Large Scale Growth of MoS2 Monolayers by Low

Pressure Chemical Vapor Deposition

Omar Salih Omar

Doctor of Philosophy

University of York
Physics
March 2018
 Abstract

Monolayers of molybdenum disulphide MoS2, a two dimensional (2D)

semiconductor with a direct band gap of 1.9 eV, have been proposed as a
candidate for next generation nanoscale electronic and opto-electronic devices.
Controlled synthesis of MoS2 monolayers is critically important since the thickness
uniformity and grain size are major concerns for the fabrication of opto-electronic
devices.
In this study, we demonstrated the growth of wafer scale uniform MoS2
monolayers on SiO2 covered silicon wafers, at a range of growth temperatures
(650 oC-850 oC) with optimum grain sizes as large as 400 μm, using low pressure
chemical vapor deposition (LPCVD). By controlling the partial pressure of the
reactant species at the growth surface and the limiting time, we can achieve
prefered monolayer growth over multilayer growth.
The MoS2 monolayer crystals follow a lognormal size distribution, consistent
with random crystal nucleation, with single crystal domains as large as 400 μm.
We estimated the thermal expansion coefficient to be (2.5±1.2) ×10-6 /oC, which is
at least double that of the bulk.
We have found film growth can be clearly classified into the reaction limited,
feed limited and desorption limited regimes. With the help of COMSOL
simulations, we have related the local growth environment such as growth
temperature, MoO2 concentration, sulphur chemical potential and growth time with
the macroscopic growth parameters such as Ar flux. In the feed limited regions, it
is the supply of Mo that is the rate limiting factor. In the desorption regions, the
growth is controlled by thermal stability of MoS2 monolayers. The growth modes
also can be used to tune the grain morphology from perfect triangles to hexagons.
Finally, we have also compared our approach with an LPCVD approach
based on MoO3 as the Mo source. MoO3 has a higher vapor pressure than MoO2
which was used in the previous approach. By tuning the the S:MoO3 ratio, we
could grow controllably planar MoS2 monolayers, vertically aligned MoS2/MoO2
and planar MoO2 crystals.

2
Table of Contents
Abstract 2

Table of Contents 3

List of figures 6
Acknowledgements 16
Dedication 17
Declaration 18

Chapter 1: Introduction 19
1.1 Transition metal dichalcogenides (TMDs) 19
1.2 Why MoS2 monolayers are interesting 23
1.3 Outline of the structure of the thesis 25

Chapter 2: Preparation of 2D materials 28

2.1 Overview 28
2.2 Top-down approaches 28
2.2.1 Mechanical exfoliation 28
2.2.2 Liquid exfoliation 30
2.2.3 Electrochemical Exfoliation 32
2.3 Bottom-up approaches 33
2.3.1 Chemical vapor deposition 33
2.3.2 Physical vapor deposition 37
2.3.3 Atomic layer deposition (ALD) 37
2.3.4 Molecular beam epitaxy (MBE) 38
2.4 Conclusion 39

Chapter 3:Experimental and Simulation Techniques 41

3.1 Characterization techniques 41
3.1.1 Optical microscope 41
3.1.2 X-ray photoelectron spectroscopy (XPS) 43
3.1.3 Photoluminescence spectroscopy (PL) 44
3.1.4 Raman spectroscopy 45
3.1.5 Second harmonic generation microscopy 47
3.1.6 Scanning electron microscopy (SEM) and energy-dispersive X-ray
(EDX) spectroscopy 49
3.1.7 Transmission electron microscopy (TEM) 52
3.1.8 Atomic force microscopy (AFM) 55
3.1.9 X-ray diffractometry (XRD) 57
3.2. CVD simulation using COMSOL 59
3.3 Summary 61
Chapter 4: LPCVD growth of continuous MoS2 monolayer films 62
Abstract 62

3
 4.1 Literature review 62
4.2 Experimental 64
4.2.1 CVD set-up, temperature dependence 64
4.2.2 Experimental procedure 65
4.3 Results and discussion 66
4.3.1 Self-limiting growth of monolayer thin films 66
4.3.2 Chemical analysis by XPS measurements 68
4.3.3 Thickness measurements by AFM 70
4.3.4 Uniformity of monolayer films using multiphoton microscopy 72
4.3.5 Phase identification by Raman spectroscopy, TEM and X-ray
diffraction 75
4.3.5.1 Raman spectroscopy 75
4.3.5.2 Transmission electron microscopy investigations 76
4.3.5.3 Synchrotron X-ray in-plane grazing angle diffraction GIIXRD 77
4.3.6 In situ annealing of the MoS2 monolayers 81
4.4 Physical property measurement 83
4.4.1 Optical properties 83
4.4.2 Electrical properties 85
4.5 Grain size distribution and nucleation density 87
4.6 Conclusions 91
Chapter 5: Growth mechanism studies with the aid of COMSOL 93
Abstract 93
5.1 Literature review 94
5.1.1 CVD process 94
5.1.2 Thermodynamics of CVD 96
5.1.3 Monolayer growth mechanisms 98
5.2 Experimental investigation 99
5.2.1 Effect of gas flow on furnace temperature profile 99
5.2.2 Starting material vapor pressure and concentration 102
5.3 Experimental results 103
5.3.1 The effect of the vertical distance between substrate and Mo source
on the monolayer film coverage 104
5.3.2 The effect of sulphur flux on the coverage of monolayers 109
5.3.3 Grain size and nucleation density temperature dependence 112
5.3.4 Temperature dependence of the initial monolayer growth rate 118
5.3.5 One minute growth (flushing growth) 120
5.3.6 The effect of sulphur chemical potential on grain morphology 123
5.3.7 The effect of the growth time on the film uniformity 127
5.4 Discussion and conclusion 129

Chapter 6: MoS2/MoO2 nanostructure growth 135

Abstract 135
6.1 Literature review 135
6.2 Experimental investigation 137

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 6.3 Results and Discussion 138
6.3.1 Effect of non-uniform MoO3 concentration on the grown film 138
6.3.2 The effect of MoO3 concentration on the film growth at constant
temperature 142
6.3.3 The effect of MoO3 concentration on the film growth at different
growth temperatures 145
6.3.4 Film characterization using XRD and TEM 146
6.4 The effect of S/MoO3 ratio on the film growth mechanisms 148
6.5 Conclusion 150

Chapter 7: Conclusion and Future Work 152

7.1 Conclusion 152
7.2 Future work 155
Appendix 158
A1: MoS2 monolayers transferring method 158
A2: Polarization resolved SHG and grain orientation (Matlab script) 158
A3: Sulphur chemical potential parameters 159

Symbols and abbreviations 160

References 165

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List of figures

Figure Page

Fig. (1:1): The transition metals and the three chalcogens that are 20
highlighted have a layered structure. Partial highlights for Co, Rh,
Ir and Ni refer to some of the dichalcogenides that form layered
structures.

Fig. (1:2): Schematics of the structural polytypes: 2H, 3R and 1T 21

stacking sequences in TMD materials. The chalcogen atoms (X)
are yellow and the metal atoms (M) are grey.

Fig. (1:3): (a,b) Atomic models showing 1H and 1T phases of 21

monolayer MoS2 respectively.

Fig.(1:4): (a): Crystallographic orientations of a MoS2 monolayer 22

(b): The shaded region bounded by dashed lines corresponds to
one primitive cell.

Fig. (1:5): (a) Bright-field TEM image of a single-crystal triangle 22

with a Mo-zigzag edge orientation. (b) Diffraction patterns from a.
The asymmetry of the Mo and S sublattices separates the 1100[ ]
{ }
diffraction spots into two families: ka = (1100), (1010), (0110) and
kb =−ka.

Fig. (1:6): Calculated band structure of (a) bulk MoS2 (b) 4 layers 23
(c) 2 layers and (d) 1 layer.

Fig. (1:7): Photoluminescence of single and bilayer MoS2. 24

Fig. (1:8): Band structure of the first Brillouin zone. Green cones 24
represent the conduction band, blue and red cones are the
spin-orbit split valence band, and the arrows are spin up and down
of the carriers.

Fig (2:1): The micromechanical cleavage technique (“Scotch tape” 29

method) for producing graphene. Top: Adhesive tape is used to
cleave the top few layers of graphite from a bulk crystal of the
material. Bottom left: The tape with graphitic flakes is then
pressed against the substrate of choice. Bottom right: Some flakes
stay on the substrate, even on removal of the tape.

Fig. (2:2): Mechanically exfoliated single layer and multilayer MoS2 30

films on SiO2/Si. Optical microscope images of single-layer (1L),
bilayer (2L), trilayer (3L), and quadrilayer (4L) MoS2 films (A–D).
Panels E–H show the corresponding AFM images of the 1L
(thickness ≈ 0.8 nm), 2L (thickness ≈ 1.5 nm), 3L (thickness ≈ 2.1
nm), and 4L (thickness ≈ 2.9 nm) MoS2 films shown in (A–D).

Fig. (2:3): Schematic illustration of the most widely used liquid 31

6
exfoliation methods. (A) Ion intercalation. Ions represented by
(yellow spheres) are intercalated between the TMD crystal layers
in a liquid environment causing crystal swelling and resulting a
decrease in the interlayer attraction. Adding energy such as shear,
ultrasonication, or thermal to the system causes bulk TMD to
exfoliate into a dispersed layers. (B) Some TMDs have ions
between their layers, these ions are represented by (red spheres).
In a liquid environment, these ions can be replaced by larger ions
(yellow spheres) weakening the interlayer attraction. After Ion
exchange, an external perturbation can exfoliate bulk TMD crystal
to layers. (C) Sonication-assisted exfoliation. A bulk TMD crystal
can be exfoliated into separated layers by sonication in a solvent
with an appropriate surface energy. The solvent stabilize the
exfoliated layers against re-aggregation and sedimentation.

Figure (2:4): (a) The electrochemical circuit used for exfoliation of 32

bulk MoS2 crystal. (b) Bulk MoS2 crystal held by a Pt clamp (c)
Dispersed MoS2 layers in Na2SO4 solution (d) Dispersed MoS2
layers in N-methyl-2-pyrrolidone (NMP) solution. (e) Schematic
description summarizing the electrochemical exfoliation
mechanisms of bulk MoS2 crystal.

Fig. (2:5): Schematics of the most common methods used to 34

deposit TMD from vapor phase. (a) Metal (M) and chalcogen (X)
powders. (b) Metal or metal oxides deposited on substrate and
chalcogen powders. (c) Metal or metal oxides deposited on
substrate and chalcogen supplied as gaseous precursor. (d) Metal
and chalcogen compounds supplied by gaseous precursors.

Fig. (2:6): a, Optical image of CVD growth of typical large-grain 35

MoS2 on a SiO2 (285 nm)/Si substrate. The image contrast has
been increased for visibility; magenta is the bare substrate, and
violet represents monolayer MoS2. b, Optical image of a
monolayer MoS2 triangle. The triangle is 123 μm from tip to tip.

Fig. (2:7): (a) and (b), Optical images of the MoS2 monolayer and 35
bilayer films, respectively. The insets are optical micrographs of
the MoS2 monolayer and bilayer films grown on SiO2/Si
substrates. The scale bars in the insets are 80 μm. (c) and (d),
AFM height profiles for typical MoS2 monolayer and bilayer films
grown on sapphire, respectively.

Fig. (2:8): (a) Schematic illustration for growing MoS2 layers by 36

MoO3 sulfurization. A MoO3 film (∼3.6 nm) was thermally
evaporated on the sapphire substrate. The MoO3 was then
converted to a MoS2 by a two-step thermal process. (b) MoS2
layer grown on a sapphire wafer. (c ) AFM thickness
measurements.

Fig. (2:9): Schematic illustration of the two-step thermolysis 37

process for the synthesis of MoS2 thin layers on insulating
substrates. The precursor (NH4)2MoS4 was dip-coated on SiO2/Si

7
or sapphire substrates followed by the two-step annealing
process.

Fig. (2:10): Schematic illustration of one growth cycle of an ALD 38

MoS2 film.

Fig. (3:1): Schematic depiction of optical reflection and 41

transmission for a nanolayer with thickness d1 and complex index
of refraction n1 deposited on an SiO2 layer characterized by
thickness d2 and index of refraction n2 that is grown on top of a Si
substrate. Nanolayers deposited on SiO2 are visible due to
interference between light rays A, B and C reflected at various
interfaces in the stack.

Fig. (3:2): Color contrast plot of calculated contrast as a function 42

of the number of layers of MoS2 ultrathin films and the illumination
wavelength for 300 nm thick SiO2/Si substrates.

Fig. (3:3): (a–m) Color optical images of 1L–15L MoS2 on 300 nm 43

SiO2/Si. The scale bars are 5 μm for images a–l and 10 μm for
image m. (n) Contrast difference values of 1L–15L MoS2
nanosheets on 300 nm SiO2/Si.

Fig. (3:4): Idealised model of Rayleigh scattering and Raman 45

stokes and anti-stokes scattering.

Fig. (3:5): (a) Raman spectra of thin (nL) and bulk MoS2 films. (b) 46
Frequencies of E 12g and A11g Raman modes (left vertical axis) and
their difference (right vertical axis) as a function of layer thickness.

Fig. (3:6): (A) Optical image of CVD-grown monolayer MoS2. (B) 49

SHG image of a polycrystalline monolayer of MoS2 of the same
area showing the grain boundaries. (C) Color coded orientation
map of the same area (D) Crystal orientation vectors of the grains
I,II,III, b and a.

Fig. (3:7): Electron ray traces through a schematic SEM column 50

with a condenser lens and a probe-forming or objective lens. Lens
distances p and q are marked for each lens.

Fig. (3:8): SEM image of an MoS2 film. Monolayer (ML), bilayer 51

(BL) and substrate (SUB) areas are marked.

Fig. (3:9): The two basic operation modes of a typical TEM 53

imaging system: (A) the diffraction mode: projecting the diffraction
patterns (DP) onto the viewing screen and (B) the image mode:
projecting the image onto the screen.

Fig. (3:10): Schematic diagram of a basic setup of AFM. 55

Fig. (3:11): Idealized forces between tip and sample surface 57

highlighting where the three imaging modes are operative.

Fig. (3:12): Schematic of Bragg Brentano XRD. 58

8
Fig. (3:13): Schematic of GIIXRD. 59

Fig. (4:1): The Elite thermal system used for growing MoS2 64
monolayers.

Fig. (4.2): Measured furnace temperature profiles for different 65

set-temperature of the tube furnace.

Fig. (4:3): Temperature profile for sublimation of starting materials 66

during a typical growth run.

Fig.(4:4): Optical images of a centimeter scale MoS2 monolayer 67

(the upper part) and the bare SiO2 covered Si substrate (the lower
purple part).

Fig. (4:5): Second harmonic generation (SHG) nonlinear 67

two-photon microscopy revealing the polycrystalline nature of
continuous MoS2 monolayers where grains with different
orientations show different color.

Fig. (4:6): XPS survey spectrum of MoS2 monolayers grown on 69

SiO2/Si substrate.

Fig. (4:7): Mo 3d and S 2s XPS spectrum of MoS2 monolayers. 69

Fig. (4:8): S 2p XPS spectra of MoS2 monolayers. 70

Fig. (4:9): AFM topography of a single crystalline MoS2 monolayer. 71

Fig. (4:10): AFM topography of a MoS2 polycrystalline film 71

showing grain boundaries (GB).

Fig. (4:11): AFM micrograph of MoS2 single crystal monolayer 71

edge.

Fig. (4:12): Height profile measurements of the crystal shown in 71

fig. (4:11); data taken along dark line.

Fig. (4:13): AFM micrograph of scratch continuous monolayer 71

MoS2 film.

Fig. (4:14): Height profile measurements of film shown in fig. 71

(4:13): data taken along dark line.

Fig. (4:15): SHG image of MoS2 monolayers grown at 800 oC 72

(grey regions) on a SiO2/Si substrate (dark regions).

Fig. (4:16): SHG images of continuous MoS2 monolayer films 73

grown at growth temperatures of (a) 700 oC, (b) 750 oC, (c) 800 oC
and (d) 850 oC.

Fig. (4:17): SHG image showing 3R-Stacking MoS2 bilayer grains 74

G1,G2, G3 and G4 (white triangles).

Fig.(4:18): SH intensity profile taken from MoS2 monolayers and 74

3R stacking bilayer, data taken along dark lines on G1, G2, G3

9
and G4 shown in fig. (4:17).

Fig.(4:19): SHG image of 2H-stacking MoS2 bilayer grains G1, G2 74

and G3 (dark triangles).

Fig. (4:20): SHG intensity from monolayer and 2H-stacking bilayer 74

MoS2, data taken along the white lines on G1,G2 and G3 shown in
fig. (4:19).

Fig. (4:21): Raman spectrum of a continuous MoS2 monolayer film 76

grown at 800 oC.

Fig. (4:22): (a) Transferred MoS2 monolayer on lacey carbon film 77

(b) HRTEM image of monolayer region (c) SAED from the same
region in (b).

Fig. (4:23): Measured GIIXRD from a MoS2 monolayer film. 78

Fig. (4:24): Zoom in scan of (100) and (110) MoS2 monolayer 78

planes at seven different points on the sample.

Fig. (4:25): Lattice constant variation as result of growth induced 80

strain in the polycrystalline monolayer film.

Fig. (4:26): GIIXRD peak splitting from differently strained grains. 80

Fig. (4:27): GIIXRD of a MoS2 monolayer at room temperature and 81

an annealing temperature of 800 oC.

Fig. (4:28): High resolution (100) and (110) MoS2 monolayer 82

planes at different annealing temperatures.

Fig. (4:29): MoS2 monolayer lattice expansion as a function of 83

temperature extracted from (110) plane.

Fig. (4:30): Lorentzian peak fitting of MoS2 monolayer PL spectra 84

showing neutral exciton and trion peaks.

Fig. (4:31): Schematics of FET used for electrical 86

characterizations.

Fig. (4:32): Drain-source current Ids as a function of back-gate 86

voltage Vbg at different drain-source bias voltage Vds=0.1V,Vds=0.5
V and VDS=1 V.

Fig. (4:33): SHG image of MoS2 monolayer film grown at 800 oC. 88

Fig. (4:34): (a) X-polarized SHG image of the film shown in fig. 89
4.33 (b) Y-polarized SHG image of the film shown in fig. 4.33.

Fig.(4:35): Orientation color map of the film shown in fig. (4:33). 89

Fig. (4:36): SHG image showing two opposite grains labelled G1 90

and G2.

10
Fig. (4:37): Polarization-resolved SHG image showing two 90
opposite grains G1 and G2.

Fig. (4:38): SHG profile across the white arrow shown in fig. 4:37. 90

Fig. (4:39): Grain size distribution of a MoS2 polycrystalline film. 91

Fig. (5:1): Substrate in the reaction zone of the CVD system. 95

Fig. (5:2): Typical Arrhenius plot, dependence of growth rate on 96

temperature.

Fig. (5:3): Two possible nucleation routes for growing MoS2 99

monolayers (a) MoS2 monolayer cluster (b) molybdenum
oxysulphides (MoOxS2-y, y≥x) nanoparticles.

Fig. (5:4): Typical tube furnace surface temperature for a growth 100
temperature of 1073 K and a flow rate of 100 SCCM.

Fig. (5:5): Main (MoO2) zone temperature profile at a growth 100

temperature of 1073 K under different flow rates.

Fig. (5:6): Sulphur zone temperature profile at a growth 101

temperature of 1073 K under different flow rates.

Fig. (5:7): Cross sectional temperature profile of the MoO2 zone at 101
a growth temperature of 1073 K under 200 SCCM.

Fig. (5:8): Cross Sectional temperature profile of the sulphur zone 101
at growth temperature 1073 K under 200 SCCM.

Fig. (5:9): Cutline along which the data is taken. 102

Fig. (5:10): Typical position of the substrates with respect to MoO2 104
powder in the reaction boat.

Fig. (5:11): MoS2 monolayer coverage as a function of 105

MoO2-substrate distance.

Fig. (5:12): Photograph of a continuous monolayer film (right part) 105

and bare substrate (left part) of a typical sample grown at an angle
48o.

Fig. (5:13): Isolated grains from left side of sample shown in fig. 105
5:12.

Fig. (5:14): Continuous monolayer film from right side of sample 105
shown in fig. 5:12.

Fig. (5:15): MoO2 concentration as function of MoO2-substrate 106

vertical distance at a growth temperature 1073 K and an Ar flow
rate 200 of SCCM.

Fig. (5:16): Cross-section of MoO2 concentration profile inside the 106

reaction boat at a growth temperature 1073 K and an Ar flow rate
of 200 SCCM.

11
Fig. (5:17): Upper: ZX-plane, sulphur concentration profile at a 107
growth temperature 1073 K and a flow rate of 200 SCCM. Below:
cutline (red) along the tube.

Fig.(5:18): Sulphur concentration profile taken along the cutline 107

shown in fig. (5:17).

Fig. (5:19): Left: Ar velocity profile in the XY-plane at the center of 108
the furnace tube. Right: red cutline.

Fig. (5:20): Ar velocity profile taken along the cutline shown in 108
figure 5:19 (right).

Fig. (5:21): Monolayer coverage as a function of Ar flow rate for a 109

growth temperature 1073 K.

Fig. (5:22): Left: typical sulphur flux at the inlet of the growth 110
region at a growth temperature of 1073 K and a flow rate 200
SCCM. Right: cutline along which data for different flow rates is
taken.

Fig. (5:23): Sulphur convective flux under different flow rates, and 110
at a growth temperature 1073 K. Data is taken along the red
cutline shown in fig 5:22,right.

Fig. (5:24): Close-up of the sulphur flux profile at the reaction zone 111
as a function of Ar flow rate.

Fig (5:25): Average sulphur flux entering the reaction zone as a 111
function of Ar flow rate.

Fig. (5:26): Sulphur concentration profile as a function of Ar flow 112

rate. Data is taken along the red cutline shown in fig. 5:22.

Fig. (5:27): Colored orientation map of a polycrystalline film grown 113

at 850 oC.

Fig. (5:28): Colored orientation map of a polycrystalline film grown 113

at 800 oC.

Fig. (5:29): Colored orientation map of a polycrystalline film grown 113

at 750 oC.

Fig. (5:30): Colored orientation map of a polycrystalline film grown 113

at 700 oC.

Fig. (5:31): Colored orientation map of a polycrystalline film grown 113

at 650 oC.

Fig. (5:32): Grain size distribution at a growth temperature of 850 114

o
C.

Fig. (5:33): Grain size distribution at a growth temperature of 800 115

o
C.

12
Fig. (5:34): Grain size distribution at a growth temperature of 750 115
o
C.

Fig. (5:35): Grain size distribution at a growth temperature of 700 115

o
C.

Fig. (5:36): Grain size distribution at a growth temperature of 650 116

o
C.

Fig. (5:37): Nucleation density as a function of growth temperature 116

with the inset fitted with an Arrhenius equation.

Fig. (5:38): Sulphur flux in the reaction zone at different growth 118
temperatures.

Fig. (5:39): Arrhenius plot of MoS2 monolayer growth for 15 120

minutes.

Fig. (5:40): Average sulphur flux used in the one minute growth. 120

Fig. ( 5:41): Arrhenius plot for one minute growth. 121

Fig. (5:42 a-f ): Are the gray style images of monolayers grown in 122
one minute at different growth temperatures between 700 oC and
1000 oC respectively.

Fig. (5:43): Shape evolution of MoS2 monolayers grown between 125

700 oC and 1000 oC.

Fig. (5:44): Depiction of MoS2 monolayer nucleus in the early 125

growth stage showing Mo and S zigzags.

Fig. (5:45): Sulphur chemical potential as a function of growth 127

temperature and vapor pressure indicating the morphology of
MoS2 monolayers at each growth temperature range.

Fig. (5:46): Optical images of (a) continuous MoS2 monolayer film 128
grown at 15 min., (b) continuous monolayer film grown at 20 min.
The initial stages of bilayer growth is shown in the dark circles. (c)
bare 300 nm SiO2/Si substrate.

Fig. (5:47): Optical micrograph of a MoO2 grain (b) the 129

corresponding Raman spectra and (c) Raman spectra for MoO2
from the literature.

Fig. (6:1):(a) Possible growth processes of MoS2 by the reaction of 136

MoO3-x and S. (b) The Mo–O–S ternary phase diagram, in which
the labelled arrows indicate reaction pathways for the CVD growth
of MoS2 from MoO3 precursors.

Fig. (6:2): Rectangular duct for MoS2 growth. 137

Fig. (6:3): Duct with pillars for MoS2 growth. 137

Fig. (6:4): Optical image of MoS2-MoO2 film grown using a duct 138

13
with (a) walls (b) pillars.

Fig. (6:5a): Edge of the sample MoS2 monolayers. 139

Fig. (6:5b): 1 mm from the edge, MoO2 crystals. 139

Fig. (6:5c): 2 mm from edge, MoO2 crystals and vertically aligned 139
MoS2/MoO2.

Fig. (6:5d): 2.5 mm from edge, MoS2 film and vertically aligned 139
MoS2/MoO2.

Fig. (6:5e): 3 mm from edge, mostly vertically aligned MoS2/MoO2. 139

Fig. (6:5f): 4 mm from edge, vertically aligned MoS2/MoO2. 139

Fig. (6:6a): Edge of the sample, vertically aligned MoS2/MoO2. 140

Fig. (6:6b): 1 mm from edge, vertically aligned MoS2/MoO2. 140

Fig. (6:6c): 2 mm from edge, vertically aligned MoS2/MoO2. 140

Fig. (6:6d): 4 mm from edge, vertically aligned MoS2/MoO2. 140

Fig. (6:7): MoO3 concentration profile on the substrate when using 141
a duct with (a) walls sticking coefficient of one, (b) wall sticking
coefficient of zero (c) pillars.

Fig. (6:8): Cross-sectional view of MoO3 concentration profile in 141

the center of the reaction zone when using duct with (a) wall
sticking coefficient of one (b) wall sticking coefficient of zero (c)
pillars.

Fig. (6:9): MoO3 concentration profile on the substrate in the three 141
cases :duct with walls with a sticking coefficient of one and zero,
and a duct with pillars.

Fig. (6:10): Duct walls (a) before and (b) after deposition. 142

Fig. (6:11): SEM images of MoS2-MoO2 films grown at T=650 oC 143

and MoO3-substrate vertical distances of (a) 1 cm (b) 2 cm (c) 3
cm (d) 4 cm and (e) 5 cm.

Fig. (6:12): Calculated MoO3 concentration on the substrate as a 144

function of vertical distance at a growth temperature of 650 oC.

Fig. (6:13): SEM images of MoS2-MoO2 films grown at (a) 650 oC 145
(b) 700 oC (c) 750 oC and 800 oC.

Fig. (6:14): MoO3 concentration profile at different growth 146

temperatures on a substrate placed at a vertical distance 5 cm.

Fig. (6:15): XRD data for a film grown at T=650 oC. 146

Fig. (6:16): XRD data for a film grown at T=700 oC. 146

14
Fig. (6:17): XRD data for a film grown at T=750 oC. 147

Fig. (6:18): XRD data for a film grown at T=800 oC. 147

Fig. (6:19): (a-b): TEM image of typical MoS2/MoO2 crystals 148

showing MoS2 planes at the edges.

Fig.(6:20): MoS2-MoO2 structures grown at 650 oC, (a): MoS2 148

monolayers, (b): vertically aligned MoS2/MoO2 crystalts, (c): planar
MoO2 crystals.

Fig. (6:21): Film composition at different S:MoO3 ratios at a growth 149

temperature of 650 oC.

Fig. (6:22): Partially covered planar MoS2 monolayers, vertically 150

aligned MoS2/MoO2 grains and nucleation of the vertically aligned
MoS2/MoO2 grains.

