growth and characterization of two-dimensional materials ... · finally, the growth of...

Growth and Characterization of

Two-Dimensional Materials and Their

Heterostructures on SiC

SUBMITTED IN FULFILMENT OF THE REQUIREMENTS FOR THE DEGREE OF

DOCTOR OF PHILOSOPHY

Jonathan BradfordB.AppSc.(Phys)/B.Maths., B.Sc.(Hons)(Phys)

School of Chemistry, Physics and Mechanical Engineering

Science and Engineering Faculty

Queensland University of Technology

2019

Growth and Characterization of Two-Dimensional Mate-rials and Their Heterostructures on SiC

Jonathan Bradford (Email: [email protected])Principal Supervisor: Prof. Nunzio Motta

Associate Supervisor: Dr. Jennifer MacLeod

External Supervisor: Dr. Mahnaz Shafiei

School of Chemistry, Physics and Mechanical Engineering

Science and Engineering Faculty

Queensland University of Technology

Statement of Original Authorship

The work contained in this thesis has not been previously submitted to meet requirements

for an award at this or any other higher education institution. To the best of my knowledge

and belief, the thesis contains no material previously published or written by another

person except where due reference is made.

Signature:

Date:

i

24/10/2019

QUT Verified Signature

Abstract

Two-dimensional (2D) material heterostructures represent an elegant approach to de-

signing and synthesizing nanoscale materials with properties that can be selected to suit

the desired application by making informed choices of the constituent layers. Utilizing

graphene grown epitaxially on SiC as a platform for 2D material heterostructure growth

could facilitate a route to large scale assembly of 2D material devices since the SiC sub-

strate itself can be used as a device substrate. This work explores scalable and transfer-free

growth of 2D material heterostructures based on epitaxial graphene on SiC, and the use of

SiC as a substrate for 2D materials beyond graphene. We focus on the growth of three 2D

material systems on SiC: (1) Lateral heterostructures of graphene and hexagonal boron-

nitride (h-BN); (2) van der Waals heterostructures of transition metal dichalcogenides on

epitaxial graphene; (3) vertically aligned nanosheets of molybdenum disulfide (MoS2).

The synthesis of lateral heterostructures of graphene and hexagon boron nitride (h-

BN) has attracted attention due to the ability to broadly tune the electronic properties of

the hybrid monolayer. Conventional synthesis methods for lateral heterostructures rely

on chemical vapor deposition on metal surfaces, and application of the heterostructure in

devices requires transfer of the material onto a suitable substrate during which the sample

is susceptible to damage or contamination. Herein, a transfer-free synthesis method

to produce lateral heterostructures of graphene and h-BN by chemical conversion of

epitaxial graphene on 6H-SiC(0001) is demonstrated. X-ray photoelectron spectroscopy

measurements confirm the substitution of graphene with h-BN, while scanning tunneling

microscopy reveals that the h-BN domains reside in-plane with graphene forming an

interface along the zig-zag direction. Raman spectroscopy measurements provide insight

into the reaction mechanism in which h-BN is substituted for graphene at defect sites

thereby reducing the defect density in the lateral heterostructure.

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Van der Waals heterostructures of MoS2 and WS2 with epitaxial graphene/SiC have

been synthesized due to the practical appeal of their complementary electronic and op-

tical properties. Chemical vapor deposition (CVD) of MoS2 multilayers on epitaxial

graphene are shown to form either concentric multilayers or spiral structures depending

on nucleation density. MoS2 spiral growth is driven by formation of screw dislocation

defects, and four different types of spiral structures were observed and characterized

by the number of screw dislocations and their chirality. Monolayer and few layer WS2

was synthesized by sulfurization of WO3 thin films deposited on epitaxial graphene by

electron-beam physical vapor deposition. This method allows the growth temperature to

be significantly reduced compared to conventional CVD, however achieving long-range

uniformity remains a challenge.

Finally, the growth of vertically-standing MoS2 nanoflakes on vicinal and on-axis 4H-

SiC substrates was investigated. In both cases the MoS2 flakes exhibit three preferential

orientations, aligning with the 〈112̄0〉 substrate directions due to strain minimizationof the MoO2 intermediate phase. While MoS2 grown on vicinal SiC substrates exhibit

strict vertical alignment, scanning electron microscopy and Near-Edge X-ray Absorption

Fine Structure (NEXAFS) measurements indicate that the vertical orientation of MoS2

grown on on-axis SiC varies. Photoemission spectroscopy and NEXAFS measurements

indicate the presents of defects and disordered edges. We exploit these defects for NO2

gas detection.

This work establishes the relevance of epitaxial graphene as a platform for scalable

synthesis of two-dimensional materials and their heterostructures, and is expected to

facilitate a route towards large-scale synthesis of novel devices directly on-chip.

iv

Keywords

Epitaxial Graphene, h-BN, MoS2, WS2, Transition Metal Dichalcogenides, Van der Waals

heterostructure, Lateral heterostructure, XPS, STM, NEXAFS, Raman spectroscopy

v

Acknowledgments

First and foremost, I owe my deepest gratitude to my supervisors Prof. Nunzio Motta,

Dr. Jennifer MacLeod and Dr. Mahnaz Shafiei, all of whom have offered guidance and

advice throughout the entire project and have always been available to help when I have

needed it. I could not have asked for a more supportive supervisory team.

The data reported in this thesis were obtained at the Central Analytical Research

Facility (CARF) operated by QUT’s Institute for Future Environments. Access to CARF

is supported by generous funding from the Science and Engineering Faculty (QUT). I

would also like to acknowledge the support of all CARF staff for providing access and

assistance with the laboratory equipment. In particular I would like to thank Dr. Josh

Lipton-Duffin, Dr. Peter Hines, Dr. Llew Rintoul, Dr. Jamie Riches.

Part of this research was undertaken on the SXR beamline at the Australian Syn-

chrotron, part of ANSTO. Thank you to Dr. Dongchen Qi for doing PES and NEXAFS

measurements on my samples, and for providing assistance with the NEXAFS analysis.

This work was performed in part at the Queensland node of the Australian National

Fabrication Facility, a company established under the National Collaborative Research

Infrastructure Strategy to provide nano and micro-fabrication facilities for Australia’s

researchers.

I also acknowledge the financial support from the Australian Government’s Research

Training Program and QUT Excellence Top Up Scholarship.

Thank you to all of the past and present group members that I have had to pleasure of

working with for all of your support. In particular, thank you to Dr. Iolanda Di Bernado,

Dr. Nima Khoshsirat and Dr. Maryam Abyazisani all of whom have provided tremendous

support during the project and helped me to keep things in perspective.

vii

Thank you to my family and friends for their patience and support, and providing

much-needed distractions from my work on occasion. Above all, thank you to Ben Blunt

for being my biggest supporter and motivator, and for reminding me to take care of myself

during the past few years.

viii

List of Publications

• J. Bradford, M. Shafiei, J. MacLeod and N. Motta. Transfer-free Synthesis of Lat-eral Graphene-Hexagonal Boron Nitride Heterostructures from Chemically Con-

verted Epitaxial Graphene. Advanced Materials Interfaces. 2019, 1900419.

https://doi.org/10.1002/admi.201900419

Publications not presented in this thesis:

• M. Abyazisani, J. Bradford, N. Motta, J. Lipton-Duffin and J. MacLeod. Adsorp-tion, deprotonation and decarboxylation of isophthalic acid on Cu(111). Langmuir.

2019. https://doi.org/10.1021/acs.langmuir.8b04233

• N. Khoshsirat, J. Bradford, M. Shahbazi, M. Shafiei, H. Wang and N. Motta. Effi-ciency Enhancement of Cu2ZnSnS4 Thin Film Solar Cells By Chromium Doping.

Solar Energy Materials and Solar Cells. 2019. https://doi.org/10.1016/j.solmat.2019.110057

• M. Abyazisani, J. Bradford, N. Motta, J. Lipton-Duffin and J. MacLeod. Adsorp-tion and Reactivity of Pyridine Dicarboxylic Acid on Cu(111). The Journal of

Physical Chemistry C. 2018. https://doi.org/10.1021/acs.jpcc.8b04858

• M. Shafiei, J. Bradford, H. Kahn, C. Piloto, W. Wlodarski, Y. Li and N. Motta.Low-Working Temperature NO2 Gas Sensors Based on Hybrid Two-Dimensional

SnS2-Reduced Graphene Oxide. Applied Surface Science. 2018.

https://doi.org/10.1016/j.apsusc.2018.08.115

• F. Ali, N.D. Pham, J. Bradford, N. Khoshsirat, K. Ostrikov, J. Bell, H. Wang and T.Tesfamichael. Tuning the Amount of Oxygen Vacancies in Sputter-Deposited SnOx

films for Enhancing the Performance of Perovskite Solar Cells. ChemSusChem.

