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Stoichiometric Hydrogenated Amorphous Silicon
Carbide Thin Film Synthesis Using DC-Saddle Plasma
Enhanced Chemical Vapour Deposition
By
Behzad Jazizadeh
A thesis submitted in conformity with the requirements for the degree of Master of Applied Science
Graduate Department of Electrical and Computer Engineering
University of Toronto
Copyright © 2013 by Behzad Jazizadeh
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Stoichiometric Hydrogenated Amorphous Silicon Carbide Thin Film Synthesis Using DC-Saddle Plasma Enhanced Chemical
Vapour Deposition
Behzad Jazizadeh Master of Applied Science
Graduate Department of Electrical and Computer Engineering
University of Toronto 2012
Abstract
Silicon carbide is a versatile material amenable to variety of applications from electrical
insulation to surface passivation, diffusion-barrier in optoelectronic and high-frequency devices.
This research presents a fundamental study of a-SiC:H films with variable stoichiometries
deposited using novel technique, DC saddle-field plasma-enhanced chemical-vapour deposition,
a departure from conventional RF PECVD commonly used in industry. DCSF PECVD is an
alternative technique for low temperature large area deposition. Stoichiometric a-SiC:H obtained
by fine-tuning precursor gas mixture. Annealing up to 800oC showed no significant change in
elemental composition; particularly indicating thermal stability at stoichiometry. Ellipsometry
showed wide range of optical gaps whose maximum surpasses values reported in literature.
Refractive index measured and change in values studied as function of increasing carbon content
in the films. Also attainment of very smooth surface morphology for stoichiometric a-SiC:H
films reported. Surface roughness of 1 nm rms demonstrated for films grown at temperature as
low as 225oC.
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Acknowledgments
Firstly, I would like to acknowledge and thank my supervisor, Professor Kherani, for his support
and advice as well as for providing the opportunity to partake in research studies in the APD
Labs and Dr. Ali Badar Alamin Dow for partially supervising me and guiding me in all the steps
throughout the research. I am also grateful to them for their trust in me from the start and their
belief in me all the way. I would like to express my deepest gratitude to Dr. Davit Yeghikyan and
Dr. Tome Kosteski for their assistance with the DC Saddle-Field PECVD system, which is the
corner stone of my research. I am also extremely grateful to Zahid Chowdhury and Pratish
Mahtani for their help in instructing me on system operation, discussing various research topics,
and teaching me many characterization techniques. My appreciation goes to Dr. Rana Sodhi for
carrying out XPS measurements and to Edward Xu for assisting me with annealing experiments.
I am also thankful to my colleagues and lab mates John Zhu, Jack Yang and Dave Jeong for their
advice and companionship. My sincere thanks also go to my dear friends who have made my life
in Toronto extremely enjoyable. Finally, I would like to thank Mom, Dad, Behrouz and
Banafsheh for believing in me in every single step.
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Table of Contents
Abstract ........................................................................................................................................... ii
Acknowledgments.......................................................................................................................... iii
Table of Contents ........................................................................................................................... iv
List of Tables ................................................................................................................................. vi
List of Figures ............................................................................................................................... vii
List of Symbols and Acronyms....................................................................................................... x
CHAPTER 1: Introduction ............................................................................................................. 1
CHAPTER 2: Background .............................................................................................................. 4
2.1. Properties .............................................................................................................................. 4
2.2. Applications ......................................................................................................................... 5
2.3. Deposition techniques .......................................................................................................... 6
2.4. Motivation ............................................................................................................................ 7
2.4.1. Stoichiometric a-SiC:H .................................................................................................. 7
2.4.2. DC Saddle Field PECVD .............................................................................................. 8
2.4.3. State-of-the-art ............................................................................................................... 9
CHAPTER 3: Experimental .......................................................................................................... 11
3.1. DC Saddle Field PECVD ................................................................................................... 11
3.2. X-ray Photoelectron Spectroscopy (XPS) .......................................................................... 16
3.3. Spectroscopic Ellipsometry ................................................................................................ 20
3.4. Profilometry ....................................................................................................................... 23
3.5. Atomic Force Microscopy .................................................................................................. 25
3.6. Scanning Electron Microscopy (SEM) .............................................................................. 27
4. CHAPTER 4: Results and Analysis .......................................................................................... 29
4.1. Synthesis of Unity Stoichiometry a-SiC:H Films .............................................................. 29
4.2. Post-Deposition Annealing of Silicon Carbide Films ........................................................ 33
4.3. X-ray Photoelectron Microscopy (XPS) Analysis ............................................................. 37
4.3.1. Introduction ................................................................................................................. 37
4.3.2. Operating Principle of XPS ......................................................................................... 38
4.3.3. Pre-setting Experiments ............................................................................................... 40
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4.3.4. Variation in Temperature............................................................................................. 41
4.3.5. Variation in Location ................................................................................................... 41
4.4. Results and Observations ................................................................................................... 42
4.5. Analysis .............................................................................................................................. 45
4.6. Spectroscopic Ellipsometry ................................................................................................ 57
4.8. Atomic Force Microscopy .................................................................................................. 62
Chapter 5: Concluding Remarks ................................................................................................... 64
5.1. Conclusion .......................................................................................................................... 64
5.2. Issues and Constraints ........................................................................................................ 66
5.3. Future work ........................................................................................................................ 67
References ..................................................................................................................................... 69
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List of Tables
Table 1 A summary of the synthesis experiments as a function of the variation in methane mole fraction, along with corresponding gas mixture by parts and by partial pressure. ....................... 30
Table 2 shows mean thickness values acquired from profilometry measurements. Thickness values are in nanometer (nm) range and the variation is 5% of the (mean) value in the table. .... 31
Table 3 is a summary of compositional measurements acquired from XPS scans. ...................... 39
Table 4 a summary of the fixed deposition parameters for the synthesis of hydrogenated amorphous silicon carbide ............................................................................................................ 40
Table 5 A reference summary of the binding energies for chemical bonds involving Carbon, Silicon, Hydrogen and Oxygen [11]. ............................................................................................ 45
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List of Figures
Figure 1 Schematic diagram of the pentode configuration of the DC-SF PECVD system. Modulation of the proximity of the plasma to the substrate, controlled via the distance between the semi-transparent cathode and the substrate, affects the deposition rate, film properties in general, and surface roughness. .................................................................................................... 12
Figure 2 Schematic diagram of the triode configuration of the DC-SF PECVD where the cathode, substrate and chamber are at ground potential. The remote plasma provides for a low ion energy film growth environment at the substrate. .................................................................. 13
Figure 3 Photograph of the grids, chamber shield and holder following an a-SiC:H deposition run and before KOH cleaning. ............................................................................................................ 14
Figure 4 Photograph of the DC-SF PECVD Deposition System. The deposition chamber is evident from the center to the far left of the photograph. ............................................................. 15
Figure 5 the band diagram at thermal equilibrium shows the shift in the vacuum levels (uniform Fermi levels at equilibrium). The kinetic energy refers to the energy detected by the spectrometer with reference to the vacuum level [39]. ...................................................................................... 17
Figure 6 The XPS tool. ................................................................................................................. 19
Figure 7 An illustration of the mechanism of X-ray Photoelectron Spectroscopy which is based on photoelectric effect................................................................................................................... 20
Figure 8 Photograph of the SOPRA Spectroscopic Ellipsometer. ................................................ 22
Figure 9 Tencor Alphastep 200 profilometer. ............................................................................... 24
Figure 10 (a) A plot of the force between the tip and surface as a function of vertical distance. (b) An illustration of the operating principle of the AFM; as the cantilever deflects the reflection of the laser beam is altered thus providing the vertical difference. ................................................... 26
Figure 11 Nanoscope Dimension 3100 AFM tool. (a) The sample holder is shown and the laser beam source on top. (b) The control unit of the Nanoscope Dimension 3100. ............................ 27
Figure 12 Evolution of the a-Si1-xCx:H film growth/deposition rate as a function of the increasing methane mole fraction χ in the precursor gas mixture. The dashed line is a guide for the eye. ... 31
Figure 13 (a) Carbon mole fraction x and silicon mole fraction 1 – x in the a-Si1-xCx:H films as a function of the methane mole fraction χ in the precursor gas mixture. (b) The rate of increase in the carbon content x with respect to the methane mole fraction χ. ............................................... 33
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Figure 14 the temperature profile of the furnace during annealing. The total annealing sequence time was 375 minutes with the sample at 800oC for 4 hours. ....................................................... 34
Figure 15 (a) Comparison of C (x) and Si (1 – x) mole fractions for As-Deposited and Post-Deposition Annealed samples of a-SiC:H. (b) Bar chart representing percentage relative change for carbon and (c) for Silicon. ....................................................................................................... 36
Figure 16 The oxygen content in the films as a function of the methane mole fraction. ............. 43
Figure 17 Shows the difference between the predicted model (dashed line) for increase in carbon content as a function of the increase in methane partial pressure and the actual trend (triangular symbol). ........................................................................................................................................ 44
Figure 18 The XPS scans (a) before Ar+ etch and (b) after Ar+ sputtering. ................................. 46
Figure 19 The XPS scan deconvoluted for (a) Carbon, (b) Silicon and (c) Oxygen. ................... 47
Figure 20 The XPS scan for four consecutive methane mole fractions, (a) χ = 67%, (b) χ = 71%, (c) χ = 75% and (d) χ =78%. The corresponding increase in the Carbon count is evident from the four diagrams. ............................................................................................................................... 48
Figure 21 The individual peaks appropriately deconvoluted. ....................................................... 51
Figure 22 XPS profile for the samples with methane mole fractions of (a) χ = 91% and (b) χ = 92%. Carbon content is x = 0.43 and x = 46, respectively. .......................................................... 52
Figure 23 The deconvoluted peaks for (a) C1s scan and (b) Si2p scan for χ = 91% CH4 mole fraction. ......................................................................................................................................... 53
Figure 24 the deconvoluted peaks for (a) C1s scan and (b) Si2p scan for χ = 92% CH4 mole fraction. ......................................................................................................................................... 53
Figure 25 The XPS profile for O1s scan for (a) χ = 91% and (b) χ = 92%. ................................. 54
Figure 26 The complete XPS profile of the stoichiometric a-SiC:H. This corresponds to the methane mole fraction χ = 93%. ................................................................................................... 55
Figure 27 All deconvoluted peaks for (a) C1s, (b) Si2p and (c) O1s scans. The peaks correspond to the χ = 93% or the 14 to 1 methane to silane partial pressure. ................................................. 56
Figure 28 The optical bandgap (Tauc gap) as a function of Methane Mole Fraction, χ. The increasing trend of the bandgap is shown in two line segments, (A) and (B), corresponding to the range of χ from 65% to 85% and from 85% to 95%, respectively. .............................................. 57
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Figure 29 An illustration of the increasing trend of optical (Tauc) gap as a function of percent increase in Carbon content of the a-SiC:H films. The trend, as shown, is best approximated by an exponential regression fit (R2 > 0.96); this proves the.................................................................. 59
Figure 30 The change in refractive index, n, as the carbon content x increases. The values are obtained for wavelengths of (a) λ = 248 nm, (b) λ = 365 nm and (c) λ = 633 nm. ...................... 61
Figure 31 (a) rms surface roughness as a function of increasing methane mole fraction. (b) The surface morphology of the stoichiometric a-SiC:H film obtained at χ = 93%. The rms surface roughness of this sample was measured to be 1.014nm. .............................................................. 63
Figure 32 An overview and comparison of achievements in optical bandgap for hydrogenated amorphous silicon carbide films synthesized using different deposition techniques. .................. 66
x
List of Symbols and Acronyms
SiC Silicon Carbide
a-SiC:H Hydrogenated Amorphous Silicon Carbide
a-Si Amorphous Silicon
a-C Amorphous Carbon
RF Radio Frequency
PECVD Plasma Enhanced Chemical Vapour Deposition
x Carbon content
1 - x Silicon content
DC Direct Current
DCSF DC Saddle Field
eV Electron Volt
MV.cm-1
Mega Volts per centimetre
CVD Chemical Vapor Deposition
PLD Pulsed Laser Deposition
LAD Laser Ablation Deposition
HWCVD Hot Wire CVD
ECRCVD Electro Cyclotron Resonance CVD
SiH4 Silane
CH4 Methane
C2H4 Ethylene
C4H10 Butane
KOH Potassium Hydroxide
HF Hydrofluoric Acid
VHF Very High Frequency
MW Microwave
Gen. 1 First Generation
AC Alternative Current
XPS X-ray Photoelectron Spectroscopy
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AFM Atomic Force Microscopy
SEM Scanning Electron Microscopy
IPA Isopropyl Alcohol
DI Deionized
Si - C Silicon Carbon bond
UHV Ultra High Vacuum
SE Spectroscopic Ellipsometry
R2 Standard Error
n Refractive Index
λ Wavelength
α Absorption coefficient
χ Methane Mole Fraction
E Energy
Eg Bandgap
UV Ultra Violet
Å Angstrom
TEM Transmission Electron Microscopy
3D Three Dimensional
kV Kilo Volt
RCA Radio Corporation of America (a method of cleaning in microfabrication)
oC Degrees centigrade
V Volts
mA Milli Amperes
sccm Standard Cubic Centimeters per Minute
rms Root Mean Square
PL Photo-luminescence
XRD X-ray Diffractometry
1
CHAPTER 1: Introduction
Silicon Carbide (SiC), a versatile semiconducting material exhibits a host of interesting
properties which include high mechanical hardness, electrical insulation and chemical inertness.