15
Acknowledgements
I would like to thank all those that helped me complete this PhD project. A special
thanks goes to my supervisor professor Jun Yuan for his supervision , helpful
discussions ,invaluable feedbacks and encouragement. Without his guidance and
persistent help this thesis would not have been possible.
I am particularly grateful to Ben Dudson,my thesis advisor panel,advice and
comments given by Ben has been a great help in completing this project.
A special thank you to York nanocenter staff: Ian Wright, Leonardo Lari and
Jon Barnard for doing a brilliant job in keeping the microscopes working, for
teaching to use many of the research tools used during my PhD work.
I am particularly grateful for the assistance given by Dr. Andrew Pratt for
X-ray photoelectron spectroscopy measurements and data analysis.
Special thanks also to professor Kevin O'Grady and his group for XRD
measurements.
My sincerely thank Dr. Stuart Cavill and to diamond light source beam 107
staff for the in-plane XRD measurements.
Special thanks to professor Chuanhong and his his group in China for the
Raman and photoluminescence measurements and data analysis.
I would like to thank the technicians Dave Coulthard, Mark Laughton, Adam
Stroughair and Neil Johnson for helping to assemble the CVD system and keep it
in working order.
My deepest thanks to the Kurdistan regional government for the financial
support.
A final thank to my beloved parents, brothers, sisters and my wife Awaz
Badry who have been supportive all times and have always encouraged me to aim
high.

16
Dedication

To my beloved parents, brothers, sisters and my wife Awaz

Badry

17
Declaration
I declare that this thesis is a presentation of original work and I am the sole
author. This work has not previously been presented for an award at this, or any
other, University. All sources are acknowledged as References

18
Chapter 1: Introduction

Two dimensional materials can be defined as the “materials in which the atomic
organization and bond strength along two dimensions are similar and much
stronger than along a third dimension”[1]. With recent advanced techniques almost
all layered three dimensional materials can be exfoliated to make atomically thin
layers. Particularly, the family of van der Waals solids is one of the most popular
sources of 2D materials, in which the strong in plane bonds provide stability for
atomically thin layers and weak interlayer van der Waals forces offer the feasibility
of exfoliation [1], [2].
Silicon is the most common material that has been employed in the
electronics devices industry. However, in the era of Nanoelectronics devices,
silicon is facing size limitations. This motivated researchers to quest for novel
materials that could overcome those shortcomings. The family of 2D materials
could have a range of electronic properties; they could be metals, semimetals,
insulators and semiconductors with a band gap from ultraviolet to infrared. This
makes 2D materials a good candidate to replace silicon [2]. Due to the high
mechanical strength, carrier mobility and thermal conductivity, graphene appeared
to be one of the most promising 2D materials that might have potential applications
in nanoelectronics, optoelectronics, energy harvesting, and biosensing fields [3].
However, graphene, as a semimetal, has a zero band gap, limiting its applications
in nanoelectronics [4], [5]. With this consideration, other graphene-like materials
such as transition metal dichalcogenides (TMDs) have recently emerged and
differently to graphene, many of them have the advantage of being
semiconductors and have a sizable band gap, thus making them appealing for
nanoelectronics applications [6], [7].

1.1 Transition metal dichalcogenides (TMDs)

Layered transition metal dichalcogenides have the stoichiometry formula MX2. The
letter M stands for the transitional metal atom and X is the chalcogen atom (for
example, M = Ti, Zr, Hf, V, Nb, Ta, Re; X = S, Se, Te (see figure 1:1). These
materials belong to family of van der Waals solids with the transition metal atoms
sandwiched between two chalcogen layers. M-X are strongly bonded by covalent

19
bonds, while the sheets are held together by weak van der Waals forces which
make these materials easy to be exfoliated to a single and few layers [8].
Different transition metal dichalcogenides possess very different electrical
properties depending on polytype and number of electrons in the d-shell of the
transitional metal atom: some are insulators such as HfS2, while the others could
be semiconductors such as MoS2, WS2, MoSe2 and WSe2, or semimetals such as
WTe2 and TiSe2, or even true metals such as NbS2 and VSe2. Among 2D transition
metal dichalcogenides, MoS2, WS2, MoSe2 and WSe2 are most widely studied in
nanoelectronic applications since they are semiconductors with a sizable band
gap, and have a high chemical and thermal stability [2], [6], [8].

Fig. (1:1): The transition metals and the three chalcogens that are highlighted have a
layered structure. Partial highlights for Co, Rh, Ir and Ni refer to some of the
dichalcogenides that form layered structures [6].

Bulk MoS2 can be found in different polytypes (different stacking sequences). The
most identified polytypes are 2H which is mostly found in naturally grown material
and 3R found in synthetic material where the letters represent hexagonal and
rhombohedral respectively, and the numbers represent the number of layers in the
stacking sequence. The 2H phase has a AbA BaB ( capital letter stands for the
sulphur atoms and the lower-case letters for molybdenum atoms) stacking
sequence with trigonal prismatic coordination and the 3R phase follows a AbA
CaC BcB stacking sequence with trigonal prismatic coordination [8]. A 1T phase
with stacking ABC ABC has been found in other TMDs such as TiS2 [6], [8]. Figure
1:2 represents different stacking sequences in TMD materials.

20
 Fig. (1:2): Schematics of the structural polytypes: 2H, 3R and 1T stacking sequences in
TMD materials. The chalcogen atoms (X) are yellow and the metal atoms (M) are grey
[9].

Monolayer of MoS2, a sheet of S-Mo-S sandwich layer, has two phases: 1T with
stacking AbC and 1H with stacking sequence AbA. It has been found more stable
than 1T [10].
For MoS2 monolayers, 1H has trigonal prismatic coordination that belongs
D3h point group and the 1T has octahedral coordination belongs to the D3d point
group. Figure 1:3 shows both monolayer polytopes [6].

Fig. (1:3): (a,b) Atomic models showing 1H and 1T phases of monolayer MoS2
respectively [6].

21
The crystallographic orientations of MoS2 monolayers are depicted in figure 1:4a
and the primitive unit cell is shown in figure 1:4b. Figure 1:5 a,b shows a
bright-field TEM image and diffraction patterns of a single-crystal triangle grown by
CVD. As seen in the schematic in figure 1:5b, the lattice of monolayer MoS2 is
divided into molybdenum and sulphur sublattices, which reduces the hexagonal

[ ]
lattice from six-fold to three-fold symmetry. As a result, the six 1100 diffraction

{
spots belong to two distinct families ka = (1100), (1010), (0110) } and k = −k
b a [11].

Fig.(1:4): (a): Crystallographic orientations of a MoS2 monolayer (b): The shaded region
bounded by dashed lines corresponds to one primitive cell.

Fig. (1:5): (a) Bright-field TEM image of a single-crystal triangle with a Mo-zigzag edge
orientation. (b) Diffraction patterns from a. The asymmetry of the Mo and S sublattices
[ ] { }
separates the 1100 diffraction spots into two families: k a = (1100), (1010), (0110) and
kb =−ka [11].

22
1.2 Why MoS2 monolayers are interesting
As bulk MoS2 is thinned down, the band structure changes with the number of
layers. The sizable band gap of MoS2 is one of the most important features that
make MoS2 to be a promising candidate for 2D electronic devices. The MoS2 band
gap size varies from 1.2 eV which is an indirect band gap associated with bulk
MoS2 to a 1.9 eV direct band gap for a single layer [12]. An example of the
calculated band structure of MoS2 is shown figure 1:6 [13]. In bulk MoS2, the
indirect band gap occurs at the Γ-point. The conduction band states at the Γ-point
arise from the hybridization between pz orbitals on S atoms and the d orbitals on
the Mo atoms. The conduction band states at the K-point mainly occur due to the
localized d orbitals on Mo atoms. The Mo atoms are sandwiched in the middle of
S-Mo-S layers and they are fairly isolated from the interlayer coupling therefore the
K-point is not influenced by a decreasing layer number. However, the Γ-point is
strongly affected by interlayer coupling effect since the stats at this point are due to
combinations of the pz orbitals on the S atoms and the d orbitals of the Mo atoms.
Thus, when the layer number decreases, the energy at the Γ-point increases, while
the conduction band at the K-point remains stable as shown in figure 1:6 [13].

Fig. (1:6): Calculated band structure of (a) bulk MoS2 (b) 4 layers (c) 2 layers and (d) 1
layer [13].

Such a transition from an indirect to a direct band gap semiconductor has

been confirmed by photoluminescence (PL) measurements. Bulk and multilayer
MoS2 exhibit a negligible photoluminescence peak due to the indirect band gap.

23
However, a bright photoluminescence peak has been detected from direct band
gap (1.9) eV single layer MoS2 as shown in figure 1:7. This emerging PL in the
monolayers in the visible range make them a promising candidates for future
photodiodes [14].

Fig. (1:7): Photoluminescence of single and bilayer MoS2 [12].

With the band gap in the visible range, the monolayers can also be used for
solar energy harvesting. MoS2 monolayers as active layers have been integrated
with silicon to make MoS2-Si heterostructures. A maximum efficiency of 5.23% has
been achieved from such heterostructures [15].
MoS2 monolayers and 3R bulk material have broken inversion symmetry
(figure 1:2) and they exhibit spin orbit splitting at the top of the valence band.
These two properties make them favorable candidates for spin based devices. The
band structure of the first Brillouin zone of a MoS2 monolayer is shown in figure
1:8. The points (valleys) located at the corners (K and -K) have equal energies and
different well separated momentums. This makes selective excitation of carriers
with various combinations of valley and spin index possible [16], [17].

Fig. (1:8): Band structure of the first Brillouin zone. Green cones represent the
conduction band, blue and red cones are the spin-orbit split valence band, and the
arrows are spin up and down of the carriers [16].

24
 The broken inversion symmetry in MoS2 monolayers can also be used for
converting mechanical energy to electricity. A measured piezoelectric coefficient of
e11 = 2.9 × 10–10 C m−1 of the monolayers makes them an excellent material for
atomically thin piezoelectric devices [18].
Another promising application of MoS2 monolayers is expected to be in the
field of electronics. The calculated room temperature phonon limited mobility in the
MoS2 monolayer is found to be ∼410 cm2 V−1s−1 [19]. Field effect transistors (FET)
fabricated from single layer MoS2 exhibit a room temperature on/off ratio of 108,
and a mobility over 200 cm2V-1s-1 with very low standby power dissipation [7].
Mechanically, MoS2 monolayers are flexible, with a Young's modulus of
270±100 GPa, which is comparable to that of steel. Such exceptional mechanical
properties make this material suitable to be integrated in flexible electronic devices
[20].

1.3 Outline of the structure of the thesis

The aim of this thesis is to develop a reliable chemical vapor deposition approach
for growing high quality, uniform wafer scale MoS2 monolayers with the largest
possible grain size and exploring their intrinsic optical, electrical and structural
properties.
Chapter one covers a brief introduction of two dimensional materials with
special attention given to the structural, electrical and optical intrinsic properties of
molybdenum disulphide monolayers.
In chapter two, we mainly focus on the synthesis approaches established for
2D materials. It includes top-down approaches such as mechanical exfoliation,
chemical exfoliation and electrochemical exfoliation. The bottom up approaches
cover chemical vapor deposition, physical vapor deposition, atomic layer
deposition and molecular beam epitaxy.
In chapter three, we present the characterization and simulation tools that
have been employed in this project. We introduce X-ray photoelectron
spectroscopy (XPS) that has been used for chemical analysis of MoS2
monolayers. Microscopic tools such as scanning electron microscopy (SEM),
transmission electron microscopy (TEM), scanning transmission electron
microscopy (STEM) and atomic force microscopy (AFM) that have been used in
monolayer imaging, studying the crystalline nature of the monolayer and their
thickness measurements are discussed. We also covers Raman spectroscopy for

25
qualitative thickness measurements and evaluating growth induced strain in the
monolayers. Second harmonic generation (SHG) was used for confirming the
uniformity of the monolayers and studying the grain size distribution of the
monolayers as function of growth conditions. X-ray diffractometry (XRD) and
in-plane grazing incidence angle X-ray are presented for studying the crystalline
nature of the monolayers. Finally, the CVD simulation software COMSOL is
presented and the fluid dynamics module, chemical engineering, and heat transfer
module are outlined.
Chapter four presents a low pressure chemical vapor deposition method for
growing MoS2 monolayers on SiO2 covered Si. In this chapter we took advantage
of the low vapor pressure of molybdenum dioxide as the Mo source to establish an
approach for producing uniform MoS2 monolayers on a wafer scale. The
production of such large scale monolayers is essential in the practical world of
optoelectronic devices. Under optimized conditions, we grew a uniform wafer scale
polycrystalline monolayer film with a grain size up to 400 μm. Different techniques
such as XPS, AFM, TEM, PL, Raman spectroscopy, SHG and XRD were used for
studying the chemical composition, thickness measurements, uniformity and
intrinsic structural, optical and electrical properties of the as grown monolayers.
In chapter five we further modified the growth conditions of the approach
presented in chapter four. We report the dependence of MoS2 film growth as a
function of growth conditions such as temperature, MoO2 concentration, sulphur
flux and carrier gas flow rate. We also studied the grain size distribution as a
function of growth temperatures to optimize the growth of largest possible grains
within a continuous polycrystalline film. We employed COMSOL software to
simulate the concentration distribution at the surface of the substrate with a view to
understanding the reaction conditions at the growth front under different growth
conditions. We also investigated the growth rate of isolated grains for better
optimizing the growth regimes in our CVD process. Finally, we present an
approach for tuning the morphology of the grains as a function of growth
temperatures.
In chapter six, another approach based on using MoO3 as a Mo source is
presented. MoO3 has a much higher vapor pressure than MoO2 at the same
growth temperature. As a result of this, different structures such as laterally
aligned MoS2 monolayers, vertically aligned MoS2/MoO2 crystals and laterally

26
aligned MoO2 are produced. COMSOL is used for fluid dynamics simulations and
the ratio of Mo/S for each growth regime is comprehensively analysed and the
growth mechanisms of the film growth is discussed.
Finally, chapter seven summarizes the conclusion for the whole project as
well as an outlook of potential future research.

27
Chapter 2: Preparation of 2D materials

2.1 Overview
The fabrication methods of transition metal dichalcogenides (TMDs) can be
classified into two types of general approaches: top-down and bottom-up. In
top-down approaches, techniques such as mechanical exfoliation, liquid exfoliation
and electrochemical exfoliation are used to thin down bulk crystals to few layers
and monolayers. In bottom-up approaches, starting materials in the form of gases
and powders are used to grow multilayers and monolayers using chemical vapor
deposition CVD, physical vapor deposition, atomic layer deposition and molecular
beam epitaxy.

2.2 Top-down approaches

2.2.1 Mechanical exfoliation

The nature of bulk layered materials has been studied for more than 150 years
[21]. Isolation of monolayers dates back at least to a few years ago (2004) when A.
Geim, and his colleagues isolated graphene (a monolayer of graphite) using
mechanical exfoliation by Scotch tape. This technique is the starting point for much
monolayer research in the world [22].
In mechanical exfoliation, an adhesive tape is used to peel off a thin layer of
graphite from the bulk crystal. Then the freshly cleaved nanosheet is brought into
contact with an appropriate substrate (SiO2) for further thinning and investigation
[22]. As well as exfoliation of graphene, the technique has also been widely used
in exfoliating other layered materials such as transition metal dichalcogenides [7],
hexagonal boron nitride [23] and black phosphorus [24]. What makes the isolation
of a monolayer possible is that the adhesion between the substrate and the bottom
layer of the bulk material can be stronger than the interlayer coupling of the bulk
material. Therefore, for exfoliation, of different materials different substrates can be
used [25]. For example, the flake size of MoS2 using SiO2 as a substrate is about
one order of magnitude smaller than the graphene size exfoliated using the same
substrate. To overcome this problem, gold is being used instead of SiO2 for

28
producing MoS2, as the sulphur’s affinity to gold is higher and MoS2 strongly binds
to gold [25].
The exfoliation technique is straightforward, inexpensive and reliable for
producing high quality monolayers for research purposes (for studying optical,
electrical and structural properties), however, the exfoliation technique is time
consuming and the lateral size of flakes is very small, in the range of a few
microns. Also, the number of layers in the exfoliated material cannot be precisely
controlled, as the exfoliation yield comprised of monolayers, multilayers and even
bulk material. Finally the technique is unreliable for mass production [22]. Figure
2:1 shows the micro cleavage technique used for graphene exfoliation [26]. Figure
2:2 shows optical images of monolayers and few layer of MoS2 deposited on a
SiO2/Si substrate using Scotch tape together with their corresponding thickness
measurements using AFM [27].

Fig (2:1): The micromechanical cleavage technique (“Scotch tape” method) for
producing graphene. Top: Adhesive tape is used to cleave the top few layers of graphite
from a bulk crystal of the material. Bottom left: The tape with graphitic flakes is then
pressed against the substrate of choice. Bottom right: Some flakes stay on the
substrate, even on removal of the tape [26].

29
 Fig. (2:2): Mechanically exfoliated single layer and multilayer MoS2 films on SiO2/Si
substrate. Optical microscope images of single-layer (1L), bilayer (2L), trilayer (3L), and
quadrilayer (4L) MoS2 films (A–D). Panels E–H show the corresponding AFM images of
the 1L (thickness ≈ 0.8 nm), 2L (thickness ≈ 1.5 nm), 3L (thickness ≈ 2.1 nm), and 4L
(thickness ≈ 2.9 nm) MoS2 films shown in (A–D) [27].

2.2.2 Liquid exfoliation

The liquid exfoliation of 2D materials can be categorized into direct and ion
intercalated sonication of bulk 2D crystals in an appropriate liquid. Figure 2:3
represents the schematic of the two approaches [28]. In a direct sonication, a bulk
crystal or powder of the 2D material is dispersed into a solvent with surface
tension in certain ranges, then the mixture is sonicated to produce dispersed mono
and multilayers, finally the product is centrifuged to remove unexfoliated residues
[29].
The suitability of many solvents for liquid exfoliation of TMDs has been
evaluated and it has been found that only solvents with a surface tension close to
40 mJ/m2 are suitable for exfoliation of monolayers and multilayers. Several TMDs
such as MoS2, MoSe2, WS2 and NbS2 have been exfoliated by direct sonication
and the two solvents, N-methyl-pyrrolidone (NMP) and isopropanol (IPA), resulted
in better yields [29]. TMDs have also been exfoliated in water and in this case a
surfactant such as sodium cholate is added to the mixture to prevent the
re-aggregation of the exfoliated material, as the water does not provide enough
surface energy to keep the exfoliated material dispersed [30].
In the ion intercalation approach, atomically thin TMDs layers can be
prepared by intercalating bulk TMDs crystal with ions such as lithium and then

30
exposing them to water [31]. Typically, a bulk TMD is first submerged in lithium for
more than 24 hours and then the intercalated TMD is exposed to water [32]. The
water strongly reacts with lithium between the layers to release hydrogen and
separate bulk the TMD into layers [32]. Such methods can be employed to
produce relatively large quantities (grams) of single layer TMDs [33]. However,
only small size flakes up to a fraction of a micrometer could be produced.
Moreover, the structural change (in the case of MoS2) (from trigonal prismatic
2H-MoS2 to octahedral 1T-MoS2) can result from the lithium interaction. This could
make MoS2 lose its pristine semiconducting properties. To reverse the phase
change and restore semiconducting properties of chemically exfoliated MoS2, the
samples need to be annealed at 300 oC [31], [33].

Fig. (2:3): Schematic illustration of the most widely used liquid exfoliation methods. (A)
Ion intercalation. Ions represented by (yellow spheres) are intercalated between the
TMD crystal layers in a liquid environment causing crystal swelling and resulting a
decrease in the interlayer attraction. Adding energy such as shear, ultrasonication, or
thermal to the system causes bulk TMD to exfoliate into a dispersed layers. (B) Some
TMDs have ions between their layers, these ions are represented by (red spheres). In a
liquid environment, these ions can be replaced by larger ions (yellow spheres)
weakening the interlayer attraction. After Ion exchange, an external perturbation can
exfoliate bulk TMD crystal to layers. (C) Sonication-assisted exfoliation. A bulk TMD
crystal can be exfoliated into separated layers by sonication in a solvent with an
appropriate surface energy. The solvent stabilize the exfoliated layers against
re-aggregation and sedimentation [28].

31
2.2.3 Electrochemical Exfoliation
Two dimensional materials nanosheets can also be prepared by electrochemical
exfoliation in which a DC bias voltage is applied between a 2D crystal and a Pt
electrode in an electrolyte (Na2SO4) solution. Initially, the 2D crystal is wetted by
applying a small voltage, and then the voltage is increased for exfoliation. Flakes
can be dissociated from the bulk material and become suspended in the
electrolyte [34]. The mechanisms of this technique can be understood as follow:
when applying positive voltage to the working electrodes, radicals OH- and O- or
SO4-2 are produced around the 2D crystal as a result of water oxidation. Such
radicals intercalate the 2D crystal and weaken the van der Waals forces between
the layers. Gases such as O2 and SO2 are released as a result of second oxidation
causing more weakening in the interlayer forces and finally flakes are detached
from the bulk crystal by the erupting gas bubble [34]. Figure 2:4 summarizes the
schematic production and mechanisms of this technique [34].

Figure (2:4): (a) The electrochemical circuit used for exfoliation of bulk MoS2 crystal. (b)
Bulk MoS2 crystal held by a Pt clamp (c) Dispersed MoS2 layers in Na2SO4 solution (d)
Dispersed MoS2 layers in N-methyl-2-pyrrolidone (NMP) solution. (e) Schematic
description summarizing the electrochemical exfoliation mechanisms of bulk MoS2
crystal [34].

32
2.3 Bottom-up approaches
2.3.1 Chemical vapor deposition
For mass production and integrating the 2D materials into industrial applications,
new techniques that enable control over the layer number, crystal quality, lateral
size of grains, etc. must be introduced. For this purpose, many bottom up
techniques such as CVD, PVD, ALD, MBE…, etc. have been proposed [35]. Such
processes utilize one or two starting materials to grow two dimensional materials.
In the case of TMDs the starting materials in form of the powder are vaporized,
carried by a carrier gas to the substrate where the reaction takes places and the
film grows. Through optimization of growth parameters such as temperature of
starting materials and their concentrations, reactor design, the species can be
uniformly delivered to the substrate which in principle permits uniform growth [35].
The most adopted scenarios for the CVD growth of TMDs are summarized in
figure 2:5. We will now focus on the growth of MoS2 as the most widely grown
TMD using approaches shown in figure 2:5. In the case of (a) the chalcogen
powder, mostly sulphur with high purity is placed at the upstream of a furnace tube
in a region where the temperature reaches above its melting point. The Mo source
is also in the form of powder such as molybdenum trioxide MoO3. It is placed at the
center of the heating zone where the temperature is high enough for Mo source
sublimation. The S and MoO3 vapors are carried by an inert gas (Ar or N2) to the
substrate placed downstream a few centimeters away from the Mo container. On
the substrate, and sometimes within the carrier gas atmosphere, the reaction takes
place and a film grows, with the reaction byproducts carried by the carrier gas out
of the furnace. The most common substrates used are insulating materials such as
SiO2/Si, quartz, mica or noble metals such as gold foils [34], [36]–[41]. Using this
method, films up to centimeter scale as well as isolated crystals with edge length
up to few hundreds of microns can be grown as shown in figure 2:6 [11]. Figure 2:7
is an optical image of continuous monolayers and bilayers with their corresponding
AFM thickness measurements, grown using MoCl5 and sulphur as the starting
materials [41].

33
Fig. (2:5): Schematics of the most common methods used to deposit TMD from vapor
phase. (a) Metal (M) and chalcogen (X) powders. (b) Metal or metal oxides deposited
on substrate and chalcogen powders. (c) Metal or metal oxides deposited on substrate
and chalcogen supplied as gaseous precursor. (d) Metal and chalcogen compounds
supplied by gaseous precursors [35].

34
 Fig. (2:6): a, Optical image of CVD growth of typical large-grain MoS2 on a SiO2
(285 nm)/Si substrate. The image contrast has been increased for visibility; magenta is
the bare substrate, and violet represents monolayer MoS2. b, Optical image of a
monolayer MoS2 triangle. The triangle is 123 μm from tip to tip [11].

Fig. (2:7): (a) and (b), Optical images of the MoS2 monolayer and bilayer films,
respectively. The insets are optical micrographs of the MoS2 monolayer and bilayer films
grown on SiO2/Si substrates. The scale bars in the insets are 80 μm. (c) and (d), AFM
height profiles for typical MoS2 monolayer and bilayer films grown on sapphire,
respectively [41].

Cases (b) and (c) in figure 2.5 are a two-step process CVD, in which MoO3 film
with the required thickness is first deposited on the substrate using physical vapor
deposition techniques such as e-beam evaporation and then sulfurized at high
temperature using sulphur powder or H2S gas. Employing this approach, one can
grow wafer scale MoS2 thin films, however the uniformity of the grown films is still
an issue, monolayer, bilayer and trilayers coexist on the substrate of the same
growth run [42], [43]. Similarly, instead of MoO3, Mo metal has been predeposited
on the substrate and sulfurized to get MoS2 films. Despite large scale films that can
be prepared using this method, some residual Mo atoms tend not to react with
sulphur affecting the semiconducting properties of the final product. Additionally,

35
as Mo metal has a much higher melting point (2610 oC) compared to MoS2 growth
temperatures (typically 650-850 oC), Mo atom migration is suppressed at these
growth temperatures which in turn affects the grain size of the grown film [44]–[46].
An example of two step growth using MoO3 is summarized in figure 2:8; the same
growth process can be applied to Mo-based growth as well.

Fig. (2:8): (a) Schematic illustration for growing MoS2 layers by MoO3 sulfurization. A
MoO3 film (∼3.6 nm) was thermally evaporated on the sapphire substrate. The MoO3
was then converted to a MoS2 by a two-step thermal process. (b) MoS2 layer grown on
a sapphire wafer. (c ) AFM thickness measurements [42].

The Mo and S sources can also be supplied in the form of gas precursors at the
furnace inlet as in the example of metalorganic chemical vapor deposition
(MOCVD) (see Fig. 2:5, case d). In this approach, molybdenum hexacarbonyl
(MHC) as Mo source and diethyl sulphide (DES) as sulphur source are diluted in
H2 and Ar carrier gases. Wafer scale films with a controlled number of layers can
be produced. The disadvantage of this approach is the growth time is quite long
time (26 hrs) for each run and the grain size is very small (about 10 microns) as
well as that the precursors are highly toxic and they need special precaution during
the growth [47].
Dip coating is another MoS2 CVD growth method that has been used. The
process starts by immersing an insulating substrate in ammonium thiomolybdate
(NH4)2MoS4 diluted in dimethylformamide (DMF). Then the substrate coated with
(NH4)2MoS4) is annealed at 500 oC in an Ar-H2 mixture to remove the residual

36
solvent, NH3 molecules, and other byproducts are dissociated from the precursors.
Finally, by sulfurization in 1000 oC, MoS2 film is obtained. The approach can be
used to grow large scale multilayer MoS2 films, however the uniformity of the film is
not under control because of the difficulties of coating uniform (NH4)2MoS4) films at
the beginning of the process [48]. Additionally, the self-assembly of the precursor
on the substrate during dip-coating could lead to the growth of different
morphologies such as MoS2 nanowires rather than films [49]. A typical procedure
of the dip coating approach is shown in figure 2:9 [48].

Fig. (2:9): Schematic illustration of the two-step thermolysis process for the synthesis of
MoS2 thin layers on insulating substrates. The precursor (NH4)2MoS4 was dip-coated on
SiO2/Si or sapphire substrates followed by the two-step annealing process [48].