2018. https://doi.org/10.1002/cssc.201801541

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Conference Presentations

• The 9th Vacuum and Surface Science Conference of Asia and Australia. Sydney,Australia (oral presentation)

• NanoSEA 2018 “Nanostructures self-assembly and Nanomaterials”. 2018. Car-queiranne, France (oral presentation)

• QUT-CSIRO Joint Laboratories Workshop. 2018. Brisbane, Australia (poster pre-sentation)

• International Conference on Nanoscience and Nanotechnology. 2018. Wollongong,Australia (oral presentation)

• Wagga 2018 – The 42nd Annual Condensed Matter and Materials Meeting. 2018.Wagga Wagga, Australia (oral presentation)

• Nanostructures for Sensors, Electronics, Energy and Environment. 2018. Brisbane,Australia (poster presentation)

• Nanotechnology and Molecular Science HDR Symposium. 2017. Brisbane, Aus-tralia (oral presentation)

• Materials Research Society Spring Meeting. 2017. Phoenix, USA (oral presenta-tion)

• Wagga 2017 – The 41st Annual Condensed Matter and Materials Meeting. 2017.Wagga Wagga, Australia (poster presentation)

• Joint 13th Asia Pacific Physics Conference and 22nd Australian Institute of PhysicsCongress. 2016. Brisbane, Australia (poster presentation)

• 5th International Symposium on Graphene Devices. Brisbane, Australia (posterpresentation)

• Nanotechnology and Molecular Science HDR Symposium. 2016. Brisbane, Aus-tralia (oral presentation)

x

Table of Contents

Abstract iii

Keywords v

Acknowledgments vii

List of Publications ix

Nomenclature xv

List of Figures xxvii

List of Tables xxix

1 Introduction 1

1.1 Background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

1.1.1 The Library of 2D Materials . . . . . . . . . . . . . . . . . . . . 2

1.1.2 Heterostructures of 2D Materials . . . . . . . . . . . . . . . . . . 4

1.2 Research Problem . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 6

1.3 Thesis Outline . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

2 Literature Review 9

2.1 Epitaxial Growth of Graphene on SiC . . . . . . . . . . . . . . . . . . . 9

2.1.1 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12

2.2 Lateral h-BN/Graphene Heterostructures . . . . . . . . . . . . . . . . . . 13

xi

2.2.1 Tuneable Electronic and Magnetic Properties . . . . . . . . . . . 14

2.2.2 Synthesis Methods . . . . . . . . . . . . . . . . . . . . . . . . . 18

2.2.3 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 30

2.3 Transition Metal Dichalcogenide/Graphene van der Waals Heterostructures 31

2.3.1 Chemical Vapor Deposition of Transition Metal Dichalcogenides 34

2.3.2 Synthesis of TMD/Graphene/SiC Heterostructures . . . . . . . . 36

2.3.3 Electronic and Optoelectronic Properties . . . . . . . . . . . . . 41

2.4 Knowledge Gap and Research Questions . . . . . . . . . . . . . . . . . . 47

3 Methodology and Experimental Details 51

3.1 Methodology and Research Design . . . . . . . . . . . . . . . . . . . . . 51

3.1.1 Preparation of Epitaxial Graphene on SiC . . . . . . . . . . . . . 53

3.1.2 Chemical Conversion of Graphene to h-BN . . . . . . . . . . . . 53

3.1.3 Chemical Vapor Deposition of Transition Metal Dichalcogenides 54

3.1.4 Transition Metal Oxide Thin Film Sulfurization . . . . . . . . . . 55

3.2 Characterization Techniques . . . . . . . . . . . . . . . . . . . . . . . . 55

3.2.1 Scanning Probe Microscopy (SPM) . . . . . . . . . . . . . . . . 56

3.2.2 Photoemission Spectroscopy . . . . . . . . . . . . . . . . . . . . 60

3.2.3 Near-Edge X-ray Absorption Fine Structure (NEXAFS) . . . . . 64

3.2.4 Raman Spectroscopy . . . . . . . . . . . . . . . . . . . . . . . . 66

3.3 Instrumentation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

3.3.1 Omicron Ultrahigh Vacuum (UHV) System . . . . . . . . . . . . 69

4 Transfer-free Synthesis of Lateral Graphene–Hexagonal Boron Nitride Het-

erostructures from Chemically Converted Epitaxial Graphene on SiC 71

4.1 Abstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72

4.2 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72

4.3 Results and Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . 74

xii

4.4 Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 82

4.5 Experimental Details . . . . . . . . . . . . . . . . . . . . . . . . . . . . 82

5 Growth and Characterization of Transition Metal Dichalchogenide/Epitaxial

Graphene van der Waals Heterostructures 85

5.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85


5.2.1 MoS2 Chemical Vapor Deposition . . . . . . . . . . . . . . . . . 87

5.2.2 WO3 Thin Film Sulfurization . . . . . . . . . . . . . . . . . . . 88


5.3.1 Chemical Vapor Deposition of MoS2 . . . . . . . . . . . . . . . 89

5.3.2 WS2 Synthesis by Thin Film Sulfurization . . . . . . . . . . . . . 99

5.4 Chapter Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111

6 Chemical Vapor Deposition of Vertically Standing MoS2 Nanosheets on SiC 113

6.1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 113


6.2.1 MoS2 Synthesis . . . . . . . . . . . . . . . . . . . . . . . . . . . 114

6.2.2 Material Characterization . . . . . . . . . . . . . . . . . . . . . . 115

6.2.3 Gas Sensing . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 116


6.3.1 MoS2/SiC Characterization . . . . . . . . . . . . . . . . . . . . . 116

6.3.2 Growth Mechanism . . . . . . . . . . . . . . . . . . . . . . . . . 123

6.3.3 Gas Sensing Properties . . . . . . . . . . . . . . . . . . . . . . . 127

6.4 Chapter Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 128

7 Conclusions and Future Work 129

7.1 Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 129

7.2 Future Work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 130

xiii

A Supplementary Information for Chapter 4 133

A.1 Apparent Step Height of the h-BN/graphene Interface . . . . . . . . . . . 133

A.2 Raman Spectroscopy Additional Data . . . . . . . . . . . . . . . . . . . 134

B WS2/graphene/SiC High Resolution XPS 135

C CVD Synthesis of WS2 137

D MoS2/SiC Additional NEXAFS Data 139

E Vertical MoS2 Gas Sensors Operating at Elevated Temperatures 141

References 143

xiv

Nomenclature

Abbreviations

AFM Atomic Force Microscopy

ARPES Angle-resolved Photoemission Spectroscopy

CVD Chemical Vapor Deposition

DFT Density Functional Theory

DOS Density of States

FET Field Effect Transistor

FWHM Full Width at Half Maximum

EG Epitaxial Graphene

HREELS High Resolution Electron Energy Loss Spectroscopy

HRTEM High Resolution Transmission Electron Microscopy

IMFP Inelastic Mean Free Path

LBL Layer-by-layer

LDOS Local Density of States

LEED Low Energy Electron Diffraction

LEEM Low Energy Electron Microscopy

MBE Molecular Beam Epitaxy

MOCVD Metal-Organic Chemical Vapor Depostion

NEXAFS Near-Edge X-ray Absorption Fine Structure

PL Photoluminescence

PVD Physical Vapor Deposition

xv

QFEG Quasi-freestanding Epitaxial Graphene

SAED Selected Area Electron Diffraction

SD Screw Dislocation

SDD Screw-disolocation-driven

SEM Scanning Electron Microscopy

SPA-LEED Spot Profile Analysis Low Energy Electron Diffraction

SPM Scanning Probe Microscopy

STM Scanning Tunneling Microscopy

STS Scanning Tunneling Spectroscopy

TEM Transmission Electron Microscopy

TMD Transition Metal Dichalcogenide

UPS Ultraviolet Photoemission Spectroscopy

XPS X-ray Photoemission Spectroscopy

VBM Valence Band Maximum

xvi

List of Figures

1.1 Classification of various 2D materials according to their bandgap and the accord-

ing optical wavelength. The shaded grey bars indicated the variability in reported

values. Reproduced from [10], with the permission of Nature Publishing Group. 3

1.2 Representation of 2D materials using building blocks, and their ability to be

stacked accordingly into artificial van der Waals heterostructured materials. Re-

produced from [55], with the permission of Nature Publishing Group. . . . . . . 5

1.3 Trade-off between quality and cost for common methods of graphene production,

and their suitability towards specific applications. Reproduced from [93], with

the permission of Nature Publishing Group. . . . . . . . . . . . . . . . . . . 7

2.1 (a) Atomic model of the bottom-up growth of graphene on Si-terminated SiC

(0001). At sufficiently high temperature Si atoms sublimate resulting in layer-by-

layer growth of graphene passivated by a buffer layer partially covalently bonded

to the substrate; (b-e) LEED patterns indicating the evolution of the SiC (0001)

surface through (1 × 1) (b), (√3 × √3) (c) and (6√3 × 6√3) reconstructionsbefore graphene formation (d),(e) [14]; (f-i) AFM and LEEM images showing

the improvement of the graphitized SiC surface morphology for UHV grown

graphene (f),(g) and graphene grown at atmospheric pressure (h),(i). Reproduced

from [111], with the permission of Nature Publishing Group. . . . . . . . . . . 11

2.2 Atomic models of monolayer graphene and hexagonal boron nitride. The brown,

green and grey spheres represent carbon, boron and nitrogen atoms, respectively. 13

xvii

2.3 Calculated band structures of graphene doped with h-BN with concentrations of

(a) 0%, (b) 25%, (c) 50%, (d) 75%, and (e) 100%. The calculated bandgap is

plotted as a function of h-BN concentration in (f). Reproduced from [75], with

the permission of RSC Publishing. . . . . . . . . . . . . . . . . . . . . . . . 15

2.4 (a) Interface energies of different terminations of graphene embedded in h-BN

plotted as a function of chemical potential. Blue, purple and red represent B-

terminated zigzag, armchair, and N-terminated zigzag interfaces respectively.

