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Studies of diamond and diamond-like carbon nucleation and growth on non-diamond substrates by microwave plasma-enhanced chemical vapor deposition

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S tu d ie s o f d iam o n d a n d d iam ond-like c a rb o n n u cleatio n a n d
g ro w th on n o n -d iam o n d s u b s tra te s by m icrow ave p lasm a
e n h an c ed chem ical v a p o r d ep o sitio n
O ng, Tiong Peng, Ph.D .
Northwestern University, 1991
UMI
300 N. Zeeb Rd.
Ann Arbor, Ml 48106
NORTHWESTERN UNIVERSITY
STUDIES OF DIAMOND AND DIAMOND-LIKE CARBON
NUCLEATION AND GROWTH ON NON-DIAMOND
SUBSTRATES BY MICROWAVE PLASMA ENHANCED
CHEMICAL VAPOR DEPOSITION
A DISSERTATION
SUBMITTED TO THE GRADUATE SCHOOL
IN PARTIAL FULFILLMENT OF THE REQUIREMENTS
for the degree
DOCTOR OF PHILOSOPHY
Field o f Materials Science and Engineering
by
Tiong P. Ong
~
EVANSTON, ILLINOIS
DECEMBER 1991
^
A B STR A C T
Studies o f diamond and diamond-like carbon nucleation and growth on
non-diamond substrates by microwave plasma enhanced chemical vapor
deposition
by
Tiong P. Ong
T his thesis encompasses the study of diam ond nucleation and
growth on copper, steel, and quartz surfaces.
W hile chem ical vapor
growth o f diamond on diamond surfaces is well ch aracterized , the
nucleation and growth o f diamond on non-diamond substrates are very
com plicated. Yet in practical applications, there is the need to form
diamond films on various surfaces. The three substrates-copper, steel, and
quartz-had been selected for detailed study because of their unique and
different properties.
Unlike silicon (the most studied substrate for diamond grow th),
copper does not form a carbide. Since its lattice constant is close to that of
diamond, it is pivotal to understand how diamond can be grown on copper
surfaces. Diamond nucleation and growth had been attempted on C u ( lll)
and (100) implanted with carbon at room and elevated temperatures. The
implantation conditions were : ion d o se = lx l0 18 c n r2, energy=60-125 keV.
T he carbon layer im planted at elevated tem perature consisted o f
turbostradc graphite islands (20-30 p m size) with its c-axis perpendicular
to the surface of Cu. In the case of room temperature implantation, the
carbon layer was embedded inside the Cu. It was found that diamond
preferentially nucleated at the edges of the graphite islands.
ii
A simple
lattice model, in which the
< 1 ll> d iam o nd
was parallel <0 0 0 1 >grephite and
the < 1 1 0 > diamond was parallel to the < 1 1 2 0 >graphite. was constructed.
Iron at high temperatures tends to crack hydrocarbon gas to form
a layer of soot on its surface, thus m aking direct diam ond growth
impossible. The use o f
200
A silicon buffer layer proved to be effective in
inhibiting the catalytic effect of iron and in preventing carbon species from
diffusing into the bulk. A novel concept of diamond com posite film s
consisting of diamond crystallites, hydrogenated amorphous carbon and/or
fluorocarbon polymer was invented for coatings on steel surfaces. The
com posite film s adhered well to the substrates even upon imposing a
scratch load of
68
newtons. They could also be bent up to ten degrees and
still remained chemically inert and impermeable to salt solutions.
Finally, quartz was chosen as a substrate so that the optical
properties of smooth diamond film could be studied. A pulsed microwave
plasma technique was invented to deposit smooth diamond films on quartz.
It was found that as the time duration of each cycle (At) and the final
temperature (Tf) decreased, the film surface roughness decreased yet the
num ber of process cycles to form a continuous diamond film (np)
increased. Adherent diamond films as smooth as 50 A had been deposited
using At=20 sec.,Tj=500 °C, and np=20 cycles. The transmission of the
film s were well over 60% in the range of 0.6-2 ^m .
T he optical
absorption edge was 0.225 pm .
Approved:_______________________
Professor R. P. H. Chang
iii
ACKNOW LEDGEM ENTS
First of all, I would like to acknowledge the invaluable guidance
and vision of my thesis advisor. Prof. Robert P.H. Chang. Thanks also to
my thesis examination committee members, Prof. S.A. Barnett, Prof. Y.W.
Chung and Prof. P. Stairs.
I am also greatly indebted to Dr. F.R. Chen, Dr. W.A. Chiou, and
Dr. F.L. Xiong for their assistance in a portion o f the microscopy work.
The help of my colleagues in Prof. Chang research group throughout this
entire work, in particular Kevin Grannen who had spent his precious time
to proofread this thesis, is sincerely appreciated.
I would like to thanks my beloved wife, Felly, for her patience and
inspirations which made the completion of this work possible. Thanks also
to my family for their encouragement and support over the years. This
work is dedicated to my late father, Selamat.
The support o f the Departm ent of Energy, O ffice o f Naval
Research, and EPRI is greatly appreciated. Thanks also to the generous
gift of the Lee foundation over the last three years. I greatly acknowledge
the use of the M aterials Research Center, C enter for E ngineering
Tribology, Analytical Chemistry, and Basic Industrial Research Laboratory
facilities.
iv
TABLE OF CONTENTS
ABSTRACT
ACKNOWLEDGEMENTS
TA BLE O F CONTENTS
LIST OF TABLES
LIST OF FIGURES
1.
2.
ii
iv
v
vii
viii
INTRODUCTION
1.1 Brief history of metastable diamond growth
1.2 Properties and technological applications
1.3 Fundamentals of metastable diamond growth
1.3 .1 The role of atomic hydrogen
1.3.2 Diamond nucleation and growth
1.4 Surface issues related to diamond nucleation
and growth
1 .4.1 Surface modification of single crystal
copper
1.4.2 Surface effects of transition metals on
diamond nucleation and growth
1 .4.3 Surface roughness of diamond films
1.5 Summary of thesis
EXPERIMENTAL
2.1 Microwave plasma reactor
2.2 Capacitively coupled radio frequency plasma
reactor
2.3 Thermal evaporator
2.4 Crystal growth by Bridgman technique
2.5 Film growth techniques
2.5.1 Nucleation and growth of diamond on
carbon implanted single crystal copper
2.5.2 Growth of diamond composite films
2.5.3 Growth of smooth diamond films
2.6 Characterization techniques
2.6.1 Raman spectroscopy
2.6.2 Scanning and transmission electron
microscopy
2.6.3 X-ray diffraction and back-reflection
Laue
v
1
1
4
11
11
18
26
26
27
29
33
37
37
37
41
44
51
51
52
56
58
58
58
59
2.6.4
2.6.5
2.6.6
2.6.7
2.6.8
2.6.9
3.
Scanning tunneling and atomic force
microscopy
UV/Vis/near IR spectroscopy
Tribotester
Three-point bending tester
Scratch tester
Chemical test
RESULTS AND DISCUSSION
3.1
Nucleation and growth of diamond on carbon
implanted single crystal copper
3.1.1 Microstructure of the implanted
carbon layer
3.1.2 Diamond nucleation
3.1.3 Graphite edge chemistry
3.1.4 The epitaxial relationship between
diamond and carbon implanted copper
3.2
Growth of diamond composite films on iron
surfaces
3.2.1 Three-point bending test
3.2.2 Scratch test
3.2.3 Chemical test
3.3
Growth of smooth diamond films
3.3.1 Effect of cycle duration and temperature
on diamond grain size
3.3.2 Raman spectroscopy
3.3.3 Transmission electron microscopy
3.3.4 Effect of film surface roughness on
optical transmission spectra
3.3.5 Tribotest
60
62
62
65
67
68
71
71
71
83
100
109
112
112
123
127
130
130
132
135
146
156
4.
CONCLUSIONS
159
5.
REFERENCES
172
184
VITA
vi
LIST OF TABLES
Page
Table
9
I
The wonderful properties o f diamond and its major
competing materials (after Bachmann, 1990).
II
The deposition conditions for a-C:H and fluorocarbon
polymer films growth.
55
III
The measured and calculated d spacings o f the
diffraction pattern shown in Figure 3.1.8.
82
IV
Transmission electron diffraction data for CVD diamond
film, the reported values o f natural cubic diamond
(ASTM 6-675) and hexagonal graphite-2H (ASTM 23-64).
(Hitachi H-700, 200 kV)
vii
137
LIST OF FIGURES
Figure
page
1.1
The patent activities in the area diamond films collected
from around the world since 1963 (Source : Photonics
Spectra, January 1969, p. 123).
3
1.2
Comparative diagram of thermal conductivity versus
temperature for different types of materials. The shaded
region indicates the temperature range from -25 to
125 °C (Geis et al., 1968).
5
1.3
Electron velocity as a function of electric field for
Si, GaAs, InP, and diamond (Geis et al, 1968).
6
1.4
The theoretical maximum allowable voltage of a
transistor versus cutoff frequency for Si, GaAs, Ge,
and diamond (Geis et al., 1988).
7
1.5
The projected global market for diamond thin film by
1996 (Bachmann and Messier, 1969).
10
1.6
The equilibrium phase diagram of carbon and the
regions where diamond crystals are typically grown
(Bachmann, 1990).
12
1 .7
A schematic diagram showing the effect o f atomic
hydrogen etching on diamond and graphite growth
rate (Badzian and Devries, 1988).
14
1.8
The surface structure of diamond with no hydrogen
termination (Anthony, 1990).
15
1.9
The surface structure of hydrogen terminated diamond
(Anthony, 1990).
17
viii
page
Figure
1.10
The polycyclic aromatic hydrocarbon radicals
known to be the graphite precursor (Anthony, 1990).
19
1.11
The proposed hydrocarbon cage compounds for
twinned diamond crystals:
(a) cubo-octahedron, (a') adamantane ;
(b) twinned cubo-octahedron, (b’> bicyclo [2 .2 .2 ] octane,
(b”) tetracyclo [4 .4 .0 . 13,9. 1 4,8] dodecane ;
(c) decahedral-Wulff-polyhedron, (c’) hexacyclo
[5 .5 . 1 . 1 2 ,6 1 8,12.o3,11.q5,9] pentadecane ;
(d) icosahedron, (d’) dodecahedrane
(Matsumoto and Matsui, 1983).
21
1.12
The preferential nucleation of diamond crystals
on the induced scratches on silicon
(Ong and Chang, unpublished).
22
1.13
A schematic diagram showing the role of atomic
hydrogen during diamond growth (Anthony, 1990).
24
1.14
A plot of the light scattering coefficient (asc)versus
refractive index of sample for three different
surface roughness values. The dotted line represents
the values of ctsc for o = X/100 and A/10 in the case
of n = 2.4 (such as the case in diamond)
(Filinski, 1972).
32
2.1
2.45 GHz microwave plasma CVD apparatus for
growing diamond films.
38
2.2
A schematic diagram of the microwave plasma
CVD reactor.
39
2.3
A schematic diagram of the capacitively coupled
RF (27.15 MHz) plasma reactor for growing a-C:H
and fluorocarbon films.
40
ix
Page
Figure
2.4
The configuration of the capacitively coupled RF
electrodes for : (a) a-C:H and (b) fluorocarbon
polymer films growth.
42
2.5
The JEOL JEE4C thermal evaporator for depositing
a-Si films.
43
2.6
The NRC vertical furnace system for growing single
crystal copper using the Bridgman technique.
45
2.7
A schematic diagram of the graphite mold dimension
for single crystal Cu growth.
46
2.8
The programmed temperature inside the furnace
during the Cu shots melting.
48
2.9
A simplified diagram of the Bridgman technique for
growing single crystal Cu.
49
2.10
The temperature history of the furnace during single
crystal Cu growth.
49
2.11
The back-reflection Laue pattern of Cu( 111)
(W tube, 20 kV, 15 mA, collection time = 2.5 minutes).
50
2.12
The polymerizing and etching conditions of
fluorocarbon plasmas as a function of fluorine
to carbon ratio (Cobum and Winters, 1979).
55
2.13
A schematic diagram of the pulsed plasma processes.
57
2.14
A schematic system diagram of the STM Nanoscope
II unit. (Source : Digital Instruments Inc.)
61
2.15
A schematic system diagram of the AFM Nanoscope
II unit. (Source : Digital Instruments Inc.)
63
2.16
A schematic diagram of a block-on-ring tribotester.
64
x
page
Figure
2.17
A schematic diagram of a three-point bending tester.
66
2.18
A schematic diagram of the Revetest scratch adhesion
tester.