2.3.2 Physical vapor deposition

The growth of MoS2 monolayers using physical vapor deposition has also been
reported in a few works. This is a single component process in which high purity
MoS2 powder is evaporated in a thermal system at high temperature (900 oC) with
the vapor carried away by an inert gas to the substrate placed in a cooler region
(600 oC) on which the physical deposition happens and an MoS2 film grow. The
technique has successfully been used to produce micron-sized monolayers.
However, there is no reports about scaling up the method as further work is
needed for optimization of the process [50].

2.3.3 Atomic layer deposition (ALD)

ALD is also one of the common techniques that has been used for growing TMDs.
The technique has multiple steps as shown in figure 2:10. The steps are as
follows: (1) introducing a sublimed Mo source to the reaction chamber where the
substrate is placed; (2) purging the chamber with an inert gas; (3) introducing the
sulphur source to the reaction chamber; (4) purging the chamber again by an inert

37
gas to remove the reaction byproducts. The mentioned steps represent one cycle
of ALD. Such cycles are repeated based on the required thickness to be deposited
[51].

Fig. (2:10): Schematic illustration of one growth cycle of an ALD MoS2 film [51].

Although the technique can be used to grow large scale MoS2 films, the uniformity
of the continuous monolayers on a wafer scale has not been reported yet. Beside
the ALD grown MoS2 films exhibit poor optical properties attributed to the
amorphous nature of the as-grown films. To improve the crystalline quality of the
as-grown films, further annealing in a sulphur rich environment is needed [51].

2.3.4 Molecular beam epitaxy (MBE)

The MBE technique has been widely used in growing TMD monolayers such MoS2
[52], MoSe2 [53] and WS2 [54] as well as TMD heterostructures like MoS2/h-BN
[55], MoTe2/MoS2 [56] and HfSe2/MoS2 [57].
In this technique the transitional metal source and chalcogen are
co-evaporated on a heated substrate using Knudsen effusion cells and e-beam
evaporators [53]. The growth rate and the morphology of the grown 2D material is
determined by source material fluxes [56].
There are several advantages of MBE over other deposition techniques, it
can control the elemental deposition rates, and switch easily from one source
material to another during deposition which enables the technique to have precise

38
composition control for heterostructure growth [56]. Another advantage of MBE is
that the growth of the films can be in-situ monitored using tools such as reflection
high energy electron diffraction (RHEED) [53].
However, the challenge in using MBE for growth of TMD heterostructures is
that the chalcogens are volatile either in elemental form or as small molecules.
This causes a low sticking coefficient on the growth-substrate, especially on inert
van der Waals substrates. Therefore, MBE generally requires low growth
temperature regimes and chalcogen rich conditions [56]. Finally, as the films are
grown in low temperature regimes, the grains have a limited grain size (<200 nm)
[54].

2.4 Conclusion
We have presented the main techniques that have been employed in the
deposition of TMDs and TMDs heterostructures.
Initially, researchers focused on depositing TMDs using top-down techniques
for fundamental research purposes and studying the intrinsic opto-electronic and
structural properties. The field of 2D TMDs was started by micromechanical
exfoliation of graphene in 2004 and then this technique become one of the most
widely used in exfoliating other 2D materials such as MoS2, WS2, h-BN ...etc. The
technique is simple and does not need sophisticated tools to produce monolayers
and multilayers. However, it is time consuming and for mass production it is not
reliable.
To scale up the production, other exfoliation techniques such as liquid
exfoliation have been tried. Beside the success that has been achieved in
producing large amount of 2D materials, the technique failed to produce pure
monolayers and there is always a mixture of monolayers and multilayers in the
yield. The problem of reaggregation of the produced monolayers still needs to be
tackled.
Top-down techniques, such as CVD, PVD, ALD and MBE, are offering an
alternative for growing 2D materials on large scales that are essential for industrial
applications. Rapid development in the deposition of 2D materials, driven by
optimizing the growth conditions and using different starting materials, has been
reviewed.
Graphene and other 2D materials with grain sizes of several hundreds of
microns have been grown using CVD. Continuous films on relatively large areas

39
have been achieved. Heterostructures of different kinds are successfully grown
using MBE.
Although remarkable successes have been achieved in depositing 2D
materials, the grown films are polycrystalline with relatively small grain size, and
there is a lot of grain boundaries that affect the opto-electronic and mechanical
properties of the film. The growth of films with large grain size is still a challenge.
More understanding about the film growth mechanisms and optimum growth
conditions still need further investigations.
The continuity and uniformity of the grown films will also affect the potential
applications of the films. More work is needed to produce continuous uniform
wafer scale films. This can be done through more investigations about the nature
of the starting materials, substrates and trying different growth regimes.
All in all, the 2D materials are becoming more and more interesting and
potential applications in future microdevices have already appeared on the
horizon. Through the available techniques, a variety of monolayers and
heterostructures have been produced. However, despite the intense research work
towards controlled deposition of 2D materials, wafer scale growth, uniformity and
grain size remain challenging issues. The chemical vapor deposition (CVD)
technique has shown great potential to grow large area TMDs. Nevertheless, mass
production for industrial applications is still at the very early stages which requires
more efforts to achieve the goal.

40
Chapter 3:Experimental and Simulation Techniques

In this chapter we introduce the characterization and simulation tools that have
been employed during the course of this project. Different techniques have been
used for characterization of MoS2 mono and multilayers. The following is a list of
techniques: optical microscopy, X-Ray photoelectron spectroscopy (XPS),
photoluminescence spectroscopy (PL), atomic force microscopy (AFM), Raman
spectroscopy, second harmonic generation (SHG) microscopy, scanning electron
microscopy (SEM), transmission electron microscopy (TEM) and X-ray
diffractometry (XRD). For CVD simulations we used COMSOL multiphysics 5.2a.

3.1 Characterization techniques

3.1.1 Optical microscope

Optical microscopes are inexpensive, rapid and nondestructive tools for
characterization large areas of 2D materials. As a result of light interference
between the 2D layer and underlying substrate, even single layer of graphene on
SiO2 covered Si show detectable contrast under optical microscope [58]. Figure
3:1 is a depiction of 2D/SiO2/Si stacking [59].

Fig. (3:1): Schematic depiction of optical reflection and transmission for a nanolayer
with thickness d1 and complex index of refraction n1 deposited on an SiO2 layer
characterized by thickness d2 and index of refraction n2 that is grown on top of a Si
substrate. Nanolayers deposited on SiO2 are visible due to interference between light
rays A, B and C reflected at various interfaces in the stack [59].

Under normal incidence the intensity of the reflected light from 2D/SiO2/Si stacking
can be calculated as follow [59]:

41
 i(φ1+φ2) +r2e−i(φ1−φ2) +r3e−i(φ1+φ2) +r1r2r3ei(φ1−φ2) ∣2
R(n) = ∣∣ er1e
i(φ1+φ2) +r1r2e−i(φ1−φ2) +r1r3e−i(φ1+φ2) +r2r2r3ei(φ1−φ2 ∣ 3:1

where
no −n1 n1 −n2 n2 −n3
r1 = no +n1
, r2 = n1 +n2
, r3 = n2 +n3
3:2
2πdi ni
are the relative indices of refraction and φi = λ
are the phase shifts induced by
changes in the optical path. no, n1, n2 and n3 are the refractive indices of air, 2D
material, SiO2 and Si respectively.
The reflectivity of the bare substrate is given by:
′2ei(φ2) +r3e−i(φ2) ∣2
R(n1 = 1) = ∣∣ rei(φ2) +r′2r3e−i(φ2) ∣
3:3
no −n2
where r′2 = no +n2
is the relative index of refraction at the interface between air and

the dielectric thin film.

Then the contrast between the the substrate and 2D material is given by:
R(n=1)−R(n)
C ontrast = R(n=1) 3:4

Using optical microscopy, the contrast of the 2D material depends on the thickness
of the underlying SiO2, 2D material thickness and camera filter used for imaging.
Therefore, optical microscopy is a rapid tool for distinguishing monolayers from
multilayers and confirming the uniformity and continuity of the grown film [60].
Figure 3:2 is a typical example of the calculated contrast as a function of MoS2
layer number when using 300 nm SiO2 covered Si as substrate [60].

Fig. (3:2): Color contrast plot of calculated contrast as a function of the number of layers
of MoS2 ultrathin films and the illumination wavelength for 300 nm thick SiO2/Si
substrates [60].

42
Figure 3:3 (a-m) color optical micrograph of exfoliated 1L-15L MoS2 on a 300 nm
SiO2/Si substrate. Figure 3:3 (n) is the measured contrast between MoS2 and
substrate as a function of layer number [61].

Fig. (3:3): (a–m) Color optical images of 1L–15L MoS2 on 300 nm SiO2/Si. The scale
bars are 5 μm for images a–l and 10 μm for image m. (n) Contrast difference values of
1L–15L MoS2 nanosheets on 300 nm SiO2/Si [61].

3.1.2 X-ray photoelectron spectroscopy (XPS)

X-ray photoelectron spectroscopy (XPS) is a non destructive surface analysis
technique that can be used to study the surface chemistry of materials. It
measures the binding energy of the materials under investigation. Such
measurements can be employed to find the chemical composition of sample
constituents. The XPS spectra can be obtained by exposing the sample to a
monochromatic soft X-ray with a certain wavelength. The incident X-ray removes
core electrons which are then collected and analysed. Since the mean free path
(inelastic scattering length) of the emitted electrons in solids is very short, the
detected electrons originate from the top few atomic layers, making XPS a surface
sensitive technique. The energy of the ejected electron represents the binding
strength of the core level which is an indicator of the presence of a certain
element. The emitted electrons have measured kinetic energies given by [62]:
K E = hυ − (Φ + B E) 3:5

43
where KE is the kinetic energy of the emitted electron, h is Planck's’ constant , Φ is
the work function BE is electron binding energy.
Furthermore, the atoms’ core level binding is very sensitive to the chemical
environment of the corresponding atom. For an atom in two different chemical
states, the binding energy for the same core level will be different. Such variation
in the binding energy results in a shift in the position of XPS peaks. This effect is
called a chemical shift and can be used to study the chemical status of the
elements in the sample [62]. The spectral intensity for the core level transition is
proportional to the quantity of that element in the sample.

3.1.3 Photoluminescence spectroscopy (PL)

In photoluminescence spectroscopy, we use the optical emission spectrum of
semiconductors to study their electronic structure.
When a semiconductor is exposed to radiation with photon energy greater
that the band gap of the semiconductor, electrons are excited to the conduction
band leaving holes in the valence band. When the electrons return to the ground
state, they recombine with holes and emit radiation with a certain wavelength that
is characteristic of the semiconductor under study [63].
Experimentally, a laser beam with an appropriate photon energy is focused
by an objective lens on the sample surface and the emitted radiation is recorded
after filtering the incident beam. The emission spectrum extracted provides
valuable information about the type of semiconductor whether it has a direct band
gap or an indirect band gap, the semiconductor quality and the presence of
impurities and defects [63].
As we explained in chapter one, an MoS2 monolayer is a direct band gap
semiconductor and, due to the valence band splitting in MoS2, there are two
excitonic transitions from the valence band to the conduction band called A and B.
The observed PL peaks of the two excitons are at 1.85 eV and 1.98 eV
respectively [13].
In this project, we employed PL measurements for a non-destructive study of
the intrinsic optical properties of our samples and evaluating the crystalline quality
of the as-grown monolayer films.

44
3.1.4 Raman spectroscopy
When photons interact with matter, they can be transmitted, scattered or absorbed.
The scattering processes have three categories depending on the frequency of the
out-coming photons (νo) which one shown in Figure 3:4. When the photons have
the same frequency νo to the incident ones, it is called Rayleigh scattering (elastic
scattering). When the scattered photons have a different frequency (ν0±νv), the
process is called Raman scattering (inelastic scattering). The intensity of Raman
scattering is very weak, in the range of 103 to 106 times lower than the intensity of
Rayleigh scattered photons which means the sample needs to be exposed to a
laser light in the Raman spectroscopy [64], [65].

Fig. (3:4): Idealised model of Rayleigh scattering and Raman stokes and anti-stokes
scattering [64].

Raman spectroscopy is a non-destructive technique based on the analysis of the

inelastically scattered photons from the medium, produced by the interaction of the
photon with the atomic vibrations that are also called phonons. Experimentally, the
shift in frequency between the incoming and inelastically scattered photons is
measured. The Raman effect was first observed in molecules by C.V. Raman in
1930, and nowadays the technique is one of the most widely used in analysing
molecules and crystals [66].
Raman spectroscopy is a very powerful characterization tool in the field of 2D
materials. Due to its sensitivity to symmetric carbon-carbon bonds it has been
used to identify different carbon based materials: graphene, graphite, single-wall
carbon nanotube (SWCNT), multi-wall carbon nanotube [67]. The G band of 1580

45
cm-1 corresponds to planar sp2 C-C carbon in graphene, graphite and carbon
nanotubes. The D band at 1350 cm-1 is a defect induced band corresponding to
sp2 carbon rings and its intensity is proportional to the presence of defects in the
sample and it has been standardized for defect detection in graphene [68].
In our project, Raman spectroscopy is employed to assess the presence of
MoS2 crystals grown by CVD and to extract a qualitative thickness measurement
as well as the growth induced strain in as grown films. For this purpose, we focus
on analysing the most intense Raman modes detected from MoS2 crystals. For
1 1
bulk material, the so called the E 2g and the A1g modes are located around 380

cm−1 and 405 cm−1 respectively. The former is due to in-plane vibrations of the
atoms, while the latter results from the out of plane vibrations. It has been shown
1
that for exfoliated MoS2 layers as the MoS2 is thinned down, the frequency of E 2g
1
increased while that of A1g decreased as shown in figure 3:5. The decrease in the

latter one is attributed to a decrease in the interlayer van der Waals interaction that
causes weaker restoring forces in the vibrations. The increase in the former one
might arise from either long range Coulombic forces or stacking induced interlayer
coupling. Therefore, this anomalous behavior is used as an indicator in identifying
the number of layers present in a sample [69].

1
Fig. (3:5): (a) Raman spectra of thin (nL) and bulk MoS2 films. (b) Frequencies of E 2g
1
and A1g Raman modes (left vertical axis) and their difference (right vertical axis) as a
function of layer thickness [69].

In addition to thickness measurements, it has been found that the positions of

these two peaks are very sensitive to the presence of strain in the film. In the case

46
of MoS2 monolayers, a considerable shift of -2.1% strain in the E 12g mode, and a

smaller, but observable shift of -0.4 % strain of the A11g mode has been detected

[70].

3.1.5 Second harmonic generation microscopy

When the intensity of the incident light is relatively low, interaction between light
and matter can be assumed to be linear, and the induced polarization by the
electric field is given by [71]:
P (t) = ε0 χ(1) E(t) 3:6

where P is the polarization, ε0 is the permittivity of free space, χ(1) is known as

the linear susceptibility and E is the electric field.

When the amplitude of the applied field is of the order of the characteristic
atomic electric field strength of 5.14×1011 V/m, the linear relationship breaks down
and the electric polarization depends on higher powers of the electric field [71]:
P (t) = ε0 [χ(1) E(t) + χ(2) E 2 (t) + χ(3) E 3 (t) + ...] 3:7
The quantities χ(2) and χ(3) are known as the second and third order nonlinear
optical susceptibilities, respectively.
The second-order polarization:
2 2
P (t) = ε0 χ(2) E (t) 3:8
gives rise to second harmonic generation SHG. In the SHG process, a laser beam
with an electric field of [71]:
E (t) = E e(−iωt) + cc 3:9
is incident upon non centrosymmetric crystals with second order susceptibility χ(2)
is non-zero, the nonlinear polarization created in such a crystal is given by [71]:
2 2
P (t) = 2ε0 χ(2) EE * + (ε0 χ(2) E e−2iωt + cc) 3 : 10
We see that the second-order polarization consists of a contribution at zero
frequency (the first term) and a contribution at frequency 2ω (the second term).
According to the driven wave equation 3:10, this latter contribution can lead
to the generation of radiation at the second-harmonic frequency.
MoS2 monolayers, odd number of the 2H [72] and 3R [73] polytypes possess
D3h symmetry. With the D3h symmetry, the second-order nonlinear susceptibility

47
tensor has nonzero elements of χ(2)′ ′ ′ =− χ(2)′ ′ ′ =− χ(2) =− χ(2) = χ(2) where x′y ′z ′
yyy yxx x′x′y ′ x′y ′x′

are crystalline coordinates [71].

The relationship between the incident electric field of the pump laser and the
coherently generated nonlinear polarization can be expressed as follows [74]:
(2) ︿ ︿ 2
I 2ω ∝ ୲e
︿
2ω . (χ ′ .e ω ).e ω୲ 3 : 11
x y ′z ′

where ︿
e ω and ︿
e 2ω are the polarization vectors for the beams at the fundamental
and the harmonic frequencies.
If the incident laser radiation is polarized along the x-direction, then x- and
y-polarized SHG emission from a MoS2 monolayer is given by [75]:
x 2
I 2ω ∝୲χ(2) .cos(3Θ)୲ 3 : 12
and
y 2
I 2ω ∝୲χ(2) .sin(3Θ)୲ 3 : 13

Here Θ is the angle of the crystal orientation, i.e. the angle between the orientation
of incident polarized radiation and one of the mirror planes of the MoS2 monolayer.
The orientations of each individual grain therefore can be determined as
follow:

√
y
1 I2ω
Θ = 31 tan− x
I2ω
3 : 14

MoS2 monolayers have three fold rotational symmetry, SHG without phase
information can not distinguish between opposite crystal orientations and the
measurable difference between grain orientations is from 0 to 30 degrees i.e. due
to the reflection symmetry of the MoS2 lattice, the six-fold symmetry present when
using SHG [75]. Equation 3:14 has an angular resolution of about 1o, which is
comparable to that achieved by TEM methods [76].
Figure 3:6(a) is an optical image of a polycrystalline MoS2 monolayer film,
with the corresponding (b) polarization resolved SHG image and (c) color coded
orientation map of the same film and (d) select regions of the sample showing the
actual crystal orientation vector of the grains [75].

48
 Fig. (3:6): (A) Optical image of CVD-grown monolayer MoS2. (B) SHG image of a
polycrystalline monolayer of MoS2 of the same area showing the grain boundaries. (C)
Color coded orientation map of the same area (D) Crystal orientation vectors of the
grains I,II,III, b and a [75].

In this project, we used SHG for confirming the uniformity of our MoS2
monolayer films. Polarization resolved SHG is used for finding the grain size
distribution within the polycrystalline film.

3.1.6 Scanning electron microscopy (SEM) and energy-dispersive

X-ray (EDX) spectroscopy
The scanning electron microscope (SEM) is a powerful instrument that permits the
observation and characterization of materials on a nanometer (nm) to micrometer
scale with resolving power of 1-5 nm [77].
The story of the scanning electron microscope and every other electron
microscope started in 1926 when H. Busch (1926) studied the effect of electric and
magnetic fields on the trajectories of charged particles in axially-symmetric fields
[78] and the first prototype instrument was built by Knoll (1935) [79].
The major components of an SEM are the electron column and the control
console. In the column there is an electron gun for generating electrons and
electromagnetic lenses for controlling the electron path down to the sample [77].
Two different kinds of guns are used in SEM. Thermionic emission sources
rely on high temperature to generate electrons such as a tungsten hairpin or
lanthanum hexaboride LaB6. The other is a field emission cathode, a sharp wire
fashioned into point a (100 nm or less in radius) and supported by a tungsten
hairpin [77].
The generated electrons accelerate to energies in the range 0.1-30 keV
(100-30,000 electron volts) and pass through the condenser lense which
de-magnifies the beam. The beam finally gets focused on the sample by an
objective lens. The electron ray path is shown in figure 3:7.

49
 Fig. (3:7): Electron ray traces through a schematic SEM column with a condenser lens
and a probe-forming or objective lens. Lens distances p and q are marked for each lens
[77].

The focused electron beam interacts with sample atoms elastically or inelastically.
This interaction derives many types of first and secondary emissions: elastically
scattered electrons leaving the sample (sometimes via a process called
"backscattering") provide an important class of information for SEM imaging.
Simultaneously with elastic scattering, the incident electrons lose energy and
transfer it in different ways to the target atoms giving rise to useful imaging signals
such as secondary electrons (SE), visible light (chathodoluminecence, or CL),
X-ray photons and Auger electrons [77].
The sample surface morphology of the sample can be determined by
analysing the SE. The SE signal results from ionization of near surface sample
atoms. More specifically SE are electrons ejected from the conduction band in the
case of metals and the valence band in the case of semiconductors and insulators
and their energy lies between (0-50) eV [77].
Backscattered electrons (BSEs) are beam electrons that are reflected from
the sample at different depths by elastic scattering and their energy is very close to
the incident beam. The scattering intensity of BSEs is proportional to to the atomic
number of atoms existing in the sample. Heavy atoms give rise to a strong BSE

50
signal, therefore the BSE spectrum can be used for identification of sample
composition [77].
In X-ray photon generation, the incident beam interacts with the tightly bound
inner shell electrons of the target atoms. As a result, the inner shell electron gets
ejected from the atom, leaving the atom in an excited state. The atom can relax to
the ground state in two different ways: 1) via an Auger process in which the
difference in shells energies is transferred to a valence electron ejecting it from the
atom with a specific kinetic energy; 2) via X-ray generation where the difference in
the shell energy is released as a characteristic X-ray photon with a sharp energy
peak. By analysing the generated X-ray the elemental composition of the sample
can be estimated [77].
For MoS2 monolayers, SEM is one of the powerful tools for quick
characterization of CVD samples on their growth substrates. Using the SE signal,
a clear contrast between the monolayer and substrate and monolayer and
multilayer can be observed. There is a weak interaction between the incident
beam and 2D material monolayers and few layers because the depth of the
interaction volume of secondary electrons in the 2D materials monolayers and few
layers are much greater than their layer thickness. Therefore, the contrast of 2D
materials with layer number might not be a result of direct interaction between the
beam and 2D material. However, it is suspected that 2D materials monolayers and
few layers attenuate a portion of the SE signal generated in the underlying
substrate. As a result of this, increasing the thickness of the 2D material will result
in a decrease in the number of the secondary electrons reaching the detector and
thus a reduction of the SE signal over 2D material layers [80]. Figure 3:8 shows
SE image of an as grown MoS2 film showing monolayer, bilayer and bare substrate
regions [43].

Fig. (3:8): SEM image of an MoS2 film. Monolayer (ML), bilayer (BL) and substrate
(SUB) areas are marked [43].

51
In addition, energy dispersive X-ray (EDX) could be the first checking point of the
elemental composition of the grown film.

3.1.7 Transmission electron microscopy (TEM)

Transmission electron microscopy (TEM) is an electron based technique used for
characterization of nanomaterials with resolution of 0.1 nm or better. As in the case
of SEM, electrons in TEM are generated by either thermionic emission or field
emission using (LaB6) crystals and fine tungsten needles [81]. The generated
electrons are accelerated by a high voltage (100-400 kV). Because of this high
voltage, the electron wavelength is decreased dramatically, essential for the
formation of high-resolution images [81].
Generally the image formation system in TEM consists of several lenses
(electromagnetic lenses) and apertures (see figure 3:9). A condenser lens is
located above the sample and it is used to produce parallel or convergent beam
illumination of the sample under study. The parallel beam is usually used in TEM
and convergent beam in the STEM. The objective lens sits just below the
specimen to collect the electrons that pass through the specimen and are directed
onto a focal plane and an image plane. The objective lens can be considered as
the most important part of the TEM as microscope resolution and the image quality
is largely determined by this lens. Under the objective lense there are intermediate
lenses which allow either image mode or diffraction mode to be formed and control
of the magnification in the case of image mode and change the camera of length in
the case of diffraction mode. There are also projection lenses that project the
image and diffraction pattern onto an imaging system [81].
Regarding the apertures, the condenser aperture controls the fraction of the
electron beam that hits the specimen and thus controls the intensity of illumination.
The objective aperture is placed in the back focal plane of the objective lens, the
place at which diffraction patterns are formed. This allows one to use the objective
aperture to select specific diffracted beams in diffraction contrast imaging. The
selected area aperture is located in the image plane of the objective lens and
allows selection of an area of interest in the conjugate sample plane. The basic
image operation systems are shown in figure 3:9 [81].

52
 Fig. (3:9): The two basic operation modes of a typical TEM imaging system: (A) the
diffraction mode: projecting the diffraction patterns (DP) onto the viewing screen and (B)
the image mode: projecting the image onto the screen [81].

In our work, we used TEM to study the crystallinity of our samples using selected
area diffraction patterns (SAED) as well as bright field images of large areas of the
film, high resolution transmission electron microscopy (HRTEM) for atomic
resolution structural imaging and scanning transmission electron microscopy
(STEM) for annular dark field or Z-contrast imaging [81].
As the electron beam passes through a crystalline sample the electrons are
diffracted by the atomic planes nearly parallel to the incident beam. The diffraction
is described by Bragg’s law [81]:
nλ = 2dsin(θ) 3 : 15

53
where n is an integer called diffraction order, λ is the wavelength of the incident
electrons, d is the separation between atomic planes and θ is the angle between
the incident beam and the atomic planes. In the SAED, a selected area aperture is
used to select an area of interest, then the intermediate lens is focused on the
back focal plane of the objective the place where diffraction patterns are formed.
In bright field (BF) imaging the diffracted beams are blocked using an
objective aperture and the transmitted beam is used for image formation and the
image contrast depends on the amplitude of the transmitted beam. In dark field
(DF) imaging, the transmitted beam is blocked and the diffracted beams are used
for image formation, resulting in so called diffraction contrast imaging.
For HRTEM, the image formation is the result of the interference between the
transmitted and diffracted beams in the image plane. This type of imaging is called
phase contrast imaging. As the phase variations can not be detected directly, the
phase modulation must be transformed into an amplitude modulation. This is done
by the interference of the transmitted beam with suitably phase delayed diffracted
beams. The phase difference in the sample exit wave changes into changes in the
amplitude of the wave. By taking the effect of the spherical aberration and lens
defocus only, the origin of the contract in HRTEM imaging can be described by a
contrast transfer function (CTF):
C T F = sin[πΔf λm2 + 21 πC s λ3 m4 ] 3 : 16

where ∆f is the defocus of the microscope, λ is the wavelength of incident

electrons, Cs is the spherical aberration and m is the spatial frequency.
In a special case when the macroscope is set to Scherzer defocus [81]. :
0.5
Δf Sch = 1.2(C s λ) 3 : 17
the point resolution of the microscope is given by:
1 3
rSch = 0.66C s4 λ 4 3 : 18
In STEM, the optics are different than in TEM. The electron beam is scanned over
the sample using deflection coils, such that the incident beam is always parallel to
the optical axis of the machine. During the scan, the STEM detectors register the
scattered or transmitted beams at each probe position for a selected scattering
angle. There are different types of detectors when using STEM mode. The
electrons scattered at 0-10 mrad are registered using a bright field detector. For
registering scattered electrons, Annular Dark Field (ADF), Medium-Angle Annular
Dark Field (MAADF) and High Angle Annular Dark Field (HAADF) detectors are

54
used. HAADF imaging is also called Z-contrast imaging as the scattered electrons
detected are mostly from Rutherford-type scattering with contrast proportional to
atomic number Z.

3.1.8 Atomic force microscopy (AFM)

AFM is a particular type of scanning probe microscopy that relies on the
interatomic forces between the tip and surface. AFM uses a tip typically made of
silicon or silicon nitride and with a typical radius of tens of nanometers and is
attached to the end of a small cantilever. By bringing the tip close to the sample
surface, the forces between tip and surface cause deflection of the cantilever.
There are different kinds of forces that the cantilever may experience, such as
repulsive mechanical force, van der Waals force or magnetic force [82] .
In AFM, the tip scans the sample surface by using a piezoelectric scanner.
The piezoelectric scanner moves the sample precisely in the x, y, and z directions
as the piezoelectric materials used can change their shape in response to the
applied voltage. A photo detector is used to detect the deflection of the cantilever
during the scanning. This is done by measuring the displacement of a laser beam
reflected from the back side of the cantilever. AFM generates an accurate
topographic map of the features on the sample surface by using a feedback loop
to control the height of the tip above the sample [83]. A schematic diagram of the
basic setup of AFM is shown in figure 3:10 [84].

