtained with the photon helicity of the incident X-rays parallel (μ1) and antiparallel (μ2) to the direction of magnetization. XMCD spectrum is defined as the ratio of (μ1 2 μ2)/(μ1 1 μ2). The general line shapes and positions of the features in the O K-edge XANES and XMCD spectra in Fig. 8
4A are consistent with the features of ZnO nanostructures that were obtained in our earlier study (Singh et al., 2014). In O K-edge XANES, different spectral

FIGURE 8–4 (A) Normalized O K-edge and (B) Zn L3,2-edge XANES spectra with photon helicity of incident X-rays parallel (μ1) and antiparallel (μ2) to the direction of magnetization for ZnO
C:NW and ZnO-NW. Upper inset in (A) displays magnified O K-edge XANES spectra; lower panels in (A and B) display O K-edge and Zn L3,2-edge XMCD spectra of ZnO
C:NW and ZnO-NW. (C) C K-edge XANES spectra of ZnO
C:NW. Upper inset magnified view C Kedge near-edge spectra of ZnO
C:NW. Lower panel displays C K-edge XMCD spectra of ZnO
C:NW. Wang, Y.F., et al., 2018. Origin of magnetic properties in carbon implanted ZnO nanowires. Sci. Rep. 8, 7758:1
13. Copyright Nature Springer publications. Reproduced with permission.

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features arise due to electron excitation of O 1s states to 2pσ (along the bilayer) and 2pπ (along the c-axis) states, and their intensities are approximately proportional to the density of the unoccupied O 2p-derived states (Chiou et al., 2004a,b). The enhancements of those features in ZnO
C:NW represent the increases in local density of states (DOSs) that are caused by the defects and unpaired/dangling O 2p-derived states in the surface region of the ZnO nanorods (Chiou et al., 2004a,b). This is because the implanted C atoms reduce the number of unpaired/dangling O 2p bonds as the interstitial defects of Ci form C
O bonds in the surface region. It is evident from O K-edge XANES that the density of the O 2p-derived states is lower in the surface region at/above the conduction-band minimum (ECBM) or Fermi level (EF) in ZnO
C:NW than that of ZnO-NW. The XMCD spectra in the lower inset in Fig. 8
4A is weak but confirmed magnetic moment at O sites and caused the imperfect alignment of spin moments of unpaired/dangling O 2p-derived states in the surface region of ZnO-NW (Singh et al., 2014). The intensity of the O K-edge XMCD spectrum of ZnO
C:NW is lower than that of ZnO-NW, revealing no clear spin moment of the O 2p-derived states in ZnO
C:NW. This result is highly consistent with the above claim that the d0 magnetic behavior of ZnO-NW is closely related to the number of unpaired/ dangling O 2p-derived states. In Zn L3,2-edge XANES spectra shown in Fig. 8
4B, different features are primarily associated with the transition of Zn 2p electrons to Zn 4d/s-derived states (Chiou et al., 2004a,b). The variation in the line shapes of the Zn L3,2-edge XANES of ZnO
C:NW can be caused by the change in the electronic structures upon the formation of Zn
C bonds around Zn sites when C atoms are implanted (Kucheyev et al., 2005). No clear difference of both ZnO
C:NW and ZnO-NW are observed in the XANES and XMCD spectra of Zn 4d states as shown in lower panel in Fig. 8
4B, indicating the possibility of Zn
d orbital d0 magnetism in ZnO
C:NW and ZnO-NW. Even there is a possibility that implanted C atoms are substituted at VO or VZn sites to form Zn
C bonds in the first and second shells around Zn atoms in ZnO
C:NW. The C K-edge XANES features in Fig. 8
4C are associated with the C 1s!2p (π ) and 1s!2p (σ ) transitions, respectively (Chuang et al., 2014; Wang et al., 2015). The lower panel also displays the corresponding C K-edge XMCD spectrum that was obtained with the photon helicity of the incident X-rays parallel (μ1) and antiparallel (μ2) to the direction of magnetization. The difference between μ1 and μ2 intensities was magnified in the upper inset of Fig. 8
4C. The lower panel in Fig. 8
4C displays the C K-edge XMCD spectra [(μ1 2 μ2)/(μ1 2 μ2)] of ZnO
C:NW, clearly revealing that the C 2p-derived states of the C atoms affect the magnetic behavior in ZnO
C:NW. As shown in the lower panel in Fig. 8
4C, a magnetic moment, that was associated with the C 2p-derived states in ZnO
C:NW, was also observed. The intensity of the C K-edge XMCD features in the range 280
292 eV, typically attributed to C 2p (π )-derived states, but no clear XMCD feature of 2p (σ ) states was observed in the region 292
300 eV. The C 2p (π )-derived states are clearly responsible for the magnetism of C sites in ZnO
C:NW. The spectral intensities of the XMCD features of O and C K-edge are opposite in Fig. 8
4A and C, respectively. This result also can be explained by the different projected spin contributions of the O 2p- and C 2p-derived states in ZnO
C:NW, which cause the magnetic moment of the implanted C atoms to align antiparallel to that of the