The dotted line plots the magnetism per unit perimeter as change equilibrium

graphene crystal shape changes with chemical potential. The crystal shapes are

illustrated in (b)-(f) and the arrow lengths are proportional to the magnitude of

magnetism. Adapted with permission from [127]. Copyright (2011) American

Chemical Society. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

2.5 STM image of an atomically resolved h-BN/graphene domain boundary ex-

hibiting a 0.1 nm height difference (line profile inset) due to the difference in

the density of states. Reproduced from [77], with the permission of American

Association for the Advancment of Science. . . . . . . . . . . . . . . . . . . 19

2.6 STM image a h-BN/graphene domain boundary showing perturbations to the

moiré pattern due to misfit dislocations (left panel). The right panel shows an

atomically resolved misfit dislocation (MD) caused by the lattice mismatch of

graphene and h-BN. Adapted with permission from [144]. Copyright (2014)

American Chemical Society. . . . . . . . . . . . . . . . . . . . . . . . . . 21

2.7 Schematic illustration of the patterned regrowth procedure for h-BN/graphene

lateral heterostructures with controlled domain shapes and sizes. Reproduced


2.8 (a),(b) 40 nm × 40 nm STM images of polymerized products of monomers 1and 2, respectively, as shown schematically in (c). Adapted with permission

from [161]. Copyright (2015) American Chemical Society. . . . . . . . . . . . 24

xviii

2.9 Mechanism of the graphene to BN conversion reaction. (a) OH functionalized

C atom in graphene; (b) Nitrogen substitution in graphene; (c) BN domain

nucleation; (d) Extension of the BN domain; (e) Merging of multiple BN do-

mains into a singular irregularly shaped region; (f) Totally converted h-BN film.

Reproduced from [169], with the permission of Nature Publishing Group. . . . . 27

2.10 (a) Raman mapping of the graphene 2D peak of a graphene/h-BNC/h-BN hybrid

film (scale bar 10 μm). The colour spectrum represents the 2D peak intensity;

(b) Carrier mobility vs. on/off ratio for h-BNC, graphene, and MoS2 FETs.


2.11 STM images showing the evolution of the epitaxial h-BNC film from h-BN to

graphene by high temperature annealing. (a) h-BN layer with a (5 × 5) super-structure; (b)-(d) h-BNC layer after annealing at 1250◦C, 1350◦C and 1450◦C

respectively, the light regions correspond to h-BN and the dark regions indicate

graphene domains; (e) graphene with R0◦ after annealing at 1600◦C. The (6√3×

6√3)R30◦ structure is marked in red; (f) and (g) high resolution images of h-

BN/graphene boundary showing an atomically sharp interface in (f). Adapted

with permission from [172]. Copyright (2015) American Chemical Society. . . . 29

2.12 Evolution from h-BN to h-BNC to graphene on 6H-SiC (0001). h-BN is de-

posited on the SiC surface and partially decomposed and replaced by graphene

domains by high temperature annealing. At sufficient temperatures the h-BN

layer is completely replaced by graphene and a (6√3 × 6√3)R30◦ interface

layer is formed beneath the R0◦ graphene layer. Adapted with permission from

[172]. Copyright (2015) American Chemical Society. . . . . . . . . . . . . . 30

2.13 Atomic structure of transition metal dichalcogenides when viewed along (a) the

c-axis, and (b) the a-axis. The purple and yellow spheres represent transition

metal atoms (e.g. Mo or W) and chalcogen atoms (S, Se or Te), respectively;

(c),(d) Calculated band structures of monolayer and bulk MoS2 and WS2, re-

spectively. (c) and (d) reproduced from [22], with the permission of Nature

Publishing Group. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 33

xix

2.14 Four growth routes of monolayer TMDs grown by CVD on SiO2 obtained through

controlling the mass flux of the metal precursor and the growth rate. Reproduced


2.15 STM images of MoS2 grown by deposition and sulfurization of Mo on graphene/SiC

[200]. (a) 0.55 ML coverage of MoS2 coexisting with bilayer regions, Mo

particles and Mo6S6 aggregates; (b) 0.85 ML coverage with a higher density

of bilayer regions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 37

2.16 (a) STM images of the spiral growth of MoS2 at a graphene/SiC step edge; (b)

Atomic resolution STM image of MoS2 showing the 3.1Å periodicity; (c)-(f)

Growth mechanism for MoS2 spirals at a step edge. Reproduced from [201],

with the permission of AIP Publishing. . . . . . . . . . . . . . . . . . . . . 39

2.17 (a) Schematic represtation of the CVD system and growth parameters for WS2

synthesis on epitaxial graphene; (b) AFM image of a WS2 monolayer; (c-e)

SEM images showing increasing coverage with the reaction time, and extension

of the lateral dimensions of WS2 domains. Reproduced from [211], with the

permission of IOP Publishing. . . . . . . . . . . . . . . . . . . . . . . . . . 41

2.18 (a) STM image monolayer MoS2 on epitaxial graphene; (b) Nine STS spectra

measured different locations on the MoS2 monolayer in (a); (c) STM image of

multilayer MoS2 on epitaxial graphene; (d) Layer dependent STS measurements

of MoS2 taken at the positions indicated in (c). [201]. (a) and (b) were adapted

with permission from [202]. Copyright (2016) American Chemical Society. (c)

and (d) were reproduced from [201], with the permission of AIP Publishing. . . 42

2.19 (a) E12g and A1g Raman peaks for MoS2 and graphene/MoS2 devices; (b) Trans-

fer curves for graphene/MoS2 photodetectors under different illumination inten-

sities. Reproduced from [230], with the permission of Nature Publishing Group. 45

2.20 Tuneable photoresponse of MoS2/graphene/SiC heterostructure [194]. . . . . . 46

xx

2.21 Device fabrication and wavelength-dependent photoresponse of WS2/graphene/SiC

van der Waals heterostructures. Graphene regions are masked with hydrogen

silsesquioxane (a) before WS2 deposition (b). This is followed by spin-coating

with PMMA (c) and etching with potassium iodide or hydrofluoric acid to re-

move the mask and expose the graphene regions (d). Finally gold electrodes are

deposited on the exposed graphene to form an ohmic contact (e) and the final

device structure is shown in (f). The photoresponse at three difference excitation

wavelengths is shown in (g). Reproduced from [232], with the permission of

RSC Publishing. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 47

3.1 (a) Optical microscopy image of hydrocarbon bubbles (bright spots) in trans-

ferred stacks of graphene and h-BN (scale bar is 20 μm; (b) SEM image of the

contamination which appears as dark contrast; (c) Cross-sectional TEM image

of a clean area showing individual graphene and h-BN layers (scale bar is 2 nm).


3.2 Schematic of the experimental setup for the chemical conversion of graphene to

h-BN. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54

3.3 Schematic of the experimental setup for chemical vapor deposition of transition

metal dichalcogenides on epitaxial graphene. . . . . . . . . . . . . . . . . . . 55

3.4 Schematic of the WS2 synthesis by thin film sulfurization procedure. . . . . . . 56

3.5 (a) Schematic of a tunneling junction between an STM tip and a surface; (b)

Energy level diagram of the tunneling junction. EFt and EFs are the Fermi

levels of the tip and sample, respectively, Evac is the vacuum energy level, ρ

is the density of state, φ is the work function, d is the distance between the tip

and sample, and V is the bias voltage [234]. Originally published by The Royal

Society of Chemistry. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57

3.6 (a) (200 × 200) nm2 STM image of the SiC surface after graphitization in UHV(U = −1.61 V, I = 0.7 nA); (b) (6.7 × 6.7) nm2 atomic resolution STM imageof epitaxial graphene on SiC (U = −0.1 V, I = 0.4 nA). . . . . . . . . . . . . 58

3.7 Force-distance curve showing the different modes of AFM [238]. . . . . . . . . 60

xxi

3.8 Energy level diagram of the x-ray photoelectron transitions. The subscripts

s and XPS denote quantities for the sample and the instrument, respectively

Reproduced from [239], with the permission of the John Wiley and Sons. . . . . 61

3.9 Inelastic mean free path (IMFP) of electrons (pink band) with a given kinetic

energy. The data points represent experimental data for the IMFP of common

semiconductor materials. Reproduced from [240], with the permission of the

John Wiley and Sons. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62

3.10 (a) Survey, (b) C 1s and (c) Si 2p spectra measured by XPS from epitaxial

graphene on SiC. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63

3.11 (a) UPS valence band spectrum of epitaxial graphene on SiC; (b) Valence band

spectrum of epitaxial graphene near the Fermi level, indicated by the dashed box

in (a). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63

3.12 Electron transitions following photoabsorption of a soft x-ray. Reproduced from

[242], with the permission of the Royal Society of Chemistry. . . . . . . . . . 64

3.13 Origin of spectral features observed in NEXAFS. Reproduced from [242], with

the permission of the Royal Society of Chemistry. . . . . . . . . . . . . . . . 65

3.14 Energy level diagram showing the transitions involved in Raman scattering [245]. 67

3.15 Raman spectrum of Epitaxial graphene/SiC after removal of the SiC substrate

structure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 68

3.16 Photograph of the Omicron UHV System. . . . . . . . . . . . . . . . . . . . 69

xxii

4.1 (a) (500×500) nm2 STM image of epitaxial graphene/SiC (U=-2.0 V; I=1.0 nA);(b) Atomic resolution STM image of epitaxial graphene showing the graphene

lattice superimposed on the (6√3× 6√3)R30◦ moiré pattern (U = −0.2 V; I =

1.2 nA). The (6√3×6√3)R30◦ and graphene unit cells are marked in green and

red, respectively; (c) Surface morphology after the chemical conversion reaction

of graphene to h-BN (U=-0.2 V; I=0.4 nA); (d) Atomically resolved image of

the h-BN/graphene interface (U=-0.02 V; I=0.6nA); (e) Line profile across the

h-BN/graphene interface and adjacent step edge along the direction indicated

in the inset; (f) Proposed cross-sectional model of the lateral h-BN/graphene

heterostructure on SiC based on the profile in (e). Brown, blue, green and grey

spheres represent carbon, silicon, boron and nitrogen atoms, respectively. . . . . 75