69
2.19
A schematic diagram of the acoustic emission versus
scratch load obtained in a typical adhesion scratch
tester. An abrupt increase in the acoustic emission at
load Per indicates the initial stage of film adhesion failure.
70
3.1.1
Optical micrograph of carbon layer implanted into
Cu(100) surface at 820 °C, 1 xlO *8 dose, and 70 keV.
72
3.1.2
Rutherford backscattering/channeling spectra of
virgin and carbon implanted Cu(100). The conditions
of the ion implantation are : 820 °C, 1 xlO 18 dose,
and 70 keV. (Courtesy of Dr. C. W. White, Oak Ridge
National Laboratory, Oak Ridge, Tennessee).
73
3.1.3
Rutherford backscattering/channeling spectra of
virgin and carbon implanted Cu(100). The
conditions of the ion implantation are : room
temperature, 1 xlO 18 dose, and 70 keV.
(Courtesy of Dr. C. W. White, Oak Ridge National
Laboratory, Oak Ridge, Tennessee).
74
3.1.4
X-ray diffraction of carbon implanted C u ( lll).
The copper signal is obtained by the normal
6/26 scan using 20 kVand 5 mA X-ray Cu K a
line. For the case o f carbon signal, the 6 scan
is carried out using 40 kV and 20 mA.
76
xi
page
Figure
3 . 1 .3
Transmission electron micrographs (bright field)
of carbon layer implanted onto Cu(l 11) taken
using Hitachi H-7CK) at 200 kV : (a) low
magnification , (b) high magnification
(Ong et al., 1991).
77
3.1 .6
Transmission electron diffraction of carbon
layer shown in Figure 3.1.5 in the region which
does not include the fiber-like features.
(Ong et al., 1991).
78
3.1.7
A schematic diagram of (a) ideal and (b) turbostratic
graphite lattices (Hoffman et al., 1991).
79
3 .1 .8
Transmission electron diffraction of carbon layer
shown in Figure 3.1.5 in the region which include
the fiber-like features.(Ong et al., 1991).
81
3.1.9
Atomic force micrograph of the graphite layer cracks
on Cu(l 11). The implantation was done at 820 °C.
84
3.1.10 Scanning electron micrographs of diamond
crystals nucleated on : (a) virgin and
(b) 820 °C C-implanted Cu(l 11) crystals.
85
3.1.11 Scanning electron micrographs of diamond
crystals nucleated on Cu(l 11) which is implanted
with carbon at :(a) 820 °C and (b) room temperature.
87
3.1.12 Raman spectrum of diamond crystals grown
on Cu (111) implanted with carbon at 820 °C.
The diamond deposition conditions are :
3% CF4 , 0.7% 6 2 in 200 seem H2 at 33 T on
and 800 °C. The spectrum was collected using
488 nm Ar laser line.
88
xii
page
Figure
3.1.13 SEM micrographs of diamond crystals grown
on Cu(l 11) implanted with carbon at 120 keV,
lx lO 18 ion dose, and 820 °C.
(0.5% CH 4 , 0 .6 % 0 2 in 2 0 0 seem H 2 at 33 Torr
and 800 oC).
89
3.1.14 SEM micrographs of the twinned diamond crystals
on Cu(l 11) implanted with carbon at 120 keV,
1x1018 ion dose, and 820 °C, shown along with
its proposed precursors (Matsumoto and Matsui, 1983):
(a) twinned cubo-octahedron, (a’) bicyclo [2 .2 .2 ] octane,
(a”) tetracyclo [4 .4 .0 . 13.9 . 1 4,8 ] dodecane ;
(b) decahedral-WulfT-polyhedron, (b’) hexacyclo
[5 .5 . 1 . 1 2 ,6 . 1 8 , 1 2 .0 3 , 1 1 .0 5 ,9 ] pentadecane ;
(c) icosahedron, (c’> dodecahedrane
(0.5% CH 4 , 0.6% O 2 in 200 seem H 2 at 33 Torr
and 800 °C).
90
3.1.15 The Raman spectrum of samples shown in
Figure 3.1.13 and 3.1.14.
91
3.1.16 Preferential nucleation of diamond crystals on the
edges of graphite islands on Cu(l 11) : (a) low
magnification , (b) high magnification. The circle
patch of the graphitic islands were formed by C
implantation through a Ta shadow mask at the
following conditions : 75 keV, 1x10*8 cm _2 dose,
and 840 OC.
93
3.1.17 Scanning electron micrographs showing the distinct
difference in diamond nucleation density on the
(a) prism and (b) basal planes of HOPG crystals.
94
3.1.18 Scanning electron micrographs (close-up views) o f
diamond crystals on the (a) prism and (b) basal
planes of graphite crystals.
95
xiii
Figure
F»gc
3.1.19 Scanning tunneling microscope image of the basal
plane of perfect graphite lattice :
(a) the atomic image
(b) the power spectrum of the 2D Fourier Transform.
The image was obtained using 16.8 mV bias and
1.8 nA setpoint current.
%
3 . 1 .20 Scanning tunneling microscope image o f basal plane
of perfect graphite lattice near the edge of a step
after lm in diamond growth in 3% CF4 , 0.7 % 6 2
and H 2 plasma environment at 33 Torr and 800 °C :
(a) top view, (b) tilted view.
98
3.1.21 Atomic force microscope image o f basal plane of
perfect graphite lattice near the edge of a step after
20 min diamond growth in 3% CF4 , 0.7 % O 2 and
H 2 plasma environment at 33 Torr and 800 °C :
(a) top view, (b) tilted view.
99
3.1.22 Probabilities of producing methane and acetylene
by reaction of graphite with atomic hydrogen.
(Balooch and Olander, 1975).
102
3.1.23 Results of atomic H etching of HOPG crystals in
microwave plasma (H 2 : 20 seem ; He : 20 seem ;
power : 100 W ; 10 Torr ; 100 - 400 ° C ) :
(a) STM image before etching ; (b) STM image
after 1 min. etching (b) STM image after 3 min.
etching ; (d) AFM image after 5 min. etching.
104
3.1.24 (a) A schematic diagram of the crystallographic
orientation of hexagonal lattice in graphite ;
(b) The relation between the hexagonal etch pits and
twin planes in graphite lattice (Thomas, 1965).
106
xiv
Figure
page
3.1.25 The elementary mechanism for the hydrogenation of
graphite via the attack of the atoms on the {1 lJL }
plane ("arm-chair”) (after Tomita and Tamai, 1974).
107
3.1.26 The elementary mechanism for the hydrogenation
o f graphite via the attack of the atoms on the {10lL}
plane ("zig-zag") (after Tomita and Tamai, 1974).
106
3.1.27 A simple model of diamond nucleus on the edge
of graphite basal plane.
< lll> d ia m o n d // <0 0 0 1 >oraphite i
< 1 1 0 >diamond H the < 1 l 2 0 >graphite-
110
3.2.1
SEM micrographs of diamond crystals grown on
carbon steel substrate.
113
3.2.2
Cross sectional view of diamond composite films
on carbon steel substrate.
114
3.2.3
SEM micrographs of continuous diamond film
grown on carbon steel substrate. The white markers
indicate the spontaneous cracking of the film upon
sample cooling.
115
3.2.4
Schematic diagrams of the multilayer structure of
diamond composite films : (a) a-C:H, fluorocarbon,
and diamond on 304 stainless steel, (b) a-C:H and
diamond on 304 stainless steel, (c) fluorocarbon
and diamond on 304 stainless steel.
116
3.2.5
Plot of bending load versus sample displacem ent:
(a) bare 304 stainless steel, (b) a-C:H, fluorocarbon,
and diamond on 304 stainless steel,
(c) a-C:H and diamond on 304 stainless steel
(d) fluorocarbon and diamond on 304 stainless steel.
117
xv
page
Figure
3.2.6
Optical micrographs of a-C:H, fluorocarbon, and
diamond coated 304 stainless ste el:
(a) before and (b) after the bending tests.
119
3.2.7
Optical micrographs of a-C:H and diamond
coated 304 stainless ste e l: (a) before and
(b) after the bending test.
120
3.2.8
Optical micrographs of fluorocarbon and diamond
coated 304 stainless ste e l: (a) before and (b) after
the bending test.
121
3.2.9
SEM micrographs of scratch tracks of a-C:H
and diamond composite film coated 304 stainless
steel with the stylus load : (a) 5 N, (b) 20 N,
(c) 29 N, (d) 39 N, (e) 49 N, (0 6 8 N.
The arrow markers indicate the direction of
diamond stylus motion.
124
3.2.10 Schematic diagrams of (a) conformal and
(b) tensile cracking failure modes in thin film
after scratch adhesion tests
(after Burnett and Rickerby, 1987).
126
3.2.11 The results of four-hour-corrosion tests on :
(a) a-C:H and diamond coated carbon steel and
(b) bare carbon steel in a solution of 5 g NaCl
and 1 0 0 cc water.
129
3.3.1
131
Dependence of average diamond grain size on the
process cycle time interval At and the final
temperature Tf when the plasma is turned o f f :
(a) At = 60 min.,Tf = 800 °C, Tavg = 800 °C , one cycle
(b) At = 4 min.,Tf = 800 °C, Tavg = 600 °C, four cycles
(c) At = 45 sec., T f = 600 °C, Tavg = 415 °C, 16 cycles.
xvi
page
Figure
3.3.2
Scanning electron micrographs of a very smooth
diamond film : (a) low magnification ;
(b) high magnification grown at At = 20 sec.,
Tf = 500 °C, 20 cycles.
133
3.3.3
Surface profile of diamond film shown in Figure 3.3.2.
133
3.3.4
Raman spectroscopy of diamond thin films grown at
Tf = 800 °C : (a) 4 cycles, (b) 1 cycle,
(c) diamond powder.
134
3.3.5
(a) Transmission electron micrograph of the diamond
film (bright field : 200 kV), (b) The corresponding
transmission electron diffraction
(taken by Hitachi H-700 TEM).
136
3.3.6
High resolution TEM micrograph of a single crystal
diamond (taken using Hitachi H-9000 HREM,
Ong et al., 1990).
139
3.3.7
High resolution TEM micrograph of a single crystal
diamond oriented in a [110] direction. The lattice image
of { 1 1 1 } plane is clearly observed with lattice spacing
of 2.06 A : (a) image of the perfect lattice , (b) image
of the defected lattice (twins and stacking faults)
(Ong et. al, 1990).
140
3.3.8
A detailed HREM view of a primary twin plane
(£ = 3) in diamond crystal twin plane along [011]
projection (the twin plane is { 1 1 1 }):
(a) experimental results ( Ong et al., 1990)
(b) calculated structure (Narayan, 1990).
141
3.3.9
An HREM micrograph of fivefold twin in
diamond crystal taken along [0 1 1 ] projection
(the twin plane is { 1 1 1 })
(Ong et al., 1990).
143
xvii
page
Figure
3.3.10 The schematic diagram of five fold twinned
diamond crystal and the hydrocarbon cage
compounds proposed as the precursors :
(a) decahedral-Wulff-polyhedron,
(b) hexacyclo [5 .5 . 1 . A ® i8,12.o3,11.o5,9]
pentadecane, (c) icosahedron, (d) dodecahedrane
(Matsumoto and Matsui, 1983).
144
3.3.11 A proposed model of a fivefold twin diamond
shown along [110] projection based on 4 (111)
twins (Devries, 1987).
145
3.3.12 Optical transmission spectra o f :
(a) 1 -mm-thick quartz
(b) 0 .9 2 p m diamond coated quartz (roughness =
(c) 1 -mm-thick type Ila natural diamond (*)
(d) 1.25 p m diamond coated quartz (roughness =
(* KJocek et al., 1968).
147
200
A)
2000
A)
3.3.13 SEM micrographs showing the crystal
morphology and grain size of samples shown
in : (a) Figure 3.3.12 ( b ) , (b) Figure 3.3.12 (d).
149
3.3.14 Optical transmission spectra of 0.3 p m thick
diamond film on 1 -mm-thick quartz.
150
3.3.15 Optical transmission spectra of 0.4 p m thick
diamond film on 1 -mm-thick quartz.
151
3.3.16 Optical transmission spectra of 0.5 p m thick
diamond film on 1 -mm-thick quartz.
152
3.3.17 Optical transmission spectra of 0.8 p m thick
diamond film on 1 -mm-thick quartz.
153
3.3.18 SEM micrographs of samples shown in Figure
3.3.14 (a and a ') and Figure 3.3.17 (b and b ’).