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O atoms, possibly weakening the magnetic moment from O sites in the surface region. This phenomenon may also explain the lower panel in Fig. 8
4A, which exhibits lower magnetization in the O K-edge XMCD spectrum of ZnO
C:NW than in that of ZnO-NW. The d0 magnetic moments of ZnO-NW at O sites have been demonstrated, based on O K-edge XMCD measurements (Singh et al., 2014), and they result in the unpaired/dangling O 2p-derived states, as addressed above. Notably, the integrated XMCD intensity at the C K-edge (region of B287
292 eV) of ZnO
C:NW, presented in the lower panel of Fig. 8
4C, is approximately 2.5 times that of the integrated XMCD intensity at the O K-edge (region of B532
538 eV) in the lower panel of Fig. 8
4A, revealing that the C 2p-derived states of implanted C atoms may critically affect the net spin polarization in the surface and bulk regions of ZnO
C:NW. In particular, based on the results of the RRBS analysis, the C-depth profile indicates that most implanted C atoms in ZnO
C:NW are within the depth of 0
400 Å. These results provide evidence that not only the magnetic moments in the surface or bulk regions are determined by the population of unpaired/ dangling O 2p-derived states but also that these moments are associated with the population of C 2p-derived states [C 2p(π ) states] of the implanted C atoms particularly in the bulk region, explaining the difference between the magnetic behaviors in the surface and bulk regions of ZnO
C:NW and those in ZnO-NW. Spatially resolved microscopic and spectroscopic techniques are used to provide more information about the effect of implanted C atoms on the difference between the surface and bulk regions in ZnO
C:NW. Wong et al. (2013) demonstrate the experimental investigation in the use of CoFeB for fabricating crystalline interfaces with graphite for carbon-based spintronics. Incorporation of an a-FM alloy as the top electrode material on graphite/graphene enables studies on the aforementioned spin-filtering effect at FM/graphene interfaces. The main idea behind this was to induce a crystalline interface between the two dissimilar materials by solid-phase epitaxy driven by postdeposition annealing, where the a-FM alloy crystallizes at the heterointerface. Consequently, the requirements posed by lattice matching and surface energy compatibility for epitaxial growth of FM metals on graphene/graphite should be less stringent in this case. For this study, Wong et al. (2013) used X-ray absorption spectroscopy (XAS), and XMCD techniques. XAS and XMCD have been particularly chosen for the present studies, due to their element specificity and ability to allow for direct and separate quantitative determination of atomic spin and orbital magnetic moments (O’Brien et al., 1994; Chen et al., 1995; Thole et al., 1992; van der Laan, 1999) of Co and Fe in CoFeB/ HOPG. Fig. 8
5 shows the Co and Fe L2,3-edge XAS and XMCD spectra of CoFeB/HOPG film. The XMCD sum rules are used to extract the spin and orbital magnetic moments of Co and Fe from the spectra (Thole et al., 1992; Carra et al., 1993). According to the procedure of Chen et al. (1995) the contributions of the continuum states were simulated by a two-step background function and subtracted from the absorption spectra for computing the 2p to 3d XAS intensities μ1 and μ2 for parallel and antiparallel alignment of photon helicity and magnetization (Chen et al., 1995). The orbital-m or b and spin magnetic moments mspin were obtained from the integrals of the summed XAS

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209

FIGURE 8–5 Sum rule analysis of normalized Co- and Fe L2,3-edge XAS and XMCD spectra of annealed CoFeB/HOPG. Wong, P.K.J. et al., 2013. Crystalline CoFeB/Graphite Interfaces for Carbon Spintronics Fabricated by Solid Phase Epitaxy. Adv. Funct. Mater. 23, 4933. Copyright Wiley publications. Reproduced with permission.