4.2 (a) FFT of the lateral h-BN/graphene interface shown in Figure 4.1(d); (b) LEED

pattern of the hBN-graphene/SiC acquired with a beam energy of 150 eV. . . . . 77

4.3 (a) C 1s, (b) N 1s, and (c) B 1s XPS core level spectra of BN-substituted epitaxial

graphene on SiC. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78

4.4 (a) Comparison of the Raman spectra of epitaxial graphene (top) and hybrid h-

BN/epitaxial graphene (bottom); (b) Deconvolution of the D band after chemical

conversion of graphene (highlighted by the dotted red box in (a)). . . . . . . . . 80

5.1 Heating ramps applied to the two-zone furnace for MoS2 synthesis. The green

line applies to the downstream zone containing the MoO3 powder and epitaxial

graphene substrate, and the yellow line is applied to the upstream zone containing

the sulfur powder. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 88

5.2 Heating ramps applied to the two-zone furnace for WS2 synthesis. The green

line applies to the downstream zone containing the WO3/graphene/SiC sample,

and the yellow line is applied to the upstream zone containing the sulfur powder. 89

xxiii

5.3 (a)-(c) AFM topography images of MoS2 grown on epitaxial graphene with

increasing carrier gas flow rates; (d)-(f) STM images showing greater detail of

the MoS2 layer topography in images (a)-(c), respectively. The bias voltage and

tunneling current setpoints are (d) U = 1.50 V, I = 0.6 nA, (e) U = 0.96 V,

I = 0.4 nA, and (f) U = 0.50 V, I = 0.5 nA; (g) Atomic resolution image of

an exposed graphene region (U = 0.50 V, I = 0.5 nA); (g) Atomic resolution

image of MoS2 (U = −0.03 V, I = 0.5 nA); (h) (U = −1.25 V, I = 1.3 nA);(i) STM tunneling current image corresponding the topography image in (f). . . 91

5.4 (a) Double-arm MoS2 spiral appearing to originate from a single screw disloca-

tion (U = −1.50 V, I = 0.5 nA); (b) Closed-loop growth of MoS2 caused byopposing screw dislocations (U = 0.97 V, I = 0.4 nA). . . . . . . . . . . . . 93

5.5 Raman spectra of MoS2 multilayers (top) and spirals (bottom) grown on epitaxial

graphene. Schematic models of the vibrational modes are inset. The yellow and

purple spheres represent sulfur and molybdenum atoms, respectively. . . . . . . 94

5.6 Mo 3d (a) and S 2p (b) XPS core level spectra for MoS2/graphene/SiC. . . . . . 96

5.7 High resolution XPS C 1s and Si 2p core levels after MoS2 synthesis. (a) and

(b) show the spectra of a sample that is unaffected by MoS2 growth, and (c)

and (d) show the spectra of a different sample that shows evidence of oxygen

intercalation after MoS2 synthesis. The insets in each panel show a comparison

of the core level peak before the growth (black curve) and after the growth (red

curve) of MoS2. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 98

5.8 (a) STM image of an exposed graphene region in a non-intercalated MoS2/graphene/SiC

sample; (b) is taken in the region marked by the blue dashed box and shows the

(6√3× 6√3)R30◦ moiré pattern of epitaxial graphene. The unit cell is drawn in

green (U = −1.00 V, I = 0.7 nA). . . . . . . . . . . . . . . . . . . . . . . 98

5.9 (a),(b) XPS W 4f and O 1s core level spectra of WO3−x/graphene SiC; (c) AFM

topographic image of the WO3 surface; (d) Tip height line profile taken along

the blue line in (c). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100

xxiv

5.10 (a) STM image of monolayer WS2 on epitaxial graphene. The WS2 is outlined

by the dashed white line (U = −1.50 V, I = 1.0 nA); (b) Multilayer WS2 ona different region of the sample (U = −1.50 V, I = 0.7 nA); (c) Distributionof step heights measured on STM images of WS2/graphene/SiC ; (d) Atomic

resolution image of graphene. The unit cell is marked by the green rhombus

(U = −0.08 V, I = 0.1 nA); (e) Atomic resolution STM image of monolayerWS2 on epitaxial graphene. The unit cell of the WS2 lattice is drawn in magenta,

and the unit cell of the moiré pattern is drawn in blue (U = −1.57 V, I = 0.1nA); (f) STM image of multilayer WS2 on graphene showing the WS2 lattice

without a moiré pattern (U = −1.67 V, I = 1.2 nA). . . . . . . . . . . . . . . 102

5.11 XPS survey spectra of at each stage of the WS2 growth procedure. . . . . . . . 103

5.12 High resolution XPS core levels of WS2/graphene/SiC. (a) and (b) S 2p and W 4f

core levels, respectively, after WS2 synthesis; (c) C 1s core level of bare epitaxial

graphene (top) and WS2/graphene/SiC (bottom). . . . . . . . . . . . . . . . . 105

5.13 Valence band spectrum of WS2/graphene/SiC. The inset shows a higher magni-

fication of the spectrum near the Fermi level, with the blue spectrum for the bare

graphene sample included for comparison. . . . . . . . . . . . . . . . . . . . 106

5.14 Raman spectrum of WS2 grown on epitaxial graphene. The spectrum of the bare

epitaxial graphene is shown in the top right. . . . . . . . . . . . . . . . . . . 108

5.15 Comparison of the W 4f core level spectra at different sulfurization temperatures. 109

5.16 High resolution XPS core level spectra of partially sulfurized WOxSy. (a) Com-

parison of the W 4f core level before (top) and after (bottom) sulfurization;

(b),(c) S 2p and O 1s core level spectra, respectively, after sulfurization. . . . . . 110

6.1 Heating ramps applied to the two-zone furnace for MoS2 synthesis on 4H-SiC.

The green line applies to the downstream zone containing the MoO3 powder and

epitaxial graphene substrate, and the yellow line is applied to the upstream zone

containing the sulfur powder. . . . . . . . . . . . . . . . . . . . . . . . . . 115

xxv

6.2 (a),(b) Scanning electron micrographs of MoS2 nanosheets grown on 4◦ off-axis,

and on-axis 4H-SiC, respectively. The scale bars in (a) and (b) are 10 μm and

5 μm, respectively. The corresponding distrubtions of the flake orientations are

shown below in (c) and (d). Angles are measured with respect to the image

horizontal, and are shifted to lie in a 0-180◦ ; (e),(f) Top view and side view of

the oriented, vertically MoS2 nanosheets on 4H-SiC. The blue, brown, purple

and yellow spheres represent silicon, carbon, molybdenum and sulfur atoms,

respectively. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 117

6.3 (a) AFM topography image of vertical MoS2 on SiC; (b) AFM tip height profile

taken along the white dotted line in (a); TEM image of MoS2 flakes viewed along

the MoS2 [0001] zone axis; (d) TEM image of the MoS2 flake edge. . . . . . . 118

6.4 Raman spectrum of vertically aligned MoS2 nanosheets grown on 4H-SiC. . . . 119

6.5 High resolution Mo 3d (a) and S 2p (b) core level specta for MoS2 nanosheets . 120

6.6 (a) Representative synthetic fit to the sulfur K-edge NEXAFS spectrum; (b),(c)

Calculated sulfur K-edge absorption spectrum, and isosurfaces representing the

excited states contributing to the absorption features. The states in (c) i, ii and iii

are represented by the blue (i,ii) and cyan (iii) peaks in the fit to the experimental

data in (a), whereas the yellow, orange and green components do not necessarily

represent a particular excited state. (b) and (c) are unofficial adaptations of an

article that appeared in an ACS publication [330]. ACS has not endorsed the

content of this adaptation or the context of its use. . . . . . . . . . . . . . . . 122

6.7 Angular dependence of the sulfur K-edge NEXAFS spectra for MoS2 nanosheets

grown on (a) off-axis and (b) on-axis SiC substrates. The x-ray angle of inci-

dence is measured with repect to the sample surface as shown schematically in

(a). Intensity variations for the primary absorption peak are shown in the insets

with fits to equation (6.2) . . . . . . . . . . . . . . . . . . . . . . . . . . . 124

6.8 (a) Raman spectrum of an incompletely sulfurized MoS2 fin; (b) TEM image

of an incompletely sulfurized MoS2 (scale bar is 20 nm); (c) HRTEM image

showing the MoO2 lattice and MoS2 layers. . . . . . . . . . . . . . . . . . . 125

6.9 NO2 gas sensing response of vertical MoS2 nanosheets at room temperature. . . 127

xxvi

A.1 Apparent step height of the h-BN/graphene lateral interface. The example line

profile in (a) was taken from left to right along the green line in the inset STM

image. The apparent step height as a function of the bias voltage is plotted in (b).

All measurements were consistently taken along the peaks of the moiré pattern. . 133

A.2 Raman G band peak fitting demonstrating a reduction in the D’ peak intensity af-

ter the chemical conversion reaction which supports the analysis that the overall

defect density is lower. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 134

B.1 High resolution XPS core levels at each stage of the WS2/graphene/SiC process. 136

C.1 Raman spectrum of CVD grown WS2 (left); AFM phase image of WS2 spirals

(right). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 137

D.1 (a),(b) Sulfur K-edge NEXAFS spectra for MoS2 grown on off-axis and on-axis

4H-SiC, respectively; (c),(d) Normalised photon flux measurements recorded

during the NEXAFS measurments shown in panels (a) and (b), respectively. . . 139

E.1 NO2 gas sensing reponse of MoS2 nanoflakes at operating temperatures of (a)

150 ◦C, (b) 200 ◦C and (c) 250 ◦C; (d) Peak-to-valley sensing reponse as a

function of NO2 gas concentration at different operating temperatures. . . . . . 142

xxvii

xxviii

List of Tables

5.1 Raman spectroscopy peak data for different morphologies of multilayer

MoS2/graphene/SiC. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 95

A.1 Peak positions, widths and intensity ratios of epitaxial graphene before

and after chemical conversion to h-BN. . . . . . . . . . . . . . . . . . . 134

xxix

Chapter 1

Introduction

1.1 Background

The advent of any new technology is often examined in terms of whether it is an incremen-

tal change, usually represented by scaling of existing technologies, or disruptive change,

i.e. breakthroughs resulting in broad-base adoption of previously out-of-reach technology.