154
xviii
Figure
3.3.19 Plot of film thickness versus deposition time.
Note that the deposition time is taken as the
number of cycles times the cycle period (At).
155
3.3.20
Results of block-on-ring tribotest on bare and
diamond-coated quartz : (a) optical micrograph
of the grinding track on bare quartz, (b) surface
profile across the grinding track of sample (a),
(c) optical micrograph of the grinding track on
0.35 ftm diamond-coated quartz, and (d) surface
profile across the grinding track o f sample (c).
The small arrow indicates the materials transferred
from the steel bearing.
157
3.3.21
SEM micrograph showing the bearing materials
which are stuck on the diamond surface after
30 min tribotest.
158
xix
1.
1.1
B rief
INTRODUCTION
History o f Metostable Diamond Growth
T he first successful work metastable diamond growth, in which
diamond is synthesized in the graphite stability region, was started in 1953
at the Union Carbide Corporation by W.G. Eversole (Angus and Hayman,
1988, ref. 16)
he was later granted a patent in 1962 after m a n y
successful repetitive experiments (Eversole, 1962). Interestingly, the first
successful growth o f diamond at high pressure and high temperature
environment occured not until 1955 (Bundy et al., 1955).
Later in 1956, a
Russian group led by B. Deryagin initiated the research on low pressure
diamond growth. They started by growing diamond “w hiskers” using a
metal-catalyzed vapor-liquid-solid process. The group later studied the
epitaxial growth of diamond and formulated the theory o f re la tiv e
nucleation rates of diamond with respect to that of graphite (Deryagin and
Fedoseev, 1970).
It was not until 1968 when John Angus et al. (1968) revived the
research on metastable diamond growth in the United States. His group
was the first to report the preferential etching o f g rap h ite,
codeposits during the diamond growth, by atomic hydrogen.
w hich
He also
discovered some unusual catalytic phenomenon of boron on low pressure
diamond growth.
1
2
The LEED studies conducted later by Lander and Morrison (1966) further
emphasized the significant role of hydrogen in the metastable diamond
growth.
T heir
experim ental
results revealed the fact that a to m ic
hydrogen stabilizes the ( 1 1 1 ) diamond surface by satisfying the dangling
bonds at the outerm ost layer of carbon atom s, resulting
unreconstructed l x l LEED pattern.
in an
Studies on the beneficial role of
atomic hydrogen continued by Chauhan et al. (1976) and Fedoseev et al.
(1977). They found that atomic hydrogen enhances the growth rate of
diamond , but suppresses that of graphite.
Up to that point, the technology
of diamond growth was still not commercially feasible due to the slow
growth rate.
In the mid 1980s a group at the National Institute for Research in
Inorganic Materials in Japan, led by Nobuo Setaka (Matsumoto et al., 1982,
1985, and 1987 ; Kamo et al., 1983 ; Matsui et al. 1983), invented the hot
filament and plasma assisted ( RF, microwave, and DC ) chemical vapor
deposition techniques to grow polycrystalline diamond on non-diamond
substrates at few microns per hour growth rate. These fantastic results
sparked the interests from around the world, and from thereon the number
of research publications in this field started to grow in an immensely
accelerated rate (see Figure 1.1).
Up to this point, Japan is still the
unequivocal leader in the technological applications of diamond thin films.
3
500
400
J
300
US
JAPAN
200 H
OTVER
100
1963 -1977
1978 -1982
1983 -1987
Y«ar
Figure 1.1 The patent activities in the area diamond films collected from
around the world since 1963 (Source : Photonics Spectra, January 1989,
p. 123).
4
1 .2
Properties and Technological Applications
Diamond has a range of remarkable properties.
The technology
and applications of natural and synthetic diamond powders, compacts, and
crystals have been explored and utilized in many different fields.
The
thermal conductivity of diamond, for example, is superior to that o f any
other material
including copper (see Figure 1.2).
A t the same time,
diamond is also an excellent electrical insulator, which makes diamond such
a unique material. Therefore, the use of diamond in printed circuit board
will solve the existing heat dissipation problem in very-large-scaleintegrated (VLSI) circuits.
The inherent large bandgap of diamond (5.45 eV) is a great
advantage in that its electrical properties are relatively insensitive to
temperature and radiation damage. Its high electric field breakdown and
free electrons velocity, as shown in Figure 1.3 and Figure 1.4, makes
diam ond a serious challenger to GaAs and Si in high -frequency, highpower device applications.
The superb hardness and chemical inertness of diamond, which is
probably the best of all known materials, open a new exciting era in
protective coating technology. Combined with its high optical transparency
from IR to near UV, diamond films are excellent choices for protecting
various materials against wear and corrosion.
5
1000
*
300
Cu
Diamond
100
30
7819
GaAa
Diamond
30
Cu
GaAa
100
300
1000
Tamparatura (K)
Figure 1.2 Comparative diagram of theimal conductivity versus
temperature for different types of materials. The shaded region indicates
the temperature range from >25 to 125 °C (Geis et al., 1988).
6
VELOCITY
( l o ’ c m /a )
D IA M O N D
OaAt
10*
10'
10‘
10*
ELECTRIC FIELD | V / e m |
Figure 1.3 Electron velocity as a function o f electric field for Si, GaAs,
InP, and diamond (Geis et al, 1988).
DIAMOND
10 *
o«
MAXIMUM
ALLOWABLE
VOLTAGE
(V)
7
S i PST
10
1
10
10 *
10 *
104
CUTOFF FREQUENCY FT (GHx)
Figure 1.4 The theoretical maximum allowable voltage of a transistor
versus cutoff frequency for Si, GaAs, Ge, and diamond (Geis et al., 1988).
8
Diamond is also very transparent in the x-ray region due its low
atomic number (Z= 6 ). Therefore, x-ray windows and lithography masks
have been made and marketed recently by C rystallum e Com pany in
California. Diamond high-fidelity loudspeakers have also been fabricated
by Hitachi and Sumitomo utilizing the fact that diamond possesses the
highest sonic velocity (18.2 km/sec) among any known materials. T he
high-frequency sound distortion limit up to 50 kHz has been obtained on
alumina diaphragm coated with diamond film.
Table I summarizes some of the properties of diamond along with
its potential applications. These various potential exotic applications have
been the major driving force in attracting interests from around the world.
From the study conducted by the International Resource Development Inc.,
the global market for diamond thin films in 19% is projected to be around
1
billion dollar, with 60 % of the market will be in the area of
semiconductor (see Figure 1.5).
9
Table I.
The wonderful properties of diamond and its major competing
materials (after Bachmann, 1990).
Properties of Diamond
Victor hardness
(kg/mm2 )
10,000
Coefficient of
friction
Young's modulus
(N/mZ)
(in air)
1.2 xlO iZ
Sound propagation
velocity (km/sec)
Chemical inertness
Tiansmioance (ja n)
0.1
18.2
inert
6.11-2.5
and >6
Refractive index
14l
Band gap (eV)
5.45
Electron/bole
mobility (cm^/V-s)
1666/
1600
Dielectric constant
W
Thermal
conductivity
(W/cm-K)
16
Thermal expansion
coefficient (i/K )
Comparison with
Competitors
hardest material
known
very low in air
twice the value of
AI2 Q3 ; highest
mechanical
strength
1.4 tunes the
value of AI2 O3
at R.T. resistant to
all adds, bases
and solvent
in the IR orders of
magnifurif higher
than other
material
1.6 times the
value of Si0 2
1.1 for 41;
1.43 forGaAs;
3.0 for 0-SiC
1500/600 for
S i; 8500/400
forOaAs
11 (or Si
12.5 for OaAs
value for type 11a
Potential Applications
drill bits,poiisbing material,
cutting tools, sintered or
brazed diamond compacts
wear resistant coalings on
windows and bearings
stiff membranes for
lithography masks,
lightweight coatings for
audio devices
coatings for reactor vessels
UV-VIS-IR windows and
coatings, microwave
windows, optical filters,
optical waveguides
high power electronics, lugb
frequency semiconducting
devices, not thermistors, hot
transistors, lasers, detectors
insulating heat sinks for
electronic devices
room temp, is 4 x
the value or Cu or
0 .8 x 10-6
value at room temp.
doeetoSi0 2 value
of 0.57x10^
thermally stable substrates,
e.g. for X-ray lithography
10
By 1996, sem iconductors may take 60%
of worldwide diamond thin film market
S Millions
700
500
19M market for
. diamond thin films
□ Expaetad Ineraasa In
— diamond thin film marfcat
by 19M
________
500
400
300
200
100
0
O
JZZL
Figure 1.5 The projected global market for diamond thin film by 1996
(Bachmann and Messier, 1969).
11
1 .3
Fundamentals o f Metastable Diamond Growth
The review presented here is given for general readers only. It is
never intended to be comprehensive. Detailed reviews have been covered
elsewhere (Angus and Hayman, 1988 ; Spear, 1987 ; Badzian and Devries,
1988 ; Badzian et. al, 1968 ; Bachmann, 1990 ; Anthony, 1990).
1 .3 .1
The role of atomic hydrogen
Metastable growth of diamond crystals is a process where diamond
is nucleated at the graphite stability region.
From the carbon phase
diagram, this growth region is shown by the empty circle in Figure 1.6.
Here the presssure is well below atmosphere and the temperature is in the
range o f 800 - 1000 °C.
Diamond is thermodynamically unstable with
respect to graphite. Hence the free energy of diamond is higher than that
of graphite. In other words, for a given partial pressure of hydrocarbon
(e.g. methane), the supersaturation obtained on the diamond surface is
always lower than that on graphite (Fedoseev et al., 1984). Yet, metastable
diamond growth is not a forbidden phenomenon. Certainly it is expected
that graphite deposits are also obtained during metastable diamond growth.
However,
the etching rate of graphite by atomic hydrogen is higher
relative to diamond, making high quality diamond growth possible. The
difference in reactivity to atomic hydrogen between diamond and graphite
stems from the fact that atomic hydrogen significantly affects the nature of
12
\
•
\
\
\
Pressure, Psi x 10
•
\
\
i\
\
Diamond
Direct Metbo
»
Solvetit Mediiod "
1
Melt
\
\
dA-
Graphite
0
1000
2000
3000
2
4000
5000
6000
Temperature (K)
Figure 1.6 The equilibrium phase diagram of carbon and the regions
where diamond ciystals are typically grown (Bachmann, 1990).
13
electronic bonds in the aromatic rings of graphite, while there is practically
no effect in the case o f diamond.
This kinetic factor of preferential
gasification of graphite by atomic hydrogen is the fundamental factor in
graphitic-free diamond growth. Sustaining a superequilibrium environment
of atomic hydrogen concentration is a must in preventing the formation of
the stable graphite phase during diamond growth. This is crucial not only
for purifying diamond, but also for avoiding the poisoning effect on
diamond growth by the graphite. A schematic diagram, shown in Figure
1.7, is to illustrate the importance of atomic hydrogen on diamond growth
at low pressure.
Atomic hydrogen is also known to stabilize the sp 3 bonding on
diamond surface.
Several studies (Lander and J. Morrison, 1966 ; Lurie
and Wilson, 1977 ; Haneman, 1961 ; Chadi, 1982 ; Pandey, 1982 ; Himpsel
et. al., 1980 ; Pepper, 1968) found that the surface of diamond is known to
reconstruct in several version at temperature in the range of 900 - 1000
°C. Results from LEED experiments (Pate, 1986) reveal that the H-free
diamond surface reconstructs to 2 x2 (2 x 1 ) from lx l pattern at temperature
exceeding 950 °C. The reconstructed surface is sp2 hybridized in nature
according to the dimerized chain model (Pandey, 1982) as schematically
shown in Figure 1.8 . Although it consists of a ji bonding network, the
nature of the bonding on this reconstructed diamond surface bears no
similarity with that of graphite. Pate's work (1986) further reveals that a
lx l surface reconstruction is obtained on H-terminated (111) diamond
14
Only Methane
Growth
Only Hydrogen
F w h iw y
Methane «■ Hydrogen
Growth Etching
Growth
Race
+
Figure 1.7 A schematic diagram showing the effect of atomic hydrogen
etching on diamond and graphite growth rate (Badzian and Devries, 1988)
15
DIAMOND SURFACE WITH NO ATOMIC HYDROCEN
GRAPHITE-UKE
SURFACE
o
—
=
-
CARBON
SINCLE BONO
DOUBLE BONO
DIAMOND
LATTICE
Figure 1.8 The surface structure of diamond with no hydrogen
termination (Anthony, 1990).
16
surface.