Ð (μ1 1 μ2) and XMCD (μ1 2 μ2) spectra for calculating the values for p, q, and r (Chen et al., 1995), using the following sum rules:

Ð

morb 5 2

4qnh ð4q 2 6pÞnh and mspin 5 : 3r r

The number of holes nh was taken as 2.49 for Co, and 3.39 for Fe (Chen et al., 1995). The results of the sum rule analysis were compared to the bulk values. It is observed that the mspin of Co and Fe are, respectively, 1.59 6 0.20 and 2.00 6 0.26 µB, both representing the bulk-like values (Chen et al., 1995). On the other hand, the anomalously high q values obtained from the XMCD integrals for both edges give a remarkably high morb of 0.661 6 0.085 and 0.543 6 0.070 for Co and Fe of the annealed CoFeB film, respectively, that are comparable with bulk Co and Fe. In order to rule out the uncertainty, Wong et al. (2013) compared the orbital-to-spin ratios morb/mspin with those of bulk Co and Fe. The ratio for the “Co” in the alloy film is considerably larger than the “Co” bulk value, and, from a similar evaluation, the morb/mspin of Fe is also much larger than that of bulk Fe. Mukherjee et al. (2015) studied the ultrathin (1
3 nm) CoFeB/MgO bilayers and found that the CoFeB/MgO bilayers behave perpendicular to magnetic anisotropy that revealed similarly large morb/mspin suggesting its origin as magnetostriction. These results convinced that this study should lead to a better knowledge as well as further investigations involving a-FM for carbon-based spintronic applications.

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8.3 Conclusion and perspectives of carbon-metal/nonmetal/ composite-based spintronics The results indicate that the incorporation of metal/metal oxide in carbon-based materials enhance the magnetization. In other ways, when carbon is irradiated in other metal oxides, such as ZnO, magnetization is enhanced. Results revealed the modification of defects and the enhancement of magnetization in metal/metal oxide
carbon-based composite materials. In case of ZnO, the magnetism is also responsible due to formation of Zn
C system in the ZnO environment, when carbon is doped in ZnO films. It shows an intrinsic n-type FM behavior, with Curie temperatures well above RT. The intrinsic ferromagnetism originates from the Zn
C system in the ZnO environment. When iron and/or iron oxide is functionalized with graphene oxide, the magnetization is enhanced, which is higher than bulk iron/ iron nanoparticles. Different amounts of magnetic iron oxides were introduced into the carbon composites that varied specific Mss. Again, metal/nonmetal composites with carbonbased materials, such as CoFeB/graphite interfaces, enhance the magnetization and can be useful for spintronic applications. Several researchers/studies reported that production of high-quality carbon-based composites in a robust technique and mass scale remains the bottleneck that needs to be overcome at the first instance to enable carbon/carbon nanostructure materials’ practical applications in spintronics. Several research-studies had put forth the fact that chemically functionalizing the surface of carbon-based materials, hence to achieve a good interfacial interaction would enhance its applicability especially within the field of spintronics applications. An appropriate interaction between the carbon and metal/ nonmetal-oxide composites ensures the technological development of carbon-based composites is yet to attain a full shape; there are a number of challenges to overcome as well. Carbon/carbon nanostructure has extraordinary mechanical, thermal, and electrical properties and forms a better replacement for conventional spintronics material matrices. Thus carbon/carbon nanostructure composites with metal/nonmetal/metal oxide, ceramics, and metals possess unique properties that are remarkable for spintronics applications.