Perhaps the best example of such phenomena have occurred in the field of computing, in

which a shift from vacuum tubes to semiconductor technology revolutionized the field,

and the technology has since been scaled following Moore’s law by reducing the dimen-

sions of transistors on a chip. In Herbert Kroemer’s Nobel Prize Lecture in 2000, he put

forward what has come to be known as the Lemma of New Technology: The principal

applications of any sufficiently new and innovative technology always have been — and

will continue to be — applications created by that technology [1].

To that end, two dimensional (2D) materials have provided a plethora of exciting

avenues of exploration in terms of both fundamental physics, which may lead to fascinat-

ing new technological applications, and scaling of existing technology [2]. In particular,

various 2D materials have predicted functionality in spintronics [3], valleytronics [4] and

superconductivity [5]. Moreover, they are also expected find applications in existing

technology areas such a chemical and environmental sensing [6], catalysis [7], energy

storage [8], nanoelectronics [9] and optoelectronics [10].

1

2 CHAPTER 1. INTRODUCTION

1.1.1 The Library of 2D Materials

Undoubtedly, the rise of 2D materials as field of research can be attributed to the isolation

of graphene by exfoliation from highly ordered pyrolitic graphite (HOPG) in 2004 [11].

This sparked a surge of research into the material which exhibits some extraordinary

electronic [12–14], optical [15] and mechanical [16] properties. Immediately following

the exfoliation of graphene, researchers turned their attention to other layered crystals

which consist of monolayers held together by van der Waals interactions, akin to graphite.

In particular, focus was drawn to hexagonal boron nitride (h-BN) [17], transition metal

dichalcogenides (TMDs) [12, 18], and black phosphorus [19, 20]. These materials are also

able to be exfoliated to produce monolayers with vastly different properties to graphene.

Hexagonal boron nitride has a similar lattice structure to graphene, with boron and ni-

trogen replacing alternating carbon atoms, yet it is a wide bandgap (5.9 eV) insulator

[21]. Black phosphorus and common TMDs like molybdenum disulfide (MoS2), on the

other hand, are semiconducting with layer dependent bandgaps [22]. Beyond bulk van der

Waals materials, group IV materials such as silicon, germanium and tin have also been

explored. These materials do not form bulk sp2-hybridized layered crystals analogous

to graphite and therefore cannot be exfoliated. Nonetheless, 2D forms of the materials

were predicted to exist in a buckled honeycomb lattice and host Dirac fermions similar

to graphene [23–27], and these materials have since been synthesized by molecular beam

epitaxy on various metal surfaces [28–31], which is a requirement to stabilize the sp2-sp3

like bonds.

Thus, over the past fifteen years scientists have developed a library of 2D materials

with a variety of properties. Figure 1.1 summarizes some of the most common 2D materi-

als and classifies them based on their electronic properties [10]. Across the graphene and

monoelemental 2D material family, dichalcogenides, chalcogenides and 2D oxides energy

bandgaps can be chosen covering the entire range from metallic through to insulating.

Even among established materials such as graphene, there remains an opportunity to

manipulate the properties of the materials in the search for exotic electronic and quantum

phenomena. An example of such manipulation can be seen in the ability produce non-

dispersive flat bands and van Hove singularities in twisted graphene bilayers [32–34]

which has led to demonstrations of tuneable superconducting and correlated insulating

1.1. BACKGROUND 3

phases [35–37]. But this is still only scratching the surface. Recently, computational

exploration of 5619 layered compounds has suggested that as many as 1036 of these could

be isolated as monolayers by simple exfoliation [38]. Following further investigation of

the material properties of a subset of 258 compounds, 166 were found to be semicon-

ducting and 56 were identified as ferromagnetic or antiferromagnetic systems. Thus it is

clear that 2D material research remains a field rich with possibilities at fundamental and

applied levels.

(a)

Figure 1.1: Classification of various 2D materials according to their bandgap and the accordingoptical wavelength. The shaded grey bars indicated the variability in reported values. Reproducedfrom [10], with the permission of Nature Publishing Group.

Despite a wealth of new physics and the vast number of two dimensional materials,

it is unsurprising that there is no “one size fits all” material with properties to suit all

applications. For example, while graphene is extremely attractive for applications in nano-

electronics due to its extremely high carrier mobility, the lack of bandgap limits its use as

a logic device. It is widely acknowledged that, in order to fulfill the requirements for both

radio frequency (RF) and logic devices, it is imperative to find a way to make graphene


semiconducting [2]. While many methods have been developed to open a bandgap -for

example quantum confinement of electrons in graphene nanoribbons (GNRs) [39–43],

substrate induced bandgap opening [44–46] and chemical functionalisation [47, 48] –

successfully opening a bandgap in graphene without degrading its other properties (in

particular the carrier mobility) can be challenging. On the other hand, semiconducting

TMDs such as MoS2, WS2, MoSe2 and WSe2 offer a direct optical bandgap in the mono-

layer [22] making them an attractive material for logic devices and optoelectronic devices

[18, 49, 50]. In fact, the performance of field effect transistors (FET) with MoS2 channels

have been demonstrated to meet the International Roadmap for Semiconductors (ITRS)

2026 requirements for low operating power [9]. Nonetheless, common semiconducting

TMDs are fundamentally limited by the effective masses of electrons and holes [51], and

longitudinal optical phonon scattering [52] which may prevent achieving high switching

speeds. Beyond graphene and TMDs, other elemental 2D materials such as phosphorene

and silicene suffer from poor long-term stability [53, 54]. In many cases the properties of

different materials are complementary and it is therefore desirable to combine them in a

heterostructure in order to overcome their individual limitations.

1.1.2 Heterostructures of 2D Materials

The concept of 2D material heterostructures is straightforward: given the library of 2D

materials we can use different materials like building blocks and either stack them to

construct artificial van der Waals materials [55] (as demonstrated in Figure 1.2), or stitch

them laterally into a two dimensional patchworks [56]. Making informed choices of the

constituent materials can allow clever engineering of the properties of the heterostructure

which, in principle, could tailored to suit specific applications. This offers a versatile

technique to overcome many of the limitations of individual 2D materials.

Pioneering investigations of van der Waals heterostructures involved stacking of mono-

layer graphene onto exfoliated h-BN where a remarkable improvement in electronic trans-

port in the graphene layer was demonstrated [57]. This established h-BN as an ideal

substrate for 2D materials due to its atomically flat surface free of dangling bonds and

charge traps which act as scattering sites [57, 58]. In addition, full encapsulation of 2D

materials with h-BN allows the materials to exhibit properties close to their freestanding

1.1. BACKGROUND 5

Figure 1.2: Representation of 2D materials using building blocks, and their ability to be stackedaccordingly into artificial van der Waals heterostructured materials. Reproduced from [55], withthe permission of Nature Publishing Group.

counterparts and, moreover, can provide protection against ambient conditions to improve

long-term stability [59–66]. Measurements of interlayer tunnelling in stacks of graphene,

h-BN and TMDs also allow us to envisage tunnelling field effect transistors in next

generation nanoelectronics [67–69]. Lateral heterostructures of 2D materials is another

promising route to combining complementary properties of 2D materials. Structural

analogues graphene and h-BN have been combined in a lateral heterostructure to produce

a hybrid monolayer film with tuneable electronic properties [70–75], and structured into

atomically thin electronic devices [76, 77]. Similarly, lateral epitaxy of TMDs with a

relatively small lattice mismatch have been adopted to produce WS2–WSe2 and MoS2–

MoSe2 monolayer p-n junctions [78–80].

Adopting various approaches to combining the properties of 2D materials has been

demonstrated to improve device performances across traditional applications of 2D ma-

terials including nanoelectronics [81, 82], optoelectronics [81, 82], energy storage [8,

83], and sensing [6]. In addition, following Kroemer’s Lemma of New Technology,

many predictions have been made into potential uses of 2D materials heterostructures

in emerging technology in quantum optics [84], exploiting valleytronics for quantum

computing [85], or designing high-Tc superconductors [5, 55].


1.2 Research Problem

In recent years research into fabrication of 2D material heterostructures has been exten-

sive, yet there are still challenges that need to be addressed. Synthesis of 2D materials,

and in turn their heterostructures often entails a trade-off between quality, scalability,

and cost. This is demonstrated in Figure 1.3 using graphene as an example. Mechanical

exfoliation of 2D materials from their bulk counterparts, such as graphene exfoliated from

graphite, enables extremely high quality layers to be isolated; however this approach is

not scalable and is typically only suitable for proof-of-concept applications in research.