Fig. (3:10): Schematic diagram of a basic setup of AFM [84].

Depending on the distance between the tip and sample surface, three types of
AFM operating modes can be distinguished: contact, non-contact and intermediate
contact.

55
 In contact mode AFM, the tip is brought very close to the sample, on the
order of a few angstroms. At such distances the electronic clouds of the tip and the
sample’s atoms strongly repel each other due to Pauli-exclusion. This repulsive
force causes the cantilever to be bent as it passes over the sample features during
the scanning. The feedback loop ensures the actual cantilever deflection stays
constant by comparing the actual bending with a given setpoint; if they do not
coincide, the control system modulates the signal applied to the piezoelectric
scanner so that the scanner retracts or extends in order to bring the deflection
back to the setpoint [85].
In non-contact AFM, a cantilever vibrating at its resonant frequency is placed
above the sample at a distance of tens to hundreds of angstroms. The force
exerted by the sample features changes the amplitude of vibration of the cantilever
with these changes used to extract the topographical data. Analogously to the
contact mode, a constant amplitude is preserved by the feedback loop by
changing the tip sample distance [85] .
The Intermittent contact mode, also called AC or tapping mode, is similar to
the non-contact mode except that the tip is brought closer to the sample. The force
is strongly repulsive when the tip touches the sample surface and attractive at the
rest of the vibration cycle of the tip. Detecting both short range repulsive and long
range attractive force increase the signal to noise ratio. This is an advantage
compared to the non-contact mode, which relies only on the attractive long range
forces [83], [85]. Figure 3:11 shows an idealized force curve between the tip and
the sample with highlighted probe-sample separation regions where different AFM
operating modes work.

56
 Fig. (3:11): Idealized forces between tip and sample surface highlighting where the
three imaging modes are operative [83].

AFM is one of the most extensively used tools for obtaining topological
images of 2D materials and their thickness measurement. Depending on the
measurements conditions, the recorded AFM values of MoS2 monolayer thickness
range from 0.6 nm to 1 nm and for graphene from 0.3 nm to 1.7 nm [86].

3.1.9 X-ray diffractometry (XRD)

X-ray diffraction (XRD) is a non-destructive technique that can be used to
study the atomic structure of materials, and characterize the crystallographic
structure, grain size and strain in polycrystalline films [87].
In XRD experiments, an X-ray beam strikes a crystal surface and is scattered
into different directions according to the crystal orientation. The scattered beams in
a direction that satisfies the Bragg conditions (equation 3:15) add constructively
and can be detected by an XRD detector. The beams in a direction that do not
satisfy Bragg conditions interfere destructively and cancel each other [87].
A typical X-ray diffractometer consists of an X-ray source, a sample stage,
and a detector. We are going to present the two configurations that have been
used in this project. The first one is Bragg Brentano geometry in which the
detector, the X-ray source and the sample are moved during the scanning in such

57
a way to guarantee the detector is always at 2θ and the sample surface is always
at θ to the incident X-ray beam. In this configuration only the crystallographic
planes which are nearly parallel to the sample surface can be detected. The
schematic diagram of this Bragg Brentano configuration is shown in figure 3:12.
The disadvantage of this configuration is that the penetration of the X-rays is large
and as a result there is intense signal from the substrate and a weak signal from
the surface. Therefore, it is not suitable for characterization of monolayer films
[88].

Fig. (3:12): Schematic of Bragg Brentano XRD [88].

In the case of monolayer films it is more convenient to use in-plane grazing

incidence angle X-ray diffraction (GIIXRD) to maximize the signal from the
monolayers. The X-ray beam is incident on the sample at very small angles (β~0)
that are larger than the total internal reflection angle of the sample usually 1o or
less and the detector is placed in a horizontal plane with respect to the sample
surface to detect the diffracted beams from crystallographic planes which are
perpendicular to the sample surface [88]. The schematic diagram of GIIXRD is
shown in figure 3:13.

58
 Fig. (3:13): Schematic of GIIXRD [88].

In this project, we used both types of XRD: θ-2θ scans were performed using
Rigaku SmartLab diffractometer and in plane grazing angle incident X-ray
diffraction GIIXRD using synchrotron radiation at Diamond Light Source.

3.2. CVD simulation using COMSOL

In our project the CVD simulations are performed by a commercial software
COMSOL Multiphysics 5.2a. The COMSOL software has a powerful interactive
environment for modeling and solving different kinds of scientific and engineering
problems based on partial differential equations.
In our CVD simulations, the following coupling was investigated:
1- Momentum transfer - compressible Navier-Stokes equations
2- Mass transfer - convection and diffusion
3-Energy transfer - heat convection and conduction
Using the above couplings we correlate the carrier gas velocity to the
pressure and temperature of the CVD system and the reactant concentration to
the carrier gas velocity, pressure and temperature of the system.
The Navier-Stokes equations are the equations governing the fluid dynamics.
When using these equations along with a set of boundary conditions such as the
dimension of the inlets, the outlets and the walls of the system, the velocity,
pressure, temperature and the density of the fluid can described inside the system.
These equations are time-dependent continuity equations for the mass,

59
momentum and energy conservation respectively. In the case of a compressible
Newtonian fluid, the momentum equation is [89]:
ρ( მu
მt
+ u.∇u) =− ∇p + ∇.(μ(∇u + (∇u)T − 23 μ(∇.u)I) + F 3 : 19

where ρ is the fluid density, u is the velocity vector field, and µ is the fluid dynamic
viscosity, T is transpose matrix and I is identity matrix. The left-hand side term
corresponds to the inertial forces and the first term on the right hand side
describes the differential pressure effect, the second term viscous forces and the
last is external forces applied to the fluid.
Equation (3:19) is solved together with the continuity equation which
describes the relationship between the temporal rate of change of mass density
and the divergence of the mass flux [89].
მρ
მt
+ ∇.(ρu) = 0 3 : 20
The COMSOL fluid dynamics module is used to solve the above equations in our
CVD system to describe the velocity profile of the carrier gas.
We also used mass transport in our CVD simulations to study the
concentration profile of the reactant material inside the CVD reactor. The mass
transport is coupled with velocity profile and temperature gradient along the tube
furnace.
The chemical reaction engineering module of COMSOL is used for this
purpose by solving the diffusion-convection equations [90]:
∇.(− Di ∇ci ) + u.∇ci = Ri 3 : 21
N i =− Di ∇ci + uci 3 : 22
where Di is the diffusion coefficient of the ith species, ci and Ri and Ni are the
corresponding contraction, reaction rate and flux respectively.
Heat transfer is also very important in the CVD simulations since temperature
gradient can affect the density of the fluid, which can affect the fluid flow and mass
transfer. The heat transfer can be described through the following equation [90]:
ρC p u.∇.T + ∇.q = Q 3 : 23
q =− kf ∇T 3 : 24

where Cp is the specific heat of fluid at constant pressure, T is the temperature of

the system, Q are any other heat sources (chemical reaction), kf is the thermal
conductivity of the fluid, and q is the heat flux.

60
3.3 Summary
We have presented the characterization and simulation tools used during the
course of this project. We have introduced the principles of visualising 2D
materials under an optical microscope and explained how this simple tool provides
valuable information about the 2D material thickness, uniformity and continuity. For
the chemical analysis, the principle of XPS was introduced and its operation has
been explained. Raman spectroscopy provides a qualitative MoS2 thickness
measurements and strain analysis in the CVD grown MoS2 monolayer
polycrystalline films. Its working principle was presented. The origin of the Raman
modes shifting as a function of the MoS2 thickness was explained and examples
from the literature were given. Another thickness measurement method, using
atomic force microscopy, was presented and the different operational modes are
summarized. Second harmonic generation (SHG) microscopy was used for
checking the monolayer film uniformity on a large scale and studying the grain size
and orientation distribution as a function of experimental conditions. The physics
behind SHG and polarization resolved SHG was outlined. The basic principles of
the scanning electron microscopy were explained, with special attention given to
the imaging by secondary electrons and the generation of X-rays. The imaging
modes of TEM that have been used in this project such as dark field imaging,
bright field imaging, HRTEM and STEM annular dark field imaging with the origin
of contrast in each imaging mode were presented. For studying the crystalline
nature of MoS2 monolayers, two types of XRD, conventional XRD and GIIXRD
were introduced. Finally, for CVD simulations, COMSOL software with a fluid
dynamic module, a chemical engineering module, and a heat transfer module were
presented.

61
Chapter 4: LPCVD growth of continuous MoS2 monolayer
films

Abstract
We have employed low pressure chemical vapor deposition (LPCVD) as our
growth method for MoS2 monolayers. Molybdenum dioxide (MoO2) and sulphur (S)
have been used as starting materials to achieve the growth of wafer scale uniform
MoS2 monolayers with crystal sizes up to 400 microns on SiO2/Si substrates.
X-ray photoelectron spectroscopy (XPS) has confirmed the chemical
composition of the monolayers. The thickness of the monolayers have been
determined quantitatively using atomic force microscopy and the result is further
corroborated by second harmonic generation and Raman spectroscopy. The
uniformity of the films is verified using optical and nonlinear two photon
microscopy.
Grazing incidence in-plane x-ray diffraction (GIIXRD) is used to study the
growth induced strain in the as grown monolayer films as well as the monolayer
thermal expansion coefficient.
The crystalline quality of the monolayers has been confirmed by transmission
electron microscopy (TEM) and GIIXRD. Finally, the optical and electrical
properties of the monolayers have been evaluated by photoluminescence
spectroscopy and a field effect transistor respectively.

4.1 Literature review

There have been few approaches about growing continuous MoS2 monolayers
films on large scale (centimeter scale) via sulphurization of MoO3, MoCl5 and
Mo(CO)6. The growth of the continuous monolayer is being controlled via the
partial pressure of the reactant materials which plays the major role on the final
product composition and its uniformity [91]. The role of the reactant partial
pressure will be discussed comprehensively in the following chapters. In the
following paragraphs we will discuss the most frequent approaches that have been
used by classifying them according to the Mo source being employed.

62
 Regarding the molybdenum oxides, MoO3 is the most extensively used. Jing
Zhang et al. (2014), successfully grew MoS2 monolayers on a centimeter scale
using MoO3. In their approach, the partial pressure of MoO3 is controlled by
changing the evaporation temperature of MoO3 and growth temperature of the
substrate. Although, the growth of a continuous monolayer film is successfully
managed, the grown monolayer grains have relatively small sizes up 600 nm [40].
Similarly, Jaeho Jeon et al. (2014) used MoO3 to grow MoS2 monolayers, bilayers
and trilayers. However, as a result of the local variation in MoO3 concentration, the
substrate was not fully covered with film and the grain size was still in the range of
nanometers [40]. Finally, Shanshan Wang et al. (2015) used the substrate
geometry to better control MoO3 concentration in their CVD growth, however the
films produced are not uniform and the grain size was few microns [92]. Despite
the relatively large amount of MoS2 monolayers grown using MoO3, those
approaches are still facing scalability issues.
To overcome scalability and have better control on the deposition, recent
reports on TMD growth have focused on using alternative precursors as Mo
sources, primarily transition metal halides such as MoCl5 or transition metal
carbonyls such as Mo(CO)6.
Using MoCl5, the CVD growth of MoS2 monolayers and multilayers on a
large scale is well controlled [41]. In contrast to the well defined morphological
shapes of MoS2 monolayers that can be produced using MoO3, the monolayer
produced using MoCl5 are continuous and do not have well defined morphology,
and the monolayers have poor crystalline quality [93].
The growth of continuous monolayer films has been achieved by MOCVD
using Mo(CO)6 and C4H10S. Using this approach, the growth of continuous
monolayers is well controlled through the growth time. However, for monolayer ,
the growth time is quite long (26 hr) which might not be desirable from the
production point of view [47]. The grain size of the produced monolayers is only
few microns (up to 10 microns) [47]. Depending on the growth conditions, the
grown film might contain undesirable carbon compound such Mo2C or MoOC [94]
[95]. The carbon impurities from starting material may be another issue for
opto-electronic applications [91].
All in all, several attempts have been taken for growing MoS2 monolayers on
a centimeter scale using different CVD approaches. A remarkable success in the

63
film growth driven by controlling the partial pressure of reactants has been
established, however, the uniformity, continuity and grain size of these films is still
an issue and there is no standardized approach being presented for reproducible
growth.
In this chapter, we are introducing an LPCVD approach using MoO2 as a new
starting material. MoO2 has a relatively low vapor pressure compared to other Mo
sources. We take and advantage of such a low vapor pressure to establish a
standard approach that can be used to grow wafer scale, continuous, uniform,
MoS2 monolayers with the optimum grain size.

4.2 Experimental

4.2.1 CVD set-up, temperature dependence

The MoS2 monolayers have been grown using a TSH 16/75/450 single zone tube
furnace (from Elite Thermal Systems Limited) with a maximum operating
temperature of 1600 oC and a heating zone that is 45 cm in length, integrated with
Eurotherm 2416 PID and 2116 over-temperature controllers. The gas flow through
the tube furnace was controlled by an MKS Mass Flow Controller (MFC) Type
1179A with a maximum flow of 1000 Standard Cubic Centimeters per Minute
(SCCM) and calibrated for argon and argon-hydrogen mixture. The MFC was

controlled by a MKS PR4000 control box. The exhaust gas was passed through a
sealed glass bottle half filled with water to remove particulates and prevent
contaminated gas travelling up the exhaust. The setup of the furnace is shown in
figure 4:1.

Fig. (4:1): The Elite thermal system used for growing MoS2 monolayers.

64
The temperature distribution was measured along the tube from the center to both
ends when the furnace was under vacuum (10 mbar) in order to know the
temperature range that might be used in the experiments. For this purpose, a 1 m
long, 0 → +1100 °C, K type thermocouple (supplied by RS) was used, with
both-ends of the tube closed (without the flow of the carrier gas). The temperature
profile of the furnace is shown in figure 4:2. It has a symmetrical form as expected.
The central region has a constant heated zone at the designated temperature, with
a width of 40 cm providing a relatively uniform temperature in the growth region
where the Mo source and substrate are placed. The change to this temperature
profile in the presence of the flow will be simulated using COMSOL and the result
will be presented in chapter 5.

Fig. (4.2): Measured furnace temperature profiles for different set-temperature of the
tube furnace.

4.2.2 Experimental procedure

In a typical two-zone CVD growth procedure, 300 mg of high purity MoO2 powder
(99% pure from Sigma Aldrich) was placed in a quartz boat. A SiO2 covered Si
wafer (5×3 cm2 by dimension) was cleaned with acetone and ultrasonicated for 10
minutes and then placed face down on the MoO2 container. The MoO2 container
with the substrates were then placed at the center of the 75 mm diameter tube
furnace. Another boat containing 5-10 mg of sulphur powder (99.98% pure from
Sigma Aldrich) was placed 42 cm away from the MoO2, in a region where the
temperature reaches 200 oC. Prior to the growth, the furnace was flushed with

65
Argon gas (1000 SCCM) for about 30 minutes. Then the furnace was pumped
down to 10 mbar and the center was heated to 800 oC with a heating rate of 15 oC
/minute in an argon flow (200 SCCM). After keeping at 800 oC for 10-20 minutes,
the furnace was naturally cooled down to room temperature. The temperature of
the sulphur source also rises, roughly in line with the temperature profile at the
centre of the tube furnace, reaching an average of 200 oC during the 15 minutes
when the sample is in the growth temperature. The temperature profiles during the
growth process are summarized in figure 4:3.

Fig. (4:3): Temperature profile for sublimation of starting materials during a typical
growth run.

4.3 Results and discussion

4.3.1 Self-limiting growth of monolayer thin films

Figure 4:4 shows a low-resolution optical photograph of an as grown centimeter
scale MoS2 monolayer film (the upper purple colored part) and the bare 300 nm
SiO2/Si substrate (the lower part). There is a clear contrast for color of the area
covered by the monolayer film (blue) compared to the color of the substrate (light
pink). The color change is consistent with the expected light interference between
the monolayer film and the 300 nm SiO2 covered Si substrate [58], [59], [96], as we
have explained in section 3.1.1. For a growth time close to 15 minutes, the whole
substrate was found to be covered by a uniform monolayer thick film.

66
 Fig.(4:4): Optical images of a centimeter scale MoS2 monolayer (the upper part) and the
bare SiO2 covered Si substrate (the lower purple part).

Fig. (4:5): Second harmonic generation (SHG) nonlinear two-photon microscopy

revealing the polycrystalline nature of continuous MoS2 monolayers where grains with
different orientations show different color (see section 4:5 for details).

As the contrast between monolayer and multilayers of MoS2 thin film can be
quite subtle, we turn to nonlinear optical microscopy for additional confirmation as
the SHG signal is both sensitive to layer thickness as well as the in-plane
orientation of the crystallites [97]. Figure 4.5 shows a snapshot of our
centimeter-sized monolayer. This shows that the film indeed consists of a
continuous monolayer. The pseudo-color texture is due to its polycrystalline
nature, as grains with different orientations showing different color (see section

67
3.1.5 of Chapter 3). This suggests that the film-growth, once started, can
proceeded very rapidly but is self-limiting with very little evidence for pronounced
secondary- or multilayer growth.

4.3.2 Chemical analysis by XPS measurements

Elemental composition and bonding in the CVD grown MoS2 monolayers was
examined with X-ray photoelectron spectroscopy (XPS). The data was taken at a
base pressure of 2.5×10-9 mbar using an Al X-ray anode. The XPS source is the
XM1000 MkII monochromator/x-ray source from Omicron (Scienta Omicron). The
analyser is an Omicron EA125 hemispherical energy analyser. Survey scan was
conducted at 0.5 eV per step, with a dwell time of 0.5 second at each step.
Detailed scans were conducted at 0.1 eV per steps, with 1 second at each step,
and an average 10 cumulative sweeps. For data analysis, the software XPSPeak
4.1 was used with the background subtraction type set to Shirley. Five elements
are present in the spectra acquired. An example of the survey scan is presented in
figure 4:6: Mo and S from the monolayer MoS2, as well as Si, O from the substrate
and C due to residual contamination from the samples having been in air.
In Figure 4:7 a detailed scan performed over the Mo 3d peaks and adjacent
S 2s peak is presented. Clear Mo peaks corresponding to 3d 3/2 and 5/2 states
are present with intensity ratio of 3:2 and an energetic position relative to the S 2s
peak as expected. In our samples, the binding energy of the molybdenum doublet
Mo 3d3/2 and Mo 3d5/2 appear at energies of 232.7 eV and 229.6 eV, and the
sulphur S 2s1/2, S 2p1/2 and S 2p3/2 at 226.85 eV, 163.6 eV and 162.4 eV
respectively. This is in good agreement with literature values for characteristic
band positions of fully transformed MoS2 grown by sulfurization of MoO3 [98], [99]
confirming complete conversion of MoO2 to MoS2 in our CVD growth process.
Figure 4:8 zooms into the S 2p peak region where a clear doublet structure is
observed with a good 1:2 ratio, as expected for emission from the corresponding
2p1/2 and 2p3/2 levels.

68
Fig. (4:6): XPS survey spectrum of MoS2 monolayers grown on SiO2/Si substrate.

Fig. (4:7): Mo 3d and S 2s XPS spectrum of MoS2 monolayers.

69
 Fig. (4:8): S 2p XPS spectra of MoS2 monolayers.

4.3.3 Thickness measurements by AFM

For direct thickness measurements, a JEOL JSPM-5200 atomic force microscope
was used. Figure 4:9 shows an AFM topography map of a single crystal region of
a MoS2 monolayer film. Figure 4:10 is a topographical measurement in a
polycrystalline region of the same film showing a grain boundary. The monolayer
thickness measurements were performed on the edges of the films. The edge of
an isolated crystalline film is shown in figure 4:11 with the corresponding height
profile displayed in figure 4:12. For a continuous film, the film is first scratched with
a sharp tip and the measurements were then taken at the edges of the scratch as
shown in figure 4:13 with the corresponding height profile displayed in figure 4:14.
The height difference between the substrate and the crystal edge is found to be
0.9 nm and 0.85 nm for the isolated crystal and the continuous polycrystalline film
respectively. These values (0.85-0.90 nm) agree well with the results from
mechanically exfoliated MoS2 monolayers reported in the literature [13].
To confirm that our films are uniform, other techniques such as Raman
spectroscopy and second harmonic generation microscopy will be discussed in the
next paragraphs as AFM is time consuming and continuous films have to be
scratched, i.e. damaged.

70
 Fig. (4:9): AFM topography of a single Fig. (4:10): AFM topography of a MoS2
crystalline MoS2 monolayer. polycrystalline film showing grain
boundaries (GB).

Fig. (4:12): Height profile measurements

Fig. (4:11): AFM micrograph of MoS2 of the crystal shown in fig. (4:11); data
single crystal monolayer edge. taken along dark line.

Fig. (4:13): AFM micrograph of a scratched Fig. (4:14): Height profile measurements
continuous monolayer MoS2 film. of the film shown in fig. (4:13): data taken
along the dark line.

71
4.3.4 Uniformity of monolayer films using multiphoton microscopy
Due to the absence of the inversion symmetry in the MoS2 monolayers, 3R
stacking bulk MoS2 and odd layers of 2H stacking, they exhibit very strong second
harmonic generation (SHG) [73]. Therefore, this spectroscopic property can be
employed for checking the uniformity of MoS2 films and their stacking order. For
this purpose, a confocal Zeiss LSM 780 multiphoton microscope was used for
measuring second harmonic generation from MoS2 monolayers. The
measurements were performed in a reflection geometry using normal incidence
excitation. The pump radiation was supplied by a mode-locked Ti:sapphire
oscillator operating at an 80-MHz repetition rate. The pulses were of 140-fs
duration and centered at a wavelength of 800 nm. Using a 10X objective and 20X
with numerical apertures of 0.3 and 0.8 respectively. Figure (4:15) shows a SHG
image of an incomplete coverage of MoS2 monolayers on SiO2 covered Si
substrate. For this film the growth time was 5 minutes which is shorter than 15
minutes for a complete coverage. There is a clear contrast between the substrate
(dark regions) and the monolayers (grey regions). This is because SiO2 is
amorphous and does not produce SHG.
Figure 4:16 (a-d) shows SHG maps of continuous polycrystalline monolayers
films grown at temperatures of 700 oC,750 oC , 800 oC, 850 oC with a growth time
of 15 minutes. From the SHG contrast in the images, we can confirm that the films
are uniform and the slightly dark lines represent the grain boundaries [76].

Fig. (4:15): SHG image of MoS2 monolayers grown at 800 oC (grey regions) on a
SiO2/Si substrate (dark regions).

72
 Fig. (4:16): SHG images of continuous MoS2 monolayer films grown at growth
temperatures of (a) 700 oC, (b) 750 oC, (c) 800 oC and (d) 850 oC.

Moreover, to stress that our films are monolayers we have shown a 3R

stacking bilayer in figure (4:17). In this figure, there is a continuous background of
monolayers. On top of that, there is a secondary nucleation of 3R stacking (bright
triangles). In such a stacking there is constructive interference between SHG
generated from the top and bottom layers which is why the top layer appears to be
about 2.5 times brighter (the profile is shown in figure 4:18). In the case of 2H
stacking bilayers as shown in figure 4:19, there is also a continuous monolayer
background with triangles (dark ones) that have 2H stacking. In this case there is
destructive interference between SHG generated from the top and bottom layers
and the two beams completely cancel each other (profile is shown in figure 4:20).
Theoretically, the total SHG from the stacking region of bilayers can be
approximated as follows [97]:

73
 I t (α) = I 1 + I 2 +
√I I cos(3α )
1 2 4:1

where I t , I 1 and I 2 stand for the SH intensity in the stacking region, top grain and
bottom grain respectively. α is the stacking angle between the two grains.
However, the measured SH from the stacking regions is slightly different due
to the absorption and intensity variation (∼10-15%) among different flakes [97].

Fig.(4:18): SH intensity profile taken

from MoS2 monolayers and 3R
Fig. (4:17): SHG image showing 3R-Stacking stacking bilayer, data taken along dark
MoS2 bilayer grains G1,G2, G3 and G4 (white lines on G1, G2, G3 and G4 shown in
triangles). fig. (4:17).

Fig. (4:20): SHG intensity from monolayer

and 2H-stacking bilayer MoS2, data taken
Fig.(4:19): SHG image of 2H-stacking along the white lines on G1,G2 and G3
MoS2 bilayer grains G1, G2 and G3 (dark shown in fig. (4:19).
triangles).

The synthesized MoS2 monolayers show unprecedented uniformity over a

o
range of growth temperatures (700-850 C). Beside the uniformity of the
monolayer, the as-grown MoS2 thin films are continuous and uniform across an

74
area of wafer scale. We have taken the data from different growth runs and our
approach proved to be reproducible. We can consider our approach to be the
starting point for mass production of MoS2 monolayers.

4.3.5 Phase identification by Raman spectroscopy, TEM and X-ray

diffraction
4.3.5.1 Raman spectroscopy
Raman spectroscopy is one of the most important characterization tools in the field
of the TMDs that has been used in qualitative thickness measurements and
identifying strain in the as grown films. MoS2 exhibits two characteristics Raman
1 1
peaks which are known as E 2g and A1g modes corresponding to the in plane

vibrations of Mo and S atoms and out of vibrations of S atoms. The separation

between these two peaks is very sensitive to the thickness of MoS2 film [69]. As
the thin film thickness decreases, the separation between the peaks also
1 1
decreases such that the E 2g peak red-shifts (phonon softens), while the A1g peak

blue-shifts (phonon stiffens), therefore such shifts in the Raman peaks can be
used to determine the layer number of MoS2 [69]. However, the real situation is
complicated by the presence of growth induced strain. The separation between the
two Raman peaks is affected by the presence of strain within the monolayer,
1
especially the position of E 2g , as well as the laser wavelength that has been used

for the the measurements [100], [101]. In our case, using a SiO2 covered Si
substrate, the thermal expansion coefficient of MoS2 (10-5 /oC) is roughly 1000
times greater than that of silicon dioxide (5.6 10-8 /oC). Our CVD-grown MoS2 on
SiO2/Si at elevated temperatures always suffers from a native tensile strain
induced by interactions between the MoS2 film and the substrate. In some cases
the strain is strong enough to cause fracture of the monolayer, especially in fast
cooling rates during CVD growth [102]. The native strain in our samples will be
quantitatively analysed using XRD measurements in the following section.
1
For MoS2 monolayers, the position of the E 2g peak is at about 384 cm-1 and
1
the A1g peak at about 405 cm-1. There is a slight variation in this position of the two

peaks depending on the method that has been used for depositing monolayers,
the presence of the strain and the laser power and wavelength that has been used
for characterization [41], [69], [101], [103].

75
 In our case, Raman measurements were carried out under ambient
conditions at room temperature using a JY Horiba HR800 micro-Raman system. A
532 nm wavelength laser beam with power of 0.5 mW and spot diameter of 1 μm
was focused on the sample. Figure (4:21) is a typical Raman spectrum of a
1 1
continuous film grown at 800 oC. The E 2g peak is located at 383.6 cm-1 and A1g at

404.8 cm-1, with separation about 21 cm-1. These measurements agree well with
values reported for CVD grown MoS2 monolayers [41]. However, this separation is
slightly larger than the one recorded for the counterpart mechanically exfoliated
monolayer (separation 18 cm-1) confirming the presence of strain in the CVD
grown monolayers [101], [104].