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974. Djayaprawira, D.D., et al., 2005. 230% room-temperature magnetoresistance in CoFeB/MgO/CoFeBCoFeB/ MgO/CoFeB magnetic tunnel junctions. Appl. Phys. Lett. 86, 092502. Dong, X., et al., 2007. Magnetic properties of mesostructured ç-Fe2O3/carbon composites by a Co-casting Method. Chem. Mater. 19, 3484
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7685. Karpan, V.M., et al., 2007. Graphite and graphene as perfect spin filters. Phys. Rev. Lett. 99, 176602. Kim, D.Y., et al., 2015. Ferromagnetism of single-crystalline Cu2O induced through poly(N-vinyl-2-pyrrolidone) interaction triggering d-orbital alteration. J. Phys. Chem. C 119, 13350
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9

Challenges and emerging direction of carbon nanostructure materials in magnetism and spintronics Spintronics, or spin-based electronics, is an emerging field of technology, which exploits an electron’s “spin” and its associated magnetism rather than, or in addition to, its charge. Spintronics holds the promise for future information technologies. Devices based on manipulation of spin are most likely to replace the current silicon complementary metal oxide semiconductor devices that are based on manipulation of charge. The challenge is to identify or design materials that can be used to generate, detect, and manipulate spin. Since the successful isolation of graphene and other two-dimensional (2D) materials, there has been a strong focus on spintronics based on 2D materials due to their attractive properties, and much progress has been made, both theoretically and experimentally. Carbon and carbon nanostructure materials in spintronic are hot topic in the field of research exploiting the influence of the electron spin on the electrical conduction (Fert et al., 2012). It is mainly known for the “giant magnetoresistance” (GMR) (Baibich et al., 1988; Binash et al., 1989) and the large increase of the hard disk capacity obtained with the read heads based on GMR. However, the research on spintronic has also revealed many other interesting effects and is now developing along promising novel directions. The spin orbital coupling interaction in graphene is very weak and there is almost no nuclear magnetic moment, and the electron spin transfer process is easier to control. Therefore the unique magnetic characteristics of graphene make it ideal for the preparation of spintronic devices (Raes et al., 2016; Kamalakar et al., 2015; Barone et al., 2008; Dankert et al., 2017; Avsar et al., 2015). The physical basis of spintronic is the influence of the electron spin orientation on the electrical conduction in magnetic materials: the conductivity can be much larger for electrons having, for example, their spin aligned with the magnetization (Mott, 1936; Fert et al., 1968). The GMR (Baibich et al., 1988; Binash et al., 1989; Fert et al., 2006) exploits this spin dependence in magnetic multilayers composed by a stack of ultrathin layers (a few nm) with, alternately, layers of magnetic (iron for example) and nonmagnetic metals (copper or chromium for example). A magnetic field, by aligning the magnetizations of all the magnetic layers, makes that there is an electrical short circuit by half of the electrons, which have the “good” spin orientation in all the magnetic layers. Without magnetic ordering by an applied field, the short-circuit effect does not exist, and the electrical resistance is much larger. The large reduction of the electrical resistance by a magnetic field has been called “GMR”. The GMR is Magnetism and Spintronics in Carbon and Carbon Nanostructured Materials. DOI: https://doi.org/10.1016/B978-0-12-817680-1.00009-3 © 2020 Elsevier Inc. All rights reserved.