Experim ental results confirm that each carbon atom on the
surface bonds to one hydrogen adatom (see Figure 1.9). The bond is 1.09
A long and sp 3 hybridized (Vidali et al., 1983). This H-stabilized sp 3
bonded surface ensures the conditions for continuing diamond growth and
prevents the graphite formation which is very favorable in the case of the
reconstructed surface.
17
OIAMOND SURFACE WITH ATOMIC HYDROGEN
(HI o -
HYDROGEN
CARBON
|
ATOMIC
DIAMOND
LATTICE
Figure 1.9 The surface structure of hydrogen terminated diamond.
(Anthony, 1990).
18
1 .3 .2 Diamond N udeation and Growth
Based on theory of nucleation (Venables et al., 1984), the creation
of a new phase originates from the formation of a nucleus of critical size.
The stability of this new phase depends on the size of this nucleus. The size
of this nucleus must exceed a critical value so that the surface energy
contributions will be insignificant com pared to the volum e energy
contributions. The competition between the volume and surface energies
dictates whether the newly created nucleus will be stable enough to grow
further.
According to this theory, the nucleation density depends on
temperature, supersaturation, and the nature of the substrate surface.
In the metastable synthesis of diamond it is imperative to maintain a
deposition condition in which subcritical diamond nuclei can continue to
grow above critical size necessary for achieving further stable growth. Yet
at the same time, it is absolutely necessary to suppress the subcritical
graphite nuclei from achieving that particular favorable growth condition.
The presence of atomic hydrogen prevents the formation o f polycyclic
aromatic hydrocarbon (PAH) radicals (Figure 1.10) which are known to be
the graphite precursors. The discovery of incorporating atomic hydrogen
to the following reaction :
Heat and atomic Hydrogen
CH 4 + H2 ..................................>Diamond + 2 H2
( 1 . 1)
19
Figure 1.10
The polycyclic aromatic hydrocarbon radicals known to
be the graphite precursor (Anthony, 1990).
20
is a tremendously important contribution to the success o f metastable
diamond growth . Atomic hydrogen also significantly reduces the surface
energy o f diamond nuclei by chemisorbing on diamond surface. This
thereby reduces the critical diamond nuclei to only a few atoms. Several
hydrocarbon cage compounds have been proposed by M atsumoto and
M atsui (1983) as the diamond embryos. The lattice structures o f these
compounds along with diamond are shown in Figure 1.11. All o f these
cage compounds are saturated hydrocarbons. The adamantane structure
consists of the smallest diamond skeleton and is believed to be the embryo
for a "perfect" diamond crystal.
The other four cage com pounds are
suggested to be the embryos for twinned (cubo octahedron, D -W ulffpolyhedion, and icosahedron) diamond crystals.
Diamond nucleation on diamond substrate can be relatively easy to
achieve and results in single crystal film (homoepitaxy) when optim um
growth conditions are employed.
Diamond does not readily nucleate,
however, on non-diamond substrates such as Si, Mo, W, Au, Cu, SiC>2 , etc.
Pretreatment of the substrate surfaces is found to be necessary. It is only
when the substrates are roughened by diamond or SiC pow der can
enhanced diamond nucleation be obtained, as shown in Figure 1.12. In
contrast, rubbing the substrate (e.g. Si) with plastic, a diamond-tipped
scriber or aluminium has hardly any effect at all on the nucleation density
(Chang et al., 1988). Similarly, silicon substrates roughened with various
plasm a and chemical etching treatment result in no enhancem ent of
(O')
(b‘)
(h")
(«')
Id')
JO Q ©
Figure 1.11
The proposed hydrocarbon cage compounds for twinned
diamond crystals : (a) cubo-octahedron, (a') adamantane ; (b) twinned
cubo-octahedron, (b’) bicyclo [2 .2 .2 ] octane, (b”) tetracyclo
[4.4.0.13 -9 .14-8] dodecane ; (c) decahedral-Wulff-polyhedron,
(c’> hexacyclo [5.5.1.12 A 18 . 12.03.11.O5.9] pentadecane ;
(d) icosahedron, (d’> dodecahedrane (Matsumoto and Matsui, 1983).
22
Figure 1.12 The preferential nucleation of diamond crystals on the
induced scratches on silicon (Ong and Chang, unpublished).
23
diam ond nucleation. Iijim a et al. (1990) recently showed from high
resolution electron microscopy experiments that by rubbing the substrates
with diamond powder, enhancement of diamond nucleation comes from the
fact that some diamond residues are left on the surface o f the substrates
even after the best surface cleaning procedure is performed. Therefore,
the process is essentially homoepitaxial growth in nature.
The current method of CVD diamond growth em ploys hydrogen
concentration in excess o f 95 %. The generation of a superequilibrium
concentration of atomic hydrogen can be easily achieved by hot filament
and plasma (RF, microwave, and DC) methods.
This atomic hydrogen
saturates the entire diamond surface as depicted schematically in Figure
1.13.
However, atomic hydrogen may collide with the surface and grab a
surface hydrogen to form an H 2 molecule resulting in a vacant site which
can be refilled again by an atomic hydrogen from the surrounding gas
(Figure 1.13 c). Occasionally a carbon radical may also hit the vacant site
and form a sp 3 carbon-carbon bond. This carbon radical is usually
generated by a reaction o f hydrocarbon molecules with atomic hydrogen.
The exact nature of the radical which contributes to diamond growth is still
a mystery.
Some theoretical and experimental results hint that methyl
and/or acetylene radicals are the important players. The mechanism
proposed by Tsuda etal. based on quantum chemical calculations (Tsuda et
al., 1986) utilizes methyl cation as the precursor for diamond growth. The
mechanism involves two steps : ( 1 ) diamond ( 1 1 1 ) surface is covered with
24
H H HHH
HH
HHH
DIAMOND
<S?X j3©
' \ /
H H H .
H H H H
iH f
S
H H H ' H H
H H H
DIAMOND
H HCC:<
HY
X
I
I I I
H
H H
H
I
I I
I
DIAMOND
Figure 1.13 A schematic diagram showing the role o f atomic hydrogen
during diamond growth (Anthony, 1990).
25
m ethyl groups ; ( 2 ) three neighboring m ethyl
g ro u p s
in teract
spontaneously via methyl cation to form diamond bonding structure.
The
major flaw of this proposed mechanism is that it does not support the
experimental observations. The important role of atomic hydrogen does
not come into the proposed scheme.
In addition, the population of the
methyl cations in the plasmas and hot filament environment where diamond
is usually grown is rare (Boenig, 1982).
The Tsuda et al. mechanism was later challenged by Frenklach and
Spear (1988). They hypothesized that the main monomer growth species is
acetylene.
The proposed scheme accounts for two m ajor alternating
reaction steps, namely the activation of surface carbon by hydrogen atom
abstraction and the addition of monomers (acetylene molecules ) formed in
the gas phase. There is indeed strong evidence that acetylene is the most
stable gaseous species in high temperature pyrolysis (Duff and Bauer, 1962
; Khandelwal and
Skinner, 1981), combustion (Glassman, 1986), and
plasma (Boenig ,1982) environments.
The m echanism considers the
superequilibrium of atomic hydrogen as the key factor in the initiation and
propagation of diamond reaction synthesis.
In general, the mechanism
agrees in large extent to the various reported experimental observations
listed as follow s:
1)
Gas activation by some means (plasmas or hot filament) is required.
However, the quality of the film is relatively independent of the
26
method of activation.
2)
Starting material is not important. Methane, ethane, propane,
methanol, ethanol, acetone etc. have been used to grow diamond
films.
3)
Superequilibrium concentration of atomic hydrogen is a must for
high quality diamond growth. Small amount of oxygen enhances
both the growth rate and quality of the grown films.
4)
The predominance population of acetylene in diamond synthesis
environment is undoubtable.
5)
There is an optimum temperature range where diamond growth rate is
highest. This can be accounted by the competitive effect of diamond
and graphite growth.
1 .4
Surface Issues Related to Diamond Nucleation and
G row th
1 .4 .1
Surface M odification o f Single Crystal Copper
The unique properties o f diamond make it one of the most
desirable material for high power electronic devices (Geis et al., 1988).
Yet, the technology of high quality single crystal diamond film growth thus
far is only limited to homoepitaxy (diamond on diamond) (Nakazawa et al.,
1987).
Diamond nucleating on non-diamond substrate is very difficult due
to the extremely high surface energy of diamond ( Y (lll) ~ 5.3 J/m 2 ;
Y(100) ~ 9.2 J/m2) (Field, 1979). The heteroepitaxial diamond growth is
certainly a challenge that is currently being actively pursued.
An attempt
27
to grow epitaxial diamond film on single crystal nickel substrate (lattice
constant (a= 3.5239 A , c.f. a<jiamond= 3.5671 A ) had been carried out
(Belton and Schmieg, 1989) with no success. The reason is that nickel
preferentially turns the carbon gas species into soot or graphitic-like
carbon structure (Banaijee et al., 1961 ; Trimm, 1963). Koizumi et. al.
(1990) reported that epitaxial diamond films can be grown on single crystal
c-BN substrate (a = 3.616 A ). However, a large single crystal of this
substrate material is even harder to come by than diamond itself.
The
largest c-BN crystal that can be synthesized so far is only 100 /<m. Copper
(a = 3.6148 A ) is considered to be the next best candidate for the substrate.
Unlike Ni, Mo, W, Fe, Si etc., copper is not known to form carbide when
it interacts with carbon. Nonetheless, the surface energy of copper, as of
most other materials, is very small compared to that of diamond
( y c u (IOO)
= 2.08 J/m 2) ( Richter et al., 1985). It is therefore very interesting to
carry out some initial studies on the surface preparation of this material to
alter the surface energy of this material for growing single crystal diamond
films.
1 .4 .2
Surface Effects o f Transition Metals on Diamond
Nucleation and Growth
There are several obstacles that one has to confront in attempting to
deposit diamond films on metal substrates which contain transition elements
such as iron, cobalt, nickel, or platinum. One of the major problems is the
widely known catalytic effect of those elements in forming coke from
28
carbon containing gases at elevated temperature (Banarjee et al., 1961 ;
Trimm, 1983).
The deposits range from polycyclic arom atic rings to
disordered carbon in structure and usually in the form o f laminar graphite
(Presland and Walker, 1969), non-oriented carbon (D erbyshire and
Trim m , 1975), and fibrous carbon (Baker et al., 1972) film s. The
formation of these deposits significantly affects the nucleation of diamond
particles on such metal substrates. Thus far no satisfactory explanation has
been given to account for the catalytic effect of these materials, although
one can speculate that it might be due to the high electron affinity which
originates from the unfilled d-band election shell in the metals. Moreover,
the rate of carbon diffusion in iron is known to be so high
~
1 0 -8
( D c ..> fc at a o o ° c
cm^/sec ) in such a way that the carbon species from the diamond
forming plasma environment are consumed significantly by the material.
Differences in thermal properties between the substrate and the
deposited film can introduce severe stress at the film substrate interface. It
is commonly known that the thermal expansion coefficient of most metals,
including transition elements, far exceeds that o f diamond at the growth
temperature giving rise to spontaneous spallation failure between the film
and the substrate.
In addition, diamond films alone cannot provide a
suitable protective coating for sheet metal owing to the stiffness of diamond
(E=1054 GPa) (Mcskimin and Andreatch, 1972). It is pivotal to examine
this issue, since in today's technological world sheet steel or stainless steel
constitutes a significant portion of total metal production owing to its vast
29
demand for various products.
1 .4 .3 Surface Roughness o f Diamond Films
A s mentioned previously, diamond possesses very good optical
transparency ranging from infrared to ultraviolet wavelength.
Combined
with its high strength, diamond films are excellent protective coating
materials for optical elements. However, the grown films on non-diamond
substrates are polycrystalline in nature with relatively large grains (1-50
/<m), resulting in very rough surfaces.
It is well known that material
imperfections can affect the optical spectra of a sample in many ways. The
effects o f surface roughness (Beckmann and Spizzichino, 1963 ; Berreman,
1967 ; Bennet and Portens , 1969 ; Beaglehole and Hunderi, 1970 ; Hunderi
and Beaglehole, 1970), surface layer damage (Thomas, 1960), and optical
inhomogeneities of materials (Bruevich, 1970) on their transmission and
reflection spectra have been well documented.
From the elem ents of optical theory, specular reflection from a
perfectly smooth surfaces is due to the constructive interference of
wavelets coming from the first few Fiesnel zones. Provided the surface is
su fficien tly
large to cover the Fresnel zones, the interference is
constructive in a specific direction governed by Snell's law o f reflection.