It is a stable material at high temperatures and can be synthesized in both amorphous and
crystalline forms. Hydrogenated amorphous silicon carbide (a-SiC:H) has an amorphous
morphology within which the hydrogen atoms primarily serve to passivate the dangling bonds in
the structure and thereby reduce the density of dangling bonds and serve to relax the network.
Amorphous SiC is an important technological material. Through the controlled variation of the
relative concentration of its individual elements amorphous silicon carbide exhibits a range of
structural, mechanical, electrical and optical properties.
A great deal of attention has been given to amorphous SiC for a variety of optoelectronic and
microelectronic applications, as well as for protective coatings and hard-wear materials
applications. One important reason for this wide utility of the material is that a-SiC:H can be
easily tuned to have a continuum of varying stoichiometric compositions and corresponding
gradual change in its structural properties.
The characteristics of a-SiC:H can be varied from that of semiconducting amorphous silicon (a-
Si) at one extreme to that of insulating amorphous carbon (a-C) at the other end of the continuum
[1, 2] through variation of the constituent elements.
Moreover, through appropriate annealing treatments amorphous silicon carbide can be
crystallized – a transformation that renders a further change in its properties which can be
desirable. Electrical conductivity and optical transparency improve after crystallization [3].
2
Increases in the electrical breakdown field strength and in the thermal conductivity are among
the progressive modifications which result from the crystallization of hydrogenated a-SiC.
Hydrogenated amorphous silicon carbide films have been prepared using a variety of techniques
including rf PECVD [4 – 8], electron cyclotron resonance [9], rf sputtering [10], and high
temperature chemical vapour deposition [1]. In the context of rf PECVD, a-Si1-xCx:H films have
been studied extensively, leading to the successful preparation of stoichiometric (x = 0.5)
amorphous silicon carbide.
Recently, the dc saddle-field PECVD technique has been used for the preparation of
hydrogenated amorphous silicon carbide as a wide bandgap passivation layer on crystalline
silicon surface [11] albeit the carbon content was relatively dilute and far from stoichiometry (x
= 0.5, where C to Si ratio is unity).
The DC saddle field (DCSF) PECVD technique, a direct field method, offers potentially
significant advantages over existing oscillating field techniques. The DCSF technique provides a
saddle-field through an appropriate configuration of DC biased electrodes and accordingly a
number of distinct advantages follow. Key advantages include its scalability for uniform large
area deposition making it amenable for mass production, and its ability to separate the plasma
regime from the growth regime thereby providing fine control over the growth conditions and
thus the potential of rendering high quality interfaces, films and devices.
The objective of the current research is to systematically study the effects of the variation of gas
mixture on the structural and optical properties of hydrogenated amorphous SiC, while using the
novel Direct Current Saddle Field Plasma-Enhanced Chemical Vapour Deposition (DC-SF
3
PECVD) system, with the ultimate objective of approaching a stoichiometry of unity (Si:C =
1:1).
Further, the overall advantage in using the DC-SF PECVD in relation to the often utilized RF
PECVD will be examined, contrasting the unique aspects of the films deposited with the DC-SF
PECVD – in particular, as an industrially viable low-temperature large-area deposition system
for the fabrication of a-SiC:H.
The current document is organized as follows: Chapter 2 provides background information on
the properties of amorphous silicon carbide. This chapter also includes an overview of the
stoichiometric dependent properties of amorphous silicon carbide.
Chapter 3 describes the experimental apparatus of DC saddle field PECVD system, its unique
characteristics and also various characterization techniques and tools used in this research. A
brief overview of X-ray Photoelectron Spectroscopy, Profilometry, Spectroscopic Ellipsometry,
Atomic Force Microscopy and Annealing Furnace is given.
Chapter 4 presents the experimental results and analyzes the gradual modulation of the
stoichiometry of silicon carbide as a result of varying the methane mole fraction in the precursor
gases and deposition conditions. A discussion of the elemental properties of a-SiC in as-
deposited and annealed films is presented.
In closing, Chapter 5 provides a summary of the overall experimental results their significance
and emerging conclusions. , A discussion of the research limitations, in terms of tools and
material restrictions, is given and potential future research directions are presented. A summary
of the research contributions is also provided.
4
CHAPTER 2: Background
2.1. Properties
Hydrogenated amorphous silicon carbide (a-Si1-xCx:H) has excellent physical and chemical
properties, namely it is thermally stable and thermally conductive, chemically inert, wear and
corrosion resistive. Further a-Si1-xCx:H has superior passivation and insulation properties; the
former can be tuned by doping [12]. In addition its masking and protective coating
characteristics are well documented [13]. Silicon carbide is also a non-oxide ceramic with high
hardness and oxidation resistance not available in other ceramic materials [14]. Hydrogenated
amorphous silicon carbide is further known as an important class of organic-inorganic cross of
glasses with moisture resistivity and low sensitivity to moisture-induced fractures, hence a great
choice for silica-based etch stop layers [15]. It has high optical transmittance and wide tuneable
band-gap [16]. Hydrogenated amorphous silicon carbide has also been an amorphous material of
interest at a fundamental level [17], given this alloy system’s range of possible elemental
compositions and associated variation in properties.
The presence of carbon in this binary semiconductor has created the variability characteristic.
The atomic ratio of carbon in the (C+Si) compound, can hence vary and tune the characteristics
of the a-SiC:H for a specific purpose, such as its optical (e.g. band gap) and electrical (current-
voltage relation) properties. In general the properties of the material can be modified gradually
from the amorphous silicon side of the alloy to the amorphous carbon side of it; for instance the
bandgap of the alloy can vary from approximately 1.75 eV from amorphous silicon to roughly 4
eV for that of amorphous carbon while the electrical conductivity changes from being a
semiconductor on the a-Si:H side to an insulator on the a-C:H side [1]. Electrically, important
among other semiconductors such as Si or GaAs, SiC can have electrical field breakdown as high
5
as 3.0 MV.cm-1 and saturated carrier velocity as high as 2.0x107 cm.s-1 at high fields [18].
Hydrogenated amorphous silicon carbide alloys that possess low carbon content have extensively
been studied and their transparent as well as protective coating properties have been determined
[19].
2.2. Applications
Hydrogenated amorphous silicon carbide has been of interest in both
industrial applications as well as a material of fundamental interest in general and in particular as
a non-oxide amorphous semiconducting material. The wide range of applications is due to its
variable alloying that yields a remarkable combination of mechanical, electrical, chemical and
optical properties.
The potential of a large range of tuneable properties makes a-SiC:H a compelling material of
choice that could provide significant performance enhancements in a spectrum of applications.
Further the simplicity of synthesis, by merely varying precursor gas composition, has promoted a
host of basic material studies in the quest of appropriate combination of properties. As an
example, its wide optical bandgap makes a-SiC:H amenable as a window layer while providing
surface passivation in crystalline silicon solar cells for high conversion efficiency devices [12],
[20].
The wide bandgap property also makes a-SiC:H very useful when applied to electro-
photographic detection. Thin films of a-SiC:H due to its high tolerance to X-ray radiation and
transparency to X-rays have attracted extensive interest in the development of X-ray masks for
X-ray photolithography and proximity printing applications [16, 21]. Their outstanding chemical
and mechanical stabilities make them very useful as protective coatings for a variety of sensing
6
application in automotive, power plant and aerospace industry; they have been shown to have a
better performance as protective layers compared to amorphous silicon oxide and amorphous
silicon nitride [22]. Corrosion and moisture resistive properties of a-SiC:H thin films along with
their chemical inertness has made them suitable for harsh, corrosive and humid environments,
specifically as protective layer for humidity and other chemical gas sensors. Their potential as a
buffer layer integrated within piezoelectric material based devices open the possibility of high
frequency sensors for high temperature applications [24 – 26].
2.3. Deposition techniques
A variety of deposition techniques have been proposed for the deposition of crystalline and
amorphous (hydrogenated) silicon carbide films. Most of these techniques are slight variations of
the main processing technologies, namely chemical vapour deposition (CVD), pulsed laser
deposition (PLD), laser ablation deposition (LAD) and sputtering techniques. Among these
techniques, CVD is the most frequently utilized family of techniques, which at high deposition
temperatures produces the crystalline phase and at low temperatures yields the amorphous phase;
further, in the latter case the as-deposited samples are hydrogenated and if subsequently annealed
thermally then unhydrogenated amorphous and crystalline forms are synthesized [23].
Plasma-Enhanced chemical vapour deposition (PECVD) is the most commonly used process
from the CVD family for the deposition of a-SiC:H. There are other studies that report on the
deposition of thin film a-SiC:H using Hot Wire Chemical Vapour Deposition (HWCVD) and
Electron Cyclotron Resonance (ECR) CVD. The HWCVD approach, besides permitting low
substrate temperatures, features resistive heating of the ‘hot’ wire and accordingly catalytic
decomposition of the precursor gases on the wire as a parameter to control the ion-free
deposition of a-SiC:H with implications in the tuning of its bandgap [20]. ECRCVD on the other
7
hand features separate control over the plasma regime and the substrate peripheral/growth-region
conditions [9]. The PECVD technique however not only features low temperature processing in
general but is also the most widely used process in industry. The deposition temperature used for
this technique is often in the range of 200oC and 500oC [27, 29]. The use of precursor gases in
this system assume a manifold of configurations, ranging from pre-mixed gases introduced
through shower heads to separate introduction of gases in the reaction chamber. The precursor
gases usually consist of silane SiH4 and a hydrocarbon such as methane CH4, ethylene C2H4 and
butane C4H10.
Radio Frequency (RF) PECVD is one variation of the PECVD itself, which is used in dual [28]
and single frequency modes corresponding to the utilization of two and one RF plasma sources,
respectively. Other variations of PECVD sources include the integration of Electron Cyclotron
Resonance [19]. The processes using pulsed laser deposition and magnetron sputtering are
among other deposition techniques used for thin film deposition; these processes use ceramic
SiC targets as the deposition source [29, 30].
2.4. Motivation
2.4.1. Stoichiometric a-SiC:H
Amorphous hydrogenated silicon carbide has certain mechanical, electrical, optical and chemical
properties that when appropriately combined are very attractive in various industrial, commercial
and academic applications and studies. Among the aforementioned characteristics the most
interesting physical properties of a-SiC:H are hardness, optical gap and chemical inertness, all of
which can be controlled, among other parameters, by hydrogen content and stoichiometry of C to
8
Si [23]. Hence, the motivation of this study is to the stoichiometric synthesis of silicon carbide
using the novel DC Saddle-Field PECVD technique.
The focus of this research is to obtain, to the greatest accuracy possible, the 1:1 C to Si
stoichiometric ratio in a-SiC:H thin films by means of a low temperature inexpensive industrially
viable process. Films of a-SiC:H with stoichiometries close to 1:1 have proven to be chemically
stable and inert to a great extent. Further, this is extremely interesting in wet chemical etching of
silicon with KOH and HF, where the use of hydrogenated a-SiC as a passivation layer is etch-
resistive [31].
As the amount of carbon is increased in the a-SiC:H to nearly 50% of the silicon carbide alloy
system, the optical bandgap is also increased [9]. This characteristic along with the etch resistive
property of near stoichiometric a-SiC:H has the potential of this material playing an important
role in the development of the field of silicon micromachining [32]. Other studies have suggested
a-SiC:H thin films to have the largest hardness values near C to Si stoichiometry of unity and
hence a very promising candidate for applications in harsh operating conditions[10, 33].