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used to read the magnetic inscriptions on the hard disks of today (Parkin, 2002; Chappert et al., 2007), and the possibility of reading smaller inscriptions has led to a considerable increase (three orders of magnitude) of the capacity of the hard disks (Parkin, 2002; Chappert et al., 2007). The discovery of the GMR in 1988 kicked off an intense search of other phenomena also related to the influence of the electron spin on the electrical conduction. New effects have been found, and this domain of research is now called spintronics (Tsymbal et al., 2011), sometimes described as a new type of electronics exploiting both the charge and the spin of the electrons. An example of very active field of research is the study of the spin transfer phenomena (Slonczewski, 1996; Albert et al., 2000; Rippart et al., 2004). In a spin transfer experiment, one manipulates the magnetization orientation of a magnet without applying any magnetic field—the usual way—but by a transfusion of spin angular momentum from a spin-polarized current. This can be used, for example, to reverse the magnetization (Albert et al., 2000), and this will be used soon in the next generation of magnetic memories called spin-transform torque random access memory (STT-RAM). The STT-RAM, in contrast with the semiconductor RAM of today, is nonvolatile, and it does not need any electrical power to maintain the memory alive (Chappert et al., 2007). This will probably lead to a significant reduction of the energy consumption by the computers ad servers. In another regime the spin transfer can be used to generate oscillation in the radio wave frequency range (Rippart et al., 2004). The spin transfer oscillators are very promising of applications in telecommunications. The research in spintronics extends today in many promising directions. Spintronics with semiconductors aims at combining the potential of conventional semiconductors with the potential of spintronics. Spintronics with graphene, carbon nanotubes, or organic molecules has revealed the advantage of carbon-based materials on metals and semiconductors in term of long spin lifetime and long spin diffusion length (Hueso et al., 2007). The recent results on graphene are promising for the relay of conventional electronics in the so-called beyondCMOS perspective and open the road, for example, to “spin only logic circuits” for a novel type of computer technology. Emerging directions are also single-electron spintronics, one of the way to quantum computing, and neuromorphic spintronics in the direction of bioinspired computers.

References Albert, F.J., et al., 2000. Spin-polarized current switching of a Co thin film nanomagnet. Appl. Phys. Lett. 77, 3809 3811. Avsar, A., et al., 2015. Air stable transport in graphene contacted, fully encapsulated ultra-thin black phosphorus-based field-effect transistors. ACS Nano 9 (4), 4138 4145. Baibich, M.N., et al., 1988. Giant Magnetoresistance of (001) Fe/(001) Cr magnetic superlattices. Phys. Rev. Lett. 61, 2472 2475. Barone, V., et al., 2008. Magnetic boron nitride nanoribbons with tunable electronic properties. Nano Lett. 8 (8), 2210 2214.

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Binash, G., et al., 1989. Enhanced magnetoresistance in layered magnetic structures with antiferromagnetic interlayer exchange. Phys. Rev. B 39, 4828 4830. Chappert, C., et al., 2007. The emergence of spin electronics in data storage. Nat. Mater. 6, 813 823. Dankert, A., et al., 2017. Electrical gate control of spin current in van der Waals heterostructures at room temperature. Nat. Commun. 8, 16093. Fert, A., et al., 1968. Two current conduction in nickel. Phys. Rev. Lett. 21, 1190 1192. Fert, A., et al., 2006. Spin transport in magnetic multilayers and tunnel junctions. In: Mills, F., Bland, J. A.C. (Eds.), Nanomagnetism: Ultrathin Films and Nanostructures. Elsevier, Amsterdam, pp. 153 226. Fert, A., et al., 2012. Challenges and Emerging Directions in Spintronics. MEMS, Paris, France, pp. 1 2. Hueso, L.E., et al., 2007. Transformation of spin information into large electrical signals using carbon nanotubes. Nature 445, 410 413. Kamalakar, M.V., et al., 2015. Long distance spin communication in chemical vapour deposited graphene. Nat. Commun. 6, 6766. Mott, N.F., 1936. The electrical conductivity of transition metals. Proc. Roy. Soc. A 153, 699 718. Parkin, S.S.P., 2002. Applications of magnetic nanostructures. In: Maekawa, S., Shinjo, T. (Eds.), Spin Dependent Transport in Magnetic Nanostructures. Taylor and Francis, pp. 237 279. Raes, B., et al., 2016. Determination of the spin-lifetime anisotropy in graphene using oblique spin precession. Nat. Commun. 7, 11444. Rippart, W.H., et al., 2004. Direct current induced dynamics in Co90Fe10/Ni80Fe20 point contacts. Phys. Rev. Lett. 92, 027201. Slonczewski, J.C., 1996. Current-driven excitation of magnetic multilayers. J. Magn. Mater. 159, L1 L7. Tsymbal, E.Y., et al., 2011. Handbook of Spin Transport and Magnetism. CRC Press, edited by.