The intensity o f light reflections in all other directions is zero due to
destructive interference.
However, the reflected and transm itted light
30
wavelets from rough surfaces and grain boundaries are scattered in various
directions leading to the so called diffuse scattering.
A very simple criterion for rough surface based on optical ray
theory was suggested by Raleigh (see Beckmann and Spizzichino, 1963).
This well known " Raleigh criterion" states that a surface is considered
smooth if
Ah < X / 8 sin y
(1.1)
where :
Ah = height o f surface irregularities
X = wavelength
Y = grazing angle (a function of refractive index of material)
Therefore, for the case of electromagnetic waves in the visible region ( X ~
6000 A ),
Ah must be less than 800 A in order for a surface to be
considered smooth, assuming normal incident waves (i.e. sin y ~1).
Filinski (1972) had formulated the effects of surface roughness on
the reflection and transmission spectra of a sample. The concept o f constant
phase plane roughness (i.e. fluctuations in the plane of constant phase) was
used in determining how much light will be reflected or transmitted into
the directions other than those which are expected for an ideally smooth
and perfectly constant phase plane.
Assuming a Gaussian distribution of
31
surface roughness, the transmission of a slab o f sample after being
modified for the surface roughness factor (neglecting multiple reflections)
is given in Equation (1.2) and (1.3).
Ir = lo d -R s )2 expC-aeffd)
(1.2)
oteff = d o + <»sc = oto +(2/dX2* o (n-no)/Ko)2
(1.3)
where :
Ir
transmission for rough surface
Io
incident light intensity
Rs
reflection coefficient for smooth surface (Fresnel coefficient)
o e ff
effective absorption coefficient
ao
absorption coefficient from smooth surface
Osc
absorption coefficient from rough surface
d
sample thickness
o
r.m.s. value of surface roughness ( I
Ah
surface irregularities height
A h 2 )l/2
vacuum wavelength
n
refractive index of sample
no
refractive index of surrounding
Equation (1.3) reveals that the effective absorption coefficient strongly
depends on a and n of a sample (i.e. Oeff is proportional to (on/X)2), as
schematically plotted in Figure 1.14. For a 1 p m thick sample with n equal
to 2.4 (e.g. diamond), there is a factor of ~ 2.75 difference in the intensity
32
n0)a \
I
$
77
1213U15 10
J « 5 6 810
refractive index o f sample —
►
Figure 1.14 A plot of the light scattering coefficient (a*c)versus
refractive index of sample for three different surface roughness values.
The dotted line represents the values of a K for o = X/100 and X/10 in the
case of n = 2.4 (such as the case in diamond). (Filinski, 1972).
33
of light transm ission for a = X/100 and o = X/10.
Due to the strong
exponential dependence of the intensity of the transmitted beam on this
absorption coefficient due to scattering, it requires diamond film s with
such a high value of n to be extremely smooth in order to minimize the
incoherent light scattering which leads to surface haziness and poor
transparency.
Therefore, it is desirable to obtain diamond films with
nanocrystallites.
1.5
S u m m ary o f T hesis
This thesis investigates in full details the three surface issues related
to diamond nucleation and growth previously described, namely the surface
modification of single crystal copper for diamond heteroepitaxial growth,
the surface effects of transition elements, and the surface roughness of
diamond films.
The outcome of the studies is not only fundamentally
important to the science of metastable diamond growth, but also very
crucial for the technological applications of diamond films in a variety of
fields.
The three substrates - copper, steel, and quartz - have been
selected for detailed study because of their unique and different properties.
Initial study on the modification of single crystal copper surface
for diamond heteroepitaxial growth is carried out by implanting copper
with carbon at medium energy (~ 60-125 keV).
The implantation was
carried out at elevated and room temperatures to obtain carbon profile
primarily on the surface and approximately few hundred angstroms below
34
the surface respectively.
The idea is to modify the surface energy o f
copper and at the same time to provide seeds for diamond nuclei to
chemically attach on the copper surface epitaxially. The structure and
morphology of the implanted carbon in relation to diamond nucleation will
be examined in detail by various techniques such as x-ray diffraction,
scanning electron microscopy (SEM), transmission electron microscopy
(TEM ), scanning tunneling microscopy (STM ), and atom ic force
microscopy (AFM). Epitaxial relationship between the diamond and the
underlying layer will be assessed. The studies hopefully would provide
some preliminary clues to the route for diamond heteroepitaxial growth on
single crystal copper.
A novel concept of composite films consisting o f d ia m o n d
crystallites, hydrogenated amorphous carbon (a-C:H) and/or fluorocarbon
polymer is introduced for coatings on metal surfaces. A buffer interfacial
layer such as amorphous silicon based film is deposited on transition metal
surfaces prior to CVD diamond growth This is to isolate the catalytic
effect o f the elements and to prevent the rapid carbon diffusion into the
metals.
The composite films would not only be adherent, but also
relatively flexible, making them suitable for sheet m etal coating
applications (Ong and Chang, 1991). The a-C:H and fluorocarbon films,
which possess the diamond-like (Koidl an Oelhafen, 1967) and teflon-like
(d’Agostinos et al., 1990) properties respectively, are used in the composite
structure because they are relatively flexible compared to diamond. The
35
composite films consist o f a non-continuous diamond film covered by a
layer of a-C:H and/or fluorocarbon film.
The idea is to com bine the
superb properties of diamond with the somewhat more flexible nature of
the carbon films.
By combining these films, one is able to o b ta in
multilayer film s with very unique properties, which are flexible, wear
resistant, chemically inert, corrosion resistant and hard. Among all of these
desired properties for protective coating purposes, diamond crystallites
function as strength/hardness reinforcement for the com posite structure.
Interestingly, the surface morphology of the com posite film is very
smooth.
This is very desirable from the aesthetic standpoint.
In the
present study, growth of diamond composite films on sheet steel and
stainless steel are carried out.
In order to obtain smooth polycrystalline diamond film s, one
should realize the fact that the final diamond grain size is determined by
the heterogeneous nucleation rate. It is widely known from the classical
theory of nucleation (see Venables et al., 1984) that there are two distinct
optimum temperatures for crystal nucleation and growth. The temperature
for optimal nucleation is known to be lower than that for growth.
By
depositing diam ond at low temperature, the nucleation rate can be
optim ized relative to growth rate.
However, one cannot maintain the
tem perature too low due to the preferential formation of am orphous
carbon phase. T o maintain a low average temperature, a plasma discharge
is operated in a pulsed mode. In this way the diamond grains size can be
36
tailored (Ong and Chang, 1989).
Growth of optical diamond films on
quartz substrates will be presented in this thesis.
2.
2 .1
EXPERIMENTAL
M icrowave Plasma Reactor
M icrowave plasma chemical vapor deposition system shown in
Figure 2.1 and 2.2 was used to grow diamond films. The discharge was
excited by a 1 kW microwave generator (Conversion technology P I 000)
operating at 2.45 GHz frequency. The microwaves were confined by a set
o f waveguides attached to the system. The system was equipped with an
isolator,a power monitor,a three-stub tuner,and a plunger. The microwaves
w ere allow ed to interact with the quartz reactor (Quartz S c ie n tific
Incorporated, diam eter = 25 mm , length = 0.83 m) in a water-cooled
applicator.
Gas flowrates were controlled by mass flowmeters (Sierra
Instruments) and the pressure was monitored through a by-pass valve
installed between the chamber and the mechanical pump (Sargent Welch
1397).
T em perature was measured by a K -type chrom el/alum el
thermocouple with the junction installed just right beneath the graphite
substrate holder inside a quartz tube holder (diameter^ 6 mm).
2 .2
Capacitively Coupled Radio Frequency Plasma Reactor
Figure 2.3 illustrates the radio frequency (RF) stainless steel
cham ber used to deposit a-C:H and fluorocarbon polymer films.
The
system was equipped with a diffusion pump (Varian M-4) and a mechanical
pump (Sargent Welch 1397) capable of bringing the system pressure below
37
38
Figure 2.1 2.45 GHz microwave plasma CVD apparatus for growing
diamond films.
39
( ch 4 / cf4
CD»H
2.45 GHz
M icrowm Oooancor
3-Stab
3-Port Circulator
Dummy
w ^ .
MiarowtTcs
1. Baclc-to-air valve
2. By-pass valve
3. Butterfly valve
□
Macbmucal
Pomp
Figure 2.2
ThraoconpU
Oanf*
A schematic diagram of the microwave plasma CVD reactor.
40
115 AC
IIS AC
<3 O
2115MHl
115 AC
Figure 2.3 A schematic diagram of the capacitively coupled
RF (27.15 MHz) plasma reactor for growing a-C:H and fluorocarbon
films.
41
10-6 Torr. The pressure was monitored by an ionization gauge tube (MDC)
and a digital gauge controller (Perkin Elmer) in the <1 mTorr range and a
therm ocouple gauge (Varian) in the 1-400 mTorr range.
M atheson
flowm eters were used to control gas flowrates. The configuration of
electrodes made of two graphite disk was schematically shown in Figure
2.4.
In order to confine the plasma just within the two electrodes, a pair
of concentric ceramic magnets (INDOX 1 from Dexter Magnetic Materials
; O.D.= 2.086”, I.D. = 0.692”, thickness = 0.3”) was used. The plasma was
excited by a 500 W RF power generator (Comdel CPS-501 A) operating at
27.15 MHz and tuned by a match box (Nye Viking).
2 .3
Therm al Evaporator
A thermal evaporator shown in Figure 2.5 was used to deposit
amorphous silicon films which served as buffer layers for diamond growth
on steel and stainless steel substrates. The system (JEOL JEE4C) was
available at the scanning electron microscopy center which is part of the
M aterials Research Center facilities. Pieces of cleaved silicon were
mounted in a tungsten boat which was clamped to the power electrodes.
The substrates were placed on top of a glass slide approximately 10 cm
right beneath the boat. The system was pumped down to ~10 -5 Torr using
a diffusion pump prior to film deposition. The desired film thickness,
measured by Tencor Step profllometer on the steps created on the glass
slide, was in the order of
200
A which appeared blue to slight purple in
42
(a)
T
^ ^ P la a m a
Cathode Sheath
rTnrrr«n
Graphite
Magnet
(b)
Cathode Sheath
JS S S S s”
& §§§$ Plasma
__________
Subatrata
Graphite
Magnet
Figure 2.4 The configuration of the capacitively coupled RF electrodes
f o r : (a) a-C:H and (b) fluorocarbon polymer films growth.
Figure 2.5
a-Si films.
The JEOL JEE4C thermal evaporator for depositing
44
color on the substrates.
2 .4
C rystal Growth by Bridgman Technique
Single crystal Cu was first acquired from Monocrystal company in
Ohio in the form of cylindrical rod ( length : 1/2" ; diam eter : 1/2 " ;
orientation : (100)).
The rod was cut using diamond isomet machine into
approximately 1/8 " thick disks.
Later in the work, the crystal was grown
using the Bridgman crystal growth technique. The work was carried out at
the M aterials Research Center's Crystal Growth facility (N orthw estern
University) using the NRC 2804 vertical furnace system (M aterials
Research Corporation) shown in Figure 2.6. The raw materials were Cu
shots (99.999 % ; 4-6 mm in diameter) purchased from the Aesar group of
Johnson Matthey Inc.
First of all, approximately 250 grams of Cu shots were etched
with 50% H N 0 3 in H 2 O solution for ~10 sec and then rinsed with d e ­
ionized (DI) H 2 O for ~ 2 minutes. They were then blown dry with N2 gun
before being immediately poured into an Ultra Carbon graphite crucible
(see Figure 2.7 for the dimensions) and loaded into the growth chamber.
The system was equipped with a mechanical and a diffusion pump which
could provide a base pressure of ~2xl0** T on. The furnace temperature
was controlled by a Barber Colman microprocessor based temperature
controller series 570. The K-type thermocouple was inserted in the middle
Figure 2.6 The NRC vertical furnace system for growing single crystal
Cu using the Bridgman technique.
46
1 caper
T
3
5.7'0.5"
Figure 2.7 A schematic diagram of the graphite mold dimension for
single crystal Cu growth.
47
of the furnace. The Cu shots were first melted inside the graphite crucible
under ~ 1 atm of flowing purified argon gas.
This was carried out to
obtain a homogeneous polycrystalline Cu rod.
The temperature of the
furnace was programmed as shown in Figure 2.8.
The single crystal
growth was then carried out by initially placing the graphite crucible
outside the hot zone as schematically shown in Figure 2.9. It was then
slowly brought down by a motor at a speed approximately 0.24 mm/min
into the furnace. The temperature history of the furnace is shown in Figure
2 . 10.