Fig. (4:21): Raman spectrum of a continuous MoS2 monolayer film grown at 800 oC.

4.3.5.2 Transmission electron microscopy investigations

The crystal quality of the CVD grown MoS2 was characterized at the atomic level
using aberration-corrected transmission electron microscopy (TEM). Figure 4:22(a)
shows a continuous MoS2 film that was transferred to a TEM grid following a poly
(methyl methacrylate) (PMMA) assisted method (see Appendix A1 for more
details). The continuity of the transferred film indicated the high quality of the
grown sample; folds and holes observed on the TEM specimen were caused by
the transfer process. Figure 4:22(b) is an HRTEM image of the same monolayer
region. Figure 4:22(c) is the corresponding selected area diffraction patterns

76
(SAED) with [001] zone axis. It can be seen that the monolayer exhibits high
quality single crystalline nature with hexagonal lattice symmetry. Using SAED
diffraction spots, we have found the value of the MoS2 monolayer lattice constant
(a=3.20 Å). This value agrees with the ones reported in literatures [41], [105].

Fig. (4:22): (a) Transferred MoS2 monolayer on lacey carbon film (b) HRTEM image of
monolayer region (c) SAED from the same region in (b).

4.3.5.3 Synchrotron X-ray in-plane grazing angle diffraction GIIXRD

The crystalline structure and quality of our films is further verified using in-plane
grazing incidence X-ray diffraction (GIIXRD). The measurements were done using
synchrotron radiation (with a wavelength=0.65255 Å) and an incidence angle of
0.2 degree on the beam line l07 at the Diamond Light Source in Harwell
Campus/(Didcot, OX11 0DE, UK). The scanning was performed at fixed substrate
angle using an in-plane rotating detector. The X-ray beam was defocused on an
area of 200 μm2 and data were taken from different parts of the monolayer film.
The measured data is compared with that calculated for the 1H phase (fig. 4:23).
The peaks (100) and (110) have been detected on the different parts of the film
and the positions of these peaks coincides with the calculated 1H phase. A zoom
in scan of the two planes (100) and (110) of seven points on the sample are
plotted in figure 4:24.

77
 Fig. (4:23): Measured GIIXRD from a MoS2 monolayer film.

Fig. (4:24): Zoom in scan of (100) and (110) MoS2 monolayer planes at seven different
points on the sample.

78
 The XRD peak broadening could be due to strain in the film, grain sizes and
thin film thickness [106]. In our case the grain size is several hundreds of
micrometers, and in the grazing incidence, the film can be assumed to be infinitely
thick. Therefore the broadening in the XRD peaks can only be attributed to the
strain and instrumentation. Using synchrotron monochromatic radiation, the
instrumentation broadening is minimized and can be neglected.
We have investigated the growth induced strain in the films quantitatively
focusing on the shift of the peak (100) that is taken from seven different positions
of the film from different growth runs as shown in figure 4:24. The strain can be
calculated as follows [106]:
Δd
d
=− Δθ × cot(θ) 4:2

where d is the lattice spacing and Δθ is the XRD peak broadening (FWHM).
For a hexagonal system the lattice spacing is given as:
2 2 2
1
2 = 34 ( h +hk+k
a2
)+ l
c2
4:3
d

where a is the lattice constant, (hkl) are Miller indice, and c is the lattice constant
perpendicular to the basal plane.
In the case of monolayer l=0, therefore eq. 4:3 is reduced to:
2 2
1
2 = 34 ( h +hk+k
a2
) 4:4
d

and :

a= 2
√3
d √h2 + hk + k2 4:5

Δd√h + hk + k
2 2 2
Δa = 4:6
√3
Dividing equation 4:6 by 4:5, we get:
Δa Δd
a
= d
4:7
Δa
a
represents the strain in the monolayer or the error in the lattice constant.
Figure 4:25 shows the lattice constant measurement at seven different
points on the monolayer film. The variations in the lattice constant range between
3.17 to 3.156 Å. Such a variation in the lattice constant values indicate that the film
is under nonuniform strain. The error bars in figure 4:25 represent the strain within
a specific point, as the XRD spot was 200 μm2, these error bars also represent the
strain variation in a 200 μm2 film area. The strain within these spots ranges
between 0.5% (point 1) up to 1% (point 6). These values of strain agree with those
reported in the literature using Raman spectroscopy [104].

79
 Fig. (4:25): Lattice constant variation as result of growth induced strain in the
polycrystalline monolayer film.

We propose that this strain is induced by natural fast cooling of the samples
at the end of the growth process. As we have mentioned, the thermal expansion
coefficient of the underlying substrate (silica) is roughly 1000 times smaller than
MoS2 monolayers. A natural cooling process from the growth temperature (800 oC)
could induce a significant strain in the film as a result of this mismatch in the
thermal expansion coefficient. The resulted strain is not uniform across the film as
shown in figure (4:25).
When the data is taken from neighboring grains that are strongly differently
strained, a splitting the in peaks is expected. Figure 4:26 is an example of such an
expected case. As we can see the peaks (100) and (110) are clearly splitted.

Fig. (4:26): GIIXRD peak splitting from differently strained grains.

80
4.3.6 In situ annealing of the MoS2 monolayers
We have also studied the effect of temperature on the as grown monolayers. For
this purpose we have tested the samples under 25, 100, 200, 300, 400, 500 and
800 oC in an ultra-high vacuum (UHV) to see how the film is affected by annealing.
Figure 4:27 shows GIIXRD measurements of an in-situ annealed MoS2 monolayer
film. As we can see the film preserve its crystalline nature up to 800 oC and it is
stable under this temperature only for a few minutes (t⋍5 minutes).
Figures 4:28 show the details of the diffraction profile zoomed in on the
(100) and (110) reflections of MoS2 monolayer planes at different annealing
temperatures.

Fig. (4:27): GIIXRD of a MoS2 monolayer at room temperature and an annealing

temperature of 800 oC.

81
 Fig. (4:28): High resolution (100) and (110) MoS2 monolayer planes at different
annealing temperatures.

The thermal expansion coefficient of MoS2 monolayers on SiO2/Si has been

extracted from lattice constant expansion as a function of temperature as shown in
figure 4:29. At room temperature the film is under tensile strain induced by the
growth temperature (800 oC). The lattice constant at room temperature is found to
be 3.167 Å. When the fim is heated to 100 oC, the strain is relaxed, as can be seen
by the sharp drop in the lattice constant value which is 3.161 Å. Such relaxation is
expected as there is a weak interaction between the film and substrate.
Furthermore, this relaxation confirms that the film was under tensile strain.
Similarly, strain release has been observed when transferring as grown monolayer
to a fresh substrate [104].
By heating the film, the lattice constant increased linearly to reach a value of
3.170 Å at a temperature 800 oC, i.e. nearly reached the value of the growth
temperature.

82
 Fig. (4:29): MoS2 monolayer lattice expansion as a function of temperature extracted
from (110) plane.

The experimental data were fit with polynomial:

2
xT = xo + x1 T + x2 T 4:8

Here T is the temperature (°C) and the ratio x1/xo gives the principal linear
coefficient of thermal expansion in the a direction [107].
The fit through the experimental points gives the temperature dependence
of the lattice constant a :
6 9 2
a = 3.160 + 7.845 × 10− T + 7.22 × 10− T 4:9
The principal linear coefficients of thermal expansion is (2.5±1.2) x10-6 /oC. This
value is slightly greater than the one reported for bulk MoS2 powder (1.9 × 10-6/°C)
[107].

4.4 Physical property measurement

4.4.1 Optical properties
Photoluminescence (PL) measurements, were performed on MoS2 monolayer
samples (both continuous and isolated films) . All measurements were carried out
under ambient conditions at room temperature using a JY Horiba HR800
micro-Raman system.

83
 A 532 nm wavelength laser beam with power of 0.5 mW and spot diameter of
1 μm was focused on the sample. A typical example of PL measurements is
shown in figure 4:30. There are two pronounced peaks known as A exciton and B
exciton centered at 1.80 eV and 1.96 eV respectively. These two peaks are
associated with direct transitions from the lowest conduction band to the spin-orbit
split valence bands (see Fig. 1.6). This result agrees with PL measurements
recorded from mechanically exfoliated MoS2 monolayers [13]. The full width at half
maximum (FWHM) width of A exciton is 71 meV. This is even narrower than the
mechanically exfoliated MoS2 monolayers on SiO2 substrate and confirms the high
quality of our films [108].
The A exciton peak is comprised of transitions due to neutral exciton (Ao) and
trions (A-) generated in the monolayer [109]. The PL spectra of single-layer MoS2
fitted with three Lorentzian functions as shown in figure 4:30, with the trion peak
centered at 1.78 eV and the neutral exciton peak at 1.81 eV [110].

Fig. (4:30): Lorentzian peak fitting of MoS2 monolayer PL spectra showing neutral
exciton and trion peaks.
The position of the excitonic peaks can be affected by growth induced strain
in the film. As we explained in the previous sections there is non uniform strain
across the polycrystalline film.
The physical properties of MoS2 monolayers are affected by the underlying
substrate in the MoS2/substrate heterostructure. It has been reported that there is
a charge transfer from substrate to the monolayer and the amount of the charge
transfer depends on the type of substrate. For example, the trion peak for

84
MoS2/SiO2 is found to be higher than that for MoS2/STO, suggesting that charge
injection to the MoS2 is substrate dependent. The trion intensity is further reduced
when taking PL from freestanding monolayers [110], [111].
The presence of the trion peak in our films confirms the n-type doping of the
MoS2 by charge transfer effects from the SiO2 substrate and this result is in good
agreement with previous reports [109], [111], [112].
A PL quantum yield (QE) of 1% has been recorded for an as deposited MoS2
monolayer [113]. Such a low PL yield is mainly due to the presence of defects
especially sulphur defects which act as nonradiative recombination centers and
significantly quench the emission [114], [115]. This situation has been resolved by
post deposition treatments and the QE can be increased to near unity (i.e. 100%)
by treating monolayers with a non oxidizing organic superacid:
bis(trifluoromethane) sulfonimide (TFSI) [116]. The mechanisms by which the acid
passivates MoS2 defects is not quite clear. The propagation of the acid can remove
the surface contamination such as H2O and hydroxyl groups that adsorbed to the
surface from the environment. Such cleaning opens the active defect sites to
sulphur adatoms on the surface. The presence of sulphur adatoms and clusters on
the MoS2 surface has previously been confirmed by scanning tunneling
microscopy and the acid facilitates the reaction between the adatoms and active
defect sites. The increase in the S/M ratio in the monolayer after treatment is
confirmed using XPS indicating that the sulphur defects have filled [116].

4.4.2 Electrical properties

Figure 4:31 shows a schematic diagram of a field effect transistor FET used for
electrical characterization and figure 4:32 is the IV characteristics of the FET
transistor. The FET was fabricated as follows: the MoS2 sample was covered with
PMMA by spin coating at 4000 rpm, subsequent baking at 170 °C for 2 minutes
followed by electron-beam lithography. After that, the sample was developed in AR
600-55 solution (a mixture of methyl isobutyl ketone and Isopropyl alcohol in the
proportion of 1:3 according to the Raith company) for 2 minutes, and
photographically fixed for 20 seconds in IPA (isopropyl alcohol). 5 nm of Ti and 45
nm of Au were evaporated by electron beam deposition (AXXIS Kurt J. Lesker
company)with lift-off in acetone.The back gate FET devices were then fabricated.
We used a Keithley 4200 semiconductor analyzer to test the electrical
characteristics at room temperature. The IV characteristics of the device at

85
different Vds are shown in figure 4:32. We extract the device mobility using the
equation [7]:
dI L
μc = ( dV ds ) × ( W CV ) 4 : 10
bg ds

Here μc is the carrier mobility, L is the channel length, W is the channel

εO εr
width, C = d
is the per unit area capacitance of the oxide layer between the
back gate and the channel, ε0= 8.85x10-12 F/m is the dielectric constant of vacuum,
εr= 3.9 is the dielectric constant of SiO2, d=300 nm is the thickness of SiO2 [7].
Taking L=2 µm, W=26 µm in to equation (4:10), one obtains a mobility which is
about 0.1 cm2V-1s-1.

Fig. (4:31): Schematics of the FET used for electrical characterizations.

.

Fig. (4:32): Drain-source current Ids as a function of back-gate voltage Vbg at different
drain-source bias voltage Vds=0.1V,Vds=0.5 V and VDS=1 V.

86
The mobility of our films is comparable to the mobility extracted from mechanically
exfoliated films [112], [117] as well as CVD grown films [118]. The mobility of the as
grown monolayers is quite low and far below the theoretically estimated value of
410 cm2V-1s-1 [19]. There are several factors affecting the carrier mobility in MoS2
monolayers and other 2D materials. The main ones are phonon scattering,
Coulomb scattering due to the presence of impurities, and substrate roughness. In
addition, large contact resistance induced by the Schottky barrier at the
MoS2-metal interface could also be another factor [9], [119]. The effect of these
factors should be reduced to improve the carrier mobility to reach a level required
for practical realistic applications. Several attempts have been made to improve
the mobility of the as grown monolayers. A MoS2 FET is capped with a high
dielectric material like HfO2 as a backgate, increasing the mobility to 200 cm2V-1s-1.
Such an increase is primarily due to the dielectric screening of Coulomb scattering
of charged impurities. This figure of mobility is comparable to that of silicon thin
films and graphene nanoribbons. This advancement in the mobility, could allow
MoS2 monolayers to be the backbone in future electronics [7]. Using metal
contacts with low work function such as scandium, improves contacts with thin
MoS2 flakes, resulting in high carrier mobility compared to the counterpart metals
with higher work function [119]. However, to reach the theoretical value of the
mobility, intensive research on transistors is still needed.

4.5 Grain size distribution and nucleation density

Growing uniform and continuous 2D monolayer crystals with the largest possible
grain size is still one of the main challenges in the field of 2D materials. As the
physical properties of 2D materials are thickness dependent (as explained in
Chapter One), the uniformity of the grown film is crucial. Similarly, without film
continuity, making electronic devices on the film will be a difficult task. Additionally,
the presence of grain boundaries in the polycrystalline film affects its physical
properties such as a non uniform the PL yield and carrier mobility across the film
[11], [120].
Using SHG we have shown that our films are uniform and continuous and
now we discuss the grain size within the polycrystalline film. For this purpose, we
have employed polarization resolved SHG for mapping grain size, orientation and
boundaries. This technique has an angular resolution of of ±1o with measurable
angles from 0o to 30o due to the reflection symmetry in MoS2 [76]. Due to the loss

87
of phase information, the orientation (Mo-S) and (S-Mo) cannot be determined i.e.
the technique cannot distinguish between grains in opposite directions unless the
grains are grown next to each other [75], [76], [121]. For neighboring opposite
grains, the SHG generation interfere destructively at the grain boundary and the
boundary appears as a dark line and this will be shown later in this paragraph.
The experimental procedure for polarization resolved SHG is similar to that
explained in 4.3.4, with the addition of a linear polarizer for measuring X and Y
components of linearly polarized SHG. Figure 4:33 is a SHG image of the film
under investigation. Figures 4:34(a) and 4:34(b) are the corresponding X-polarized
and Y-polarized SHG images respectively. The orientation map of the grain sizes
obtained using a Matlab script (Appendix A2) and equation (3:14) can be mapped
over all intensities as shown in figure 4:35.
To evaluate the grain size distribution, we have considered grains to have
elliptical shapes and the area of grains is calculated on this basis. The gains within
continuous film have irregular shapes as shown in figure 4:35 and the ellipse is the
closest shape.

Fig. (4:33): SHG image of MoS2 monolayer film grown at 800 oC.

88
 Fig. (4:34): (a) X-polarized SHG image of the film shown in fig. 4.33 (b) Y-polarized
SHG image of the film shown in fig. 4.33.

Fig.(4:35): Orientation color map of the film shown in fig. (4:33).

Figure 4:36 is a typical SHG image of neighboring opposite grains labeled as

G1 and G2 with a dark line between them (destructive SHG interference at the
grain boundary). Figure 4:37 shows a polarization resolved SHG image of the
same area. As we can see the two grains have the same contrast and have not
been resolved as two different grains. However, the presence of the dark boundary
confirms that they are opposite grains. Taking data along the white arrow in figure
4:37, the intensity dramatically drops at boundaries as presented in figure 4:38.

89
Fig. (4:36): SHG image showing two Fig. (4:37): Polarization-resolved SHG
opposite grains labelled G1 and G2. image showing two opposite grains G1
and G2.

Fig. (4:38): SHG profile across the white arrow shown in fig. 4:37.

The grain size distribution versus frequency of occurring is shown in figure

4:39. The grain size ranges between 2600±515 μm2 and 79740±9670 μm2 with an
average of 13837±1614 μm2 and an average nucleation density of 82±9 mm-2.

90
 Fig. (4:39): Grain size distribution of a MoS2 polycrystalline film.

4.6 Conclusions
The establishment of a scalable synthesis approach for growing MoS2
monolayers is an important fundamental step towards the technology requirements
in the manufacturing world. However, despite the intense research efforts towards
controlled deposition of 2D MoS2 monolayers, wafer-scale synthesis with optimum
grain size, uniform and continuous films is still a challenging issue.
In this chapter we have introduced a scalable LPCVD approach using MoO2
and S as starting materials for growing wafer scale uniform MoS2 monolayers with
grain sizes up to 400 microns.
Compared with MoS2 monolayer growth using MOCVD which is only
achievable at low temperatures and therefore imposes restrictions on the the grain
size, our approach is reproducible in a range of growth temperatures and hence
our grain size at an optimum growth temperature 800 oC can be 40 times greater
than that achieved by MOCVD [47].
In the case of using MoO3, the substrate is always located downstream. This
means that variation in the concentration depends on the CVD system size and
the growth might not be reproducible in different systems. While in our case the
substrate is placed face down on the MoO2 container, and so there is no variation
in the MoO2 concentration in the case when using different systems. Also, the

91
MoO3 concentration is temperature dependent therefore the reproducibility might
not be achievable at different growth temperatures. In our case, the substrate and
the MoO2 are in the same heating zone and the growth proved to be reproducible
at different growth temperatures. Regarding the grain size, our grains are almost
600 times greater than those achieved using MoO3 [40].
Our films have high crystalline quality, as confirmed by TEM, GIIXRD, Raman
and PL measurements compared to the poor crystalline quality of MoS2
monolayers produced using MoCl5 [93].
We have used a range of different techniques to characterize our films. XPS
is used to confirm that a complete conversion to MoS2 occured as result of
reaction between MoO2 and S in the reaction zone. The AFM measurements
confirm the monolayer nature of the film, and this was further confirmed by Raman
and PL measurements.
A SHG microscopy investigation proved that the monolayers are uniform over
a range of growth temperatures. The crystalline quality of the films was studied
using TEM and GIIXRD and the results indicated high crystalline quality. Finally,
the carrier mobility was measured using a prototype FET and the mobility value
found to be in the range that has been extracted from CVD and exfoliated
monolayers.
Our GIIXRD results indicate that there is a native strain in the as-grown
monolayer films. This strain is growth induced strain as a result of the thermal
expansion mismatch between the monolayers and the underlying SiO2 substrate.
The value of the strain is found be in agreement with the ones reported in
literatures [104]. GIIXRD is further used to study the thermal expansion of the
monolayers by in situ annealing. The monolayers were found to be stable up to
800 oC under UHV conditions and the thermal expansion coefficient is comparable
with that reported for bulk MoS2 [107].
All in all, we have established an LPCVD approach for growing MoS2
monolayers on a wafer scale. The approach is proven to be scalable and
reproducible in a range of growth temperatures. The monolayers have
unprecedented size when grown at optimum temperature.

92
Chapter 5: Growth mechanism studies with the aid of
COMSOL

Abstract
In this chapter we will present a mechanistic study of the LPCVD approach used
for growing a reproducible MoS2 monolayer on a wafer scale. We report the
dependence of the MoS2 film growth as a function of growth conditions such as
MoO2 concentration, sulphur flux, growth temperature, sulphur chemical potential
and growth time. We will use COMSOL to simulate the concentration distribution at
the surface of the substrate with a view to understanding the reaction conditions at
the growth front.
The growth of continuous monolayer films was found to depend on the
vertical distance between the substrate and MoO2 powder. We also found that full
monolayer film coverage can be obtained when the MoO2 vapor pressure is equal
to the saturation vapor pressure. A sulphur flux of at least 7×10-6 mol/m2s is
required for growing continuous monolayer films.
Our results suggest that continuous reproducible monolayer films can be
grown at temperature range 650 oC - 850 oC with the maximum possible grain size
at temperatures 800 oC - 850 oC and a minimum nucleation density is obtained in
the same range of growth temperatures.
We further studied the growth rate of a single monolayer grain and we have
found that growth occurs in the reaction limited regime at 650 oC ≤ T< 800 oC, and
that the mass transport or feed limited regime happens at 800 oC ≤ T ≤ 900 oC. A
desorption limited regime occurs at 900 oC ≤ T ≤ 1000 oC. Beyond 1000 oC, rapid
desorption occurs before the reaction of the starting materials and no growth is
observed.
From a flush growth study (one minute growth) of the individual grains at
temperatures of 700 oC - 1000 oC, it is determined that the growth rate could be as
large as 710±260 μm2/s at a growth temperature of 900 oC. The optimum growth
temperature was found to be 800 oC. Beyond this growth temperature, the growth
of additional layers is observed. The optimum monolayer growth time at 800 oC is
15 minutes beyond which bilayers start to initiate.

93
 Finally the morphology of the grains can be tuned using different growth
temperatures: triangles with bent edges at a temperature T ≤ 700 oC to perfect
triangles at 800 oC ≤ T ≤ 900 oC and perfect hexagons at 900 oC ≤ T≤ 1000 oC.

5.1 Literature review

5.1.1 CVD process

In chemical vapor deposition, the film growth rate, nucleation density, grain size,
grain morphology, and thickness can be controlled via several parameters such as
the composition and chemistry of the precursors and their partial pressure in the
CVD system, the pressure and the design of the CVD system and the growth
temperature [122] .
To understand the effect of different parameters on the film growth, the model
shown in figure 5:1 (a substrate placed in the reaction zone of the CVD tube) is
proposed. In this model, the mass transport of the precursors through the
boundary layer depends on the mass diffusion. Therefore, there is a concentration
gradient of the precursors [122]. The flux (F1) of the precursor from the bulk gas
stream to the substrate can be written as follows [122]:
F 1 = hG (C 1 − C 2 ) 5:1
D
hG = δ 5:2
Assuming the special case when the chemical reaction is of first order:
F 2 = kS C 2 5:3
where F 1 and F 2 are the fluxes in stream and on the substrate respectively, hG and
kS are the mass transfer rate constant and the reaction rate constant respectively,
and C 1 and C 2 are the concentrations in the stream and on the substrate
respectively. D is the diffusion constant of the precursor and δ is the thickness of
the boundary layer.

94
 Fig. (5:1): Substrate in the reaction zone of the CVD system [122].

Under steady conditions, the flux F1 is equal to the the flux on the substrate F2.
Then the relationship between C 1 and C 2 obtained as :
C1
C2 = k 5:4
1+ h S
G

When discussing the effect of the growth temperature on a CVD process,

one can classify the growth process into three regimes: at relatively low
temperatures kS ≪ hG , the film growth is controlled by the kinetics of the
chemical reactions of the precursors. Even if there is an overabundance of the
reactants in the reaction zone, the growth is very slow. In this regime the growth
rate is weakly affected by the partial pressure of the precursors and is strongly
temperature dependent (growth rate ∝T3/2 ) [123].
At moderate temperatures kG ≪ kS , the growth is controlled by feed or mass
transport of the reactant. The growth rate can be approximated as follows:
mP
Growth rate = u0.5 T Ps
5:5

where u is the velocity of the carrier gas in the CVD tube, T is the growth
temperature, P is the partial pressure of the reactant and P s is the total pressure
of the CVD system.
In the case of atmospheric pressure chemical vapor deposition (APCVD)
[123], the growth rate depends on the diffusion of reactants through the boundary
layer to the substrate while in the case of low pressure chemical vapor deposition,
the boundary layer over the substrate is relatively thin and the growth is controlled
by feed of the reactants into the CVD system [123]. In mass transport or feed

95
controlled regime, the growth is weakly temperature dependent and the optimum
growth rate can be achieved [123].
At higher temperatures, the system is in the desorption regime. In this
regime, the desorption and reaction of the starting materials are competing and the
growth rate is reduced. Further increase in the temperature, a quick desorption of
the precursors occurs before the reaction and the film growth cannot be achieved
[123]. Each of the above mentioned growth regimes can be described by an
Arrhenius equation [124]:
−E a
Growth rate = Ae kT 5:6
where A is a constant, E a is the relevant activation energy and k is Boltzmann
constant. A typical Arrhenius plot for the growth rate displaying different growth
regimes is shown in figure 5:2.

Fig. (5:2): Typical Arrhenius plot, dependence of growth rate on temperature [123].

In this chapter, we are going to present our own Arrhenius plot analysis of the
LPCVD of MoS2 monolayers to demonstrate how we employed the effect of the
mentioned CVD parameters for establishing a standardized protocol for growing
MoS2 monolayers on a wafer scale.

5.1.2 Thermodynamics of CVD

The chemical reaction can be described by the Gibbs free energy of the reactants
and products [125]. Gas A reacting with B to form C, as shown in the following
equation [126]:
A(g) + B (g) ↔ 2C (g) 5:7

96
 At time t=0, A=1 mol, B=1 mol and C is absent. At any other time, t, A=1-x
mol, B=1-x mol and C=2x. The Gibbs free energy, G, of the system becomes
[125]:
G = (1 − x)μA + (1 − x)μB + 2xμC 5:8
where μA is the chemical potential of A, μB is the chemical potential of B and μC is
the chemical potential of C.
The derivative of G with respect to x can be positive, zero or negative. In the
case of a positive derivative, the reaction in eq. 5:7 would increase the Gibbs free
energy of the system which is not allowed. The product C is unstable and
decomposes to A and B.
dG
dx
= 2μC − μA − μB > 0 → μA + μB < 2μC 5:9

A zero derivative means the reaction is in equilibrium:

dG
dx
= 2μC − μA − μB = 0 → μA + μB = μC 5 : 10

and a negative derivative means the energy of the system would decrease and the
reaction would fe forward.
dG
dx
= 2μC − μA − μB < 0 → μA + μB > 2μC 5 : 11

For ideal gas behaviour, the chemical potential at a given temperature and
pressure can be defined as follows [125]:
μ(T , P ) = μ(T , P ref ) + kT ln( PP ) 5 : 12
ref

where T is temperature in K, P is the gas partial pressure, Pref is the atmospheric

pressure and k is the Boltzmann constant. In the case of MoS2, a more realistic
empirical formula exists and is used in section 5.6.4.
In the case of MoS2, the stable monolayer can be described as follows [127]:
μM o + 2μS = μM oS 2 5 : 13

The allowed values of the sulphur chemical potential are as follows [128]:
Hf
2
≤ μS ≤ 0 5 : 14

where Hf= 2.61 eV is the formation energy of MoS2. From equation 5:14, we can
see that large values of µS correspond to sulphur-rich conditions, whereas small
values of µS relate to Mo-rich conditions. The maximum chemical potential µS is
zero, corresponding to conditions at which sulphur condenses as bulk. The lower
limit for µS is −1.30 eV. Below this limit MoS2 monolayers are reduced to metallic
body-centered cubic (bcc) Mo [128], [129].