2.4.2. DC Saddle Field PECVD
Direct Current Saddle Field Plasma Enhanced Chemical Vapor Deposition (DC-SF PE-CVD) is
a novel deposition technique used in this research to produce hydrogenated amorphous silicon
carbide in its near stoichiometric ratio of 1:1. DC-SF is introduced as an improved alternative to
the DC Diode for PECVD processes [34] as well as incorporating other features such as remote
plasma, very low pressure plasmas (lower than RF PECVD), simplicity of operation (dispensing
away with impedance matching issues commonly found in RF systems), and a large parameter
9
space through the use of a multiplicity of semi-transparent electrodes [35]. Furthermore the DC
SF technique, besides giving separate control over plasma parameters on the one hand and local
substrate conditions on the other hand, is a scalable approach for large area production. The
former was, until recently only available through the ECR PECVD; DC-SF provides a
considerably more cost-effective alternative to ECR plasma CVD for industrial applications [36].
2.4.3. State-of-the-art
Research studies have been carried out to produce hydrogenated amorphous silicon carbide thin
films with a range of properties and the dependence of these properties on processing parameters
have been investigated. Although a large number of these studies discuss the importance of the
ability to modulate the elemental composition and in particular the attainment of unity
stoichiometry, very few studies examine the elemental composition of the films synthesized and
even fewer report on the synthesis of unity stoichiometry films. Further, there is no studies to-
date of stoichiometric silicon carbide films using the DC Saddle Field technique, the objective of
the research presented here.
Thin films of a-SiC:H have been prepared by Kuhman D. et. al. (1989) [1] by the employment of
silane (SiH4) and a series of hydrocarbon gases, together as precursors, systematically introduced
to the process, using (rf) PECVD technique. Bandgap, refractive index and mechanical abrasion
resistance and their dependence on the elemental composition have been studied. In another
work, the effect of methane (CH4) to silane (SiH4) ratio and other deposition parameters such as
deposition temperature and total pressure on the quality of surface passivation of a-SiC:H thin
films have been studied by Martin I. et. al. (2001) [4]; incorporation of carbon was shown to
have improved passivation properties. In addition to the above, RF PECVD has extensively been
10
used for the deposition of a-SiC:H thin films of various elemental compositions by Pascual E. et.
al. (1995) [2],Wuu D. S. et. al. (1999) [5], Wang Y. et. al. (2000) [6] and Choi W.K. and
Gangadharan S. (2000) [7] for various fundamental materials studies.
Kaneko T. et. al. (2002) [8] prepared near stoichiometric a-SiC:H films using two different RF
PECVD processes which used different precursors, methyltrichlorosilane and monomethylsilane
each containing both carbon and silicon. The plasma power was the main parameter by which the
composition of the films was determined. Using a different technique, ECRCVD, Cui J. et. al.
(2001) [37] prepared near stoichiometric thin films of a-SiC:H. The increased carbon fraction
and wider optical bandgap was achieved as a result of increasing RF bias voltage. Examining
post-processing follow RF PECVD synthesis, G. De Cesar used a laser treatment was by to
improve the stoichiometry as well as the crystallinity of the as-deposited a-SiC:H films. A
fundamentally different technique, non-reactive RF sputtering was employed by Goranchev B.
et. al. (1986) [10]; in this a silicon target was used as the source in a gas mixture of argon and
methane. Unhydrogenated variation of a-SiC thin films has also been produced to near
stoichiometric ratio using the PLD technique by Tabbal M. et. al (2007) [33]. These films are
synthesized at relatively higher temperatures than those produced using CVD, typically in the
range of 400oC and 950oC, and are subsequently annealed at higher temperatures to improve its
crystallinity.
11
CHAPTER 3: Experimental
3.1. DC Saddle Field PECVD
The majority of the synthesis research on hydrogenated amorphous silicon carbide, as mentioned
previously, has been performed using the Radio Frequency (RF) PECVD technique. The current
work however focuses on using the new DC saddle-field (DCSF) technique for the deposition of
a-SiC:H. The DCSF system not only uses DC power, rather than the commonly used AC power
to generate the plasma, but it is also a departure from the traditional DC diode systems in which
a simple dual-electrode setup is used.
The DC Saddle field approach employs a DC electric field which in a basic configuration is
symmetric along the axis of the plasma deposition chamber (Figure 2), hence causing electrons
to oscillate within the saddle field not unlike a marble oscillating in a bowl under the influence of
a symmetric gravitational force. The saddle-field effectively extends the path of the electrons
[38], which in turn leads to ionizations and excitations that result in a stable plasma. Inside the
plasma chamber the positively biased anode electrode is located at the center, while the other two
cathode electrodes are grounded and located on both sides of the anode. The arrangement for the
two cathodes could be symmetric or asymmetric with respect to their distance from the anode.
The substrate holder is located beyond the cathode. Semi-transparent conducting grids, such as
stainless steel, are typically used for both the anode and cathodes; this allows for both the
oscillation of the electrons and the passage of precursor gas species along the plasma chamber
and thus to the substrate. These decomposed gaseous species and radicals impinge onto the
substrate and hence lead to film growth. The main advantage of the DC Saddle Field over
conventional DC diode plasma techniques is the higher collision probability between the
12
electrons and gas species, which contributes to the activation of the gas species [35]. Further, not
only does the DC Saddle Field method yield stable plasma at lower pressures than the
conventional RF technique but it also provides the capability of separating the growth and
plasma regime through the use of the semitransparent cathode, hence enabling essentially
independent control of the parameters affecting the two regions.
Figure 1 Schematic diagram of the pentode configuration of the DC-SF PECVD system.
Modulation of the proximity of the plasma to the substrate, controlled via the distance between the semi-transparent cathode and the substrate, affects the deposition rate, film properties in
general, and surface roughness.
The choice of electrode locations relative to each other and the number of electrodes used lead to
different configurations for the DC Saddle Field approach. These parameters along with the
distance between the electrodes and the extent of their transparency can be varied to give greater
flexibility and more control over growth conditions. In particular, there are two configurations,
the triode configuration and the pentode arrangement. The former is also sometime referred to as
the shielded triode configuration and is deemed suitable for larger area deposition and yields low
surface roughness film owing to a larger separation between the cathode and the substrate
13
(Figure 2). The latter however features higher deposition rate but also can suffer from higher
surface roughness (Figure 1). This research utilizes the triode arrangement in order to separate
the growth region from the plasma and in particular to definitively establish a remote plasma
which by using the largest separation from the substrate and the cathode. The remote plasma is
generated about the semi-transparent anode with the semi-transparent cathode on the substrate
side and the solid cathode on the other side. The electric field driven ions and radicals are
transported to the substrate passing through the transparent electrodes. These ions are propelled
by the sheath established at the cathode while the radicals are driven by diffusion.
Figure 2 Schematic diagram of the triode configuration of the DC-SF PECVD where the
cathode, substrate and chamber are at ground potential. The remote plasma provides for a low ion energy film growth environment at the substrate.
Sample substrates were cut from silicon wafers with the following parameters
• 100 mm diameter
• n-doped float-zone wafer
• 525 µm thickness
14
• 1-5 Ω cm resistivity
• <100> orientation, single side polished.
Sample pieces of 25 mm square crystalline silicon were used to deposit the a-SiC:H films and
accordingly characterized using XPS (X-ray Photoelectron Spectroscopy), Profilometry, AFM
(Atomic Force Microscopy) and SEM (Scanning Electron Microscopy).
Figure 3 Photograph of the grids, chamber shield and holder following an a-SiC:H deposition run and before KOH cleaning.
Prior to each a-SiC:H film deposition, a three step cleaning process was used to prepare the
silicon samples. In the first step, samples were immersed in acetone for 10 minutes with the
beaker placed inside an ultrasonic bath. The samples were then transferred into another beaker
containing isopropanol alcohol (IPA) also in an ultrasonic bath for 10 minutes. The final
cleaning step involved immersing the samples in a deionized (DI) water ultrasound bath for 10
minutes. Silicon samples were then blow dried using a high pressure nitrogen gun.
15
Figure 4 Photograph of the DC-SF PECVD Deposition System. The deposition chamber is evident from the center to the far left of the photograph.
Typically the samples were loaded into the chamber and placed under a vacuum within 5
minutes of the last cleaning and drying step. Every effort was made to ensure that each
deposition run was carried out under identical systematic chamber conditions. Prior to each
deposition, a pre-cleaned stainless steel shield was placed on the interior cylindrical wall of the
chamber. Furthermore, pre-cleaned anode and cathode grids were also installed in the chamber.
The samples were placed on a pre-cleaned substrate holder. Following each deposition, the
stainless steel shield, grids and substrate holder were then cleaned by immersing them in a hot
bath of KOH; Figure 3 shows the shield and grids following a typical deposition. Potassium
hydroxide serves as an etchant to remove any unwanted deposits of amorphous silicon carbide.
The cleaned set was subsequently rinsed with hot water followed by cleaning with DI water.
They were then dried in an oven, at least over-night prior to being placed in the chamber for
another deposition. A photograph of the DC-SF PECVD System is shown in Figure 4.
16
A summary of the deposition parameters used for the synthesis of the hydrogenated amorphous
silicon carbide films are given and discussed in Section 4.3.3 of Chapter 4.
3.2. X-ray Photoelectron Spectroscopy (XPS)
X-ray photoelectron spectroscopy (XPS) is used for elemental and compositional analysis of the
silicon carbide films (see Figure 6). The data collected by the XPS tool for the carbon, silicon
and oxygen content was then analyzed to quantitatively determine the composition of the thin
films of a-SiC:H deposited using the DCSF tool. XPS also provides quantitative data regarding
the types of chemical bonds; hence detailed XPS scans can reveal information regarding the
nature of the bonding in alloys such as hydrogenated amorphous silicon carbide.
As it might appear from the name of this characterization tool, the working physics of XPS
comes from the principle of the photoelectric effect. XPS uses X-ray as the source for short
wavelength electromagnetic radiation to bombard the surface of the sample. An electron detector
and analyzer then perceive photoelectron emission induced from the surface. The system then
measures simultaneously the kinetic energy and the number of electrons in the detected emission.
The collection of this measured data results in the XPS spectrum. The XPS spectrum is used to
determine the composition of the material, where the X-ray induced photoelectrons have
characteristic energies corresponding to the atoms in the sample. The photoelectron energies are
measured and analyzed to determine the elemental and chemical compositions. This is obtained
from an energy relation that calculates the electron binding energy of the photoelectrons from the
17
measured kinetic energy (Ekinetic) and known values of the X-ray wavelength and work function
of the spectrometer (Øspectrometer). Both the sample holder and the spectrometer are connected
together electrically and in turn connected to ground. This prevents the accumulation of surface
charges. The band diagram at thermal-equilibrium is shown below and equation (3.1), given
below, is derived directly from the band diagram.
∅ (3.1)
Figure 5 the band diagram at thermal equilibrium shows the shift in the vacuum levels (uniform
Fermi levels at equilibrium). The kinetic energy refers to the energy detected by the spectrometer with reference to the vacuum level [39].
The elements are identified based on their binding energies, as each element has its own binding
energy, resulting from the XPS scan. The binding energy itself brings two more interesting
characteristics, the peak intensity and the shift. The intensity of a given binding energy
18
corresponding to an element makes up for the quantitative analysis of the element within the
film. Higher counts and larger area for a peak corresponds to a more abundant element within the
compound. Further, the chemical bond configurations are determined from the shifts in the peaks
[40]. The varying magnitudes of the shifts explain the variation in chemical bonds and their
different types of bonds with other elements. In other words different bonding configurations
with other elements cause the detected binding energy to shift from the main known peak for that
element. The reason for the shift is the Coulomb interactions of the core electrons between the
two elements; the interactions cause the shift in binding energy that can vary from 0.1 to 10 eV
in magnitude. As an example, the shift from 99.2 eV to 100.5 eV corresponds to Si electrons in
pure Si atoms and Si electrons from Si - C bonds, respectively, indicative of the small difference
in binding energies owing to the change in Coulombic fields [41]. A list of the binding energies
for different bonds is listed in table 1, section 4.1. A sample XPS survey scan and an example of
the deconvolved C 1s peak are shown below. The peaks of the XPS spectrum are an illustration
of the electron configuration corresponding to the core electrons (1s, 2s, 2p, etc.). These peaks
are then correlated with elements of the same binding energy. The extracted XPS spectrum is a
superposition of peaks from the various types of bonds and it is possible to deconvolve the entire
spectral range in order to determine different chemical bonding configurations corresponding to
different elements within the samples. Also in order to cancel any offsets resulting from
asymmetry due to the noisy background, the Shirley background was applied to the recorded
data [42].