Back-reflection X-ray Laue technique was used to orient the grown
single crystal Cu The sample was prepared by cutting the rod using an
ELOX series 300 wire cutter to obtain a flat half circle disk approximately
3 mm in thickness.
It was then etched with 50 % HNO 3 solution prior to
the Laue examination.
Figure 2.11 shows an example of the Laue pattern
of Cu crystal which had been oriented in the < 1 1 1 > direction.
The crystal
rod was then cut to the desired orientation using the wire cutter.
The samples were then first polished using 600 grit SiC paper.
Tw o alternative methods of subsequent polishing procedures were used :
1.
- 5.0 /<m and 0.3 p m AI2 O3 polishing in DI water
- cleaned with dish washing soap
- ultrasonically cleaned in DI water for ~ 2 min and blown dry
- annealed at 500 °C for 1 hour in 10*6 Torr vacuum
48
1500
1200*C
0 .2 h r
U
o
1000
w
u
a
ab
S.
E
H
500
0 .8 hr
u.O
0.4
1 hr
0.8
1.
2.0
Tim e (h r)
Figure 2.8 The programmed temperature inside the furnace during the
Cu shots melting.
49
Figure 2.9 A simplified diagram of the Bridgman technique for growing
single crystal Cu.
1300
1200 *C
9 kr
u
1000
5
I
I
300
T tao (hr)
Figure 2.10
Cu growth.
The temperature history of the furnace during single crystal
50
Cu (111)
Figure 2.11 The back-reflection Laue pattern of Cu(l 11)
(W tube, 20 kV, 15 mA, collection time = 2.5 minutes).
51
2.
electropolishing with H3 PO 4 : H2 O = 1 : 2 ; 2-3V using a Cu
electrode with the sample was placed on the anode.
X-ray
6
scan technique (with detector fixed at 50.50 and 43.32
degrees for the case of ( 1 0 0 ) and ( 1 1 1 ) oriented crystal respectively) was
then used with Cu K a line source to evaluate the orientation of the crystal
after the polishing steps. A diffraction peak at
6
=24.3 and 21.15 degrees
were obtained for < 1 1 1 > and < 1 0 0 > oriented crystals respectively. These
values are in good agreement with the expected values for ( 1 1 1 ) and ( 2 0 0 )
peaks of Cu ( 6 lexp(ill) = 25.23 degrees ; Oexp(2 0 0 ) = 21.66 degrees) within
one degree.
2 .5
Film Growth Techniques
2 .5 .1
Nucleation and Growth of Diamond on Carbon Implanted
Single Crystal Cu
Carbon ion implantation experiments were carried out by Dr. C.W.
White at Oak Ridge National Laboratory in Tennessee. Three implantation
conditions were used : (a) 60-75 keV, 820 °C , (b) 60-75 keV, room
temperature , and (c) 125 keV, room temperature. The ion dose used was
lx lO 18 cm *2 for all of the experiments.
The im plantations were also
carried out through a Ta shadow mask (diameter = 254 pm ) both at 820 °C
and room temperature. Rutherford backscattering (RBS)/ channeling
52
analysis with 2 MeV He+ and 160 degrees scattering angle were also
perform ed by Dr. W hite on the Cu crystals before and after the ion
implantations.
Pure and carbon implanted single crystal Cu (100) and (111), as
well as highly oriented pyrolytic graphite (HOPG) acquired from Digital
Instruments Inc., were used as substrates for diamond growth.
Prior to
the growth, the top surface layer of HOPG was peeled off using a scotch
tape. Typical deposition conditions are as follows : 3% CF 4 or 0.5% CH4
and 0.7% O 2 in H 2 ; total flowrate = 200 seem ; pressure = 33 Torr ;
temperature = 800 °C ; microwave power = 340 W.
2 .5 .2
G ro w th o f D iam ond C om posite Film s
Clean mirror-polished carbon steel and 304 stainless steel substrates
were used in the experiments. The substrates were first rubbed with 1/4
ftm diamond powder on a soft polishing cloth to shorten the induction
period for diamond to nucleate and to enhance diamond nucleation. The
samples were then ultrasonically cleaned with acetone and methanol for
approximately
10
minutes. The substrates were then coated with ~
200
A of
amorphous silicon (a-Si) which was prepared using a thermal evaporator.
The thickness of this buffer layer had been empirically optimized for the
diamond nucleation. There were two factors that govern this optimization :
the plasma environment and the substrate temperature. In the plasma
53
environment used here for diamond nucleation and growth, there was a
com bination o f sputtering, etching and deposition taking place on the
surface of the substrate. If an iron substrate surface was covered with too
thin a layer of a-Si, then a combination of sputtering and hydrogen etching
will completely remove it away before diamond paiticles could nucleate
and grow appreciably . Because the substrate was inductively heated by the
incident microwave, the substrate temperature increased in time initially.
If the barrier layer was too thick, then the carbon precursor species would
not find the nucleation sites on the iron surface prepared by diamond
powder polishing during the initial temperature rise o f the substrate. It
was during this initial temperature rise that most of the nucleation took
place. Even though the hydrogen etching of the silicon barrier took place
at all substrate temperature, at high substrate temperatures (> 600 °C ) the
silicon layer would start to oxidize due to the presence o f the oxygen in the
plasma. Once the oxide layer was formed, the hydrogen etch rate was
dropped by about an order of magnitude (Chang et al, 1982). This oxide
layer quenched the thickness reduction process of the barrier layer for
nucleation of the diamond on the iron surface.
Precursor gases were introduced into the system at the very initial
stage when the plasma was just turned on.
Diamond nucleation was
enhanced using this technique since the average substrate temperature was
low at the very beginning (Ong and Chang, 1989). Typical deposition
54
conditions were : gas mixture, CH4 : H2 : 0 2 = 1 : 99 : 0.4 ; total flow rate
= 100 seem ; pressure = 40 mbar ; microwave power = 280 W.
Hydrogenated a-C and fluorocarbon polymer films were deposited
using a parallel plate, capacitively coupled rf plasma system (see Figure
2.3). The substrates were placed on the powered, negatively self biased,
electrode (cathode) for the case of depositing a-C:H layer, allowing the
positive ion bombardment action to take place on the substrate surface
during the film growth. The ion bombardment is believed to be the
prerequisite for obtaining this metastable diamond-like phase (Koidl and
Oelhafen, 1987).
The deposition of fluorocarbon film was carried out by placing the
substrate on the anode. A mixture of CF4 and CH 4 gases was used to grow
this polymer film.
The addition of CH 4 enabled the deposition of the
polymer to occur due to the reaction between atomic H and F which would
otherwise acts as a highly efficient etchant (Winter and Cobum, 1979) to
form HF.
Generally, the plasma discharge is more suitable for etching if
the fluorine-carbon ratio is higher (d'Agostino et al, 1990) as originally
described by Coburn and W inters (1979). A boundary b e tw e e n
polymerizing and etching condition based on the F/C ratio was formulated
as shown in Figure 2.12. The deposition conditions for a-C:H and
fluorocarbon polymer films are listed in Table II.
55
• IliCM U « 4lf»f
N , i<Jhlw>
• -200
J
\
i
cr«
«■*
^ • 4 4 HIM
1
\ ITCHING
\
\
\
\
>
•100
\
\
\
\
POLYMERIZATION \
I
-------------- 1--------------1
*
2
F/C
Figure 2.12 The polymerizing and etching conditions of fluorocarbon
plasmas as a function of fluorine to carbon ratio (Cobum and Winters,
1979).
Table II.
The deposition conditions for a-C:H and fluorocarbon polymer
films growth.
R im
CH4
(seem)
Ar
(seem)
a-C:H
100
25
15
—
Fluorocarbon
CF 4
(seem)
P
(m T orr)
10
75
30-40
Power
(W)
15
< 5
56
2 .5 .3
Growth o f Smooth Diamond Films
Clean fused quartz slides ( 2 x 2 cm 2), purchased from Quartz
Scientific Incorporated, were used as substrates. They were first polished
with 1/ 2 -1 ftm diamond powder to reduce the incubation time for diamond
nucleation, and then ultrasonically cleaned with acetone, methanol, and de­
ionized water for -
1 /2
hr. The deposition conditions were : gas mixture,
C H 4 : 0 2 : H 2 = 0.3 : 0.2 : 99 ; total flow rate, 100 seem ; pressure, 40
mbar, and microwave power, 400 W. The substrate was inductively heated
by the incident microwave. No external heating source was provided.
The hydrocarbon gas mixture was fed into the reactor a few
seconds after the plasma was turned on. The discharge was maintained for
a certain time interval until the substrate temperature T f reaches 500-800
°C and then the plasma was immediately turned off. The substrate was
cooled to room temperature in vacuum (which takes about
10
minutes)
before another similar cycle was initiated. This process was repeated until
a continuous film was obtained. The time duration of each cycle (At) and
T f can be varied to change the diam ond grain size. Figure 2.13
schematically shows a diagram of the pulsed plasma processes.
57
jepositiofi
T.
f
T
Tim* (t)
Suburate
Cycle I
Subdraw
SubKrate
Cycle II
Cycle n
As At and Tf decreases, final grain size decreases, since the
average process temperature is lower.
However, the number of cycles to obtain a continuous film (n) increases, and
thus lower deposition rate.
Figure 2.13
A schematic diagram of the pulsed plasma p ro cesses.
58
2 .6
C haracterization Techniques
2 .6 .1
Raman Spectroscopy
T he carbon bonding configuration in diam ond film s was
investigated using Raman scattering spectroscopy. It involved the detection
o f inelastic light scattering due to vibrational excitations in a sam ple.
C rystalline diam ond has been shown to scatter at 1332 cm~l shift,
characterized as the first-order zone-center mode. Graphite, however, has
a sharp first-order Raman peak at 1580 cm~l. Due to the distinct signature
between the two carbon phases, Raman scattering has been widely accepted
as a powerful tool to determine the quality of crystalline diamond grown
by CVD techniques. The analysis in this study was performed using the
SPEX 1401 double spectrometer at 180 degrees back-reflection geometry
with a 488 nm argon laser light source (Spectra Physics series 2000). The
unit was controlled by a Data General computer system equiped with the
Complot plotting routines (Houston Instruments). The typical laser power
used was 200-300 mW.
2 .6 .2
Scanning and Transmission Electron M icroscopy
Hitachi H-570 (25keV) scanning electron microscope was routinely
used to examine the morphology and nucleation density o f diamond
particles. The samples were mounted on an aluminium stub using graphite
paste. The working distance typically used was 4-6 mm.
59
Hitachi H-700 (200 keV) and H-9000 (300 keV)
transmission
electron microscope were used to evaluate the crystalline microstructure of
diamond films grown on quartz. The films were etched off from their
quartz or Si substrates using dilute HF or HNO3 solutions respectively, and
then fetched by Cu grids after being rinsed in DI w ater for several
m inutes. A pproxim ately one hour prior to being loaded into the
m icroscopes, the samples were cleaned in acetone vapor for about 30
minutes. The sam ple exam inations in the H-9000 high resolution
microscope were performed by Dr. W. A. Chiou and Dr. F. R. Chen. The
author greatly acknowledges their generous help. In the case of carbon
layer implanted on Cu substrate, the film was floated off using 30% HNO 3
solution and examined using the H-700 TEM unit.
2 .6 .3
X-ray Diffraction and Back-Reflection Laue
X-ray diffraction work were performed using a Rigaku x-ray
diffractometer. The samples were mounted using a double sided scotch
tape on a piece o f aluminium plate.
The slits typically used were :
divergent slit=0.5 degree, scattering slit = 0.15 degrees, and receiving slit
= 0.5 degrees.
Back-reflection Laue camera was equipped with a goniostat placed
at a distance of 3 cm from the film plane. The Cu crystal was mounted on
the goniostat by screwing the conical end of the crystal into a concentric
60
alum inium rod. Polaroid type 57 instant film was used to record the
diffraction pattern. Typically, the current from the x-ray tube was 15 mA
and the voltage is 20 kV, with exposure time in the range of 2.5-3 minutes
to generate the Laue pattern of single crystal Cu.
2 .6 .4
Scanning Tunneling and Atomic Force M icroscopy
The Nanoscope II made by Digital Instruments Inc. was used in the
STM and AFM experiments of diamond growth on carbon-implanted Cu
and HOPG crystals. The unit, available at the Surface Science Facility, was
capable of imaging the surface topography of a sample by raster scanning a
sharp metallic tip such as Pt, which was used in the experiments, or W in
an X-Y scan across the surface of the sample and measuring the height
profile using the voltage applied to a piezoelectric material. Three basic
components comprise the Nanoscope unit : the microscope, the control
unit, and the computer workstation.