97
5.1.3 Monolayer growth mechanisms
The growth of thin films can be classified into three primary modes: Frank van der
Merwe, Volmer-Weber and Stranski–Krastanov modes [91]. The Volmer–Weber
mode is a layer-by-layer growth mode, the Frank van der Merwe mode is an island
growth mode and the Stranski–Krastanov mode is a layer-plus-island growth
mode. The growth of monolayer TMDs is believed to follow the Frank van der
Merwe or Stranski–Krastanov modes [91]. In the Frank van der Merwe mode, TMD
islands of different thickness stitch to form a complete thin film. In the
Stranski–Krastanov mode, monolayer TMD domains grow and interconnect with
each other until a complete coverage and then the second layer starts to grow in
the same manner [91]. Although, the TMDs monolayer growth is still a
controversial topic, all our observations on the MoS2 monolayers for different
growth temperatures support Stranski–Krastanov mode as shown by the SHG
images (previous chapter).
Now we focus on the growth mechanisms of MoS2 monolayers when using
MoO2 as Mo source and sulphur powder as sulphur source. At high temperatures,
MoO2 sublimes in the form of (MoO2)n and (MoO3)n clusters [130] and is then
deposited on the substrate that is placed face down on the MoO2 container. Under
a sulphur-rich environment, there is a probability that those clusters might partially
or fully sulfurise before being deposited. However there is no concrete evidence
for gas phase sulfurization yet [131]. On the substrate, the sulphur and Mo react
and the MoS2 monolayer nucleus initiates either from atomic MoS2 monolayer
clusters or unsaturated molybdenum oxysulphide (MoOxS2-y, y≥x) nanoparticles.
Both nucleation processes have been confirmed ex-situ through an electron
microscopy study by Zhu et al. (2017) in APCVD of MoS2 on graphene as shown
in figure (5:3) [131]. In a similar study for deposition on a 20 nm thick SiO2
membrane, JD Chain et al. (2016) [132] have found that a Mo(S/Se)2 alloyed
monolayer starts from MoOx(S/Se)2-y nanoparticles deposited on the substrate in a
sulphur poor environment, before the complete conversion to MoS2 in the optimal
sulphur atmosphere is achieved, resulting in the nucleation and growth of
monolayers at the nanoparticle sites [132].

98
Fig. (5:3): Two possible nucleation routes for growing MoS2 monolayers (a) MoS2
monolayer cluster (b) molybdenum oxysulphides (MoOxS2-y, y≥x) nanoparticles [131].

All the above mentioned work on the nucleation of MoS2 monolayers has
been done using APCVD. All the evidence presented in the literature on the MoS2
nucleation and growth mechanisms relied on post-growth studies. It is still not
clear if the formation of MoOxS2-y nanoparticles happens in the gas phase or on the
substrate, therefore we believe more in situ experiments based on techniques
such as mass spectroscopy or TEM are still required for finding the origin of the
nucleation.

5.2 Experimental investigation

5.2.1 Effect of gas flow on furnace temperature profile

The furnace temperature profile is one of the most important parameters that affect
the CVD growth of thin films. We have presented our measured data of
temperature profile of a closed-ended furnace in (Chapter Four). We expect the
furnace temperature profile also depends on the gas flow. Here we use COMSOL’s
heat transfer in fluids module for simulation of the temperature profile under
different flow conditions. For this purpose, we use typical growth conditions of a
set temperature of 800 oC (1073 K) and a CVD tube pressure of 10 mbar to see
the effect of varying the Ar flow rate (between 0 to 1000 SCCM) on the
temperature variations in both the sulphur and MoO2 zones. Figure 5:4 is a typical
simulated furnace temperature profile file for the growth temperature of 1073 K
and a flow rate of 100 SCCM.

99
 Fig. (5:4): Typical tube furnace surface temperature for a growth temperature of 1073 K
and a flow rate of 100 SCCM.

The temperature profile as a function of flow rate in the main (MoO2) zone
and sulphur zone at a growth temperature 1073 K and different flow rates is shown
in figures 5:5 and 5:6 respectively.

Fig. (5:5): Main (MoO2) zone temperature profile at a growth temperature of 1073 K
under different flow rates.

100
 Fig. (5:6): Sulphur zone temperature profile at a growth temperature of 1073 K under
different flow rates.

A typical surface temperature profiles of MoO2 and sulphur zones at growth

temperature 1073 K and flow rate of 200 SCCM are shown figures 5:7 and 5:8
respectively. The data shown in figure 5:5 and figure 5:6 is taken along the cutline
shown in figure 5:9.

Fig. (5:7): Cross sectional temperature Fig. (5:8): Cross sectional temperature
profile of the MoO2 zone at a growth profile of the sulphur zone at a growth
temperature of 1073 K under 200 SCCM. temperature of 1073 K under 200 SCCM.

101
 Fig. (5:9): Cutline along which the data is taken.

We have seen that there is little change in the reaction zone temperature
under different flow rates while the sulphur zone temperature fall sharply when
going to higher flow rates especially if the sulphur boat is placed somewhere close
to the center of furnace tube. In experimental work this should be taken into
consideration as the sulphur vapor pressure is very sensitive to the temperature
changes. The sulphur vapor pressure as a function of temperature will be
discussed later on in this chapter.

5.2.2 Starting material vapor pressure and concentration

The vapor pressure of molybdenum dioxide in the range (1620-1860) K is given in
the following empirical correlation [133].
17873
log P =− T
+ 6.035 5 : 15

where P is the vapor pressure in (atm.) and T is temperature in K.

As there are no reliable measured data in the range of our growth
temperatures, we extrapolated Eq. (5:15) for finding the MoO2 vapor pressure.
Sulphur: The vapor pressure of sulphur is given by the following correlation (389 K
< T < 1313 K) [134].
T 3/2 T 3 T 6 T
log PP = [A(1 − T
Tc
) + B (1 − Tc
) + C (1 − Tc
) + D(1 − Tc
) ]( T ) 5 : 16
c c

where 𝑇 is the temperature of the sulphur precursor in K, A = -7.246, B = 0.187, C

= 5.271, D = -12.128, 𝑇c = 1313 K, and Pc = 18208 kPa.
Assuming the validity of the ideal-gas law for a dilute vapor, the concentration
is given by following equation:
P (T )
C (T ) = RT 5 : 17

102
where C(T) is the concentration as a function of temperature and R is the universal
gas constant.

5.3 Experimental results

In this section, we present our results on the effect of vertical distance between the
MoO2 powder and the substrate on the monolayer coverage. For this purpose we
adopt the model presented in figure 5:10. The CVD reactor setup we used for
growing uniform MoS2 monolayer on a centimeter scale is a variation from those
used by other groups. In CVD, for a uniform film growth, the substrate must be
exposed to a uniform flux of the reactant materials. One way to do this is to place
the substrate face down on the MoO2 container in a place where where the Mo flux
is uniform. By changing the MoO2-substrate vertical distance, the MoO2
concentration can be tuned such that the optimum condition for full coverage of
monolayers is obtained. To cover all MoO2-substrate distances in one go, a simple
and clever method is to place the substrate at a certain angle with respect to the
MoO2 powder.
The procedure is as follows: 40 mg of high purity MoO2 powder (99% Sigma
Aldrich) was placed in a quartz boat to act as a Mo source. SiO2 covered Si wafers
(1×1.5 cm2) were cleaned with acetone and ultrasonicated for 10 minutes and then
positioned at certain angles 15o, 20o, 28o, 48o and 66o on a MoO2 container with a
typical example shown in figure 5:10. The MoO2 container with the substrates was
then placed at the center of the 75 mm diameter tube furnace. Another boat
contained 5-20 mg of Sulphur powder (99.98% Sigma Aldrich) was placed, 42 cm
away from MoO2, in a region where the temperature reaches 200 oC. Prior to the
growth, the furnace was flushed with argon gas (1000 SCCM) for about 30
minutes. Then the furnace was pumped down to 10 mbar and the center was
heated to 800 oC with a heating rate of 15 oC /minute in an argon flow (200
SCCM). After keeping the temperature at 600-1000 oC for 10-20 minutes, the
furnace was naturally cooled down to room temperature.

103
 Fig. (5:10): Typical position of the substrates with respect to MoO2 powder in the
reaction boat.

5.3.1 The effect of the vertical distance between substrate and Mo

source on the monolayer film coverage
Figure 5:11 shows the monolayer coverage as a function of substrate vertical
distance for five different different tilt angles.
We found that the substrate is always fully covered by a continuous uniform
monolayer when the MoO2-substrate distance is less than 5 mm. Above 5 mm, the
substrate is found to be partially covered with monolayers. Based on our findings,
at vertical distances of 0-5 mm, our system is in a steady state with respect to Mo
for the growing MoS2 monolayers. Beyond 5 mm, the system is in the feed-limited
regime. Figure 5:12 is an optical image of monolayer coverage on a SiO2/Si
substrate placed at 48o. The isolated grains and continuous film are shown in
optical micrographs 5:13 and 5:14 respectively.

104
 Fig. (5:11): MoS2 monolayer coverage as a function of MoO2-substrate distance.

Fig. (5:12): Photograph of a continuous monolayer film (right part) and bare
substrate (left part) of a typical sample grown at an angle 48o.

Fig. (5:13): Isolated grains from left side Fig. (5:14): Continuous monolayer film from
of sample shown in fig. 5:12. right side of sample shown in fig. 5:12.

To have a better understanding of the growth mechanism we use a

COMSOL software to find out the concentration profile of the reactant materials in

105
the furnace tube. Figure 5:15 displays the results of the simulation of MoO2
concentration as a function of the MoO2-substrate vertical distance.

Fig. (5:15): MoO2 concentration as function of MoO2-substrate vertical distance at a

growth temperature 1073 K and an Ar flow rate 200 of SCCM.

Figure 5:16 shows a cross-sectional MoO2 concentration profile inside the reaction
boat under the growth conditions of 1073 K and 200 SCCM flow rate.

Fig. (5:16): Cross-section of MoO2 concentration profile inside the reaction boat at a
growth temperature 1073 K and an Ar flow rate of 200 SCCM.

106
 We have presented the results at a typical growth temperature of 800 oC. We
used the same design for growth temperatures between 650-1000 oC. The growth
proved to be reproducible for full coverage uniform monolayers up to a 5 mm
substrate vertical distance in the temperature window 650-850 oC. The results for
different growth temperature will be presented later on in this chapter.
The sulphur concentration profile along the furnace tube in the ZX-plane is
shown in figure 5:17. A data profile taken along cutline displayed in figure 5:17 is
shown in figure 5:18. As one can see, the concentration profile becomes fully
developed 10 cm away from the sulphur boat centered 23 cm from the inlet. There
is a small but non negligible amount of sulphur diffusing back toward the inlet.

Fig. (5:17): Upper: ZX-plane, sulphur concentration profile at a growth temperature

1073 K and a flow rate of 200 SCCM. Below: cutline (red) along the tube.

Fig.(5:18): Sulphur concentration profile taken along the cutline shown in fig. (5:17).

107
 The velocity profile of the Ar carrier gas in the reaction boat is shown in
figure 5:19 (left). The data taken along the red cutline in figure 5:19 (right) is shown
in figure 5:20.

Fig. (5:19): Left: Ar velocity profile in the XY-plane at the center of the furnace tube.
Right: red cutline.

Fig. (5:20): Ar velocity profile taken along the cutline shown in figure 5:19 (right).

The average carrier velocity inside the reaction boat is 0.1 cm/s and the width of
the substrate is 2 cm. The sulphur and MoO2 diffusion coefficients under our
growth conditions are 168 cm2/s and 92 cm2/s respectively .
We can quantify the concentration in terms of a dimensionless number called
the Péclet number (Pe), which is the ratio of the contributions to mass transport by
convection to those by diffusion [135]:
Lr U
Pe = 4D
5 : 18

108
where Lr is the length of reactor, U is the mean carrier velocity and D is the
diffusion constant.
In our case the Pe value for sulphur is 3×10-4 and for MoO2 is 5.4×10-4. If Pe
˂˂1, the diffusion is dominant and the concentration is uniform, while for Pe >>1,
there is a gradient in the species concentration [135]. For our model we have seen
that Pe is <<1 and the concentration of sulphur and MoO2 are therefore uniform.
These values of Pe are as expected since we are using LPCVD where the starting
materials have high diffusion constants.

5.3.2 The effect of sulphur flux on the coverage of monolayers

One of the ways to change the sulphur concentration and flux in the growth reactor
is by changing the carrier gas flow rate. Depending on the size of the reactor, the
flow rate can be tuned to get the optimum condition for the growth. A set of
experiments have been done using the same model explained in section 5:3. In
those experiments the growth temperature was set to be 800 oC, the sulphur was
placed at a region where the temperature reaches 200 oC, the growth time was 15
minutes and gas flow was 10, 20, 30, 40, 50 and 200 SCCM. In such conditions,
the growth rate is expected to be a function of gas flow rate only. The experimental
data of monolayer film coverage as a function of the gas flow rate is shown in
figure 5:21.

Fig. (5:21): Monolayer coverage as a function of Ar flow rate for a growth temperature
1073 K.

109
The sulphur concentration and flux is proportional to the carrier gas velocity inside
the furnace tube which is in turn related to the tube cross section and the gas flow
rate. It is better to present the monolayer coverage as a function of sulphur flux.
The left subplot of figure 5:22 shows COMSOL simulations of typical sulphur
convective flux at 200 SCCM at growth temperature of 1073 K. The right subplot of
figure 5:22 shows the red cutline along which data is taken for different flow rates
and displayed in figure 5:23.

Fig. (5:22): Left: typical sulphur flux at the inlet of the growth region at a growth
temperature of 1073 K and a flow rate 200 SCCM. Right: cutline along which data for
different flow rates is taken.

Fig. (5:23): Sulphur convective flux under different flow rates, and at a growth
temperature 1073 K. Data is taken along the red cutline shown in fig 5:22,right.

110
 The sulphur convective flux at the reaction zone inlet is magnified in the
following figure (5:24)

Fig. (5:24): Close-up of the sulphur flux profile at the reaction zone as a function of Ar
flow rate.

We also extracted the average sulphur convective flux at the inlet of the
reaction zone as displayed in figure 5:25. As we can see, the film is partially
covered up to a sulphur flux of 7×10-6 mol/m2.s indicating that the growth is in a
feed-limited regime with respect to sulphur. Above 7×10-6 mol/m2.s full coverage
monolayers is obtained and the growth became in steady state regime.

Fig (5:25): Average sulphur flux entering the reaction zone as a function of Ar flow rate.

The sulphur concentration profile as a function of flow rate at the center of

the tube furnace is shown in figure 5:26.

111
 Fig. (5:26): Sulphur concentration profile as a function of Ar flow rate. Data is taken
along the red cutline shown in fig. 5:22.

In this section, we studied the effect of the sulphur concentration and flux on
the monolayer film growth rate. We correlate both concentration and flux to the
carrier gas flow rate using COMSOL (transport of diluted species). Based on our
experimental findings a sulphur flux of at least 7×10-6 mol/m2.s is required to grow
full coverage monolayer films. Below 7×10-6 mol/m2.s, the system was in (sulphur)
feed-limited regime, i.e. the rate of providing sulphur to the reaction zone was
much lower than the rate of sulphur consumption by the reaction at the growth
front. From 7×10-6 mol/m2.s up to 10-5 mol/m2.s, a full coverage monolayer is
obtained and the system is expected to be in a steady state.

5.3.3 Grain size and nucleation density temperature dependence

We have studied the effect of temperature on the grain size distribution and
nucleation density using the same experimental procedure as explained in section
5.3. The results presented here are for temperatures of 650-850 oC that result in
the growth of reproducible continuous monolayer films. The temperature
dependent nucleation density and grain size distribution were carefully
investigated using SHG microscopy. The reproducibility of the film growth is
confirmed by several growth runs. Figures 5:27-5:31 show colored orientation
maps of continuous monolayer films grown at temperatures between 850 oC and
650 oC respectively.

112
Fig. (5:27): Colored orientation map of a Fig. (5:28): Colored orientation map of a
polycrystalline film grown at 850 oC. polycrystalline film grown at 800 oC.

Fig. (5:30): Colored orientation map of a

Fig. (5:29): Colored orientation map of a
polycrystalline film grown at 700 oC.
polycrystalline film grown at 750 oC.

Fig. (5:31): Colored orientation map of a polycrystalline film grown at 650 oC.

113
 The statistics of the grain size distribution as a function of temperature
between 650 oC and 850 oC are shown in figures 5:32-5:36 respectively. They all
appear to follow a well-defined log-normal distribution. The average nucleation
density as a function of temperature is shown in figure 5:37.
At a growth temperature of 850 oC, the grain size ranges from 3030±780 μm2
to 60560±10350 μm2 with the peak centered at 24669±12235 μm2 and the average
nucleation density of 50±7 mm-2. The figures for 800 oC are very close to that of
850 oC. At 800 oC the maximum size of gain can be as large as 79740±9670 μm2
and minimum size 2600±515 μm2 and the most probable size of 13837±1613 μm2
and the nucleation density slightly increased to 82±9 mm-2. A sharp change
happens at the growth temperature of 750 oC, the largest grain that can be grown
at this temperature 30930±6400 μm2 is less than the half size of grain grown 800
o
C. The average grain size at 750 oC was found to be 5974±706 μm2 and the
nucleation density is 180±13 mm-2. A more steep change occurs when the
temperature is reduced to 700 oC. At this point the grain size ranges from
4980±800 μm2 to 500±300 μm2 with an average of 1000±605 μm2 and nucleation
density is up to 3271±60 mm-2. At 650 oC the mean grain size reduces dramatically
to only 30±2 μm2 with the average nucleation density increased to 25794±160
mm-2.

Fig. (5:32): Grain size distribution at a growth temperature of 850 oC.

114
Fig. (5:33): Grain size distribution at a growth temperature of 800 oC.

Fig. (5:34): Grain size distribution at a growth temperature of 750 oC.

Fig. (5:35): Grain size distribution at a growth temperature of 700 oC.

115
 Fig. (5:36): Grain size distribution at a growth temperature of 650 oC.

By fitting our measured data with an Arrhenius equation, as shown in figure

−3.26±0.23
5:37, we have found that that the nucleation density N αe kT where 3.26±0.23
eV is the nucleation activation energy for a growth temperature T≤ 750 oC and
−1.33±0.123
N αe kT where 1.33±0.12 eV is the nucleation activation for growth temperature
T≥800 oC.

Fig. (5:37): Nucleation density as a function of growth temperature with the inset fitted
with an Arrhenius equation.

116
 The growth time for full coverage monolayers is 15 minute for all growth
temperatures. Based on this fact the growth rate of the film as a whole is similar for
all growth temperatures while the growth rate for single grains within the
polycrystalline film is temperature dependent. At low growth temperatures T < 800
o
C the nucleation probability is high as shown in figure 5:37 and the growth rate of
single grains is slow. While at high temperatures T ≥ 800 oC, the nucleation is
suppressed and single grain growth rate is dramatically increased.
From the kinetic theory of gases we can easily estimate the time required to
form a monolayer on the target substrate. This monolayer formation time is
inversely related to the impingement rate (flux). From our experimental
investigations, we have found that the required time for full monolayer coverage is
15 minute for all growth temperatures. The lattice constant of the 1H MoS2 phase
is 3.16 Å and each primitive unit cell contains 1 Mo atom and 2 S atoms. This
means in the case of a perfect monolayer, there are 1019 Mo atoms/ m2 and 2×1019
S atoms/m2 respectively. The Mo flux incorporated in the crystal growth is given
by:
1019 ( atoms/m2 )
F lux = growth time (s) 5 : 19

The relation between the flux and vapor pressure is given as follow:
P
F lux = 5 : 20
√ 2πRT
M

Here, P is the partial pressure of Mo consumed in the reaction, R is the universal

gas constant, M is molar mass and T is temperature in Kelvin.
Using equations 5:19 and 5:20 together with the experimentally observed growth
time, we have found that in the steady state (when the sticking coefficient is
assumed to equal to one), the required vapor pressure of the Mo source for a
monolayer growth is 10-11 atmosphere. With the help of the vapor pressure data of
different precursors, our finding suggests that choosing the right processor can be
is the key for growing TMDs monolayers. The finding suggests that there should
be a proper designing of CVD reactors for growing TMDs monolayers when having
a precursor with high vapor pressure.
Equation 5:15 shows that the MoO2 vapor pressure is ≥ 10-11 atmosphere at
growth temperatures T ≥ 800 oC. Such a low vapor pressure is essential for
growing monolayers on large scales. However, for growth temperatures T < 800

117
o
C, equation 5:15 gives a vapor pressure lower that 10-11 atmosphere, which
contradicts our experimental results indicating that extrapolation of equation 5:15
is not correct for low temperatures.
As we have seen, MoO2 provides enough flux for growing MoS2 monolayers
at a range of growth temperatures when the substrate is placed in an appropriate
place such that the whole flux generated diffuses to the substrate. The full
coverage obtained at different growth temperatures in equal growth times confirms
our point. The MoO2 vapor pressure increases exponentially with temperature,
which means that the substrate receives more Mo at higher temperatures, but at
the same time the Mo sticking coefficient decreases with raising temperature.
Therefore there is more Mo desorption from the substrate at high temperatures
such that the net Mo-flux available at each temperature is critical for MoS2
monolayer growth. In the next chapter, we will use MoO3 as precursor, which has a
much higher vapor pressure than MoO2, to study film growth beyond such critical
MoO3 vapor pressures required for growing monolayers.

5.3.4 Temperature dependence of the initial monolayer growth rate

In this section we present a systematic study of the initial monolayer growth rate
temperature dependence. Arrhenius plots under a fixed reactant concentration can
be used for this purpose. In our case, we fixed the sulphur flux in the reaction zone
by keeping the sulphur boat at the same sublimation temperature of 200 oC for all
growth temperatures (650 oC - 1000 oC). Figure 5:38 is the average sulphur flux in
the reaction zone. As we can see, the flux is slightly higher for higher temperatures
due to a slightly higher carrier gas velocity at high temperatures.

Fig. (5:38): Sulphur flux in the reaction zone at different growth temperatures.

118
 The Mo flux is also proportional to the growth temperature via equation
(5:15). However, we do not have control over it as we have control for sulphur
because the MoO2 container and the substrate are positioned in the same
temperature zone. We assume that the rates of high Mo flux at high temperatures
are compensated by the high rates of Mo desorption from the substrate. We base
our assumption on the fact that we can get a full coverage of monolayers during a
15 minute growth run for all growth temperatures as we presented in the previous
section. Therefore the amount of Mo involved in the reaction are equal for all
growth temperatures.
For comparison of data from different growth temperatures, we used the size
of the maximum possible grains observed at each growth temperatures. Figure
5:39 is the corresponding Arrhenius plot of the grain size as a function of growth
temperature.
As we can see the growth rate follows the three classical temperature
regions. At 650 oC ≤ T< 800 oC, the grain growth starts from a low level but
increase sharply with temperature. This is considered to be in the surface kinetics
limited regime with an activation energy of -2.6±0.4 eV. Although there is plenty of
sulphur in the reaction zone (as shown in figure 5:38), the grain growth (rate) is
limited by the kinetics of the chemical reactions at the surface which is strongly
temperature dependent. In this regime, the supply of the reactant materials to the
substrate is higher than the rate consumed by the reaction.
At 800 oC ≤ T ≤ 900 oC, the grain growth seems to have plateaued and the
system is in a mass transport or feed controlled regime. In this regime the growth
rate strongly depends on feed supply of the starting materials and weakly on the
growth temperature.
At 900 oC ≤ T ≤ 1000oC the growth is in desorption regime. As a result of
high desorption of the reactants, the growth rate reduces. Beyond 1000 oC, there
is a rapid desorption of the precursors from the substrate before the reaction
occurs and the growth was not observable.

119
 Fig. (5:39): Arrhenius plot of MoS2 monolayer growth for 15 minutes.

5.3.5 One minute growth (flushing growth)

We further studied the growth of single grain monolayer by reducing the growth
time to only one minute (flash growth). For this purpose we further increased the
sulphur flux to make sure the system in a rich sulphur environment and the growth
is temperature dependent. The sulphur flux for one minute growth at different
temperatures is shown in figure 5:40.

Fig. (5:40): Average sulphur flux used in the one minute growth.

120
 When using the maximum grains found on the substrate, their growth rate is
found to be 4 ±1 μm2/s, 326±53 μm2/s, 340±15 μm2/s, 710±260 μm2/s, 560±137
μm2/s and 110±21 μm2/s at growth temperatures of 700 oC, 800 oC, 850 oC, 900
o
C, 950 oC and 1000 oC respectively. The arrhenius plot for flush growth is shown
in figure 5:41. Again we can see that the monolayers are at a reaction limited
regime at T<800 oC, a transport limited regime at 800 oC ≤ T ≤ 900 oC and a
deposition limited regime at T ≥ 950 oC. These results confirm our classification for
growth regimes presented in the previous section that is based on the grain size.
The activation energy for the reaction limited regime is -4 eV much larger than -0.4
eV for the mass transport limited regime. This is as expected as the reaction
limited regime is strongly temperature dependent while the transport regime is
weakly temperature dependent.

Fig. (5:41): Arrhenius plot for one minute growth.

Figure 5:42 a-e show grey style images of monolayers grown in one minute
at growth temperatures of 700 oC - 1000 oC respectively.

121
Fig. (5:42 a-f ): Are the grey style images of monolayers grown in one minute at
different growth temperatures between 700 oC and 1000 oC respectively.

122
 From figure 5:42, we also can observe that the significant growth of
additional layers increases at growth temperatures T ≥ 850 oC even for such a
short growth run, as shown in the white circles (fig. 5:42 c-f). For growth of
monolayers free from adlayers, the optimum growth temperature would therefore
be 800 oC.

5.3.6 The effect of sulphur chemical potential on grain morphology

The precise control of the edge geometry and crystal morphology of monolayer
two dimensional materials is of particular interest for their application as catalysts
in hydrodesulfurization of fossil fuels [136]. In the case of MoS2 monolayers, edge

structures have two different kinds, the ( 1010 ) Mo zigzag edge and the ( 1010 ) S
zigzag edge (see Fig.1:4 in Chapter One). The Mo zigzag edge is found to be
catalytically active while the S zigzag equilibrium edge structure itself is not active,
and sulphur vacancies are needed for the reaction to proceed [137].
The equilibrium shape of isolated MoS2 monolayer grains may be very
different, depending on the experimental conditions of the deposition [138]. The
most common parameters that have been used for determining the morphology of
MoS2 monolayers in CVD growth are temperature [139] and concentration of the
starting materials [138]. However, precise control of the morphology has not been
achieved yet and different shapes coexist in a single experimental run.
Wang et al. (2014) [138] have found that different morphologies can coexist
in a single experimental run. They claimed that the growth of such patterns might
be due to the local variations in Mo/S ratio on the substrate. However, they did not
present a quantitative figure of the Mo/S ratio that can lead to the growth a single
type of morphology.
To better control the morphology of MoS2 monolayers, Xie et al. (2016) [139]
used the evaporation temperature of the starting materials (MoO3 and S) to
change Mo/S in the reaction zone. They found that the MoS2 domains are mainly
round, nearly round and hexagonal, and truncated triangles grew at 760 °C, 750
°C and 730 °C respectively. However, they have not presented a clear explanation
of the mechanisms of growth. In a similar work by Yang et al. (2017) [140], they
found that at low temperature MoS2 monolayers have a triangular shape. By
increasing the temperature to 850–950 °C, hexagonal MoS2 flakes are observed. A
further increase in temperature to over 950 °C, MoS2 rectangle shape monolayers
are produced. Although, both of the mentioned works use the same type of starting

123
materials (MoO3 and S), they have found different morphologies under different
growth temperatures. Therefore, we believe further understanding of the growth
conditions and growth mechanisms is still needed to obtain control over the
reproducibility of MoS2 monolayer morphology.
Here, we used the chemical potential of sulphur to finely tune the morphology
of MoS2 monolayers. We have produced bent triangles, perfect triangles and
hexagons in a reproducible manner. Figure 5:43(a-d) shows optical micrographs
of the MoS2 monolayer grains with different morphologies grown at temperatures
of 700 oC - 1000 oC respectively. The growth process of the MoS2 monolayers can
be can be summarized as follows: (1) the sublimation of sulphur powder at 200 oC
and its transport by the carrier gas to the substrate; (2) the sublimation of MoO2
powder at the growth temperatures and its diffusion to the substrate as the
substrate is directly faced down on the MoO2 source; (3) adsorption of the source
materials (S and MoO2) on the substrate (4) diffusion and reaction of the source
materials on the substrate that lead to the nucleation of a MoS2 monolayer; (5)
desorption of the reaction byproducts to the carrier gas stream.
Based on microscopic studies, in the early stages of growth the MoS2
monolayer starts as an irregular polygon with Mo and S terminations as depicted in
figure 5:44 [131]. According to crystal growth theory [141], the different growth
rates of the different crystal facets determine the final morphology of the crystal.
The growth rate of a crystal face depends on the face free energy [141]. In the

case of MoS2 facets ( 1010 ) and ( 1010 ) which are known as Mo-terminating and
S-terminating, zigzag edges are the most commonly observed and energetically
stable edges [11].