The XPS measurements utilized in this research were obtained using the Thermo Scientific K-
Alpha instrument. It uses a monochromatic aluminum K-α X-ray source, with a variable spot
size of 30 – 400 µm. The sample stage setup can hold several samples for multiple XPS
19
measurements in one loading. XPS characterization technique is a non-destructive type of
measurement of the surface chemistry and it has the capability of detecting the top 1-10 nm
surface of the material being analyzed, for surface composition analysis. X-ray Photoelectron
Spectroscopy suffers however from the following restrictions:
• It is an ultra-high vacuum (UHV) characterization tool. This is necessary since it
minimizes the error in photoelectron counts at each energy level.
• XPS is unable to detect hydrogen or helium due to their small orbital diameters.
Photoelectron catch probability is hence reduced to nearly zero. However the detector is
able to identify all other elements with atomic numbers higher than hydrogen and helium.
Figure 6 The XPS tool.
The K-Alpha instrument is also equipped with an argon ion gun that uses an argon ion beam to
sputter atoms and etch the surface of a given sample. This becomes crucial when the sample
surface being analyzed is contaminated. The argon gun in this case is used to sputter off the
surface contamination, specifically the oxygen and organic contaminants. This so called pre-
20
sputter processing step can be controlled by varying the time of sputtering and the energy level
of the ion beam impinging the surface from the gun. The argon sputtering has another advantage
which is depth profiling using XPS. In this technique XPS measurements are performed in
between two consecutive sputtering runs and the characterization is continued until the entire
thin film is etched off the (silicon) substrate.
Figure 7 An illustration of the mechanism of X-ray Photoelectron Spectroscopy which is based on photoelectric effect.
3.3. Spectroscopic Ellipsometry
Spectral Ellipsometry (SE) is another form of surface characterization which is non-destructive.
The working principle of spectroscopic ellipsometry (SE) lies behind the variation in polarization
of light incident on the surface following reflection from the sample surface. SE can be used to
resolve the properties of either single or multi-layer films and is a useful optical technique that
can be used to determine a number of dielectric properties of thin films, namely, refractive index,
extinction coefficient, thickness, absorption coefficient and optical band gap as well as other
optical constants. The modeling that is widely used for fitting the raw data is the conventional
21
Tauc-Lorenz model. In this research this model was used to carry out a regression fit of the
measured data for the hydrogenated amorphous silicon carbide thin films and thus a range of
properties were inferred. The application of the model requires defining the structure (number of
layers) and assuming initial values of the model parameters which include thickness and the
complex refractive index. In this work these requirements were applied and in particular the
following three layers were assumed:
• Native oxide film as the upper most layer represented as the thin roughness layer
• Amorphous silicon carbide film as the layer beneath
• Crystalline silicon substrate as the platform
For thickness values, the measurements from first-order approximations using profilometry can
be used for an initial guess of parameters in the fitting routine. After the first regression run a
good graphical fit is obtained with a relatively low fitting error (R2 > 0.98); the optical bandgap
as well as other optical constants are returned in the list of reported values. In this manner it is
possible to extract the refractive index and extinction coefficient (n and k, respectively) as a
function of energy and hence wavelength. The value of the optical bandgap is follows from the
equations below [43]:
√ (3.2)
(3.3)
! "# (3.4)
$" % (3.5)
22
In, Equation (3.2), the relationship between the absorption coefficient, , and the Tauc gap
() is displayed as a function of energy (E), where is a constant. The absorption coefficient
equation is displayed in Equation (3.3) as a function of wavelength (!) and extinction coefficient,
', and the wavelength can be converted to energy through Equation (3.4) using the Planck’s
constant () and the speed of light, (. Through substitution of the equations above, Equation
(3.5) is in the form of ) *+ ,, where a linear fit can be performed on the Tauc plot to
determine the Tauc gap from the X-intercept.
Figure 8 Photograph of the SOPRA Spectroscopic Ellipsometer.
The ellipsometer used for optical characterization in this research is the Lambda 9 Sopra UV-
VIS-NIR Spectroscopic Ellipsometer (Figure 8) to carry out all SE measurements on a-SiC:H
samples. The range of photon energies that was used lie between 0.5 eV and 5.5 eV. This device
has a loader arm and performs an initialization process before each analysis. After obtaining the
23
spectral data for the sample, the fitting is done manually on the raw data. It is during this manual
regression step that the Tauc-Lorentz model is used and the first set of inferred values is
obtained.
3.4. Profilometry
Profilometry is a direct characterization technique producing first-order measurements of the
thickness of thin film materials. This technique provides information about the thickness as well
as roughness by measuring the profile of the film surface. Tencor Alphastep 200 Automatic Step
Profiler (Figure 9) was used for profilometry and thickness measurements in this work. The 12.5
µm tipped stylus of the step profiler has two degrees of freedom, being able to move along the
vertical and horizontal axes. The process starts with the tip first moving downwards in a vertical
order to make contact with the surface of the sample. The tip then moves in a horizontal
direction, laterally sweeping across the sample in a straight line. The profiler tool is accurate
enough to detect vertical features ranging from 10 nm to 1 mm with a vertical resolution of 5 Å
and a horizontal resolution of 400 Å. The vertical displacement of the stylus generates an analog
signal that is converted into a digital signal to be analyzed and displayed on screen. Since the
profiler is able to detect vertical displacements, the thickness of films can be measured across a
step between the bare wafer and the film. This requires the sample to be prepared using a mask
to create a step by preventing any film growth in a selected area.
24
Figure 9 Tencor Alphastep 200 profilometer.
While the profilometry measurements are independent of optical properties of the sample, they
are sensitive to the physical and mechanical properties of the film, such as surface roughness and
hardness, respectively. In addition, the procedure may damage the sample surface in contrast to
contactless methods. Accurate thin film thickness measurements can also be determined from
cross-sectional transmission electron microscopy (TEM), but the sample preparation involves a
complex procedure. In contrast, profilometry is a fast and straightforward method that provides a
first order approximation of thickness. The trade-off of loss in accuracy can be justified by the
simplicity and quickness of profilometry measurements.
25
3.5. Atomic Force Microscopy
Atomic Force Microscopy (AFM) is considered to be a planar imaging method that is used for
the analysis of surface morphology and surface roughness. It is an improvement over the original
imaging method, Scanning Tunnelling Microscopy, which images surfaces that are either
conductive or semiconducting. AFM however provides a 3D profile of the surface on a
nanoscale and is capable of performing imaging of almost any type of surface, including
polymers, ceramics, composites, glass, and biological samples.
The working principle of AFM lies behind the reflections of a laser source onto a detector that is
also a position sensor. The reflection is due to the movement of the AFM lever and the beams are
reflected from the back of the AFM tip, as it undergoes deflections caused by the surface
roughness (see Figure 10 (a) and (b)); deflections occur as the tip touches the surface recording
small magnitudes of force between the AFM probe and the sample surface. The radius of the tip,
which is often microfabricated from silicon (Si) or silicon nitride (Si3N4), is often quite sharp
(<10 nm) but can vary from a few nanometers to tens of nanometers.
(a)
26
Figure 10 (a) A plot of the force between the tip and surface as a function of vertical distance. (b) An illustration of the operating principle of the AFM; as the cantilever deflects the reflection of
the laser beam is altered thus providing the vertical difference.
The AFM probe can be thought of as a spring which is placed on the end of the lever. According
to Hooke’s law the magnitude of the force between the AFM tip and sample depends on the
spring constant, which also can be thought of as the stiffness of the lever, and the vertical
distance z between the tip and the sample surface.
(b)
27
Figure 11 Nanoscope Dimension 3100 AFM tool. (a) The sample holder is shown and the laser beam source on top. (b) The control unit of the Nanoscope Dimension 3100.
Surface roughness and morphology in this research were measured and characterized for a-SiC:H
using Digital Instruments’ Nanoscope Dimension 3100 (see Figure 11 (a) and (b)). The
Dimension 3100 controller integrates the illuminator, power supply, and air and vacuum pumps.
It features positioning repeatability of 3µm unidirectional and 4-6 µm bidirectional on the X-Y
stage while the motorized Z stage provides accurate and automatic tip engagement and approach.
The Video Zoom Microscope also features through-the-lens illumination, Colour video camera
and focus tracking and automated engagement.
3.6. Scanning Electron Microscopy (SEM)
Cross-sectional examination of the microstructure of the films was analyzed using a scanning
electron microscope (SEM) (mode1 570, Hitachi, Hitachi Ltd, Tokyo, Japan). A scanning
(a) (b)
28
electron microscope contains an electron source which focuses a tiny beam/spot of electrons onto
the specimen. As it scans the beam across the specimen, it concurrently detects the response and
transmits it to a display. The Hitachi 570 is a general purpose, 25 kV, diffusion-pumped,
tungsten filament-based SEM.
Samples which are usually less than 2 cm on a side are mounted on a specimen stub with a
conductive paste or a metal clip. Larger pieces such as 5-10 cm wafers may be loaded with user-
supplied holders. The specimen stage is non-eucentric (i.e., the focal point does not change with
stage adjustments) and requires that the chamber be brought to atmospheric pressure for loading
and unloading. After sample loading and evacuation, a high voltage, typically 4 – 25 kV is
applied.
29
4. CHAPTER 4: Results and Analysis
4.1. Synthesis of Unity Stoichiometry a-SiC:H Films
The deposition of hydrogenated amorphous silicon carbide films using the DC Saddle Field
PECVD is performed using the precursor gases, silane (SiH4) and methane (CH4). There are two
gas bottles that are connected to a number of mixing bottles. Fine control in the composition of
the synthesized films is obtained by mixing the precursor gases first in the mixing bottles prior to
flowing the mixture into the chamber. The plasma is ignited after introducing the gas flow into
the chamber at the set pressure. The choice of the starting point for the gas mixture was made
using a traditionally used recipe for silane applicable in the deposition of hydrogenated
amorphous silicon, a-Si:H. The notation here for methane mole fraction and stoichiometric ratio
are χ, and x in Si1-xCx, respectively. Methane mole fraction is defined by the ratio of moles of
methane to the moles of the mixture of methane and silane: χ ./0./0123/0. For instance, 1 part
Methane and 1 part Silane corresponds to 4 0.50; this methane mole fraction was the starting
point of experiments. For a methane mole fraction of 0.5, the resulting stoichiometric ratio is x =
0.15 (1 – x = 0.83) in the a-Si1-xCx:H film; note, that the fact that x + (1 – x) ≠ 1 reflects the
presence of some impurities in the sample. Table 1 summarizes the range of synthesis
experiments, showing the ratio of methane to silane, the partial pressures of methane and silane
in the mixing bottles, and the methane mole fraction.
30
CH4:SiH4
By parts By Partial pressure (Torr) χ (%)
1:1 500:500 0.50 2:1 667:333 0.67
2.5:1 714:286 0.71 3:1 750: 250 0.75
3.5:1 778:222 0.78 6:1 857:143 0.86
10:1 909:91 0.91 12:1 923:77 0.92 14:1 933:67 0.93
Table 1 A summary of the synthesis experiments as a function of the variation in methane mole fraction, along with corresponding gas mixture by parts and by partial pressure.
The hydrogenated amorphous silicon carbide film growth rate as a function of the methane mole
fraction χ is shown in Figure 12. The thickness data corresponding to growth rate of Figure 12 is
given in Table 2. This table shows the mean (average) value for thickness at each methane mole
fraction (conforming to each experiment run) and the variation given in percentage of mean
value is calculated based on an average of 5 measurements on each sample. The
growth/deposition rate is observed to decrease from 1.9 nm/min to 1.5 nm/min with increasing
mole fraction χ, indicating an inversely proportional dependence on the methane mole fraction χ.
Close observation of the data in Figure 12 shows an inflection point in the growth rate at a
methane mole fraction of approximately χ = 78%, suggesting that beyond this point the
characteristics of a methane plasma dominate the growth rate and composition of the films; this
is better understood when examining the composition data presented in Figure 13 (a) below. It is
31
worth highlighting that the growth rate levels off at a practically viable 1.5 nm/min deposition
rate as the films approach stoichiometry.
Methane Mole Fraction, χ Thickness (nm)
50% 76 67% 75
71% 74
75% 72 78% 68
86% 480
91% 420 92% 450
93% 450
Table 2 shows mean thickness values acquired from profilometry measurements. Thickness values are in nanometer (nm) range and the variation is 5% of the (mean) value in the table.