The piezoelectric scanner (which
controls the motion of the tunneling tip), the head, the preamp, the base
and its support, the course adjusting screws and the stepper motor are the
basic units in the microscope itself.
Figure 2.14 shows a schematic block
diagram of the STM nanoscope unit.
The AFM is using the same control unit, display and an a ly sis
functions as the STM. The microscope consists of the scanner support, the
61
Microscope
Control Unit
Computer
W o rk sta tio n
640 * 460*8
Display
IS Bit
IS Bit
Menu
Image
Figure 2.14 A schematic system diagram of the STM Nanoscope II unit.
(Source : Digital Instruments Inc.)
62
scanner and the head. Similar to the STM, the scanner housing was
interchangeable with X-Y maximum scan ranging from 0.7 p m (A
scanner) to 75 p m (G scanner). A very sharp tip (a few atoms wide) and
micro-cantilever arm made of silicon nitride, which was mounted on a gold
coated glass substrate, was used in the experiments. The deflection of the
cantilever upon scanning the sample was projected by a laser beam from a
laser diode (5mW maximum peak output at 670 nm) which was reflecting
off the end o f the cantilever and focused into a photodetector. A system
block diagram of the unit is given in Figure 2.15.
2 .6 .5
U V /V is/near IR spectroscopy
A double-beam Perkin Elmer 330 UV/Vis/near IR spectrometer at
the Analytical Chemistry Facilities was used to m easure the optical
transmittance of the grown diamond films on quartz substrates.
Tungsten
and deuterium light sources were used. The light wave was sent from the
quartz substrate side at normal incidence.
Air was taken as the blank
reference.
2 .6 .6 T rib o te s te r
A block-on-ring tribotester, as schematically drawn in Figure 2.16,
was used to qualitatively evaluate the strength adhesion of diamond films to
quartz substrates.
This simple set-up consisted o f a 52100 steel ring
63
Microscope
Control Unit
Computer
Workstation
12 Bn
1
_ _ _ _ Signal
rrocMior
4*U
LOO* Maf.
Di m
f
Pv<cm
—
1
J
640x«SQx«
Display
I
1
---------
UIU
Manu
Figure 2.15 A schematic system diagram of the AFM Nanoscope II unit.
(Source : Digital Instruments Inc.)
64
B L O C K -O N -R IN G T R IB O T E S T E R
Load Cel
Mttd fnrfc
*DBO
‘O
»-nQ
Mineral
01 Jet
Substrata
Diamond FMm
52100
Figure 2.16
A schematic diagram of a block-on-ring tribotester.
65
(hardened to 62C Rockwell) and a plexiglass sleeve. The dimension of the
ring was 62.5 mm in O.D. and 50.8 mm in I.D., while that o f plexiglass
was 25.4 mm in I.D. A steel shaft (O.D. = 25.4 mm), driven by a two hp
motor, was used to support this ring and plexiglass assembly. The samples
(of the following dimension : 9 mm x 3 mm x 1 mm) were mounted by
self-curing resin and shaped into 11 mm x 3 mm x 12 mm size block.
Experiments were performed on bare and ~0.35 ftm diamond films coated
on quartz by applying a constant load of 18.13 lb force on the samples
against the steel ring which is rotating with a speed o f 1 m/sec. The load
was applied through a line contact of 3 mm wide. T he steel ring was
constantly lubricated by a jet of mineral oil. The total sliding distance was
1.8 km.
2 .6 .7
T h re e -P o in t B ending T ester
A three-point bending tester, which was attached to an Instron
universal testing instruments (model 1125), was used to evaluate the degree
o f bendability of diamond composite film coated steels.
A schem atic
diagram o f the system is shown in Figure 2.17. The test samples were
basically supported at both ends on the stainless steel rollers (diameter = 5
mm).
The span length, as measured from center to center o f the roller,
was 10 mm. Another similar roller was loaded at midspan between the two
supports on top of the test specimen by a screw-driven mechanism (frame
stiffness = lxlO 6 lb/in). This upper roller was attached to a compression
66
t
I
.S u b s t r a t a
S ta M a ss sta a i
RoNsr
M a tal
S u p p o rt
Fixed Metal Support
Figure 2.17
A schematic diagram of a three-point bending tester.
67
load cell with 1000 lb maximum capacity. Measurements on load versus
sample deflection due to bending were recorded by a chart recorder
connected to the load cell. The specimens were subjected to continuously
increasing loads with cross head speed moving at 200 /im/min. The sample
dimensions were : thickness = 0.60 - 0.66 mm ; width = 6.00 - 6.60 mm ;
length = 10 mm.
2 .6 .8
Scratch Tester
Several methods of measuring the adhesion between film s and
substrates had been investigated. However, scratch adhesion testing has
been widely used due to its simplicity, reproducibility, and meaningful
results. This testing technique was first introduced by Heavens (1950) to
measure the adhesion property of evaporated chromium films on glass
substrates using a smoothly rounded chrome-steel point sliding across the
film surface. Benjamin and Weaver (1960) later extended the technique to
study transparent non-m etallic film s on glass substrates. P e r r y
(1981,1983), Steinmann and Hintermann (1985), and Valli et al.(1985)
used this method to evaluate the adhesion property o f PVD and CVD
coatings on various steel and tungsten carbide substrates.
A commercial scratch test instrument (Revetest) made by LSRH of
Switzerland was used in this study to evaluate the adhesion of the diamond
composite films to the substrates. The unit was available at the Basic
68
Industrial R esearch Laboratory.
Figure 2.18 gives a s c h e m a tic
representation of the equipment. The test basically employed a diamond
stylus with 200 p m tip radius sliding across the coating surface in a
stepwise increase in load at 10 mm/min. A first indication of coating
failure either by adhesive ("flaking"), cohesive ("chipping"), or any other
mode can be immediately detected by an acoustic transducer attached to
the instrument. The load corresponding to this initial failure of the coating
w as identified as the critical load, and it was widely accepted as a
semiquantitative measurement of adhesion. The instrument used here was
equipped with a malfunctioning chart recorder. Therefore, a plot of
acoustic emission versus load such as the one schematically shown in Figure
2.19 cannot be presented in this study.
2 .6 .9
C hem ical T est
Various corrosive chemicals had been used to test the resistance and
permeability of diamond composite films. These included a mixture of
H N O 3 , HC1, and H 2 O (3:7:30), 48% HF, 96% H 2 SO 4 , 37% HC1, 71%
H N O 3 , CCI4 , CHCI 3 , and 5g NaCl in 100 cc H2 O solutions. Prior to
immersing into the solutions, the sample edges were coated with wax. This
was to ensure that the deterioration of the samples was not initiated from
the edges.
69
ConfoNar
Chart Oacordar
aatactor
Diamond atyiua
Coatad aampla
Drivan aampia
hoidar
Figure 2.18
A schematic diagram of the Revetest scratch adhesion tester.
70
Acoustic
Emission
cr
Load
Figure 2.19 A schematic diagram of the acoustic emission versus scratch
load obtained in a typical adhesion scratch tester. An abrupt increase in the
acoustic emission at load PCr indicates the initial stage of film adhesion
failure.
3.
RESULTS AND DISCUSSION
3 .1
Nucleation and Growth o f Diamond on CarbonIm planted Single Crystal Copper
3.1.1
M icrostructure o f the Implanted Carbon Layer
Figure 3.1.1 shows an optical m icrograph o f carbon la y e r
implanted into Cu (100) at the following conditions : 820 °C , lx lO 18 cm-2
dose, 70 keV. Approximately 20-30 p m size o f carbon islands are clearly
observed on the surface of the substrate. The thickness of the carbon layer
is estimated to be approximately 900 A (thickness= ion dose/1.128 x 1023
atoms/cm3). Rutherford backscattering (RBS)/channeling analysis confirms
the fact that the carbon layer is mostly on the surface o f the copper, as
shown by the shift of the surface peaks of Cu shown in Figure 3.1.2.
Notice that the scattering yield of the <100> oriented crystal is much lower
than that of the randam oriented one, which is as expected.
In the case of
copper im planted at room tem perature, the carbon is em bedded
approximately 700 A deep inside the Cu, as shown by the RBS spectra in
Figure 3.1.3. Interestingly, the scattering yields of the room temperature
implanted Cu crystal which is oriented in <100> channeling direction is
found to be very high over a wide range of distance from the surface. In
fact, the values are close to the ones obtained in a random orientation. The
exact cause of this finding is unknown, but it is very likely to be due to
lattice damage by the implantation process.
71
72
Figure 3.1.1
Optical micrograph of carbon layer implanted into Cu(100)
surface at 820 °C, 1 xlO 18 dose, and 70 keV.
73
400
DEPTH (am )
200
300
2000
1800
«
•
*
—
1600
1400
Random, Virgin
Random, Implanted
<100>, Implanted
<100>, Virgin
-100
Cu
1200
1000
800
600
400
200
1.1
1.2
1.3
1.4
1.5
ENERGY (lleV)
1.6
1.7
Figure 3.1.2 Rutherford backscattering/channeling spectra o f virgin and
carbon implanted Cu(100). The conditions o f the ion implantation are :
820 °C, 1 xlO 18 dose, and 70 keV (Courtesy of Dr. C. W. White, Oak
Ridge National Laboratory, Oak Ridge, Tennessee).
74
DEPTH (nm )
200
100
2000
•
•
•
—
Random. Virgin
Random, Implanted
<100>, Implanted
<100>, Virgin
1.3
1.4
1.5
ENERGY (HeV)
Figure 3.1.3
Rutherford backscattering/channeling spectra of virgin and
carbon implanted Cu(100). The conditions of the ion implantation are :
room temperature, 1 xlO18 dose, and 70 keV (Courtesy o f Dr. C. W.
White, Oak Ridge National Laboratory, Oak Ridge, Tennessee).
75
X-ray diffraction analysis of the sample implanted at 820 °C shows that the
carbon layer consists of graphite with the c axis of the hexagonal graphite
lattice perpendicular to the Cu surface.
The (0002) reflection o f the
graphite (6= 12.84° with detector fixed at 26.54°), as depicted in Figure
3.1.4, is found to be misoriented within 0.5° with respect to the Cu (111)
reflection. This finding is similarly found in the case o f Cu(100) substrate.
The d(ooo2) spacing is calculated to be 3.356 A which is in good agreement
with the theoretical value of 3.354 A.
T ransm ission electron microscopy (TEM ) analysis further
confirm s this highly oriented graphitic structure o f this carbon layer.
Figure 3.1.5 and 3.1.6 show the TEM micrograph and the diffraction
pattern o f the region outside the fiber-like features with the electron beam
aligned parallel to the [0001] axis of the graphite. The complete absence of
(0002) diffraction peaks further substantiates the claim that the basal plane
of the graphite film is parallel to the Cu the substrate. The streaks which
form rings around the six-spot patterns suggest that the graphite consists
o f turbostratic layers (i.e. a layer w ith nearly
orientation).
random azim uthal
Figure 3.1.7 gives a comparative diagram of an ideal and
turbostratic graphite lattices. This finding is similar to the carbon films
deposited by pyrolysis or evaporation at high tem perature on other
crystalline metal substrates (Palmberg, 1969 ; Baneijee et al., 1961 ; Karu
and Beer, 1966 ; Derbyshire, et al., 1975 ; Presland and Walker, 1969 ;
Irving and Walker, 1967 ; Lee et al., 1991).
The neighboring crystals are
76
40kV . 10
C(0002)
5
<0
Cu (222)
10000.00
0 (degrees)
Cu (111)
HI
* .i
2 0 (degrees)
Cu , 20kV, 5 mA
0.00
2 0 .0 0
4 0 .0 0
6 0 .0 0
8 0 .0 0 100.00
2 0 (degrees)
Figure 3.1.4 X-ray diffraction of carbon implanted Cu( 111).
The copper signal is obtained by the normal 0/20 scan using 20 kVand 5
mA X-ray Cu K a line. For the case of carbon signal, the 6 scan is
carried out using 40 kV and 20 mA.
77
Figure 3.1.S Transmission electron micrographs (bright field) of carbon
layer implanted onto Cu(l 11) taken using Hitachi H-700 at 200 kV : (a)
low magnification , (b) high magnification (Ong et al., 1991).