In contrast, liquid-phase exfoliation enables mass production with a low cost, but comes

with the trade-off of producing lower quality material. Nonetheless this approach is

suitable for applications in areas such as energy storage and sensing where high quality

is not a requirement, and in many cases provides and advantage. In order to produce

higher quality layers at larger scales bottom-up approaches are used such as chemical

vapor deposition (CVD) or, in the case of graphene, thermal decomposition of SiC. Large

scale production of high quality of graphene on Cu substrates can be achieved by CVD;

however, the transfer process onto a suitable device substrate is known to hinder its

remarkable electronic properties [86–90]. It is therefore desirable to develop transfer-

free approaches to 2D material synthesis onto industrially relevant substrates, as can be

done for transition metal dichalocogenides by CVD onto SiO2 and sapphire [91, 92].

To that end, epitaxial graphene growth by thermal decomposition of SiC substrates

represents a promising platform for graphene growth and subsequent heterostructure syn-

thesis. Epitaxial graphene can be produced at the wafer scale by high temperature anneal-

ing of SiC substrates, and the annealing conditions can be used to control of the number

of layers [94, 95]. Furthermore, the SiC substrate is a wide bandgap semiconductor which

is suitable as a device substrate and is compatible with existing semiconductor processing

techniques [96, 97].

The aim of this thesis is to explore epitaxial graphene on SiC as a platform for transfer-

free synthesis of lateral and vertical heterostructures of 2D materials. In particular, two

heterostructure systems are considered using epitaxial graphene on SiC as a substrate. The

first is a lateral heterostructure of graphene with hexagonal boron nitride (h-BN) which

offers the ability to tune the electronic properties almost continuously between that of

1.3. THESIS OUTLINE 7

Figure 1.3: Trade-off between quality and cost for common methods of graphene production,and their suitability towards specific applications. Reproduced from [93], with the permission ofNature Publishing Group.

graphene (semi-metallic) to h-BN (insulating), and has the potential to be developed into

atomically thin functional devices. Direct synthesis of van der Waals heterostructures of

transition metal dichalcogenides (TMDs) and graphene are also explored in the context of

developing an understanding of the crystal growth to tune the morphology from large area

mono- or multilayers, multilayer spirals, and vertical fins. Heterostructures of graphene

and TMDs enable enhanced performances in optoelectronics and chemical sensing, where

the morphology of the layers plays a critical role [98]. The methods developed in this

work represent a bottom-up approach to incorporating the three types of 2D materials

with vastly different electronic properties into heterostructures directly on a device-ready

substrate at the wafer scale. Doing so creates a pathway towards rational design and

large-scale production of a range of functional devices directly on-chip.

1.3 Thesis Outline

This thesis has been presented as a thesis by monograph, however Chapters 4, 5 and 6 are

written as standalone pieces of work. The outline of the thesis is as follows:


Chapter 2 provides an overview of the literature to develop the specific research

points of the thesis. First, graphene formation by thermal decomposition of SiC will

be discussed in the context of producing controlled, high quality graphene suitable for

subsequent heterostructure growth. Section 2.2 will establish the theoretical framework

for lateral heterostructures of graphene and h-BN, before discussing the experimental

synthesis techniques reported in the literature. Finally, in Section 2.3 van der Waals het-

erostructures of graphene and transition-metal dichalcogenides (TMDs) will be discussed

in Section 2.3. After introducing chemical vapor deposition (CVD) of TMDs as a general

technique, focus will be shifted to van der Waals heterostructure of TMDs on epitaxial

graphene/SiC.

Chapter 3 outlines the experimental approach to 2D material synthesis used in this

work, and provides details of the characterization techniques that have been applied to

study the systems.

Chapter 4 presents growth of lateral heterostructures of h-BN and graphene using

a substitution reaction of h-BN into epitaxial graphene. The lateral heterostructures

are characterized and insight is provided into the reaction mechanism. The content of

this chapter has been accepted for publication to Advanced Materials Interfaces (DOI:

10.1002/admi.201900419).

Chapter 5 demonstrates two different approaches to synthesizing van der Waals het-

erostructures of transition metal dichalcogenides MoS2 and WS2 on epitaxial graphene:

first by chemical vapor deposition and second by sulfurization of metal oxide thin films.

Chapter 6 presents the growth of vertically aligned MoS2 nanosheets by chemical

vapor deposition, and discusses the reaction mechanism and application to NO2 gas de-

tection.

Chapter 7 provides a summary of the key findings presented in this thesis, and several

avenues to extend the research further are presented.

Chapter 2

Literature Review

2.1 Epitaxial Growth of Graphene on SiC

Graphitization of SiC substrates by high temperature annealing has been demonstrated as

an effective method of producing a high quality graphene at the wafer scale [99–101].

The formation of graphene on Si-terminated SiC surfaces is illustrated in Figure 2.1.

The SiC substrate, typically 4H- or 6H-SiC (0001) (shown in Figure 2.1(a)), consists

of bilayers of silicon and carbon atoms. Si sublimation at high temperatures result in a

carbon-rich surface which rearranges to form graphene [102, 103]. On the Si-terminated

face, silicon sublimation first produces a carbon-rich interfacial layer (or buffer layer).

Although the buffer layer is structurally similar to graphene it does not exhibit properties

typical of graphene due to bonding of one third of the carbon atoms to the SiC substrate.

As the annealing time progresses silicon sublimation continues and this buffer layer is

transformed into the first epitaxial graphene layer while a new buffer is formed below

the graphene monolayer (see Figure 2.1). This process continues to form more graphene

layers as the annealing time increases. The number of graphene layers can be controlled

according to the annealing temperature and time [94, 95]. On the C-terminated face

graphitization of the SiC surface occurs without a buffer layer and growth occurs much

faster than on the Si-terminated face.

Throughout the graphene growth process the surface undergoes a series of recon-

structions depending on the specific growth conditions and the termination of the SiC

face. On the Si-terminated face, the surface begins with either a (3× 3) reconstruction or

9

10 CHAPTER 2. LITERATURE REVIEW

(√3×√3)R30◦ reconstruction. The (3×3) reconstruction occurs usually with the use of

silicon flux which deposits an adlayer of silicon atoms on top of the bulk SiC. In less Si-

rich conditions the (√3 × √3)R30◦ reconstruction is more favorable. Removing silicon

atoms by high temperature annealing results in a (6√3×6√3)R30◦ surface reconstruction

corresponding to the C-rich buffer layer, followed by the (1×1) graphene structure. Thesesurface reconstructions have been captured by Low Energy Electron Diffraction (LEED)

and Scanning Tunneling Microscopy (STM) [14, 104]. Two common problems that arise

in epitaxial graphene formation on SiC in ultrahigh vacuum (UHV) are the inhomogeneity

of graphene films and pitting of the surface. Each of these can be attributed to rapid

sublimation of Si either at step edges [105, 106] or defects. At these sites Si is more

readily sublimated, resulting in the onset of a new graphene layer before the previous

layer has completely formed. Multiple strategies have been developed to limit the rate

of Si sublimation to produce more uniform graphene layers including hydrogen etching

of the substrate prior to graphitization [107–109], growth under Si flux [110], confined

sublimation in UHV [108], and growth at atmospheric pressure [111].

Graphene growth on vicinal SiC has also been investigated due to the relevance of

off-axis 4H- and 6H-SiC wafers in commercial SiC wafer processing [96, 97]. Graphene

grown on off-axis SiC wafers is influenced by the substrate morphology, where higher

miscut angles present a higher density of step edges and narrower terrace widths [112,

113]. A high step density results in a higher Si sublimation rate during graphene growth

which produces a higher average graphene thickness on the off-axis substrates [113].

Robinson et. al. demonstrated a monotonic increase in carrier density and decrease in

mobility in graphene grown on samples with miscut angles up to 0.45◦, but interestingly

the mobility for higher miscut angles was similar that of graphene on lower angle miscut

substrates [113]. Ouerghi et. al. presented a method for producing highly uniform, large

area graphene on SiC (0001) substrates cut 3.5◦ off-axis towards [112̄0] [110]. Etching

the substrates in a hydrogen atmosphere prior to graphene growth results in an ordered

series of (0001) terraces separated by (111̄0n) nanofacets [114, 115]. Graphene nucleates

on the nanofacets before extending along the terraces allowing long-range order in the

graphene sheet.

The interfacial carbon buffer layer of epitaxial graphene on SiC also plays a substantial

role in its electronic properties. Approximately a third of the carbon atoms in buffer are

2.1. EPITAXIAL GROWTH OF GRAPHENE ON SIC 11

Figure 2.1: (a) Atomic model of the bottom-up growth of graphene on Si-terminated SiC(0001). At sufficiently high temperature Si atoms sublimate resulting in layer-by-layer growthof graphene passivated by a buffer layer partially covalently bonded to the substrate; (b-e) LEEDpatterns indicating the evolution of the SiC (0001) surface through (1 × 1) (b), (√3 × √3) (c)and (6

√3 × 6√3) reconstructions before graphene formation (d),(e) [14]; (f-i) AFM and LEEM

images showing the improvement of the graphitized SiC surface morphology for UHV growngraphene (f),(g) and graphene grown at atmospheric pressure (h),(i). Reproduced from [111], withthe permission of Nature Publishing Group.


covalently bonded to the SiC substrate which disrupt the π bands seen in graphene. The

remaining dangling bonds are responsible for n-type doping of epitaxial graphene, and

have been cited as the cause of low carrier mobilities in epitaxial graphene compared to

exfoliated or CVD graphene transferred onto other substrates [116]. In order to reduce

the substrate effects strategies to intercalate epitaxial graphene with hydrogen have been

developed. H-intercalation saturates the dangling bond of the SiC substrate and breaks

the covalent bonds to the buffer layer, thereby releasing it to form a new graphene layer

[117, 118]. Angle-resolved photoemission spectroscopy (ARPES) measurements of the

band dispersion of monolayer + buffer layer graphene on SiC before and after hydrogen

intercalation show a shift in the Dirac point from 420 eV below the Fermi level to just

above the Fermi level (extrapolated to be 100 eV) indicating a shift from n-type to p-

type doping [117]. Furthermore, a second band emerges indicating passivation of the SiC

surface transforming the buffer layer into a new graphene layer. The resulting graphene

is said to be quasi-freestanding due to the reduced substrate influence, and transport mea-

surements show an improvement in carrier mobility and device performance compared to

as-grown graphene on SiC [119–121].