Figure 12 Evolution of the a-Si1-xCx:H film growth/deposition rate as a function of the increasing
methane mole fraction χ in the precursor gas mixture. The dashed line is a guide for the eye.
1.4
1.5
1.6
1.7
1.8
1.9
0.5 0.67 0.71 0.75 0.78 0.86 0.91 0.92 0.93
Poin
t-to
-Poi
nt F
ilm g
rwot
h ra
te (n
m/m
in.)
Methane mole fraction χ
32
The carbon mole fraction x and the silicon mole fraction 1 – x in the a-Si1-xCx:H films as a
function of the methane mole fraction χ are shown in Figure 13 (a). The dashed lines represent
least squares fitting to a parabolic curve with correlation coefficient R2 for the two curves
approaching unity. It is interesting to observe that in order to obtain stoichiometric a-SiC:H we
require a methane mole fraction of approximately χ = 0.93%, or equivalently a silane mole
fraction of about 7%. This indicates that the occlusion of C is dominated by the methane plasma
and corresponding growth processes. To illustrate this further, the dependence of the rate of
increase in the carbon mole fraction x with respect to the methane mole fraction as a function of
x is shown in Figure 13 (b). The rate of C occlusion drops off rapidly with increase methane
mole fraction as the films approach stoichiometry. It is noteworthy that these films are
essentially at stoichiometric parity in C and Si albeit there is oxygen in the films in the range of 1
to 4 atomic percent as determined by XPS. The source of the oxygen is attributed to the grid and
chamber surfaces owing to the fact that the system was not baked prior to the depositions.
R² = 0.9915
R² = 0.9557
0
10
20
30
40
50
60
70
80
90
0.5 0.55 0.6 0.65 0.7 0.75 0.8 0.85 0.9 0.95
C (
x) a
nd S
i (1-
x) c
onte
nt [
%]
Methane mole fraction, χ
C
Si
(a)
33
Figure 13 (a) Carbon mole fraction x and silicon mole fraction 1 – x in the a-Si1-xCx:H films as a function of the methane mole fraction χ in the precursor gas mixture. (b) The rate of increase in
the carbon content x with respect to the methane mole fraction χ.
4.2. Post-Deposition Annealing of Silicon Carbide Films
In order to realize the stability of elemental (silicon and carbon) composition of the thin films of
hydrogenated amorphous silicon carbide deposited using DC Saddle Field PECVD, the thin film
samples were subjected to a thermal annealing process. The process involved first applying an
RCA cleaning sequence in order to remove the organic contaminants from the surface of the
films. Annealing was subsequently performed at a temperature of 800oC for duration of 4 hours.
The diagram in Figure 14 shows the ramp-soak-ramp temperature sequence in degrees Celsius as
a function of time in minutes.
R² = 0.9141
0
2
4
6
8
10
12
0.5 0.55 0.6 0.65 0.7 0.75 0.8 0.85 0.9 0.95
dx/dχ
Methane mole fraction, χ
(b)
34
Figure 14 the temperature profile of the furnace during annealing. The total annealing sequence time was 375 minutes with the sample at 800oC for 4 hours.
The a-SiC:H films showed stability in chemical composition after annealing at 800oC. Figure 15
below shows the elemental composition for carbon and silicon before and after annealing, as a
function of increasing methane mole fraction, χ.
25
175
325
475
625
775
0 121 244 366
tem
pera
ture
, oC
Time elapsed, minutes
35
10
20
30
40
50
60
70
80
0.50 0.60 0.70 0.80 0.90
C (
x) a
nd S
i (1
–x)
Methane Mole Fraction, χ
C (Annealed)
Si (Annealed)
C
Si
-30
-20
-10
0
10
20
30
Perc
enta
ge R
elat
ive
Cha
nge
-C
arbo
n
Methane Mole Fraction, χ
(a)
(b)
36
Figure 15 (a) Comparison of C (x) and Si (1 – x) mole fractions for As-Deposited and Post-Deposition Annealed samples of a-SiC:H. (b) Bar chart representing percentage relative change
for carbon and (c) for Silicon.
A visual difference between the increasing trends for Carbon and Silicon can be seen from
Figure 15 (a). Figure 15 (b) illustrates, more deeply, the difference between the increasing trends
for carbon from before and after annealing process, in terms of percentage relative change for
carbon content in elemental composition of a-SiC:H. An acceptable difference (of |10%| or less)
cannot be observed until the methane mole fraction is greater than χ = 0.78%. Although for
carbon content this is the case, the trend for Silicon content shown in Figure 15 (c) is not
representative of a meaningful trend, although the relative differences are generally within the
|10%| of the actual value, with the exception of the change at χ = 0.92%. Interestingly enough for
stoichiometric a-SiC:H however, both Figures 15 (b) and (c) show changes of less than |5%|.
This is consistent with the thermal stability of amorphous hydrogenated silicon carbide reported
by Wang et al. [44]. No trace of crystalline texture within the amorphous matrix was observed,
-30
-20
-10
0
10
20
30
Perc
enta
ge R
elat
ive
Cha
nge-
Silic
on
Methane Mole Fraction, χ
(c)
37
through the characterization of samples using XRD, Raman and SEM. One possibility for this
could come from the small thickness of a-SiC:H relative to large thickness of the substrate; and
since the measurements were performed at large incident angles it is highly probable that the X-
ray effectively penetrates the substrate instead and no crystalline information about the nano-
scale thickness of the thin film is gained. Imaging using SEM also failed to reveal a possible
crystalline texture; the fact that the surface, as reported in section 4.8, embodied a low rms
roughness and lack of contrast within the amorphous matrix made SEM ineffective.
4.3. X-ray Photoelectron Microscopy (XPS) Analysis
4.3.1. Introduction
Since the main focus of this research is to achieve near stoichiometry hydrogenated amorphous
silicon carbide films using the DC-SF PECVD technique, the elemental analysis and chemical
composition of the deposited samples is key to the characterization of the deposited films.
Furthermore the nature of chemical bonding and the determination of the abundance of the
various chemical bonds will help in gaining a deeper understanding of the as-deposited thin films
of a-SiC:H with changes in the elemental composition of the films as the deposition parameters,
(principally the precursor gas composition) are varied.
With XPS it is possible to perform a range of measurements on a given sample:
• Identify the elements in the sample
• Carry out a chemical (bonding) analysis of the elements
• Undertake quantitative characterization to determine the chemical composition
38
• Perform depth profiling of the sample using an ion gun (destructive method)
• Do non-destructive depth profiling by the angle-resolved analysis
• Also one implements surface charge neutralization, using the electron flood gun, in order
to perform chemical imaging. It is customary to use some combination of these
techniques for effective XPS analysis of a sample.
A sample XPS spectrum of a given material is a result of characteristic peaks due to the range of
elements and their chemical make-up where the peaks often overlapping each other. The peak
shift, width and intensity are keys to determining the chemical and elemental characteristics of
the material.
4.3.2. Operating Principle of XPS
The working mechanism behind an XPS instrument is shown in Figure 7 in Section 3.2. Below
the surface of the sample the electronic states of the atoms are excited by means of the x-ray
photons having a specific energy. A hemispherical analyser is utilized to filter the energy of the
ejected electrons off the surface. This is before the detector measures the intensity of the
electrons with a specific energy. The XPS spectra result in energy intensity peaks corresponding
to the resonance peaks; these peaks are reflections of the electronic structure of surface atoms, a
result of the quantized core level electrons in the elements. Although the depth of penetration for
x-rays into the sample may be large there are restrictions to the depth of the ejected electrons
from the surface. In other words the probability of electrons ejecting from depths greater than say
10 nm is small for energies of the range of 1400 eV; if these electrons undergo an energy loss
interaction, then this phenomenon creates a background signal which contributes to the overall
39
emission of electrons rather than simply well -defined peaks resulting from the primary
photoelectric effect.
The binding energies of the photoelectric lines are considered well defined in general,
corresponding to electronic states of atoms. This along with the specific chemical environment
creates the corresponding shift in energies which reflect the nature of the chemical bonds. XPS
analysis can be sub-divided into conducting and non-conducting samples. For conducting
samples, and with respect to the spectrometer the detected electron energies can be referred to
the Fermi level energy, thus providing a clear energy scale. This however is problematic for
insulating samples due to sample charging which gives rise to a retarding field. This calls for a
significant energy calibration (sometimes up to 150 eV), as without which significant energy
shifts will follow. Table 3 summarizes the compositional data extracted from XPS, for a-SiC:H
samples, as described above. The elemental compositions are listed as a function of methane
mole fraction.
Methane mole fraction, χ
Carbon content, x
Silicon content, 1 – x
Oxygen content
50 % 0.1472 0.8337 0.0191 67 % 0.2086 0.7539 0.0375 71 % 0.2393 0.7279 0.0328 75 % 0.2696 0.6936 0.0368 78 % 0.2862 0.6279 0.0289 86 % 0.3499 0.6145 0.0356 91 % 0.4306 0.5694 0.0309 92 % 0.4640 0.5360 0.0288 93 % 0.4886 0.4736 0.0379
Table 3 is a summary of compositional measurements acquired from XPS scans.
40
4.3.3. Pre-setting Experiments
In order to achieve the optimum set of the DC-SF PECVD deposition parameters for the
hydrogenated amorphous silicon carbide films, a series of experiments was performed which we
refer to as the pre-setting experiments. A number of the deposition parameters were kept
constant throughout all the experiment, such as deposition pressure, anode current, anode voltage
and gas flow rate to name a few. Table 4 lists the entire set of fixed parameters. These
parameters were categorized as constraints. They have already been optimized in general for
other similar research in the group - specifically, for previous work on the deposition of thin
amorphous films on silicon substrate.
Parameter
Value (Unit)
Heater Temperature 400 (oC)
Substrate Temperature 225-240 (oC)
Anode Current 17.5 (mA)
Anode Voltage 550 (V)
Total Mixing Gas Pressure 1000 (Torr)
Deposition Pressure 160 (mTorr)
Gas Flow 30 (sccm)
Table 4 a summary of the fixed deposition parameters for the synthesis of hydrogenated amorphous silicon carbide
41
4.3.4. Variation in Temperature
Deposition temperature was the first deposition parameter to be varied in order to determine its
effect on changes in the elemental composition of the samples. This process however was limited
to heater temperature of 400oC, a constraint imposed by the heating stage of the system.
Nevertheless, essentially no significant influences on the sample’s stoichiometry and elemental
composition were observed for heater temperature as high as 400oC. Accordingly, the heater
temperature was fixed value at 400oC for all subsequent depositions. The heater temperature of
400oC corresponds to substrate temperature (also known as film growth temperature) of, on
average, 232oC.
4.3.5. Variation in Location
The effect of the location of the samples on the sample holder was investigated and the results
were analyzed thereof. The reason for this was the fact that the sample holder has a large surface
area and the shift in location of samples may affect the thin film stoichiometry and elemental
composition. The analysis revealed that variation in the location of the samples has no effect on
their stoichiometric compositions, implying that the combination of the uniform geometry of the
electrodes in the DC-SF PECVD technique along with a sufficient separation between the semi-
transparent cathode and the substrate holder yields uniform growth conditions at the substrate
holder..
42
4.4. Results and Observations
The deposition of hydrogenated amorphous silicon carbide (a-SiC:H) using initial methane mole
fraction of χ = 50% resulted in a thin film with elemental composition of x = 0.15, 1 – x = 0.83
(the values for Si (1 – x) will rarely be used throughout the rest of this document, for the sake of
convenience and to prevent repetition). As the precursor gas mixing ratio increases towards
increasingly higher partial pressure for methane (CH4), the atomic percentage of carbon (x) starts
increasing, as expected [11]. However, the rate of increase of C (x) in Si:C ratio is considerable
as a function of the increase in methane mole fraction and, as a result, with respect to the partial
pressure of CH4. In other words as the CH4 mole fraction was increased by almost ∆χ = 17%, the
realization of C in silicon carbide shows an almost 42% rise from x = 0.15 to x = 0.21.
Subsequent experimental depositions consisted of an increase in methane mole fraction of only
∆χ = 4% each time, equivalent to an increase in methane corresponding to 2.5, 3 and 3.5 parts of
CH4 to 1 part of SiH4; these experiments resulted in carbon content of x = 0.24, x = 0.27 and x =
0.29, respectively. Another reason for the low value of methane ramping steps was the oxygen
content of the thin film - a rather drastic increase from 1.91% to 3.75% - reflecting a rather thick
oxide layer on top of the a-SiC:H thin film. The percentage rise in carbon content relative to the
content in the prior film set is therefore calculated to be approximately 15%, 13% and 6%,
respectively - which shows a sustained reduction in the increase in carbon content with
increasing methane mole fraction. One observation worth noting at this step was that the oxygen
content of the a-SiC:H thin film did not exceed 3.75%. Figure 16 shows the percentage of
oxygen content in the films as a function of the methane mole fraction used for the deposition of
the various films.