124
 Fig. (5:43): Shape evolution of MoS2 monolayers grown between 700 oC and 1000 oC.

Fig. (5:44): Depiction of MoS2 monolayer nucleus in the early growth stage showing Mo
and S zigzags.

The growth rates of advancing ( 1010 ) and ( 1010 ) facets depend on the Mo
and S concentrations in the reaction zone. In the case of a S-rich environment the

( 1010 ) facets advance faster and disappear and the final shape of the crystal will
be bent or perfect triangle as shown in figure 5:43(a) and 5:43(b) respectively. By

decreasing the sulphur concentration, the two faces ( 1010 ) and ( 1010 ) advance

125
equally and the final morphology will be a hexagon as shown in figure 5:43 c and
5:43 d respectively.
Using COMSOL simulation we have found that the average sulphur
concentration in the reaction zone is 0.016 mol/m3-0.0157 mol/m3 at growth
temperatures between 700 oC - 1000 oC. This means we are growing monolayers
under the same sulphur environment. Based on this data the grown monolayers
should have a triangular shape at all growth temperatures. However, observation
of hexagonal shaped monolayer crystals at growth temperatures above 900 oC
means that the system is in sulphur poor environment.
To understand this we correlated the MoS2 morphology to the chemical
potential of the sulphur in the reaction zone. We have used the correlation given
by Jackson et al. (2016) [142] to calculate the sulphur chemical potential as a
function of growth temperature and sulphur vapor pressure:
T −T tr μS8 T −T tr μS2 (T −(T tr −b)))2
μS (T , P ) = 21 (erf c( w
)) 8 + 21 (erf ( w
) + 1) 2 − aexp(− 2c2
) 5 : 21

where Ttr and a are represented by the polynomials in log10(P/Pa).

T tr = ∑ xT tr,i (log 10 ( PPa ))i 5 : 22

i

a = ao + a1 log 10 ( PPa ) + a2 (log 10 ( PPa ))2 5 : 23

and µS8, µS2 are pressure-corrected polynomials in T:

i
μS8 = ∑ xS8,i T + kT ln PPref 5 : 24
i

i
μS2 = ∑ xS2,i T + kT ln PPref 5 : 25
i

where Pref is 105 Pa. P is equal to the pressure in Pascals and k = 8.617×10-5 eV
K−1 is the Boltzmann constant. The values of all other parameters are given in
Appendix (A3).
Figure 5:45 shows the sulphur chemical potential as a function of sulphur
vapor pressure and growth temperature. The sulphur vapor pressure was
calculated using COMSOL. We have found that for µS=-0.834 eV , the bent,
triangles are grown, for µS=-0.995 eV, perfect crystals are observed and for µS from
-1.156 eV to -1.318 eV, hexagonal monolayers are grown. Theoretically, the
growth of hexagons is expected to happen at µS=-0.98 eV up to -0.04 eV, and the
growth of triangles starts at µS=-0.04 up to 0 eV [143].

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 We have used the temperature parameter to change the chemical potential
in our system. We have proved that the system could be switched from
sulphur-rich to Mo-rich environment. As a consequence of this, we have a control
over the morphology of the grown MoS2 monolayers.

Fig. (5:45): Sulphur chemical potential as a function of growth temperature and vapor
pressure indicating the morphology of MoS2 monolayers at each growth temperature
range.

5.3.7 The effect of the growth time on the film uniformity

As we have seen in the previous sections, our model is reproducible for growing
full coverage monolayers under a range of growth temperatures when using typical
growth time of 15 minutes. In this section we present our results when the growth
time is beyond that for a full coverage monolayers (i.e. beyond 15 minute growth).
For this purpose, we employ typical growth conditions: growth temperature 800 oC
and Ar flow rate 200 SCCM.
Figure 5:46 a-b show optical micrographs of uniform MoS2 monolayers grown
at a typical growth growth time of 15 minutes and 20 minutes respectively.

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 Fig. (5:46): Optical images of (a) continuous MoS2 monolayer film grown at 15 min., (b)
continuous monolayer film grown at 20 min. The initial stages of bilayer growth is shown
in the dark circles. (c) bare 300 nm SiO2/Si substrate.

As we can see from figure 5:46a, the film is continuous and there is no sign
of adlayers or holes in the film. Therefore the growth time of 15 minutes can be
considered as the optimized time for growing monolayers in our approach. This
observation confirms that the monolayer growth is self-limiting to an extent.
Further increasing the growth time to 20 minutes, we observed the growth of
bilayers as shown in figure 5:46b. These bilayers have not fully covered the
surface at this stage. The small triangular grains in the dark circles are the initial
stage of the bilayer growing. The nucleation of the bilayers on the continuous
background of monolayers confirms the layer-by-layer growth of our films. We did
not continue to grow the full coverage bilayers because such growth is beyond the
scope of this project. However, we can say that the growth time parameter can be
used for further bilayer growth investigations.
In the long-duration growth runs, the sulphur supply must be guaranteed,
otherwise the growth of MoO2 adlayers which is usually grown in a sulphur poor
environment is expected. Figure 5:47(a) is an optical micrograph of a
diamond-like MoO2 grain which is grown in a 30 minute run when the sulphur feed
stoke is depleted. Our Raman measurement shown in figure 5:47(b) confirms that
the diamond-like grains are MoO2 and agree with the Raman spectra shown in the
literature as in figure 5:47(c) [144].
We have shown that when using MoO2 as the Mo source, a growth time of 15
minutes is ideal for full coverage monolayers. Beyond 15 minutes, the growth of a

128
bilayer is initiated in a sulphur rich environment. If the sulphur is depleted, the
growth of MoO2 adlayers is observed.

Fig. (5:47): Optical micrograph of a MoO2 grain (b) the corresponding Raman spectrum
and (c) Raman spectrum for MoO2 from the literature [144].

5.4 Discussion and conclusion

In chemical vapor deposition, the film growth phenomena can be controlled
through several parameters such as the nature of the starting materials, transport
mechanisms in the reactor, the reaction temperature, and the reactor design. We
will discuss the effect of all the mentioned parameters one by one and explain how
we employed them in our CVD approach. Based on the mentioned parameters, we
established an unprecedented LPCVD approach for growing uniform wafer scale
MoS2 monolayers on a SiO2/Si substrate. COMSOL simulations are essential to
study quantitatively the effect of different parameters on the growth.

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 Firstly, we presented our data of furnace temperature profile in chapter four.
This data was taken under low Ar pressure with no flow conditions. Here, we
employed COMSOL (heat transfer in fluids module) to study the effect of the Ar
gas flow rate on the temperature profile of the tube furnace to see whether the flow
has an effect on the reaction zone temperature and the sulphur zone temperature.
For this purpose, we chose a growth temperature of 800 oC and flow rates of
(100-1000) SCCM. We have found that the average temperature of the reaction
zone (the place where the Mo source and substrate are placed) is only slightly
reduced by the flow rate (up to one degree), even at the highest flow rate of 1000
SCCM. Such a negligible variation of the reaction zone temperature is expected as
we use LPCVD where the density of Ar is very low. Also, the center of the reaction
zone of our system is 68 cm away from the inlet and the Ar gas gains energy from
the hot wall of the tube furnace as it passes such a long distance to the center.
Therefore, we can say that neither the evaporation rate of MoO2, nor the reaction
rates could be affected due to the temperature variations resulting from the flow
rate. However, in the sulphur zone, the situation is slightly different. The average
temperature could reduce (about 10 degrees) at the highest flow rate of 1000
SCCM. Although, we never used such high flow rates in our experimental work,
such changes in the temperature should be taken in to account when using low
growth temperatures and high flow rates.
We have introduced an experimental set-up that predominantly resulted in
the growth of MoS2 polycrystalline monolayer films under different growth
temperatures. This achievement is realized through both employing the low vapor
pressure of MoO2 as Mo source and the peculiar geometry of the reactor. As we
designed our set-up such that the substrate is placed in close proximity of the Mo
source, in this case the vapor pressure of the Mo source does matter. To
understand how the vapor pressures of the Mo sources are important, we
investigated the time needed for growing a MoS2 monolayer.
Three different candidates are available to be used as Mo source
molybdenum carbonyls such as Mo(CO)6, molybdenum salts such MoCl5;
molybdenum oxides such as MoO3 and MoO2. Mo(CO)6, and MoCl5 ; have melting
points of 150 oC and 194 oC and they can not be directly used in our approach
because they will melt away before the system reaches growth temperature which

130
is typically 800 oC. Regarding the molybdenum oxides, MoO3 has a much higher
vapor pressure than MoO2 [133].
Using MoO3, the rate of providing MoO3 to the substrate will always be higher
than the rate of consumption in the reaction leading to the formation of pure MoS2.
As result of this, the yield on the substrate will be a mixture of nonuniform MoS2,
and molybdenum suboxides. This means that the system is in the reaction limited
regime with regard to the MoO3. In this regard there are several works that have
used MoO3 in this manner that lead to the growth of MoS2 and Mo suboxides [138],
[145], [146]. We also provide our results about using MoO3 in the following chapter.
Now, regarding MoO2, our experimental results confirmed that as a result of
low vapor pressure, the rate of providing Mo to the substrate is nearly equal to the
rate of consumption at certain heights above the Mo source. By quantifying the
monolayer coverage as a function of the vertical distance between the substrate
and the Mo source, we have found that the complete coverage of monolayers can
only be achieved at certain heights. We used this height to guide us to place our
substrate horizontally on the Mo source and we have grown wafer scale uniform
monolayers as presented in chapter four.
Another key success in our approach is the simple geometry that we have
adopted. The position of the substrate relative to the Mo source is the key factor
for delivering a uniform Mo flux to the substrate. In the CVD tube, there is an
infinite number of ways to place the substrate. In our approach, we placed the
substrate at a certain angle to the Mo source for two reasons. Firstly, during the
sublimation the Mo directly diffuses to the substrate, as there are no surfaces in
between for Mo to deposit on. In this way, we avoided any reduction in the Mo flux
and guarantee a uniform Mo concentration profile to the substrate. Secondly, on
such an inclined substrate, we can study the growth as a function of Mo
concentration covering the concentration variation as a function of vertical distance
between the substrate and Mo source in one experimental run.
The CVD growth of MoS2 is a multicomponent CVD and the presence of
sulphur is as important as Mo. The sulphur flux in the reaction zone is another key
factor through which we can control the MoS2 film growth as it determines the rate
at which we provide sulphur to the reaction. To evaluate the role of sulphur flux, we
studied the effect of the flux on the monolayer coverage in a series of experiments.
There are two methods to increase the sulphur flux that reaches the reaction zone.

131
The first one is to increase the temperature of the sulphur container that would
increase the vapor pressure of the sulphur and the second one is to increase the
flow rate of the carrier gas. In our work we used the second one as we wanted the
sulphur to last during the experiment course. For this purpose, we increased the
flow rate from 10-200 SCCM and used COMSOL to correlate the sulphur flux with
the flow rate. Using sulphur flux is more convenient, as the flow rate can be
interpreted differently in furnace tubes with different sizes. We have found that a
minimum flux of 7×10-6 mol/m2.s is required for growing continuous monolayer
films. Below this value, there was an incomplete coverage of the film indicating
that the rate of delivering sulphur does not keep up with the rate of sulphur
consumption.
In light of the findings in the previous section and optimization of the growth
conditions by varying the condition for mass transport of Mo and sulphur in our
system, we next studied the effect of the growth temperature on grain size
distribution and nucleation density. We have found that full coverage monolayers
can be grown in the temperature range 650oC to 850oC, with a maximum possible
grain size up 79740±9670 μm2, and a minimum nucleation density of 82-50 mm-2
in the temperature window 800 oC-850 oC. Based on our finding, this temperature
window is therefore recommended when using MoO2 as starting material.
Furthermore, we have found that, in the case of steady state monolayer growth, a
Mo vapor pressure of 10-11 atmosphere is required for all growth temperatures.
We studied the growth of monolayers in three different regimes: the reaction
limited regime, the mass transport limited regime and the desorption limited
regime. Based in our findings, at temperatures of 650 oC ≤ T<800 oC, the growth is
in the reaction limited regime. At temperatures of 800 oC ≤T ≤ 900 oC, we have the
mass transported regime and at temperatures of 900 oC≤T≤1000 oC, growth is in
the desorption or thermodynamic controlled regime. We could not grow the
monolayers beyond 1000 oC, as the system is controlled by rapid desorption of
starting materials from the substrate before the reaction occurs.
To find the growth rate of individual grains, we limited the growth time to only
one minute and investigated the growth at temperatures of 700 oC-1000 oC. We
have found that the growth rate is proportional to the temperature up to 900 oC at
which the growth rate is 710±260 μm2/s. Beyond 900 oC the growth rate is
reduced and this expected as we are in the desorption limited regime. We also

132
 o
found that the optimum growth temperature is 800 C because at higher
temperatures, the growth of bilayers starts to appear even for a one minute
growth.
We also studied the effect of growth time on the coverage of monolayer
growth. For all temperatures, the optimum growth time that can produce full
coverage monolayers is found to be 15 minutes. When the growth time is
increased to 20 minutes, nucleation of the bilayers is observed, indicating that the
nucleation rates of monolayers and double layers are different. Our monolayer
growth explores this difference. If this difference is due to different substrates for
the nucleation (SiO2 for monolayer and MoS2 for multi-layer), then we expect that
multilayer growth follows the start of the bilayer. This seems to be observed.
Finally, we used the growth temperature and sulphur vapor pressure in the
reaction zone to switch chemical potential of sulphur for controlling the morphology
of MoS2 monolayers. We have produced bent triangles, perfect triangles and
hexagons in a reproducible manner.
From the above discussion, we can conclude that by changing the growth
parameters such as sulphur flux, MoO3 concentration, gas flow and growth
temperature we can switch our system to different growth regimes. At low
temperatures of 650 oC ≤ T < 800 oC the system is controlled by the surface
kinetics, i.e. by active species absorbed by the substrate (S and MoO2 in our
case), the reaction of MoO2 and sulphur and the mobility of the active species on
the substrate surface. The low mobility leads to short diffusion length which limits
the grain growth rate and size. At moderate temperatures of 800 oC ≤ T ≤ 900 oC,
the growth is controlled by the reactant supply (Mo and sulphur). At this
temperature window, the residence time of S and MoO2 on the substrate is very
short and their mobility and reactivity are high. The high mobility of the species
allows them to travel long distances and add at the growing edges of 2D flakes,
causing high growth rates and low nucleation probability. Both are favorable for the
growth of large grains. Further increases in the temperature to 900 oC ≤ T ≤ 1000
o
C and the growth is controlled by the thermal stability of the MoS2 monolayer,
leading the growth to be thermodynamically controlled rather than kinetic
controlled. For example, the chemical potential of sulphur becomes very low such
that the sulphur prefers to be in the gas phase rather than as a solid. The low
chemical potential of sulphur also causes the fast growing S-rich zigzag edge

133
growth rate to slow down and to be comparable to Mo-rich zigzag edges, changes
the final morphology of MoS2 grains to hexagon with Mo and S edge terminations.
All in all, in this chapter, we tried to answer five basic questions about the
CVD growth of the MoS2 monolayers. Our answers come through studying the
effect of growth parameters such as reactant concentrations, temperature, carrier
gas flow rate on the uniformity, continuity, grain size and morphology of the
polycrystalline MoS2 monolayer films. Our findings provide an unprecedented
approach for growing wafer scale uniform monolayers that could be scaled up for
industrial mass production. They also shed light on the growth mechanics
involved.

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Chapter 6: MoS2/MoO2 nanostructure growth

Abstract
In this chapter we present an alternative approach for growing MoS2 films on a
wafer scale. The approach is based on using MoO3 as the Mo source. MoO3 has a
much higher vapor pressure than MoO2 at the same growth temperature. This
enables us to tune the S:Mo ratio to a value that is not possible with MoO2. Since
MoO2 has a very low vapor pressure, the system is always under sulphur rich
conditions. By tuning the supersaturation by varying the S:MoO3 ratio, different
structures such as laterally aligned MoS2 monolayers, MoO2 and vertically aligned
MoS2/MoO2 are produced. COMSOL is used for fluid dynamics simulations and the
concentration of MoO3 and S in each growth run is comprehensively analysed and
the growth mechanisms of the film growth is discussed. The growth of different
structures is confirmed using XRD.

6.1 Literature review

Molybdenum trioxide (MoO3) is one of the most extensively studied Mo sources
and it has been used in growing MoS2 films. Before monolayer of TMDs came to
existence, this precursor had been used for growing MoS2 nanoparticles and
nanotubes [147]. In the pioneer experimental report for monolayers in 2012,
Yi-Hsien Lee et al. used MoO3 to grow laterally aligned triangular shape MoS2
monolayers and multilayers using CVD. Afterwards, MoO3 become the main Mo
source in CVD growth of MoS2 monolayers [148]. Beside growth of laterally
aligned MoS2 monolayers, MoO3 has been used in producing vertically aligned
MoS2/MoOx flakes [149].
The MoO3 vapor pressure is given by the following empirical relations [150]
for 600 oC ≤ T ≤ 795 oC :
75400
4.576log P (M oO3 )n =− T
+ 62.3 6.1
For 795 oC <T ≤ 1155 oC:
35200
4.576log P (M oO3 )n =− T
+ 24.6 6.2
where P is pressure in atm. and T is temperature in K.
When using MoO3 and S as starting material, the growth protocol is as
follows: the MoO3 and S are placed in an appropriate place where the

135
temperatures are high enough for their evaporation. Usually S is placed in the
upstream and MoO3 is placed in the downstream. The S is transported by the
carrier gas to the reaction zone and MoO3 vapor is partially reduced by sulphur to
MoO3-x and deposits on the substrate as nanoparticles [151]. The possible kinetics
as result of MoO3 and S reaction are [152]:
M oO3 + 2x S → M oO3−x + 2x SO2 6:3
7−x 3−x
M oO3−x + 2
S → M oS 2 + 2
SO2 6:4
For a value of x = 1, the reaction proceeds as follows [147]:
M oO3 + 21 S → M oO2 + 21 SO2 6:5
The growth of undesired MoO2 is possible at high MoO3 concentrations used in the
CVD as we will discuss in the rest of this chapter.
Although the growth mechanisms of MoS2 using MoO3-x and S are still not
fully understood, two possible routes are being proposed: (1) The MoO3-x species
diffuse to the substrate and further react with sulphur to form MoS2; (2) The MoO3-x
and S might fully react in the gas phase, and the resulting MoS2 clusters adsorb,
and nucleate on the substrate. The two routes are shown in figure 6:1a.
The partial pressure of MoO3 and S can be controlled by the temperature
which subsequently governs the adsorption and surface-bound reactions on the
substrate. According to the ternary Mo–O–S phase diagram shown in figure 6:1b,
poor sulphur conditions result in low-valence-state oxide or oxysulphide
nanoparticles, thus preventing the direct formation of MoS2 [153], while
sulphur-rich conditions could suppress the volatilization of MoO3, leading to a low
yield of monolayer growth. Therefore a moderate sulphur environment is
recommended during the monolayer growth [151].

Fig. (6:1):(a) Possible growth processes of MoS2 by the reaction of MoO3-x and S. (b)
The Mo–O–S ternary phase diagram, in which the labelled arrows indicate reaction
pathways for the CVD growth of MoS2 from MoO3 precursors [151].

136
6.2 Experimental investigation
The experimental work with MoO3 as one of the starting materials for growing
MoS2 monolayers started with a rectangular duct with walls as shown in figure 6:2
and a rectangular duct with pillars (see fig. 6:3). The duct is made of a
molybdenum foil with a thickness of 0.1 mm and purity of 99.99% (Sigma Aldrich).
The MoO3 powder was spread on the bottom surface of the duct and the SiO2/Si
substrate placed on the top of the duct, facing down to the MoO3 powder. The duct
height was varied between 1-5 cm and the duct was placed in the center of the
tube furnace. The sulphur powder was placed in a ceramic boat and loaded into
the upstream of the furnace in a place where the temperature reaches 200 oC.

Fig. (6:2): Rectangular duct for MoS2 growth.

Fig. (6:3): Duct with pillars for MoS2 growth.

137
6.3 Results and Discussion

6.3.1 Effect of non-uniform MoO3 concentration on the grown film

In our preliminary work, we used ducts with walls. We found that, as a result of the
walls, the MoO3 arriving of the substrate is not uniform, resulting in the non-uniform
film growth. Figure 6:4a is a photograph of the film using a duct with walls. There is
a clear color change from the center to the edges. This color change is due to
thickness changes in the film as will be shown in a series of SEM images. To
overcome the concentration nonuniformity on the substrate, we designed ducts
without walls. The duct is designed such that it leaves very thin pillars in the middle
for fixing the substrate. Figure 6:4b is a photograph of the films grown using the
duct with pillars. It is obvious from the color change that the film is much more
uniform compared to the duct with walls. The experimental conditions for both
cases are the same and as follows: the MoO3-substrate distance is 1 cm, the
growth temperature is 650 oC, and the Ar flow rate is 200 SCCM.

Fig. (6:4): Optical image of MoS2-MoO2 film grown using a duct with (a) walls (b) pillars.

Figure 6:5 (a-f) shows a series of SEM images taken from the edge of the
sample grown with the duct with walls. They clearly demonstrate thickness
nonuniformity as well as morphology changes. At the edge of the sample, isolated
MoS2 monolayer grains are observed. Toward the center, there are a number of
different MoO2 and MoS2/MoO2 structures whose crystalline structure will be
analysed later using the XRD technique.

138
Fig. (6:5a): Edge of the sample MoS2 Fig. (6:5b): 1 mm from the edge, MoO2
monolayers. crystals.

Fig. (6:5c): 2 mm from edge, MoO2 Fig. (6:5d): 2.5 mm from edge, MoS2 film
crystals and vertically aligned MoS2/MoO2 and vertically aligned MoS2/MoO2.

Fig. (6:5e): 3 mm from edge, mostly Fig. (6:5f): 4 mm from edge, vertically
vertically aligned MoS2/MoO2. aligned MoS2/MoO2.

139
 Regarding the duct with pillars, the film shows more uniformity all over the
the substrate. Only MoS2/MoO2 crystals are grown as shown in the following SEM
images (Fig. 6: (a-d)).

Fig. (6:6a): Edge of the sample, vertically Fig. (6:6b): 1 mm from edge, vertically
aligned MoS2/MoO2. aligned MoS2/MoO2.

Fig. (6:6c): 2 mm from edge, vertically Fig. (6:6d): 4 mm from edge, vertically
aligned MoS2/MoO2. aligned MoS2/MoO2.

To understand the non uniformity of the film growth, we used COMSOL

simulations for finding the MoO3 concentration profile on the substrate in both
mentioned cases. In the case of the duct with walls, MoO3 reacts with S and
deposit on the duct walls causing a concentration gradient. Because we do not
know the experimental reaction rate on the walls, we used a reactive sticking
coefficient of 1. Quantitatively, this figure of sticking coefficient may be very crude,
but qualitatively it will be helpful to understand the changes in the MoO3
concentration profiles. Figures 6:7(a-c) show the MoO3 concentration profile on the
substrate when using the duct with walls with a sticking coefficient of 1 and zero

140
and the duct with pillars, respectively. Figures 6:8(a-c) are the corresponding
cross-sectional concentration profiles. Figure 6:9 is a plot of the concentration
profile on the substrate in the three cases.

Fig. (6:7): MoO3 concentration profile on the substrate when using a duct with (a) walls
sticking coefficient of one, (b) wall sticking coefficient of zero (c) pillars.

Fig. (6:8): Cross-sectional view of MoO3 concentration profile in the center of the
reaction zone when using duct with (a) wall sticking coefficient of one (b) wall sticking
coefficient of zero (c) pillars.

Fig. (6:9): MoO3 concentration profile on the substrate in the three cases :duct with
walls with a sticking coefficient of one and zero, and a duct with pillars.

141
 The duct wall before and after deposition is shown optically in figure 6:10
(a-b) with a clear sign of deposition after the experiment run.

Fig. (6:10): Duct walls (a) before and (b) after deposition.

The configuration in which a duct or a boat with walls is placed with the
substrate facing down on the MoO3 powder is the most commonly used one in
CVD of TMDs [138], [154]–[157]. Our experimental and simulation results have
shown that this configuration will result in MoO3 concentration nonuniformity on the
substrate which in turn causes nonuniformity in the film grown. To overcome this
problem we suggest designing ducts with pillars in TMDs CVD growth.

6.3.2 The effect of MoO3 concentration on the film growth at constant

temperature
Here we study the effect of MoO3 concentration on the structure and
morphology of the grown film. The concentration is changed by changing the
vertical distance between the Mo source and the substrate. To avoid deposition on
the MoO3 container walls, we use a duct with pillars as shown in figure 6:3. In
typical experimental procedure, 5 mg of MoO3 powder was loaded in the duct and
the duct is placed in the center of the reaction zone of our furnace. A ceramic boat
containing 10 mg of sulphur was placed in the upstream of the tube furnace in a
position where the temperature reaches 200 oC. Prior to the growth, the furnace
was flushed with argon gas (at 1000 SCCM) for about 30 minutes. Then the
furnace was pumped down to 10 mbar and the center part of the tube wall was
heated to 650 oC with a heating rate of 15 oC /minute in an argon flow (at 200
SCCM). After keeping the temperature at 650 oC for 15 minutes, the furnace was
naturally cooled down to room temperature. The temperature of the sulphur source
also rises, roughly inline with the temperature profile at the centre of the tube
furnace, reaching an average of 200 oC during the 15 minutes when the sample is
at the growth temperature. Figures 6:11(a-e) show SEM images of the films grown
on substrate whose vertical distances from the MoO3 source are 1-5 cm

142
respectively. At 1-3 cm vertical distance, the film consists mostly of MoS2/MoO2
grains grown vertically. At 4 cm vertical distance, some planner MoO2 crystals start
to be observed on a continuous MoS2 multilayer background. A complete coverage
of multilayer MoS2 film is only obtained at a distance of 5 cm. The characterization
of those films will be discussed later on in this chapter.

Fig. (6:11): SEM images of MoS2-MoO2 films grown at T=650 oC and MoO3-substrate
vertical distances of (a) 1 cm (b) 2 cm (c) 3 cm (d) 4 cm and (e) 5 cm.