2.1.1 Summary

High temperature decomposition of SiC is an established method of producing graphene

layers at the wafer scale in a controlled way. In order to obtain high quality graphene it

is imperative to control the rate of Si sublimation to prevent inhomogeneity in both the

substrate morphology and graphene thickness. Both of these factors are controlled by

high temperature synthesis in an Ar atmosphere, and introducing nanofacets using off-

axis SiC wafers enable the effective elimination of step edges and allows formation of

continuous graphene sheets. By intercalation of epitaxial graphene with hydrogen the

interaction between the graphene layers and the substrate can be minimized to produce

quasi-freestanding graphene exhibiting electronic properties similar to that of pristine,

freestanding graphene.

2.2. LATERAL H-BN/GRAPHENE HETEROSTRUCTURES 13

2.2 Lateral h-BN/Graphene Heterostructures

Graphene is composed of a single layer of sp2-hybridized carbon atoms arranged in a

hexagonal lattice as shown in Figure 2.2. It has been the subject of intensive research

efforts due to its remarkable electronic, mechanical and optical properites. The existence

of π electrons in graphene gives rise to unique and remarkable electrical [12–14], optical

[15], thermal [122] and mechanical properties [16]. Graphene is a zero band gap semi-

metal and electrons propagate as massless Dirac fermions with extremely high electron

mobilities that only weakly depend on temperature. In high quality graphene films the

electronic structure allows ballistic transport over large distances at room temperature.

Graphene also exhibits remarkable thermal conductivity and mechanical strength while

remaining flexible.

Hexagonal boron nitride (h-BN) is structurally analogous to graphene, shown in Fig-

ure 2.2, sharing the same lattice structure with boron and nitrogen atoms occupying

alternating positions in the hexagonal arrangement, and a small lattice mismatch (1.7%).

Though structurally similar, h-BN and graphene have vastly different electronic properties

with monolayer h-BN exhibiting a wide bandgap of 5.9 eV. As such h-BN is considered

as an ideal substrate for 2D materials due to the absence of dangling bonds, charge traps

and surface roughness [57].

Figure 2.2: Atomic models of monolayer graphene and hexagonal boron nitride. The brown,green and grey spheres represent carbon, boron and nitrogen atoms, respectively.

Owing to the small lattice mismatch graphene and h-BN can be stitched laterally into

a hybrid monolayer. Section 2.2 will review the theoretical and experimental investiga-

tions of hybrid h-BN/graphene (h-BNG) layers, and their prospects for applications in


nanoelectronics.

2.2.1 Tuneable Electronic and Magnetic Properties

Electronic Properties

The tuneability of the electronic properties of lateral heterostructures of graphene and h-

BN have been studied by density functional theory (DFT) in various reports. A key point

of interest in each of these studies is the configurations of the graphene and h-BN domains

(i.e. graphene embedded in h-BN, and h-BN embedded in graphene), domain geometry

(quantum dots or nanoribbons), and the relative concentrations of carbon and BN.

For graphene nanostructures embedded within h-BN, alteration of the electronic struc-

ture is similar to the cases of freestanding graphene nanostructures such as quantum dots

(QDs) or nanoribbons (NRs). In freestanding structures the π electrons are essentially

confined to an infinite potential well, whereas for the hyrbrid material the potential well

is finite due to the surrounding h-BN. The case of graphene quantum dots embedded in

h-BN was studied by Bhowmick et. al. [72], and Li and Shenoy [73]. Both studies found

that a bandgap is induced jointly by quantum confinement of electrons in the graphene

QDs, and hybridization of the 2p orbitals of C,B and N. The size of the electronic gap is

inversely proportional to the size of the quantum dots, scaling with 1/√n where n is the

number of carbon atoms [72].

A bandgap can be introduced to graphene without electron confinement by doping

with h-BN. Fan et. al. calculated the band structures of graphene doped with h-BN

and found that a bandgap could be introduced to the K (K’) points of graphene [75].

Figure 2.3(a)-(e) show the calculated band structures of graphene doped with increasing

concentrations of h-BN starting from 100% graphene in (a) up to 100% h-BN in (e).

The magnitude of the gap opening at the K-point is plotted as a function of the h-BN

concentration in Figure 2.3(f). Once again, they find that the gap is widely tuneable

with the h-BN concentration, but the size of the h-BN domains isn’t observed to have a

strong influence. Instead the opening of the bandgap is caused by breaking the sublattice

symmetry due to redistribution of charge with the inclusion of h-BN. Zhao et. al. further

calculated the properties of superlattices of graphene with embedded h-BN quantum dots


[123]. After systematically studying the effects of the h-BN QD shape, size and the

superlattice structure by DFT and tight-binding calculations, they find that in all cases a

bandgap can be introduced. It is insensitive to the size of the QDs, and instead is mainly

dependent on the size of the graphene region separating them.

Figure 2.3: Calculated band structures of graphene doped with h-BN with concentrations of (a)0%, (b) 25%, (c) 50%, (d) 75%, and (e) 100%. The calculated bandgap is plotted as a function ofh-BN concentration in (f). Reproduced from [75], with the permission of RSC Publishing.

Wang et. al. used the Boltzmann transport equations in conjunction with DFT cal-

culations to calculate the transport properties of graphene with embedded h-BN domains

[74]. It was found that, in addition to the bandgap, the carrier mobility is also widely

tuneable according to the carbon content. Under a phonon scattering mechanism the

carrier mobility can range between 103 and 105 cm2 V−1 s−1 while maintaining a sizeable

bandgap between 0.38 eV and 1.39 eV. Transport in h-BNG is governed by the carbon

content and the width of the graphene domains, consistent with other works [75, 123].

The work of Wang and co-workers suggests that a field effect transistor made from h-

BNG could achieve large on/off ratios and high carrier mobility [74].

So far discussion has been limited to geometries where either graphene or h-BN

QDs are embedded within the complementary material. The electronic structure and


transport properties of alternating nanoribbons of graphene and h-BN have also been

investigated [72, 124]. The electronic properties of the embedded graphene nanoribbons

(GNRs) are dependent on the interface structure. When there is an armchair interface,

the confinement of electrons to a finite potential well defined by the h-BN results in a

gap opening. On the other hand, for a zigzag interface, opposing edges of the GNR will

have different terminations with the h-BN (B-terminated on one side and N-terminated

on the other). Due to the polarity of the B-N bond there is an intrinsic electric field

across the width of the GNR resulting in half-metallicity with spin-polarized states at

the interface. Calculated transport properties of nanostructures consisting of alternating

GNRs and boron nitride nanoribbons (BNNRs) suggest that the conductance is enhanced

by states near the Fermi level at the interface [124].

Magnetic Properties

In addition to altering the electronic properties, creating h-BN/graphene hybrids has also

been proposed as a method of inducing magnetism into light-element systems by manip-

ulating s and p electrons rather than 3d or 4f electrons in transition metals [125, 126].

Carbon atoms in graphene occupy two equivalent sublattices at alternating positions in

the honeycomb structure. At the lateral heterointerface of graphene and h-BN, imbalance

of atoms occupying each of the sublattices enables the possibility of inducing ferromag-

netism [127–129] or antiferromagnetism [126, 130].

The magnetism of carbon-doped h-BN monolayers has been studied as a function

of the interface orientation and termination [126–130]. Graphene triangles embedded in

h-BN are predicted to be ferromagnetic with a total spin equal to half the number of B

or N atoms at the interface, with the total magnetic moments of B-terminated and N-

terminated edges being antiparallel. Where mixed termination interfaces exist, the total

magnetic moment diminishes. On the otherhand, at the armchair interfaces of hexagonal

islands, the graphene domains are non-magnetic. This is illustrated in Figure 2.4, where

(a) plots the interface energies of the different terminations, and the magnitude of the

net magnetic moment as a function of chemical potential, and (b)-(f) illustrate the shapes


of the graphene domain [127]. Similar studies have investigated the magnetism of h-

BN doped graphene [131, 132]. It was found that there exist partially occupied, spin-

asymmetric mid-gap states at the interface between graphene and h-BN domains due

to the imbalance of carbon atoms occupying the two sub-lattices. The finite magnetic

moment in these systems arises from delocalization of carbon pz electrons at the interface

where C–N and C–B bonds exist [131].

Figure 2.4: (a) Interface energies of different terminations of graphene embedded in h-BN plottedas a function of chemical potential. Blue, purple and red represent B-terminated zigzag, armchair,and N-terminated zigzag interfaces respectively. The dotted line plots the magnetism per unitperimeter as change equilibrium graphene crystal shape changes with chemical potential. Thecrystal shapes are illustrated in (b)-(f) and the arrow lengths are proportional to the magnitude ofmagnetism. Adapted with permission from [127]. Copyright (2011) American Chemical Society.

Zhao et. al. experimentally investigated ferromagnetism in carbon doped h-BN

nanosheets prepared by annealing liquid-exfoliated h-BN with perylene-3,4,9,10-tetracarboxylic

acid tetrapotassium salt (PTAS) [125]. Bulk magnetization measurements revealed weak

ferromagnetism for the B-C-N nanoflakes at temperatures ranging from 2 K up to 400 K.