43
Figure 16 The oxygen content in the films as a function of the methane mole fraction.
The rate of increase in carbon content of the thin film started following a declining trend as the
methane mole fraction was raised from 67% to 78%. This suggested that for the next deposition
process runs, a higher gas mixing ratio would be required. This placed the DC Saddle Field
plasma in to a new discharge mode of operation, given the considerably larger fraction of
methane as compared to previous experiments. With the increase of methane mole fraction to χ =
86%, the resulting elemental composition showed an increase of almost ∆x = 22% in the carbon
content of the thin film.
As a result of this increase in the mole fraction of methane the new Si:C ratio achieved is 1 – x =
0.62 and x = 0.35. One important observation at this stage of carbon content is that the amount of
oxygen content incorporated within the film remains between the lower limit of 1.91% and upper
limit of 3.75%. Another important observation, discussed earlier in this section, is the rate of
00.5
11.5
22.5
33.5
44.5
5
0.50 0.60 0.70 0.80 0.90
Oxy
gen
Con
tent
, %
Methane Mole Fraction, χ
O
O (Annealed)
44
increase in carbon content versus methane mole fraction as a function of methane mole fraction
(see Figure 13(b)).
An attempt was made to predict the percentage increase of carbon content in the Si:C compound
as a result of a ramp step of 0.5 in the precursor gas mixture (partial pressure). To do this, the
rate of change of C increase as a function of increase in mixing ratio was calculated based on the
immediately preceding two previous process runs. This “worst case scenario” was used as an
estimate of the slope or rate of linear rise. Figure 17 below shows the predicted line along with
experimental results from subsequent/future deposition runs. The model provides a reasonable
approximation of the targeted carbon content, which is evidently valid within an error of ~ 12%
at its maximum spread and at its best a spread of less than 1%, is evaluated.
Figure 17 Shows the difference between the predicted model (dashed line) for increase in carbon content as a function of the increase in methane partial pressure and the actual trend (triangular
symbol).
The predictive model shown in Figure 17 serves as a guide to the selection of of methane mole
fraction for subsequent experiments in the quest for a stoichiometric compound. Based on Figure
30
35
40
45
50
55
6 6.5 7 7.5 8 8.5 9 9.5 10 10.5 11 11.5 12
Car
bon
Con
tent
x
CH4:SiH4 Partial Pressure
predictedactual
45
18, the next deposition run with χ = 92% should incorporate 0.48 of carbon content in the film;
this however proved to be lower, resulting in about 0.46 C content. At a methane mole fraction χ
of 93% we achieved films with x = 0.49 and 1 – x = 0.47. Further attempts to fine tune the
carbon content by adjusting methane mole fraction was not fruitful, as a result of excess oxygen.
4.5. Analysis
The analysis for the XPS spectrum cannot be realized unless the relative positioning of the peaks
based on the binding enery are known. Table 5 presents a summary lookup table for binding
energy peaks relevant to thin film material deposited on top of silicon substrates. It will be used
as a reference for the analysis and observations that follow.
Bond (Binding) Energy Peak
Si – Si / Si – H 99.2 ± 0.1 Si – C 100.5 ± 0.1 O – Si – C 101.8 ± 0.2 Si – Ox 103.2 ± 0.3 C – Si 283.2 C – C / C – H 284.6 ± 0.1 C – O – H 286.4 C = O 288.4 ± 0.4
Table 5 A reference summary of the binding energies for chemical bonds involving Carbon,
Silicon, Hydrogen and Oxygen [11]. From the XPS scans (see Figure 18) the carbon content is determined by calculating the
integrated intensity of the peaks for C1s to that of C1s and Si2p [45]. Further, from the XPS
scans the evolution in the structure of a-SiC:H samples can be observed with increasing methane
mole fraction. The XPS spectrum for the first deposition run is shown below. This spectrum
corresponds to the χ = 50% CH4 mole fraction, for which the Si/C ratio is 1 – x = 0.84 and C
mole fraction is x = 0.14. Deposition temperature is 400oC (heater temperature).
46
Figure 18 The XPS scans (a) before Ar+ etch and (b) after Ar+ sputtering.
From Figure 18 and based on bonding energies from Table 5, the position of peaks in the
spectrum can be found to correspond to Si2p, C1s and O1s at approximately 99, 283 and 531 eV,
respectively. Figure 18 (a) indicates major peaks visible at approximately 150 and 980 eV, which
are attributed to Si2s (associated with the substrate) and O Auger signal, respectively. The high
oxygen content, with peak area comparable to that for carbon, is attributed to surface oxidation.
For this reason argon etch process was performed using an Ar+ sputtering gun for 30 seconds and
a subsequent XPS scan was obtained. The resulting spectrum is illustrated on the right in Figure
18 (b). The removal of the oxide layer as a consequence of Ar sputtering is clearly visible with
the peak count corresponding to the O1s bond is now much lower compared to the un-etched
sample.
Figure 19 below consists of individual scans for Carbon, Silicon and Oxygen content
incorporated in the a-SiC:H thin film. The individual scans contain deconvoluted XPS profiles
for each of Carbon, Silicon and Oxygen representing various bonding information that each peak
shows. The peaks in Figure 19 (a) for carbon content are show that for the most part a single
bonding energy overlaps with the peak at 283 eV indicating the strong C – Si bonding. For Si
scan (Figure 19 (b)), 100.2 eV and 99.5 eV represent Si – C bonding with high count and Si – O
0.00E+00
2.00E+04
4.00E+04
6.00E+04
8.00E+04
1.00E+05
1.20E+05
1.40E+05
1.60E+05
1.80E+05
2.00E+05
01002003004005006007008009001000110012001300
Counts
/ s
Binding Energy (eV)
Survey
Position = 400°A, X = 53055.5 µm, Y = 55913.6 µm
0.00E+00
2.00E+04
4.00E+04
6.00E+04
8.00E+04
1.00E+05
1.20E+05
1.40E+05
1.60E+05
1.80E+05
2.00E+05
2.20E+05
01002003004005006007008009001000110012001300
Counts
/ s
Binding Energy (eV)
Survey
Position = 400°A, X = 53055.5 µm, Y = 55913.6 µm
(a) (b)
47
with a low count. The peak detected at the binding energy of almost 99.5 eV, corresponds to
Si2p3 bonding. The scan for O content (Figure 19 (c)) shows two peaks at 531.3 eV and 532.1
eV, collectively suggesting the O1s bond.
Figure 19 The XPS scan deconvoluted for (a) Carbon, (b) Silicon and (c) Oxygen.
As the Argon etching process was successful in more accurate determination of elemental
composition and other relative chemical information about the sample, it was decided to perform
this technique on all the samples from now onwards. The rest of the analysis therefore is
0.00E+00
1.00E+03
2.00E+03
3.00E+03
4.00E+03
5.00E+03
6.00E+03
7.00E+03
8.00E+03
9.00E+03
1.00E+04
1.10E+04
1.20E+04
96979899100101102103104105106107108109110
Counts
/ s
Binding Energy (eV)
Si2p Scan
Position = 400°A, X = 53055.5 µm, Y = 55913.6 µm
Si2p3
Si2p1
1300
1400
1500
1600
1700
1800
1900
2000
2100
2200
2300
526528530532534536538540542544
Counts
/ s
Binding Energy (eV)
O1s Scan
Position = 400°A, X = 53055.5 µm, Y = 55913.6 µm
O1s
O1s A
(b) (c)
1000
2000
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4000
5000
280282284286288290292294296298
Counts
/ s
Binding Energy (eV)
C1s Scan
Position = 400°A, X = 53055.5 µm, Y = 55913.6 µm
C1s(a)
48
discussed based on not the as-deposited samples, but the samples which their surfaces were
etched from 30 s to 300 s with an Argon gun; based on the amount of Oxygen hence the
thickness of the oxide layer.
Figure 20 The XPS scan for four consecutive methane mole fractions, (a) χ = 67%, (b) χ = 71%, (c) χ = 75% and (d) χ =78%. The corresponding increase in the Carbon count is evident from the
four diagrams.
Figure 20 is an illustration of the XPS spectrum for the four consecutive deposition runs in which
the methane mole fraction was sequentially increased in steps of 3% to 4% from 67% to 78%
corresponding to steps of 0.5 in the ratio of partials pressures from 1/2 to 1/3.5 for SiH4/CH4. A
monotonic trend is observed as expected in the binding energy profile for all four graphs, with
0.00E+00
2.00E+04
4.00E+04
6.00E+04
8.00E+04
1.00E+05
1.20E+05
1.40E+05
1.60E+05
1.80E+05
2.00E+05
01002003004005006007008009001000110012001300
Counts
/ s
Binding Energy (eV)
Survey
Position = 1_2 etched, X = 47186.9 µm, Y = 20326.4 µm
0.00E+00
2.00E+04
4.00E+04
6.00E+04
8.00E+04
1.00E+05
1.20E+05
1.40E+05
1.60E+05
1.80E+05
2.00E+05
01002003004005006007008009001000110012001300
Counts
/ s
Binding Energy (eV)
Survey
Position = 1_2.5 etched, X = 20368.6 µm, Y = 17951.2 µm
0.00E+00
2.00E+04
4.00E+04
6.00E+04
8.00E+04
1.00E+05
1.20E+05
1.40E+05
1.60E+05
1.80E+05
2.00E+05
01002003004005006007008009001000110012001300
Counts
/ s
Binding Energy (eV)
Survey
Position = 1_3 etched, X = 47524.7 µm, Y = 50382.7 µm
0.00E+00
2.00E+04
4.00E+04
6.00E+04
8.00E+04
1.00E+05
1.20E+05
1.40E+05
1.60E+05
1.80E+05
2.00E+05
01002003004005006007008009001000110012001300
Counts
/ s
Binding Energy (eV)
Survey
Position = 1_3.5 etched, X = 19351.8 µm, Y = 50037 µm
(a) (b)
(c) (d)
49
obvious rise in the count for Carbon, moving from lower to higher values of methane mole
fraction.
As expected the peak for silicon substrate has the highest count and the peaks for silicon
belonging to silicon carbide and Carbon within the film have the next large consequent counts. A
very slight count peak of Auger Oxygen is still (with a careful glance) visible. The Oxygen count
is still in existence with a relatively low count, as expected based on numerical percentage values
discussed in the previous section.
The set of deconvoluted peaks are also collectively shown in Figure 21. It is worth noting that
the deconvolved peaks for Oxygen were not available from the XPS measurements. The
discussion about the peaks for Silicon content is for the most part un-altered by the change in the
gas mixing ratio and the resulting change in the stoichiometric ratio. The peak corresponding to
Si2p3 is responsible for the Si – C bonding while the Si2p1 peak is thought to have been shifted
from 100.4 eV to 99.9 eV. This is generally the case with two exceptions for values of χ = 75%
and χ =78% for which the shift is approximately to 100 eV and to 99.6 eV respectively.