Index Note: Page numbers followed by “f” refer to figures. A a-C. See Amorphous carbon (a-C) Amorphous carbon (a-C), 25, 47 electrical and transport of, 52 56, 53f, 54f, 55f, 56f magnetism of, 48 52, 49f, 51f, 52f magnetoresistance of, 57 63, 58f, 59f, 60f, 62f, 64f spin field effect transistor (FET), 63 67, 65f B Ballistic electron injection, 7 Ballistic transistors, graphene in, 135 Bilayer graphene (BLG), 109 spin relaxation in, 121 123, 122f Bipolar spintronics of graphene, 137 138, 138f BLG. See Bilayer graphene (BLG) C Carbon, 23 forms of, 25 36 in spintronics, 23 25 Carbon-based composite materials, 201 carbon nanostructure-metal/nonmetal/metaloxide composites, 202 209, 203f, 204f, 205f, 206f, 209f Carbon-doped ZnO nanowires (ZnO-C:NWs), 35 36 Carbon metal oxide/sulfide composites materials, 35 36 Carbon nanomaterials, 23 carbon nanotubes, 27 28 forms of, 25 36 fullerene, 33 35 graphene, 28 31 graphene oxide, 31 33 in spintronics, 23 25

Carbon nanoparticle-based spintronics, 184 186, 185f Carbon nanoparticles, 33 35 Carbon nanosphere spintronics, 186 188, 187f Carbon nanostructure materials, challenges and emerging direction of, 213 Carbon nanostructure-metal/nonmetal/metaloxide composites, 202 210, 203f, 204f, 205f, 206f, 209f Carbon nanotube-based spintronic devices, tunnel magnetoresistance in, 93 97, 95f spin-valve devices of carbon nanotubes, 96 97 Carbon nanotubes (CNTs), 27 28, 75 magnetism of, 77 83, 77f, 80f, 82f multiwalled, 27 28, 75 76 nanostructure materials, 27 28 single-walled, 27 28, 75 76 in spintronic devices, 83 87, 84f Charged impurity scattering, 119 120 Charge transport, in magnetic tunnel junctions, 88 90, 89f Chemical vapor deposition (CVD), 29 31, 58 60 plasma-enhanced, 29 31, 47 48 CNTs. See Carbon nanotubes (CNTs) Coherent spin transport through semiconductors and interfaces, 11 12 Contact-induced spin relaxation, 120 121 Cr-doped diamond-like carbon, magnetism of, 50 51, 51f CVD. See Chemical vapor deposition (CVD) D Density functional theories (DFTs), 90 Device principles, 12 17, 13f DFs. See Dirac fermions (DFs) DFTs. See Density functional theories (DFTs) Diamond-like carbon (DLC), 25 26, 47

217

218

Index

Diamond-like carbon (DLC) (Continued) electrical and transport of, 52 56, 53f, 54f, 55f, 56f for magnetic storage disks, 67 68 magnetism of, 48 52, 49f, 51f, 52f spin field effect transistor (FET), 63 67, 65f Dilute ferromagnetic semiconductor (DMS), 47 48 Dilute magnetic semiconductor (DMS), 201 202 Dirac fermions (DFs), 123 125 DLC. See Diamond-like carbon (DLC) DRAM. See Dynamic random access memory (DRAM) Dyakonov Perel (D-P) mechanism, 117 118 Dynamic random access memory (DRAM), 15 E Edge-derived spin phenomena, 113 114 Efros Shklovskii-type VRH model, 58 60 Electrical spin transport of graphene, 123 129, 124f, 127f, 128f, 129f Electrical transport mechanism, of GO/reduced graphene oxide, 171 172, 172f Electron injection ballistic, 7 hot, 7, 8f Electron movement, manipulation of, 2f Electron spin resonance (ESR), 153 154, 183 184, 186 188 Elliott Yafet (E-Y) mechanism, 117 118 ESR. See Electron spin resonance (ESR) Exfoliation method, 29 31 F FCNs. See Fluorescent semiconductor nanocrystals (FCNs) FETs. See Field-effect transistors (FETs) Field-effect transistors (FETs) graphene in, 135, 136f spin, 63 67, 65f Fluorescent semiconductor nanocrystals (FCNs), 33 35 Fullerene nanostructure materials, 33 35 properties of, 35 Fullerene-based spintronics, 191 195, 192f, 193f, 194f, 195f