The magnetism disappeared when the carbon dopants were removed from the nanoflakes,

and concentration of magnetic metal impurities was measured to be less than 10 ppm,

thus confirming the carbon dopants as the origin of magnetism. These theoretical and

experimental investigations suggest that it may be possible to utilize heterointerfaces of

graphene and h-BN to design and fabricate half-metallic or semiconducting magnetic


nanomaterials or molecular magnets.

2.2.2 Synthesis Methods

Chemical Vapor Deposition

Driven by the success of chemical vapor deposition (CVD) for the growth of graphene

[133–135] and h-BN [136–138] on metal surfaces, the first experimental realization of

in-plane heterostructures of graphene and h-BN was also achieved by CVD. Ci et. al.

used methane (CH4) as the carbon source and ammonia-borane (NH3–BH3) as the h-BN

precursor to produce hybrid h-BN/graphene monolayers on Cu foils [139]. Consistent

with theoretical predictions [140], they observed by x-ray photoelectron spectroscopy

(XPS) that B–N bonds segregate to form isolated h-BN domains. The presence of C–N

and C–B components in the N 1s and B 1s core levels suggest the h-BN domains are in-

plane with graphene. Furthermore, the composition of the product could be tuned directly

according to the precursor concentrations supplied during CVD synthesis. Using high

resolution transmission electron spectroscopy (HRTEM), the h-BNC film was observed to

consist of a hexagonal atomic structure, and electron energy-loss spectroscopy (EELS) at

the B, C and N K edges revealed that all three elements are sp2 hybridized. The optical gap

of the hybrid h-BNG monolayers was measured in order to determine if any hybridization

effects were observed. Their results show two optical absorption edges varying from 4.48

eV to 3.85 eV and 1.62 eV to 1.15 eV as the carbon content is increased from 65 at% to

84 at%. This suggests that the domain sizes are large enough such that the graphene and

h-BN retain their own optical gap, and the shift in absorption edges can be attributed to

doping of graphene with h-BN (and visa versa).

Han et. al. later demonstrated that some control could be gained over the domain size

and shape of lateral h-BN/graphene hybrids grown on Cu foils using atmospheric pressure

chemical vapor deposition (APCVD) and alternating the precursor supplies [141]. By

tuning the growth parameters, hexagonal domains of graphene could be synthesized.

After switching off the CH4 feedstock and introducing NH3–BH3 the nucleation of h-

BN domains was observed to occur at the exposed graphene edges. The authors propose

that this is due to a higher reactivity of the exposed edges compared to the Cu surface.

However, if the h-BN growth period was extended isolated triangular domains for h-BN


would form on the bare Cu substrate. By HRTEM the graphene domains were observed

to have zigzag edges, and the h-BN was observed to have the same crystallographic

orientation as graphene except for cases where the film had torn or folded at the interface.

Further demonstration of lateral heteroepitaxy of graphene and h-BN was presented by

Liu et. al. in a study which focused on the formation of zigzag h-BN-graphene interfaces

induced by hydrogen etching monolayer graphene grown by APCVD on Cu foils [77].

The hydrogen etch process serves to prepare holes in the graphene film which are strictly

hexagonal with zigzag edges. Subsequent growth of h-BN by APCVD nucleates at the

exposed graphene edges, consistent with the work of Han et. al.. The graphene and h-

BN boundaries were identified definitively by scanning tunneling spectroscopy (STS), and

imaged by scanning tunneling microscopy (STM). The lateral heterostructure of graphene

and h-BN continues seamlessly with an atomically sharp interface. Due to a difference in

density of states between the graphene and h-BN domains the h-BN appears to be raised

by 0.1 nm, shown in Figure 2.5.

Figure 2.5: STM image of an atomically resolved h-BN/graphene domain boundary exhibitinga 0.1 nm height difference (line profile inset) due to the difference in the density of states.Reproduced from [77], with the permission of American Association for the Advancment ofScience.

CVD synthesis of lateral h-BN/graphene heterostructures has also been reported on

other metal surfaces such as Rh(111) [142], Ru(0001) [143–146], and Ir(111)[147, 148].

Gao et. al. demonstrated that alternating precursor supplies on Rh(111) led to a h-BNG

monolayer, and preferential formation of zigzag edges (75%) rather than armchair (25%)

was observed by STM and verified by DFT. On Ru(0001) substrates, preparation of a


submonolayer of graphene followed by exposure to borazine resulted in mixing of the

graphene and h-BN phases to produce a BCxN alloy separated by graphene nanoribbons

[143, 149]. Increasing the borazine dose causes segregation of graphene and h-BN into

distinct phases, and in the extreme case produces a full h-BN monolayer due to etching of

the graphene edges by hydrogen produced in the decomposition of the borazine molecule.

Sutter et. al. demonstrated that the mixing of the carbon and h-BN is caused by incorpo-

ration of carbon adatoms during borazine decomposition [143]. By exposing the surface

to a low pressure of O2 the adatoms are removed as CO, allowing distinct graphene and

h-BN phases to form.

Camilli et. al. observed that by supplying ethylene (C2H4) and borazine (B3N3H6)

precursors simultaneously in UHV allowed the formation of graphene nanodots embed-

ded in a BCN alloy [148]. The graphene dots were remarkably uniform in size with a

diameter of 1.6± 0.2 nm. By increasing the partial pressure of the ethylene precursor thedistance between neighboring dots could be decreased while maintaining their uniform

size. Lateral heterostructures of graphene and h-BN were grown by CVD on Ir(111) from

ethylene (C2H4 and NH3–BH3 using an alternating supply method [147]. Graphene and

h-BN regions could be clearly identified by STM due to their distinct moiré patterns, and

high resolution STM imaging revealed an atomically sharp zigzag boundary. In contrast

to surfaces like Ru(0001), where there is a strong coupling between h-BN/graphene and

the substrate due to hybridization of the π and d orbitals, h-BNG on Ir(111) is quasi-

freestanding. As such the electronic structure of graphene and h-BN patches could be

studied by STS and the authors found that graphene and h-BN both retain their distinct

electronic signatures. Moreover, the presence of interface states was not observed. This

is in contrast to the work done by Drost et. al. on h-BNG intercalated with Au on Ir(111)

which shows the existence of a boundary localized state at the zigzag C–B interface [150].

Similarly, h-BN/graphene interface states were observed on Cu(100) using STM and

STS by Park et. al. [151]. For zigzag terminated h-BN/graphene interfaces, boundary

states were observed to exist at around 0.6 eV above or below the Fermi level depending

on whether the interface is N-terminated or B-terminated. At the N-terminated boundary,

the interface state exists at 0.6 eV above the Fermi level, and at a B-terminated interface,

the boundary states are between 0.45 eV and 0.78 eV below the Fermi level. In both cases,

the states are localized at the h-BN/graphene interface and decay exponentially into the


h-BN and graphene regions. The presence of interface states is explained by hybridization

of the unoccupied (N-terminated) and occupied (B-terminated) π orbitals of C, B and N

at the boundary.

It is important to note that two dimensional heteroepitaxy of graphene and h-BN is

subject to strain at the interface due to the small (1.7%) lattice mismatch between the two

materials. This effect was studied by Lu et. al. for lateral heteroepitaxial h-BN/graphene

grown on Ru(0001) [144]. A pristine, atomically sharp interface was observed by STM,

and found to extend only a short distance perpendicularly from the interface before a

misfit dislocation occurs in the h-BN to reduce the interfacial strain. Figure 2.6 shows

an STM image of the interface between graphene (top right) and h-BN (bottom left)

where the two materials can be identified by their respective moiré patterns. Strain

relaxation causes a misfit dislocation which perturbs the moiré pattern, outlined in white.

An atomically resolved image of the defect is shown in the right panel. The strain is

calculated to reach a maximum of 5.9% at the sixth zigzag line of atoms from the interface

which is in good agreement with the position of misfit dislocations in the experimental

data. This phenomenon is directly analogous to the formation of misfit dislocations in the

three dimensional heteroepitaxial growth of GaAs on Si substrates [144, 152].

Figure 2.6: STM image a h-BN/graphene domain boundary showing perturbations to the moirépattern due to misfit dislocations (left panel). The right panel shows an atomically resolved misfitdislocation (MD) caused by the lattice mismatch of graphene and h-BN. Adapted with permissionfrom [144]. Copyright (2014) American Chemical Society.

Controlled patterning of the h-BN and graphene domain shapes and sizes has been

achieved by adopting a patterned regrowth scheme illustrated in Figure 2.7 [76, 77].


Patterned regrowth entails synthesis of a full monolayer of graphene (h-BN) followed

by photolithography and reactive ion etching, or focused ion beam patterning, to etch

away part of the film in a defined patterned and expose the substrate. This is followed by

CVD growth of h-BN (graphene) in turn to produce the lateral heterostructure film. This

method has the advantage of being able to selectively grow arbitrarily defined patterns of

graphene and h-BN, however, the domain size is limited to the resolution of the patterning

technique. As such there is difficulty in obtaining features down to a sub-10 nm size which

is a requirement for successfully manipulating the electron band structure. More recently

Zhang et. al. demonstrated that regions of graphene and h-BN could be predefined in

a Cu-Ni alloy [153]. Ultraviolet lithography was used to define a pattern in a Cu foil

followed by deposition of a 10 nm Ni layer by magnetron sputtering. After removing

the photoresist the graphene and h-BN were grown by CVD from CH4 and NH3–BH3

precursors. It was found that by first supplying CH4 a complete layer of graphene could

be synthesized preferentially on the Ni-free regions of the foil. Switching the precursor

supply to NH3–BH3 allowed h-BN to grow

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