1000
2000
3000
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5000
280282284286288290292294296298
Counts
/ s
Binding Energy (eV)
C1s Scan
Position = 1_2 etched, X = 47186.9 µm, Y = 20326.4 µm
C1s SiC
C1s A
C1s B
0.00E+00
1.00E+03
2.00E+03
3.00E+03
4.00E+03
5.00E+03
6.00E+03
7.00E+03
8.00E+03
9.00E+03
1.00E+04
1.10E+04
96979899100101102103104105106107108109110
Counts
/ s
Binding Energy (eV)
Si2p Scan
Position = 1_2 etched, X = 47186.9 µm, Y = 20326.4 µm
Si2p3
Si2p1(a) C1s (a) Si2p
50
1200
1400
1600
1800
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2200
2400
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526528530532534536538540542544
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Binding Energy (eV)
O1s Scan
Position = 1_2 etched, X = 47186.9 µm, Y = 20326.4 µm
1000
2000
3000
4000
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6000
280282284286288290292294296298
Counts
/ s
Binding Energy (eV)
C1s Scan
Position = 1_2.5 etched, X = 20368.6 µm, Y = 17951.2 µm
C1s SiC
C1s A
C1s B
0.00E+00
1.00E+03
2.00E+03
3.00E+03
4.00E+03
5.00E+03
6.00E+03
7.00E+03
8.00E+03
9.00E+03
1.00E+04
1.10E+04
96979899100101102103104105106107108109110
Counts
/ s
Binding Energy (eV)
Si2p Scan
Position = 1_2.5 etched, X = 20368.6 µm, Y = 17951.2 µm
Si2p3
Si2p1
1200
1400
1600
1800
2000
2200
2400
2600
526528530532534536538540542544
Counts
/ s
Binding Energy (eV)
O1s Scan
Position = 1_2.5 etched, X = 20368.6 µm, Y = 17951.2 µm
1000
2000
3000
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6000
280282284286288290292294296298
Counts
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Binding Energy (eV)
C1s Scan
Position = 1_3 etched, X = 47524.7 µm, Y = 50382.7 µm
C1s SiC
C1s A
C1s B
0
1000
2000
3000
4000
5000
6000
7000
8000
9000
10000
96979899100101102103104105106107108109110
Counts
/ s
Binding Energy (eV)
Si2p Scan
Position = 1_3 etched, X = 47524.7 µm, Y = 50382.7 µm
Si2p3
Si2p1
(a) O1s
(b) Si2p
(b) C1s
(b) O1s
(c) C1s (c) Si2p
51
Figure 21 The individual peaks appropriately deconvoluted. Each of the four iteration is based on increasing mole fraction is categorized by (a), (b), (c) and (d), where each of (a). (b), (c) and (d)
consist of C1s, Si2pand O1s scans of that iteration.
In the discussion of Carbon peaks, peaks attributed to C1s bond are prefixed as C1sA and C1sB.
It is conjectured that these peaks collectively could be regarded as being due to Organic
(contaminant) Carbon content. This would be valid if a shift of 0.4 eV from the reference C –
Organic at 285.4 eV could be realized. The minimum and maximum shifted values from the
reference are 0.4 eV and 1.6 eV, respectively. And finally the Oxygen peaks from the four
sequential graphs can collectively be considered as belonging to the oxide layer of the thin film.
1200
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526528530532534536538540542544
Counts
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Binding Energy (eV)
O1s Scan
Position = 1_3 etched, X = 47524.7 µm, Y = 50382.7 µm
1000
2000
3000
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5000
280282284286288290292294296298
Counts
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Binding Energy (eV)
C1s Scan
Position = 1_3.5 etched, X = 19351.8 µm, Y = 50037 µm
C1s SiC
C1s A
C1s B
0.00E+00
1.00E+03
2.00E+03
3.00E+03
4.00E+03
5.00E+03
6.00E+03
7.00E+03
8.00E+03
9.00E+03
1.00E+04
1.10E+04
96979899100101102103104105106107108109110
Counts
/ s
Binding Energy (eV)
Si2p Scan
Position = 1_3.5 etched, X = 19351.8 µm, Y = 50037 µm
Si2p3
Si2p1
1000
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5000
526528530532534536538540542544
Counts
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Binding Energy (eV)
O1s Scan
Position = 1_3.5 etched, X = 19351.8 µm, Y = 50037 µm
(c) O1s
(d) Si2p (d) O1s
(d) C1s
52
Figure 22 XPS profile for the samples with methane mole fractions of (a) χ = 91% and (b) χ = 92%. Carbon content is x = 0.43 and x = 46, respectively.
Figure 22 shows the complete XPS profile for χ = 91% and χ = 92% CH4 mole fraction,
respectively. The increase in Carbon content is evident in the diagrams. Another important
characteristic of the graphs above is the uneven behaviour visible close to the left end of the
profile. At binding energies close to 1220 eV, there is a jump in the profile; it is thought to
belong to Auger Carbon. This behaviour was seen in previous profiles however the intensity of it
was considerable here. This intensity has been observed to increase throughout the sequential
increase in Carbon content within the silicon carbide thin films.
The following set of figures (see Figure 23 (a) and (b)) describes the individually deconvoluted
peaks for the compound elements. In Figure 23 (a) for χ = 91% methane mole fraction the
Carbon scan contains three individual peaks namely C1s, C1sA and C1sB at binding energies
282.6, 283.8 and 284.8 eV respectively. Although they can collectively (in the convoluted form)
be considered as being correspondent to C – Si bond, the C1sB due to its proximity to the
reference value of 285.0 eV can be considered to be Carbon attributed to Organic contaminant.
As far as the area under curve is concerned the area for the Si – C bonding representing Carbon
has expectedly the largest value among all the three peaks. This makes the assumption of
0.00E+00
2.00E+04
4.00E+04
6.00E+04
8.00E+04
1.00E+05
1.20E+05
1.40E+05
1.60E+05
01002003004005006007008009001000110012001300
Counts
/ s
Binding Energy (eV)
Survey
Position = 1_10 etched, X = 51721.8 µm, Y = 23625.8 µm
0.00E+00
2.00E+04
4.00E+04
6.00E+04
8.00E+04
1.00E+05
1.20E+05
1.40E+05
1.60E+05
01002003004005006007008009001000110012001300
Counts
/ s
Binding Energy (eV)
Survey
Position = 1_12 etched, X = 16413.6 µm, Y = 22382.7 µm
(a) (b)
53
attributing C1sB to organic contaminant more probable. On the other hand, based on Figure 23
(b) Silicon peaks on the right hand side could also be thought of as a combination of Si2p3 and
Si2p both contributing to the Si – C chemical bonding.
Figure 23 The deconvoluted peaks for (a) C1s scan and (b) Si2p scan for χ = 91% CH4 mole fraction.
For the sample with χ = 92% CH4 mole fraction (Figure 24) the results for individual peaks are
quite similar to the ones obtained above for the χ = 91% mixing ratio. The C1sB peak appears to
be lower in counts compared to the case for χ = 91% ratio shown above.
Figure 24 the deconvoluted peaks for (a) C1s scan and (b) Si2p scan for χ = 92% CH4 mole fraction.
1000
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4000
5000
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7000
280282284286288290292294296298
Co
unts
/ s
Binding Energy (eV)
C1s Scan
Position = 1_10 etched, X = 51721.8 µm, Y = 23625.8 µm
C1s SiC
C1s A
C1s B
0
1000
2000
3000
4000
5000
6000
7000
8000
96979899100101102103104105106107108109110C
ounts
/ s
Binding Energy (eV)
Si2p Scan
Position = 1_10 etched, X = 51721.8 µm, Y = 23625.8 µm
Si2p3
Si2p1
1000
2000
3000
4000
5000
6000
7000
8000
280282284286288290292294296298
Co
unts
/ s
Binding Energy (eV)
C1s Scan
Position = 1_12 etched, X = 16413.6 µm, Y = 22382.7 µm
C1s SiC
C1s A
C1s B
0
1000
2000
3000
4000
5000
6000
7000
8000
96979899100101102103104105106107108109110
Co
unts
/ s
Binding Energy (eV)
Si2p Scan
Position = 1_12 etched, X = 16413.6 µm, Y = 22382.7 µm
Si2p3
Si2p1
(a) (b)
(a) (b)
54
Regarding the Oxygen content of the thin film, the following two graphs each have one peak as
illustrated for χ = 91% and χ = 92% gas mixing ratios, respectively (Figure 25). The graphs on
both sides represent peaks very close to the binding energy reference point for the hydroxide
bonding. This suggests the oxygen to come from not only the oxide naturally occurring during
the deposition process, but also resulting from a cleaning process using potassium hydroxide
(KOH). This cleaning process happens at least 24 hours prior to the scheduled deposition run, in
which all the grids, sheath and holder are cleaned to remove the films from a previous deposition
process. Both peaks show a relatively low count compared to the ones for Carbon and Silicon
content.
Figure 25 The XPS profile for O1s scan for (a) χ = 91% and (b) χ = 92%.
Figure 26 shows the complete XPS spectrometry of the hydrogenated amorphous silicon carbide
deposited using DC Saddle Field with the gas mixing ratio of χ = 93%. It shows the full profile
containing peaks representing Carbon, Silicon and Oxygen at energies 282.62, 99.54 and 531.03
eV. Besides the peak corresponding to the silicon substrate at almost 150 eV, there are other
peaks appearing at approximate locations of 245, 980 and 1220 eV along the XPS profile. These
1000
1200
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2200
2400
526528530532534536538540542544
Co
unts
/ s
Binding Energy (eV)
O1s Scan
Position = 1_10 etched, X = 51721.8 µm, Y = 23625.8 µm
1200
1400
1600
1800
2000
2200
2400
2600
526528530532534536538540542544
Co
unts
/ s
Binding Energy (eV)
O1s Scan
Position = 1_12 etched, X = 16413.6 µm, Y = 22382.7 µm
(a) (b)
55
peaks are familiar and they are thought to be attributed to (traces of) Argon (2p3, 2p and 2p1),
Auger Oxygen and Auger Carbon, respectively. From the profile, Carbon C – Si is observed as a
sharp and narrow peak, which although having a higher count than Silicon Si – C bond,
nevertheless due to its sharpness has an almost equal area under the curve to the one for Silicon
Si – C bond.
Figure 26 The complete XPS profile of the stoichiometric a-SiC:H. This corresponds to the methane mole fraction χ = 93%.
The detailed XPS profile for individual peaks for Carbon, Silicon and Oxygen are given below
(see Figure 27). Figure 27 (a) contains the single C1s peak representing a clear C – Si chemical
bond. No observations were possible to the contribution of other sources of Carbon, such as
organic contaminants. On the other hand the diagram for Silicon (see Figure 27 (b)) consists of
the two deconvoluted peaks Si2p3 and Si2p1 at 99.4 and 99.9 eV, respectively. These two peaks
are collectively considered as the Si – C chemical bond. In addition, part (c) of the set (see
Figure 27 (c)) is an illustration of two distinct peaks. The one at approximately 531 eV is
0.00E+00
1.00E+04
2.00E+04
3.00E+04
4.00E+04
5.00E+04
6.00E+04
7.00E+04
8.00E+04
9.00E+04
1.00E+05
1.10E+05
1.20E+05
01002003004005006007008009001000110012001300
Counts
/ s
Binding Energy (eV)
Survey
Position = 5h etched more, X = 29030.9 µm, Y = 25321 µm
56
assumed to have the background of the pre-cleaning contaminant from KOH, as discussed above.
Furthermore the peak tagged as O1sA is thought to be responsible for the oxide layer.
Figure 27 All deconvoluted peaks for (a) C1s, (b) Si2p and (c) O1s scans. The peaks correspond to the χ = 93% or the 14 to 1 methane to silane partial pressure.
0
1000
2000
3000
4000
5000
6000
280282284286288290292294296298
Counts
/ s
Binding Energy (eV)
C1s Scan
Position = 5h etched more, X = 29030.9 µm, Y = 25321 µm
C1s
0
1000
2000
3000
4000
5000
96979899100101102103104105106107108109110
Counts
/ s
Binding Energy (eV)
Si2p Scan
Position = 5h etched more, X = 29030.9 µm, Y = 25321 µm
Si2p3
Si2p1
800
900
1000
1100
1200
1300
1400
1500
1600
1700
1800
526528530532534536538540542544
Counts
/ s
Binding Energy (eV)
O1s Scan
Position = 5h etched more, X = 29030.9 µm, Y = 25321 µm
O1s
O1s A
(a) (b)
(c)
57
4.6. Spectroscopic Ellipsometry
The effectiveness of the variation in temperature and position of the samples, when undergoing
deposition, has already been studied [11]; and the results showed that neither an increase in
temperature (given the limits of the system), nor a change in the sample’s position on the holder
affect the values of the optical gap (Tauc gap). The optical gap values related to samples with
variable methane mole fractions were extracted from the SE absorption data and are plotted in
Figure 28. A widening effect on the optical gap is clearly observed for the samples (deposited at
400°C) with increasing methane mole fraction in the precursor gas mixture. The optical gap
ranges from a very low value of 1.06 eV to a high value of 3.77 eV for the range of gas mixture
compositions stated in Table 1.
Figure 28 The optical bandgap (Tauc gap) as a function of Methane Mole Fraction, χ. The increasing trend of the bandgap is shown in two line segments, (A) and (B), corresponding to the
range of χ from 65% to 85% and from 85% to 95%, respectively.