G Giant magnetoresistance (GMR), 1 3, 14 15, 75 76, 87 88, 213 214 GNMs. See Graphene nanomeshes (GNMs) GNR. See Graphene nanoribbon (GNR) GO. See Graphene oxide (GO) GQDs. See Graphene quantum dots (GQDs) Graphene, 103 edge-derived spin phenomena, 113 114 electrical spin transport, 123 129, 124f, 127f, 128f, 129f spin polarization, 129 into magnetic materials, making of, 105 107 doping and/or functionalization with transition metals, 107 ferromagnetism derived from hydrogenated zigzag-type pore edges graphene, 106 magnetism depending on pore edge termination by different foreign atoms, 106 graphene-ferroelectric meta-devices, 138 140, 140f magnetism in, 109 113, 110f, 111f nanostructure materials, 28 31 properties of, 31 spin detection, 115 117, 116f spin generation, 107 108 spin Hall effect, 113 114 spin injection, 115 117 spin manipulation, 108, 115 117 spin relaxation process, 117 121 charged impurity scattering, 119 120 contact-induced spin relaxation, 120 121 Hanle spin precession, 118 119, 119f single-layer graphene and bilayer graphene, 121 123, 122f spintronics, applications of, 131 138 bipolar spintronics, 137 138, 138f field-effect transistor, 135, 136f Hall effect, 135 137, 137f spin valve devices, 131 134, 132f, 134f spintronics magnetoresistance devices, 129 131 Graphene-ferroelectric meta-devices, 138 140, 140f Graphene-nano-dots spintronics, 188 191, 190f Graphene nanomeshes (GNMs), 103 104 Graphene nanoribbon (GNR), 103 104

Index

Graphene oxide (GO), 151 in-plane and out-of-plane magnetization/ magnetic anisotropy, 165, 166f magnetization of, 153 165, 155f, 156f, 157f, 158f, 160f, 161f, 162f, 163f magnetoresistance of, 165 172, 167f, 168f electrical transport mechanism, 171 172, 172f magnetic behaviors, mechanism of, 169 171 materials, 31 33 spintronics, applications of, 172 176, 173f, 174f Graphene quantum dots (GQDs), 188 190 H Hall effect of graphene, 135 137, 137f spin, 103 104, 108, 113 114 Hanle spin precession, 118 119, 119f Highly oriented pyrolytic graphite (HOPG), 109, 165 HOPG. See Highly oriented pyrolytic graphite (HOPG) Hot electron injection, 7, 8f I Information technology, 1 2 In-plane magnetization/magnetic anisotropy, 165, 166f Insulator spintronics, 3 K Kekulé structure, 24 25 L Lateral spin-valves (LSVs), 115 117 Lieb theorem, 170 LSVs. See Lateral spin-valves (LSVs) M Magnetic behaviors of GO/reduced graphene oxide, mechanism of, 169 171 Magnetic force microscopy (MFM), 110 113 Magnetic recording, 14 15, 14f Magnetic storage disks, diamond-like carbon for, 67 68

219

Magnetic tunnel junctions (MTJs), 85 application of, 86 87 fabrication of, 85 86 spin currents in, 87 93 spin and charge transport, 88 90, 89f spin polarization, 90 93, 91f tunnel magnetoresistance in, 86 Magnetite, 164 165 Magnetoresistive random access memory (MRAM), 15 17, 16f in quantum computer, 17 in spin transistor, 17 Metallic spintronics, 3 MFM. See Magnetic force microscopy (MFM) Mott VRH law, 61 MRAM. See Magnetoresistive random access memory (MRAM) MTJs. See Magnetic tunnel junctions (MTJs) Multiwalled CNTs (MWCNTs), 27 28, 75 76 magnetism of, 77 81, 77f, 80f spin polarization, 90 93 spin-valve devices of, 96 97 tunnel magnetoresistance in, 93 94, 95f N Nickel nanoparticle deposited vertical carbon nanotubes, 83 Nonvolatile memories, 15 O Ohmic injection, 5 6 Optical control of nuclear spins, 12 Out-of-plane magnetization/magnetic anisotropy, 165, 166f P PAHs. See Polycyclic aromatic hydrocarbons (PAHs) PECVD. See Plasma-enhanced chemical vapor deposition (PECVD) Perpendicular magnetic anisotropy (PMA), 165 Plasma-enhanced chemical vapor deposition (PECVD), 29 31, 47 48 PMA. See Perpendicular magnetic anisotropy (PMA) Polycyclic aromatic hydrocarbons (PAHs), 29 31