The (COMSOL calculated) concentration profile of MoO3 on the substrate is

shown in figure 6:12 as a function of vertical distance. It can be seen that the

143
concentration is uniform across the substrate. At 1 cm, the average concentration
4.41 ×10-4 mol.m-3, much higher than what is essential to grow MoS2 monolayers.
Instead, vertically aligned MoS2 /MoO2 grains are grown. At 2 and 3 cm, the
concentration is reduced to 2×10-4 mol.m-3 and 7.4×10-5 mol.m-3 respectively and
the films still have the same structure as in the previous case. At a concentration
of 2×10-5 mol.m-3 (4 cm vertical distance), no further vertically aligned grains are
seen and laterally-aligned MoO2 crystals on a continuous MoS2 film are grown. A
continuous multilayer MoS2 film is deposited at a concentration of 7.4 ×10-6 mol.m-3
(5 cm vertical distance).
We have seen that at a constant growth temperature of 650 oC, a vertically
aligned MoS2/MoO2 grains can be grown in the MoO3 concentration range from
4.41×10-4 mol.m-3 to 7×10-5 mol.m-3 and multilayer MoS2 is realized at a MoO3
concentration of 7.4×10-6 mol.m-3.

Fig. (6:12): Calculated MoO3 concentration on the substrate as a function of vertical

distance at a growth temperature of 650 oC.

144
6.3.3 The effect of MoO3 concentration on the film growth at different
growth temperatures
In this section we present the effect of temperature on MoO3 concentration and the
structure of the grown films. The experimental procedure is as presented in 6.3.2
except we kept the substrate at a vertical distance of d=5 cm and we changed the
growth temperature from 650 oC to 800 oC by increments of 50 oC. The sulphur
concentration at the substrate for all growth temperatures is estimated by
COMSOL to be 0.02 mol/m3. Figure 6:13(a-d) are SEM micrographs of the films
grown at 650 oC, 700 oC, 750 oC, and 800 oC, respectively.

Fig. (6:13): SEM images of MoS2-MoO2 films grown at (a) 650 oC (b) 700 oC (c) 750 oC
and 800 oC.

The concentration profile of MoO3 at different growth temperatures on a

substrate placed at a vertical distance of 5 cm is shown in figure 6:14 below.

145
 Fig. (6:14): MoO3 concentration profile at different growth temperatures on a substrate
placed at a vertical distance 5 cm.

The composition and morphology of the films shown in figures 6:12a-d will be
discussed in the next section as investigated using XRD.

6.3.4 Film characterization using XRD and TEM

The XRD measurements were performed using a Rigaku SmartLab diffractometer
(Cu Kα1 wavelength=1.5405 Å) and θ/2θ scan Bragg Brentano geometry (as
explained in fig. 3:12).
The results for the films grown at 650 oC, 700 oC, 750 oC and 800 oC are
shown in figures 6:15-6:18 respectively.

Fig. (6:15): XRD data for a film grown at Fig. (6:16): XRD data for a film grown at
T=650 oC. T=700 oC.

146
 Fig. (6:17): XRD data for a film grown at Fig. (6:18): XRD data for a film grown at
T=750 oC. T=800 oC.

XRD data from different growth temperatures indicate the coexistence of

MoS2 and MoO2 in the samples and there is no sign for unreacted MoO3. This
means that MoO3 is partially reduced by sulphur to MoO2 or fully reduced to MoS2.
The presence of the MoS2 (002) peak confirms that the MoS2 structure is multilayer
as no such peak should be observed in the case of monolayer growth
(experimentally it is indeed hardly detected) [48]. The sharpness of the (002) MoS2
reflection also means that the crystallites are large in the c-axis dimension. The
presence of the MoS2 (100) peak confirms that there are vertically aligned MoS2
grains. The MoS2 (100) plane is perpendicular to the basal plane (001) and in our
scanning geometry such a peak can not be detected unless there is vertically
aligned grains. The observation of a MoO2 (200) peak indicates that the MoO2
crystals are monoclinic with the orientation relationship of MoO2(100) parallel to
the substrate surface. The presence of a MoS2 (100) reflection and a MoO2 (200)
reflection suggests that the vertically aligned grains are not pure MoS2 but MoO2
wrapped by MoS2 i.e. the grain core is MoO2 and the shell is MoS2. This result
agrees with those reported in the literature [158], [159] indicating that those grains
are firstly grown as MoO2 and then the shell is sulphurized to MoS2. The
sulphurization of the MoO2 crystals is further confirmed using TEM. Figure
6:19(a-b) show TEM images of two typical vertically standing MoS2/MoO2 crystals
showing the MoS2 shell. The interlayer distance of MoS2 flakes was ~ 0.62 nm
which corresponds to the (002) crystal plane of the MoS2 nanosheets.

147
 Fig. (6:19): (a-b): TEM image of typical MoS2/MoO2 crystals showing MoS2 planes at the
edges.

6.4 The effect of S/MoO3 ratio on the film growth mechanisms

In this section, we study the effect of S/MoO3 concentration ratio on the grown fim.
For this purpose we used growth temperatures of 650 oC - 700 oC, the sulphur
concentration is set to 0.02 mol/m3 and the MoO3 is changed by varying the
vertical distance between the MoO3 and substrate.
In contrast to using MoO2 which always led to growth of MoS2 monolayers
(nothing but monolayers) as studied in the previous chapter, using MoO3 as the Mo
source, and based on our experimental result, we can classify the CVD growth
products in to three different types: planar MoS2 monolayers, vertically aligned
MoS2/MoO2 crystals and planar MoO2 crystals, as shown in figure 6:20 a-c
respectively. Our results agree with those reported in the literature [152], [160],
[161].

Fig.(6:20): MoS2-MoO2 structures grown at 650 oC, (a): MoS2 monolayers, (b): vertically
aligned MoS2/MoO2 crystalts, (c): planar MoO2 crystals.

148
 The growth of monolayers (Fig. 6:20a) is only possible at the very diluted
MoO3 concentration limit (MoO3≲8.45×10-7 mol/m3, S/MoO3=2.4×104) i.e. at a
MoO3 partial pressure of 6.4×10-8 atm. A similar value of Mo(CO)6 1.3×10-7 atm. is
found to be crucial to grow MoS2 monolayers using MOCVD [47], however the
experimental time required for MOCVD to achieve monolayers is 26 hr which is
much longer than our typical growth run of 15 minute indicating that Mo(CO)6 has
a much lower reactive sticking coefficient than MoO3 [47]. Increasing the value of
MoO3 vapor pressure to 6.4×10-6 atm (S/MoO3= 2.4×102), vertically aligned
MoS2/MoO2 crystals are observed as shown in figure 6:20b. At a high MoO3 vapor
pressure of 4×10-5 atm and in a weakly reducing atmosphere (S/MoO3=0.38×102),
the growth of planar MoO2 was observed as shown in figure 6:20c. This
observation agrees with our previous results in chapter five when using MoO2 as
Mo sources under poor sulphur conditions as shown in figure 5:47.
Figure 6:21 summarize the effect of S:MoO3 ratio on the film composition at a
typical growth growth temperature 650 oC.

Fig. (6:21): Film composition at different S:MoO3 ratios at a growth temperature of 650
o
C.

It has been assumed that the vertically aligned MoS2/MoO2 structures are
formed via mechanical collision or distortion of planar MoS2 islands [146].
However, this assumption is invalid here as we have found growth of MoS2/MoO2
on bare substrate as shown in figure 6:22.

149
 Figure 6.22 shows the edge area of a sample grown at 700 oC. The image
shows partially covered planar MoS2 monolayers, vertically aligned MoS2/MoO2
structures and the initial stages of MoS2/MoO2 structures .
As one can see in figure 6:22 (blue arrow), the MoS2/MoO2 structures first
nucleate as a nanowire and grow in the vertical direction. Our XRD data confirmed
that the shell of such vertically aligned structures is MoS2 and the core is MoO2.
Based on this, we can assume that under high MoO3 partial pressures (2×10-7
atm), those structures in the early stages are pure MoO2, and that MoO2 keeps
growing in the vertical direction, while, simultaneously the shell is sulphurized into
MoS2.

Fig. (6:22): Partially covered planar MoS2 monolayers, vertically aligned

MoS2/MoO2 grains and nucleation of the vertically aligned MoS2/MoO2
grains.

6.5 Conclusion
In this chapter we have introduced an LPCVD approach using MoO3 as the Mo
source for growing wafer scale planar MoS2 monolayers and vertically aligned
MoS2/MoO2 crystals.
In our preliminary work, we used a MoO3 container with walls, observing that
such a container design resulted in a nonuniform growth of the film on the
substrate. Such nonuniformity is due to the deposition of Mo structures on the

150
container walls causing a concentration gradient on the substrate such that the
concentration on the substrate edges is much lower than the center. The variation
of the MoO3 concentration is confirmed using COMSOL simulations. To overcome
this problem we removed the container walls leaving very thin pillars (1 mm) to fix
the substrate face down to the MoO3 powder. Such a design resulted in more
uniform films and again the MoO3 concentration profile become more uniform.
There are typically two types of designs when growing 2D materials using
powders as transition metal sources. The first one is by placing the substrate face
down on the transition metal compound container and the second one by placing
the substrate downstream in the CVD tube at a certain distance from the container.
In both cases the deposition on the walls (container walls in the first case and CVD
tube walls in the second case) is unavoidable. Therefore, for better uniform
growth, the substrate placed face down on the powder container having no walls is
recommended.
After solving the problem of the design, we studied the effect of concentration
of MoO3 and S/MoO3 ratio on the film growth. The concentration of the gaseous
species can be controlled by the temperature of S and MoO3 or by changing the
the distance between the substrate and the source powder. At a constant growth
temperature of 650 oC and a constant sulphur concentration of 0.02 mol/m3, we
tuned the MoO3 concentration by changing the vertical distance between the
substrate and MoO3 powder. We have found that the growth of vertically aligned
MoS2/MoO2 is dominant at MoO3 concentrations between 4.4×10-4 mol.m-3 and
7.4×10-5 mol.m-3 and the growth of monolayers is only realized at MoO3≲ 8.45
×10-7 mol/m3, S/MoO3=2.4×104 while the growth of planar MoO2 grains is achieved
at high MoO3 concentration 2.6×10-4 mol/m-3 and poor sulphur conditions of
S/MoO3 = 0.38×102.
We further studied the effect of growth temperature on the film growth by
fixing the substrate at a vertical distance of 5 cm above the the MoO3 powder and
changing the growth temperature from 650 oC to 800 oC. Again we tuned the MoO3
concentration from 7.3×10-6 mol/m-3, 4.8×10-5 mol/m-3, 2.6×10-4 mol/m-3 and 1×10-3
mol/m-3 and S/MoO3 ratio from 2735, 417,78 and 19 for growth temperatures of
650 oC, 700 oC, 750 oC and 800 oC respectively. Under these conditions we have
observed the growth of MoS2 multilayer films coexisting with vertically-aligned
MoS2/MoO2 structures which has been confirmed using XRD measurements.

151
Chapter 7: Conclusion and Future Work

7.1 Conclusion
Monolayer MoS2 is an ultra-thin 2D semiconductor with a direct band gap of 1.9
eV. It is proposed as a potential candidate for nanoscale electronics and
optoelectronic devices [12]. Due to its mechanical flexibility, MoS2 monolayers
could also be integrated in flexible and wearable devices. The broken inversion
symmetry monolayer and strong spin-orbit coupling also make monolayers a
promising candidate for spintronic, valleytronic and piezoelectric devices [16], [18].
For practical applications of these monolayers, scalable and controlled
synthesis of monolayers on a wafer scale is required. Thickness uniformity and
grain size are also major concerns for the fabrication of opto-electronic devices.
In this research project, we have focused on establishing an LPCVD
approach for scalable synthesis of MoS2 monolayers using MoO2 and Sulphur as
the starting materials and Si coated with 300 nm of SiO2 as the substrate and
argon as the carrier gas. When growing monolayer crystals, the starting materials
can be considered one of the most important parameters affecting the growth.
In the case of MoS2 monolayer crystals, sulphur has a very high vapor
pressure even at relatively low temperatures, therefore it can not be placed in the
growth zone with a typical temperature of 800 oC. Therefore the sulphur is usually
kept upstream in the CVD tube furnace in a place where the temperature reaches
about 200 oC and the sulphur flux toward the growth zone is controlled by the
carrier gas flow rate. This is an economical alternative to the two furnace set-ups
being reported in the literature. In case of the Mo source, the situation is different,
with the non-toxic Mo sources being Mo oxides such as MoO3 and MoO2. These
oxides have relatively low vapor pressures and need a higher temperature to
produce enough vapor for significant film growth. Therefore, they are usually
placed in the growth zone and the substrate is placed downstream a few
centimeters away from the Mo source. One of the most challenging issues in the
uniform fim growth is delivering a uniform flux of reactant materials across the
substrate. The second challenge is that of bringing the growth into a steady state
i.e. the amount of reactant delivered to the substrate must be equal to the amount

152
of reactant participating in the monolayer growth. As we have mentioned, the
sulphur flux can be controlled by varying the carrier gas flow, so our main task was
to find an appropriate Mo source. We have found among the Mo oxides MoO2 has
the lowest vapor pressure. Our experimental results confirmed that the vapor
pressure of MoO2 is low enough to achieve steady state growth in a range of
growth temperatures when adapting the design we have mentioned. The third task
was to adopt a design for growing scalable monolayer films. This calls for uniform
delivery of reactive species. In our design we have placed the substrate face down
to the Mo source such that all the generated Mo vapor diffuse relatively uniformly
to the substrate.
With the above-mentioned design and MoO2 as the Mo source, we could
grow MoS2 monolayers on a wafer scale. Our XPS measurements confirmed that
a complete conversion to MoS2 occured as result of reaction between MoO2 and S
in the reaction zone. The AFM measurements showed that the thickness of the
grown film is in the monolayer range. SHG images from different parts of the film
proved that the film is uniform. Techniques such as Raman spectroscopy, PL
measurements, TEM, XRD and GIIXRD showed the high crystalline quality of the
films.
We used GIIXRD to study the growth-induced strain in the as-grown
monolayers and thermal expansion coefficient of the monolayers. As a result of the
thermal expansion mismatch between the MoS2 monolayers and the underlying
substrate, a tensile strain up to 1% was observed in the as-grown monolayer. This
value of the strain is comparable with the ones reported in the literature [104]. By
in situ annealing of monolayers under UHV conditions, we have found the
monolayers are stable up 800 oC and their thermal expansion coefficient is found
to be comparable with that of bulk MoS2 [107].
We have tested our approach in different growth temperatures and it is found
that one can reproducibly grow uniform monolayers in the temperature window of
650 oC-850 oC.
Our second goal was to optimize the growth conditions with the help of
COMSOL simulations to produce films with the optimum grain size and to
understand the film growth mechanisms. For this purpose we have studied the
effect of different growth parameters such as MoO2 concentration, sulphur flux,
growth temperature, sulphur chemical potential and growth time on the film growth.

153
 We have related the monolayer coverage to the MoO2 concentration and
sulphur flux. We have tuned the MoO2 concentration by changing the vertical
distance between the substrate and MoO2 powder. Our results indicated that the
full monolayer coverage can be achieved at certain distances where the entire
generated Mo flux diffuses to the substrate. We also quantified the monolayer
coverage as a function of the sulphur flux, by changing the Ar gas flow rate which
subsequently changes the sulphur flux at the reaction zone. We have found that a
minimum flux of 7×10-6 mol/m2.s is required for achieving full coverage
monolayers. Below this value, the substrate was partially covered with monolayers
indicating that our system is in the feed limited-regime with respect to sulphur. At
this stage, we have optimum conditions for growing full coverage monolayers.
Next, we used the aforementioned optimum conditions to investigate the
effect of growth temperature on the grain size distribution and nucleation density.
We have found that a maximum possible grain size of up to 79740±9670 μm2 can
be grown at growth temperatures between 800 oC and 850 oC and that the
nucleation density is suppressed to only 82-50 mm-2 at the same temperature
range. Based on this finding, this temperature window is recommended when
using MoO2 as the starting material.
We further studied the effect of temperature on the film growth, and we have
found that the growth can be classified into three distinct regimes with respect to
the Mo species. At growth temperatures of 650 oC ≤ T < 800 oC, the growth is in
the reaction limited regime, i.e. the grain growth rate is temperature dependent. At
temperatures of 800 oC ≤ T ≤ 900 oC we have the mass transport regime. In this
regime the grain growth rate is weakly temperature dependent. At temperatures of
900 oC ≤ T ≤ 1000 oC, there was the desorption or thermodynamic controlled
regime. There is a high desorption rate in this regime and a large fraction of the
reactants desorb before the reaction occurs. Beyond 1000 oC, we did not observe
any monolayer growth as the growth is completely inhibited by desorption.
To find out the growth rate of individual grains, we limited the growth time to
only one minute for a range of growth temperatures from 700 oC to 1000 oC. The
growth rate was found to be proportional to the temperature and a peak of growth
rate of 710±260 μm2/s occurred at 900 oC beyond which the growth rate is reduced
as the system entered the desorption regime. This follows the classical
description of the CVD processes.

154
 For all growth temperatures, we have found the optimum growth time is 15
minutes, beyond which the nucleation of bilayers appeared on the films. This
suggests that the growth is layer-by-layer due to the different rates of nucleation
on the bare substrate and on the MoS2 monolayer. We speculate that the
tendency for monolayer growth can be enhanced for substrates interacting
strongly with MoS2.
We employed the growth temperature and sulphur partial pressure to tune
the sulphur chemical potential and subsequently tune the morphology of the
monolayer MoS2 grains. We have found that triangular shaped monolayer crystals
are grown at sulphur chemical potentials of µS=-0.834 eV to -0.995 eV, and
hexagonal shaped monolayer crystals are produced for µS=-1.156 eV to -1.318 eV.
As the non-equivalent edges of the hexagonal shaped monolayer crystals can
have different catalytic behaviours, the ability to tune the morphology also means
that we can potentially tune the chemical reactivity of these crystals.
Finally, we have also compared our approach with an LPCVD approach
based on MoO3 as the Mo source. MoO3 has a higher vapor pressure than MoO2
which was previously used. Again, we studied the effect of MoO3 concentration on
the film growth by changing the vertical distance between the MoO3 and the
substrate. Then we studied the effect of growth temperature on the film growth.
We could tune the S:MoO3 ratio, to grow planar MoS2 monolayer, vertically
aligned MoS2/MoO2 and planar MoO2 crystals.
In summary, we have established two LPCVD approaches for growing MoS2
monolayers. The first approach is used for growing uniform MoS2 monolayers on a
wafer scale and the second approach can be adopted for growing vertically
aligned MoS2/MoO2 crystals.

7.2 Future work

As we have seen we have developed a standardized approach for growing MoS2
monolayers over a range of growth temperatures. So far we have used our
approach to grow MoS2 monolayers only. By switching the chalcogen source to
selenium (Se) or tellurium (Te), we can expect to grow MoSe2 and MoTe2
monolayers which are also direct band gap monolayer semiconductors.
Furthermore, we can employ the existed model to produce more complicated
structures such as TMDs heterostructures such as MoS2-MoSe2 and MoS2-MoTe2
and TMDs alloys. 2D novel heterostructures have already been prepared which

155
led to the exploration of numerous exciting physical phenomena and novel nano
electronic and optoelectronic applications such as light emitters [162],
photodetectors [163] new generation field effect transistors [164], and memory
devices [165]. However, 2D heterostructure preparation still mostly depends on the
mechanical exfoliation and a scalable method such as CVD is required. We can
realize the growth of such heterostructures by switching between the chalcogen
sources i.e. first, growing partially-covered MoS2 monolayers and then switching
the chalcogen source to Se or Te to grow horizontally-aligned heterostructures or
growing the full coverage of film and then switching the chalcogen source to grow
vertical stacking heterostructures.
Alloying in TMDs monolayers can be used to tune the band gap of the
monolayers. The alloying can be done by annealing the as-grown MoS2
monolayers in a chalogen rich environment such as Se [166] or by supplying
different chalcogen and transitional metal sources to the CVD system [167]. It has
been reported that by selenization of MoS2, the band gap is tuned from 1.85 eV
for pure MoS2 to 1.57 eV for MoS2xSe2(1–x) [166].
Although we have grown films with optimum grain size, the film is still
polycrystalline with a large number of grain boundaries that affect the electrical
and optical properties of the film. So far we have only employed temperature to get
large grain size. Another way to get larger grain sizes and reduce the grain
boundary numbers is by epitaxial growth using hBN or graphene as a substrate. In
such epitaxial growth the MoS2 grains are expected to grow in only two prefenced
orientations (0o and 60o). At the grain boundary the grains have similar orientations
and will merge eliminating the grain boundary producing larger grains.
Another approach is growing the monolayers on other substrates rather than
SiO2/Si. SiO2 is a good substrate, with very small defects that can be responsible
for heterogeneous nucleation. Optically, the layer thickness can also be adjusted
to make monolayer crystals visible for easy inspection. Well-developed methods
exist for transferring the monolayer films to a substrate of interest. However, it
does suffer from mismatch of the thermal expansion with MoS2, resulting in
thermal cracks. Alternative substrates with smaller thermal expansion coefficients
will be useful for the production of low-stress MoS2 monolayer films. Crystalline
substrates with lattice dimensions comparable with MoS2 can also be used to
study the heteroepitaxial growth of the MoS2 and investigating the effect of the

156
lattice mismatch strain induced on the film properties. Using insulating substrates
with different dielectric constants, one can tune the optical properties of the MoS2
monolayers. As we have mentioned in the main text, the optical emission of MoS2
is excitonic, and the binding energy of the excitons depends on the dielectric
properties of the surroundings.
We believe that all the aforementioned proposals are feasible with our
current approaches. We might need a slight modification to feed different
reactants, but in principle we have developed a framework approach to achieve
them.
Our study is also helpful to our understanding of the growth mechanism
involved. Our results suggest that growth is consistent with the monolayer growth
being activated with nucleation of MoO2 or related oxide/oxysulfide particles on the
film. The growth speed of the monolayer can be very large compared with that of
the multilayer. This is the key, together with the low nucleation density on silica
surface, that allowed us to grow the monolayer MoS2 crystalline films with minimal
presence of multilayers. However, we do not know why there are such huge
differences between monolayer growth and multilayer growth. One possibility
could be attributed to the nature of oxide nanoparticles seen at the growing edge
of the crystals. Further in-situ experiments, particularly those at atomic resolution,
can shed light on the detailed chemical and physical processes at the nucleation
and growth and allow us to understand the CVD parameters from first principles so
that we can rationally design the optimal CVD process to produce desired films for
different applications.

157
Appendix

A1: MoS2 monolayers transferring method

MoS2 monolayer films grown on SiO2/Si substrate were transferred onto the
TEM grid as follows:
1: Polymethyl methacrylate (PMMA) film is spin coated on the
monolayers/SiO2/Si and then baked at 120 oC for 5 minute.
2: PMMA/Monolayer/SiO2/Si were immersed into the boiling sodium
hydroxide solution (1 mol L-1), to etch away the underneath SiO2 layer.
3: The floating PMMA film was picked up with a clean piece of glass and then
immersed in a distilled water to wash away the surface contaminations.
4: The PMMA film was lifted out by a TEM grid covered with lacey carbon film
and left it to dry in ambient temperature.
5: This TEM grid was heated at 120 oC for 5 minuet in air and then immersed
into acetone for about 24 h, to dissolve the PMMA.
6: The film is annealed at 300 oC under Ar environment to remove the PMMA
residues .

A2: Polarization resolved SHG and grain orientation (Matlab

script)

A=imread('XP1.tif'); XP1 is X-polarized SH image.

figure (1);
imagesc(A);
B=imread('YP1.tif'); YP1 is Y-polarized SH image.
imagesc(B),
figure (2);
D=uint8(A);
E=sqrt(double(D));
F=uint8(B);
G=sqrt(double(F));
H=(G./E);
I=atand(H);

158
 J=(1/3).*I;
figure (3);
imagesc (J);
axis image;
colorbar;
title('Orientation image');

A3: Sulphur chemical potential parameters [142]:

Parameter Value Unit

XS8,0 7.62×10-1 eV/atom

XS8,1 -2.457×10-3 eV/atom. K

XS8,2 -4.012×10-6 eV/atom. K2

XS8,3 1.808×10-9 eV/atom. K3

XS8,4 -3.81×10-13 eV/atom. K4

XS2,0 1.207 eV/atom

XS2,1 -1.848×10-3 eV/atom. K

XS2,2 -8.566×10-7 eV/atom. K2

XS2,3 4.001×10-10 eV/atom. K3

XS2,4 -8.654×10-14 eV/atom. K4

XTtr,0 5.077×102 K

XTtr,1 7.272×101 K

XTtr,2 -8.295 K

XTtr,3 1.828 K

a0 1.465×10-2 eV/atom

a1 -2.115×10-3 eV/atom

a2 6.905×10-4 eV/atom

b 10 K

c 80 K

w 100 K

159
Symbols and abbreviations

Symbol Description

n Refractive index

r Relative refractive index

R(n) Reflectivity

eV Electron volt

kV Kilovolt

KE Kinetic energy

h Planck constant

ν Frequency

λ Wavelength

ω Wavenumber

Φ Work function

P Polarization

ε0 Permittivity of free space

χ Susceptibility

E Electric field

θ Diffraction angle in XRD

Θ Angle between between the incident electric field of the

laser beam and MoS2 armchair

α Staking angle of MoS2 bilayers

x
I 2ω Intensity of second harmonic radiation in x-direction.

I y2ω Intensity of second harmonic radiation in y-direction.

Δf Electron microscope defocus

Δf Sch Scherzer defocus

Cs Spherical aberration aberration correction

160
m Spatial frequency

rSch Transmission electron microscope point resolution

β Grazing incident angle

ρ Density

μ Fluid dynamic viscosity

u Fluid velocity

U Mean fluid velocity

Di Diffusion constant of species (i)

ci Concentration of species (i)

Ri Reaction rate of species (i)

Ni Flux of species (i)

Cp Specific heat of fluid

kf Thermal conductivity of fluid

Q Heat sources

q Heat flux

T Temperature

p Pressure

Pref Atmospheric pressure

Ps System pressure

P vapor pressure of precursor

nm Nanometer

Å Angstrom

a Lattice constant

161
μc Carrier mobility

L Transistor channel length

w Transistor channel width

C Capacitance

Vbg Transistor back gate voltage

Vds Transistor drain-source voltage

Ids Transistor drain-source current

F1 Precursor flux from bulk gas

F2 Precursor flux on the substrate

hg Mass transfer constant

ks Reaction rate constant

C1 Concentration of precursor in stream

C2 Concentration of precursor on substrate

δ Boundary layer thickness

Ea Activation energy

k Boltzmann constant

Pe Péclet number

Lr Reactor length

μS Sulphur chemical potential

Hf MoS2 Formation energy

162
Abbreviation Description

2D Two dimensional materials

ADF Annular dark field

AFM Atomic force microscopy

ALD Atomic layer deposition

APCVD Atmospheric pressure chemical vapor deposition

BSE Backscattered electrons

CTF Contrast transfer function

CL Chathodoluminecence

CVD Chemical vapor deposition

EDX Energy- dispersive X-ray

FET Field effect transistor

FWHM Full width at half maximum

GIIXRD Grazing incidence in-plane X-ray diffraction

HAADF High angle annular dark field

HRTEM High resolution transmission electron microscopy

LPCVD Low pressure chemical vapor deposition

MAADF Medium angle annular dark field

MBE Molecular beam epitaxy

MOCVD Metalorganic chemical vapor deposition

PL Photoluminescence

PVD Physical vapor deposition

QE Quantum yield

Rheed Reflection high energy electron diffraction

SAED Selected area diffraction

SCCM Standard cubic centimeter per minute

SE Secondary electrons

SEM Scanning electron microscopy

163
SHG Second harmonic generation

STEM Scanning transmission electron microscopy

SWCNT Single wall carbon nanotube

TEM Transmission electron microscopy

TMD Transition metal dichalcogenide

UHV Ultra high vacuum

XPS X-Ray Photoelectron Spectroscopy

XRD X-ray diffractometry

164
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