The set of optical gap data was fitted with linear equations, obtaining excellent fits with R2 of
0.98 and 0.99, respectively. The increasing trend of optical gap values is rather gradual for
0
0.5
1
1.5
2
2.5
3
3.5
4
0.45 0.55 0.65 0.75 0.85 0.95
Opt
ical
Ban
d ga
p, E
g (e
V)
Methane Mole Fraction, χ
(A)
(B)
58
methane mole fractions of below χ = 80%; these mole fractions correspond to carbon contents of
x = 0.15 to x = 0.29, when the carbon atomic percent is almost doubled compared to the starting
value. During this first phase, the optical gap increases slowly from 1.06 eV to 1.40 eV, over the
range where carbon content almost doubles. The behaviour however changes drastically as the
methane mole fraction is further increased to values over χ = 85% and higher, a range over which
the carbon mole fraction corresponds to x = 0.35 to x = 0.49 - where the carbon content reaches
the stoichiometric target value of ~0.5. In this second phase, the optical gap value begins at 2.07
eV and reaches a maximum wide bandgap of 3.77 eV.
The low values of the bandgap recorded for a-SiC:H thin films deposited using DCSF-PECVD
were found not to correlate with what has been reported by Vetter, M., et al.; the minimum value
of optical bandgap obtained in this thesis (E9:3;. = 1.06 eV) for a minimum carbon content of x =
0.10 is even lower than the minimum value in the range 1.60 eV – 1.80 eV, obtained by Vetter,
M., et al. corresponding to carbon content of x = 0.0 (pure silane) [46]. Another point of evidence
for lack of correlation of results with literature is the value of optical bandgap for pure silane
deposition, to have amorphous silicon films reported by Saha, J.K., et al. to be around 1.80 eV
[64; this value is well above the 1.06 eV which is the minimum bandgap measured in this
research. We regard this band of low optical gaps as not being truly representative, attributing
the discrepancy with literature reports to the fact that these band of measurements were carried
out on very thin films (sub 100 nm). Alternative optical spectrophotometer measurements were
not possible on these samples owing to the crystalline silicon being single-side polished.
59
Figure 29 An illustration of the increasing trend of optical (Tauc) gap as a function of percent increase in Carbon content of the a-SiC:H films. The trend, as shown, is best approximated by an
exponential regression fit (R2 > 0.96); this proves the Films with increasing carbon lead to the stronger Si – C bonds substituting the Si – Si bonds.
When the carbon content increases beyond x = 0.50, as predicted in [11], the optical gap opens
up even more rapidly as a result of Si – C bonds being replaced by even stronger C – C bonds
[9]. Such strong C – C bonds at the high carbon content of x = 0.49 are believed to be the reason
why the bandgap widens to the maximum value of 3.77 eV, as anticipated that the gap would
increase to values ranging from 3 eV to 4 eV [48], and reaches a record maximum, compared to
previous work (3.7 eV) obtained at a carbon content of x = 0.71 [49], much higher than the
corresponding carbon content value (x = 0.49) reported here. The rate of change in the optical
bandgap as the methane mole fraction is raised is divided into two linear segments (A) and (B)
(see Figure 28). This is not the case when the bandgap is increased as a function of percent
carbon content – in this case the behaviour is best described by an exponential fit (see Figure 29).
It is expected that the beginning of sp2 carbon bonding will result in narrowing of the bandgap,
as the 0.65 carbon content is reached and passed [50]. It seems further that the boost of the
R² = 0.9633
0
0.5
1
1.5
2
2.5
3
3.5
4
0.10 0.20 0.30 0.40 0.50
Opt
ical
Ban
dgap
, Eg
(eV
)
Carbon Content, x
60
bandgap energy as a function of carbon content is highly dependent on the growth of the valence
band energy, as reported in the literature [51]; that as the valence band offset, <=, varies from
0.44 – 1 eV for x = 0 to x = 0.50 carbon content, the band-gap increasing trend tends to follow
the same trend from the valence-band. The fact that the valence band is shifted downwards, as
the carbon concentration increases, reflects the changes observed as the carbon concentration
increases, while the conduction band remains relatively constant.
The refractive index was observed to decrease as the carbon content increases in a-Si1-xCx:H thin
films deposited by DC Saddle Field plasma-enhanced chemical vapour deposition, as expected
from previous work [52, 53]. Figure 30 (c) shows an extreme drop of larger than 1 unit in the
values of refractive index evident as x increases beyond 0.30. The reason for this correlation can
be attributed to the growth of C – C bonds as a result of increasing carbon content, that replace Si
– C cross links throughout the network of a-SiC:H thin films . The decreasing trend saturates at
around n = 2.6, where it corresponds to the stoichiometric Si/C. While Figure 30 (c) shows an
inverse correlation between refractive index and carbon content measured at the larger
wavelength of λ = 633 nm, this trend was not observed in Figures 30 (a) and (b) - measurements
at λ = 248 nm and λ = 365 nm, respectively. In fact Figure 30 (a) shows increase in refractive
index as the carbon content grows beyond 0.30 although close to the stoichiometric value, n
drops down to reach a low value of less than 1.5 units.
61
Figure 30 The change in refractive index, n, as the carbon content x increases. The values are obtained for wavelengths of (a) λ = 248 nm, (b) λ = 365 nm and (c) λ = 633 nm.
0.51
1.52
2.53
3.54
4.5
0.10 0.20 0.30 0.40 0.50
Ref
ract
ive
Inde
x, n
Carbon Content, x
λ=248 nm
0.51
1.52
2.53
3.54
4.5
0.10 0.20 0.30 0.40 0.50
Ref
ract
ive
Inde
x, n
Carbon Content, x
λ=365 nm (b)
0.51
1.52
2.53
3.54
4.55
0.10 0.20 0.30 0.40 0.50
Ref
ract
ive
Inde
x, n
Carbon Content, x
λ=633 nm (c)
62
4.8. Atomic Force Microscopy
The surface roughness of the a-Si1-xCx:H films as a function of the methane mole fraction χ are
shown in Figure 31 (a). The rms roughness is observed to decrease almost exponentially, from
approximately 11 nm to 1 nm rms, as the films approach stoichiometry. It is certain that as the
methane mole fraction increases relative to the total gas mixture, the nature of the plasma is
altered and the methane plasma soon dominates. This likely alters the growth dynamics and
ultimately affects the surface roughness.
It is also interesting to observe that both the decrease in film roughness and decrease in film
growth rate are correlated, which is generally valid for thin film growth. The surface morphology
of the stoichiometric a-SiC:H film as measured using AFM is shown in Figure 31 (b). The
amorphous nature of the film is quite clear from Figure 31 (b), as there are no features, structures
nor grains observable.
0
2
4
6
8
10
0.65 0.7 0.75 0.8 0.85 0.9 0.95
rms
roug
hnes
s (n
m)
Methane mole fraction, χ
63
Figure 31 (a) rms surface roughness as a function of increasing methane mole fraction. (b) The surface morphology of the stoichiometric a-SiC:H film obtained at χ = 93%. The rms surface
roughness of this sample was measured to be 1.014nm.
(b)
64
Chapter 5: Concluding Remarks
5.1. Conclusion
In this study we presented the synthesis of hydrogenated amorphous silicon carbide using DC
Saddle-Field PECVD. Silane and methane were used as precursor gases at a deposition
temperature of approximately 225oC. With the variation in the precursor gas mixture and in
particular the methane mole fraction, stoichiometric a-SiC:H was achieved at a methane mole
fraction of 93% . XPS scans revealed strong Si – C bonds with high counts and no observable
shifts in bonding energies with composition up to and at the point of equal ratios of carbon and
silicon. Optical bandgap was characterized for the samples with increasing carbon content and a
noticeable increasing trend was observed. A maximum bandgap of 3.77 eV was reported, from
Spectroscopic Ellipsometry, at stoichiometric ratio corresponding to precursor gas composition
of 93% methane mole fraction; these results are in general agreement with the literature reporting
synthesis using conventional techniques. Measurements of the refractive index values disclosed a
relatively uniform decreasing trend with increasing C content, in-line with observations from
literature; specifically a drastic drop in the value of refractive index was observed for carbon
content beyond x = 0.30. The extremely low surface roughness of 1.014 nm rms was achieved at
the point of stoichiometry. These results pave the way for using DCSF-PECVD technique for the
synthesis of stoichiometric amorphous silicon carbide films at low temperatures.
Further, annealing experiments did not render any evidence of crystallization for the temperature
examined here, nevertheless a correlation in the increasing carbon content and decreasing silicon
content with increasing methane mole fraction was observed. High temperature annealing of
65
stoichiometric films maintained the relative carbon and silicon content of the a-SiC:H thin film;
further, the annealed film was thermally stable up to 800oC – the upper limit in this study.
The deposition of hydrogenated amorphous silicon carbide with varying carbon content proved a
constant and absolute increase in the optical bandgap of up to 3.77 eV at a C to Si ratio of 1 to 1,
a value slightly lower than the 4 eV bandgap of hydrogenated amorphous carbon (a-C:H),
otherwise known as diamond-like carbon (DLC). Further, upon approaching stoichiometry a
significant decreasing trend in surface roughness was observed with an rms value as low as 1.014
nm rms at stoichiometry. Both of these results exceed the results reported in the literature (higher
optical gap and lower surface roughness) where films were prepared using rf PECVD.
Optical gap values of the material prepared in this research, were compared with the results
reported in the literature and are illustrated in Figure 32. The figure shows the DCSF is able to
yield a stoichiometric amorphous silicon carbide film with an optical gap that exceeds values for
similar films reported in the literature and is remarkably close to optimal DLC value. Further,
the films in this research were synthesized at a low temperature of 225oC and using a technique
amenable to large area deposition.
66
Figure 32 An overview and comparison of achievements in optical bandgap for hydrogenated amorphous silicon carbide films synthesized using different deposition techniques.
5.2. Issues and Constraints
Relatively low deposition rate was found to be a considerable issue, specifically when a certain
minimum thickness was required. One example of this limiting factor was evident in the attempt
to characterize the crystalline morphology of the films using X-ray Diffractometry or Raman
Spectroscopy. The low thickness of the films made it very difficult and in most cases almost
impossible to extract film composition signals from the thin layers.
The reason for this is that in the process of depositing the films the plasma power was kept
constant while the methane mole fraction was increased so as to fundamentally only change one
variable in the experiment. However, with increase methane mole fraction the plasma
RF PECVD: increase in C-rich hydrocarbon gas
(C2H2) [1]
RF Sputtering: increase in RF
power [54]
RF PECVD: increase in CH4 to
98% [2]
RF PECVD: increase in
hydrocarbon to diamond-like a-C:H [31]
RF PECVD: increase in bias voltage up to
140V [9]
RF PECVD: increase in
hydrocarbon gas (CH4) [55]
This Study:
DCSF PECVD
0
0.5
1
1.5
2
2.5
3
3.5
4
1985 1990 1995 2000 2005 2010 2015
Opt
ical
Ban
dgap
, eV
Traces from literature, sorted by year of publication
67
characteristics were dominated by the methane plasma and effectively the production of carbon
radicals was much lower and hence the lower deposition rate.
Another bounding factor in having a sensible deposition rate was the utilization of the triode
configuration in the system. While this configuration was chosen over the pentode, so as to
clearly separate the growth regime from the plasma region in the chamber, and thus aided in
producing a super-smooth surface and uniform morphology it nevertheless required a longer
distance for the radicals to travel from the plasma region proper to the growth surface.
5.3. Future work
One important undertaking as an extension to the project is to increase the deposition rate of the
a-SiC:H thin films. This will not only serve to generate thicker films making characterization
easier, but also lead to the fabrication and synthesis of films with variable thickness for the next
step of device integration. In order to do this, one option is to increase the plasma potential.
Other options are to increase the gas flow rate and to potentially use the pentode configuration
instead of the triode configuration or vary the distance between the plasma region proper and the
substrate in the triode configuration.
A brief campaign of post annealing experiments was performed to test the stability of the
chemical composition under high temperature treatment of up to 800oC. The post-annealing
process can be further extended to higher temperatures of 1000oC and 1200oC, hence studying
the effect of high temperature annealing on the crystallinity of the films. In the event of being
successful in transforming to a crystalline morphology from the amorphous matrix, this would
68
open the path for the application of crystalline silicon carbide through a simple large area
synthesis technique. Further, in this context, it would also be feasible to explore laser annealing
of the film thereby locally crystallizing the film. The versatility that this would provide from a
device perspective is potentially far reaching. More advanced characterization techniques such as
TEM could be applied to samples for higher resolution imaging. In addition, XRD measurements
at lower angles could also provide more promising conclusions about the crystallinity of the thin
films.
Other efforts that would constitute advancement to this project are to use additional
characterization techniques to gain a more in-depth understanding of the properties of the films:
hardness measurements, photo-luminescence (PL), electrical characterization (resistivity,
activation energy, capacitance …).
69
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