220

Index

Q QAHE. See Quantum anomalous Hall effect (QAHE) Q-carbon, magnetism of, 51 52, 52f QSHE. See Quantum spin Hall effect (QSHE) Quantum anomalous Hall effect (QAHE), 123 Quantum computer, MRAM in, 17 Quantum dots, 12 Quantum spin Hall effect (QSHE), 123 R Reduced graphene oxide (rGO), 31 33, 151 153 magnetization of, 153 159 magnetoresistance of electrical transport mechanism, 171 172, 172f magnetic behaviors, mechanism of, 169 171 Room-temperature ferromagnetism (RTFM) defect-induced, 35 36, 201 202 S Scanning tunneling microscope (STM), 6 SDT. See Spin-dependent tunneling (SDT) Seebeck effect, 108 Semiconductor spintronics, 3 SHE. See Spin Hall effect (SHE) Signal-to-noise ratio (SNR), 115 117 Single layer graphene (SLG) spin relaxation in, 121 123, 122f Single-walled CNTs (SWCNTs), 27 28, 75 76 ferromagnet-contacted, 86 magnetism of, 78 79 spin-valve devices of, 96 97 tunnel magnetoresistance in, 94 96, 95f SLG. See Single layer graphene (SLG) SNR. See Signal-to-noise ratio (SNR) Spin accumulation, 11 Spin coherence, 11 Spin currents, in magnetic tunnel junctions, 87 93 spin and charge transport, 88 90, 89f spin polarization, 90 93, 91f Spin-dependent tunneling (SDT), 86 88, 90 Spin detection, 10 graphene, 115 117, 116f Spin field effect transistor (FET), 63 67, 65f Spin generation, in graphene, 107 108

Spin Hall effect (SHE), 103 104, 108, 113 114 Spin injection, 5, 9 10 graphene, 115 117 Spin manipulation of graphene, 108, 115 117 Spin polarization, 4 graphene, 129 in magnetic tunnel junctions, 90 93, 91f Spin relaxation, 2 5 graphene, 117 121 charged impurity scattering, 119 120 contact-induced spin relaxation, 120 121 Hanle spin precession, 118 119, 119f Spin transfer, 10 influence on magnetic tunnel junctions, 87 Spin-transform torque random access memory (STT-RAM), 214 Spin transistor, 83 MRAM in, 17 Spin transport, 2 3, 9 10 in magnetic tunnel junctions, 88 90, 89f Spintronics applications of, 14 17 fundamental aspects of, 1 insulator, 3 metallic, 3 need for, 2 3 semiconductor, 3 Spin valve, 27 Spin valve devices of carbon nanotubes, 96 97 graphene in, 131 134, 132f, 134f Spin-valve geometry, 83 84, 84f SQUID. See Superconducting quantum interface device (SQUID) STM. See Scanning tunneling microscope (STM) STT-RAM. See Spin-transform torque random access memory (STT-RAM) Superconducting quantum interface device (SQUID), 83, 184 186 SWCNTs. See Single-walled CNTs (SWCNTs) T TFTs. See Thin film transistors (TFTs) Thermoremanent magnetization (TRM), 184 186

Index

Thin film transistors (TFTs), 63 66 TRM. See Thermoremanent magnetization (TRM) Tunnel injection, 6 Tunnel magnetoresistance in carbon nanotube-based spintronic devices, 93 97, 95f spin-valve devices of carbon nanotubes, 96 97 in magnetic tunnel junctions, 86

U Ultrananocrystalline diamond (UNCD), 25 27 W Weak localization theory (WLT), 58 60 Z Zigzag graphene nano ribbons (ZGNRs), 108

221