Figure 3.1.6 Transmission electron diffraction of carbon layer shown
in Figure 3.1.5 in the region which does not include the fiber-like features.
(Ong et al., 1991).
79
0.1415 nm
! jj j 0.2456 nm
j j |i
0.3354 nm
<b)
Figure 3.1.7 A schematic diagram of (a) ideal and (b) turbostratic
graphite lattices (Hoffman et al., 1991).
80
misoriented to each other about the c-axis.
These graphite islands are
form ed by out diffusion process during cooling through the carbon
dissolution-precipitation mechanism (Oya and Otani, 1979 ; Derbyshire et
al., 1975). This mechanism is driven by the extremely low solubility of C
in Cu (0.0001% at 1100 °C ) (Bever and Floe, 1946) and the fact that
copper does not react with carbon to form a carbide compound because of
a filled d-electron shell {[A r]3d 1 0 4 s 1} (Oya and Otani, 1979).
The
preferential alignment of the graphite c-axis perpendicular to the substrate
is caused by the highly anisotropic nature of graphite lattice and the
presence of flat metal surface (Lee et al., 1991).
Formation of these
highly crystalline films is enhanced by the crystalline nature o f the copper
substrate (Kara and Beer, 1966).
In contrast, Prins and Gaigher in a recent report (1991) claim ed
that epitaxial diamond layer was obtained by this ion im plantation
technique.
Apparently, the authors mistakenly identified the TEM
diffraction peaks as those of diamond due to the very sim ilar lattice
spacings between diamond and graphite in many planes. In addition, a
slightly tilted { 1 0 1 0 } graphite pattern resembles two long {2 0 0 } and four
short {111} diamond diffraction peaks (Lee et al., 1991).
Figure 3.1.8 gives the TEM diffraction pattern o f the graphitized
carbon layer in the region which includes the fiber-like features. Table III
lists the measured d spacings from the diffraction pattern. The two
[oooij
Figure 3.1.8 Transmission electron diffraction of carbon layer shown
in Figure 3.1.3 in the region which include the fiber-like features.
(Ong, et al., 1991).
82
Table III.
The measured and calculated d spacings of the
diffraction pattern shown in Figure 3.1.8.
hkil
d spacing (A)
measured
calculated
0002
3.38
3.354
1010
2 .1 0
2.131
0004
1.69
1.677
1120
1.23
1.231
83
broad peaks near the incident electron beam are indexed to be the (0 0 0 2 )
reflection spots of graphite. These fiber-like features, approximately 1000
A wide, are nothing but the cracking network of the graphite films. The
significant difference in the coefficient o f thermal expansion between
graphite and Cu ( a a axis,graphite = 0.9 x 10-6 /K at 800 °C , c.f. a c u = 24 x
10~6 /K) (Nelson and Riley, 1945) causes the graphite film to experience
large compressive stress during cooling. Thus the films start to crack
exposing the prism plane o f the graphite. The cracking patterns have also
been imaged using AFM and they are depicted in Figure 3.1.9.
3 .1 .2
D iam ond N u d e a tio n
Diam ond nucleation had been attem pted on pure and carbon
implanted copper.
Figure 3.1.10 shows the SEM pictures o f diamond
crystals grown on virgin Cu (111) and Cu (111) implanted with carbon at
820 °C .
The diamond crystals obtained are polycrystalline and no
indication of epitaxial growth or growth with preferred orientation has
been observed in any of these samples. Notice that the graphite layer
formed on top o f Cu substrate at high temperature enhances diamond
nucleation quite significantly.
Seemingly, the graphite layers on the
surface o f the copper provide the nucleation site for the diamond.
However, diamond crystals grown on Cu are found to be smaller in the
case of implantation done at room temperature than that at 820 °C under
Figure 3.1.9 Atomic force micrograph of the graphite layer cracks
C u(l 11). The implantation was done at 820 °C.
85
Figure 3.1.10 Scanning electron micrographs of diamond crystals
nucleated on : (a) virgin and (b) 820 °C C-implanted Cu(l 1 1 ) crystals.
86
exactly the same deposition conditions. This is shown by the SEM
micrographs given in Figure 3.1.11. This is believed to be caused by the
induction period for the carbon species embedded inside Cu substrate to
com e up to the surface to provide nucleation sites for diamond.
Figure
3.1.12 shows the Raman spectrum o f the diamond crystals. The presence
o f 1332 cm*l confirms the presence of crystalline diamond.
An attempt to nucleate diamond on Cu(l 1 1 ) implanted with carbon
at 120 keV, lx lO 18 ion dose, and room temperature conditions yielded no
difference in results as the one grown on Cu(l 11) implanted at 60/75 keV.
This can be shown in by the SEM micrographs given in Figure 3.1.13,
where diamond crystals were grown using 0.5% CH4 , and 0.6% O 2 in 200
seem H 2 at 33 Torr and 800 °C . The interesting thing about this
experiment is that the implanted carbon was embedded far beneath the Cu
surface (~1500 A deep) and only traces amount of carbon was left on the
Cu surface. Prior to diamond growth, the substrate was exposed to H 2 and
O 2 (200 : 0.6) plasma for ~ 10 minutes. This was carried out in hope that
there is an ample time for the carbon species to slowly diffuses out to the
substrate surface before they started to interact with the carbon gas species
to form epitaxial diamond phase with the Cu substrate. Apparently, only
polycrystalline diamonds with the distinctive twin defects are obtained, as
shown in Figure 3.1.13 and 3.1.14. The corresponding Raman spectrum is
given in Figure 3.1.15. A strong and sharp diamond peak shift at 1332
cm *1 wavenumber is cleary detected.
87
V
V
Figure 3.1.11
Scanning electron micrographs of diamond crystals
nucleated on Cu(l 11) which is implanted with carbon a t :
(a) 820 °C and (b) room temperature.
2000
I
I
I
1750
1500
1250
1000
750
500
Wavcnumbcr (cm-1)
Figure 3.1.12 Raman spectrum of diamond crystals grown on Cu (111)
im p la n t e d with carbon at 820 °C. The diamond deposition conditions are :
3% CF 4 , 0.7% O 2 in 200 seem H 2 at 33 Torr and 800 °C. The spectrum
was collected using 488 nm Ar laser line.
89
Figure 3.1.13 SEM micrographs of diamond crystals grown on Cu( 111)
implanted with carbon at 120 keV, lxlO 18 ion dose, and 820 °C.
(0.5% CH 4 , 0.6% O 2 in 200 seem H 2 at 33 Torr and 800 °C)
(b)
Figure 3.1.14 SEM micrographs of the twinned diamond crystals on
C u(l 11) implanted with carbon at 120 keV, lxlO 18 ion dose, and 820 °C,
shown along with its proposed precursors (Matsumoto and Matsui, 1983):
(a) twinned cubo-octahedron, (a’) bicyclo [2 .2 .2 ] octane, (a”) tetracyclo
[4 .4 .0 . l3 ,9 .i4 ,8 ] dodecane ; (b) decahedral-Wulff-polyhedron, (b’>
hexacyclo [5 .5 . 1 . 1 2 ,6 .i8,12.o3,ll.o5,9] pentadecane ;
(c) icosahedron, (c’) dodecahedrane
(0.5% CH 4 , 0.6% 6 2 in 200 seem H 2 at 33 Torr and 800 °C).
91
C
3
•
•e
a
CO
C
4)
i1
i
r~
i " 't
1800 1700 1600 1500 1400 1300 1200 1100 1000
Wavenumber (cm-1)
Figure 3.1.15
and 3.1.14.
The Raman spectrum of samples shown in Figure 3.1.13
92
Figure 3.1.16 presents the optical micrographs which depict the
preferential nucleation of diamond crystals on the graphite layers formed
on the Cu crystal by ion implantation through a T a shadow mask at 840 °C.
Interestingly, the diamond crystals nucleate only near the edges o f the
"hump" structures which were previously occupied by the edges o f the
graphite islands.
In order to elucidate the effect o f graphite layers on diamond
nucleation, the author has attempted to grow diamond on HOPG crystals
(Digital Instruments). The SEM pictures given in Figure 3.1.17 show that
diamond crystals are indeed found primarily at the edges o f the graphite
sheet. A striking difference in the diamond nucleation density on the basal
and prism planes o f graphite is clearly observed. Closer viewgraphs are
given in Figure 3.1.18. Notice the fact that the diamond crystals are sitting
at the edges of the etch pits formed on the basal plane o f the graphite. This
implies that a disruption of the graphite lattice to create the edge planes is
necessary for diamond to nucleate easily on the basal plane.
Atomic force and scanning tunneling microscopy experiments
were also carried out to examine the initial stages of diamond nucleation on
graphite.
HOPG crystals (Digital Instruments) were used as substrates.
The atomic image of the graphite parallel to the basal plane obtained using
the STM is shown in Figure 3.1.19. The C-C bond distance is measured to
be approximately 1.4 A, in agreement to the literature value o f 1.42 A.
93
!
%
Figure 3.1.16 Preferential nucleation of diamond crystals on the edges of
the graphite islands on Cu(l 11) : (a) low magnification , (b) high
magnification. The circle patch of graphitic islands were formed by C
implantation through a Ta shadow mask at the following conditions : 75
keV, lx lO 18 cm-2 dose, and 8 4 0 °C.
94
Figure 3.1.17 Scanning electron micrographs showing the distinct
difference in diamond nucleation density on the (a) prism and (b) basal
planes of HOPG crystals.
95
Figure 3.1.18 Scanning electron micrographs (close-up views) of
diamond crystals on the (a) prism and (b) basal planes of graphite crystals.
96
Figure 3.1.19 Scanning tunneling microscope image of the basal plane of
perfect graphite lattice :
(a) the atomic image
(b) the power spectrum o f the 2D Fourier Transform.
The image was obtained using 16.8 mV bias and 1.8 nA setpoint current.
97
The power spectrum of the two dimensional Fourier transform o f the STM
image shown in Figure 3.1.19(b) reveals the hexagonal symmetry o f the
graphite lattice.
The spectrum is analogous to an optical or electron
diffraction pattern.
Figure 3.1.20 and 3.1.21 depict STM and AFM images of graphite
w hich had been exposed to 3% CF 4 and 0.7% O 2 in H 2 plasm a
environment at 33 Torr for one minute and 20 minutes respectively. The
diamond crystals again are found to be nucleating right at the edges o f the
basal plane o f graphite crystal. Notice that the STM image was collected in
the height mode to prevent the STM tip to crash upon encountering the
diamond.
STM.
Because diamond is an insulator, one cannot image the atoms by
An attem pt to carry out such an experim ent using the AFM
technique failed because of the limitation of the Nanoscope II unit. The
reason the AFM technique was used to image the graphite after 20 minutes
of diamond growth is that the technique produced better images on rough
surface than STM did.
98
Figure 3.1.20 Scanning tunneling microscope image of basal plane of
perfect graphite lattice near the edge of a step after lm in diamond growth
in 3% CF 4 , 0.7 % O 2 and H 2 plasma environment at 33 Torr and 800 °C :
(a) top view, (b) tilted view.
99
Figure 3.1.21 Atomic force microscope image o f basal plane o f perfect
graphite lattice near the edge of a step after 20 min diamond growth in 3%
CF 4 , 0.7 % O 2 and H 2 plasma environment at 33 Torr and 800 °C :
(a) top view, (b) tilted view.
100
3 .1 .3 . Graphite Edge Chemistry
The observed phenomena in the preferential nucleation of diamond
on the edges of the graphite lattice described previously is not surprising at
all when one considers the fact that the chemistry of the basal and prism
(edge) planes of graphite is very different.
Knowing the fact that the
bonding on the basal plane of graphite lattice is saturated with n (sp2)
bonding, the carbon-carbon bond on this plane is very strong. Each carbon
at the prism plane, however, has one dangling bond resulting in a lower
bond strength. A recent LEED study (Kelemen and Mims, 1984) reveals
that some kind o f interplanar bonding is present at the edges of graphite,
but the exact nature of the bonding is still unknown.
Many previous studies of catalyzed and uncatalyzed reactions of
graphite with oxygen (Hennig, 1964 ; Thomas, 1965 ; Evans et al., 1971 ;
Olander et al., 1972 ; Tomita and Tamai, 1974 ; Balooch and Olander,
1975 ; Yang and Wong, 1981 ; McKee, 1981) have shown strong evidence
that the carbon-carbon bonding on the edges o f the carbon sheets represent
the weakest and thus the most reactive sites in the whole graphite lattice.
Grisdale (1953) reported anisotropy factor of 17 for graphite oxidation
reactions in the a and c directions, while some other workers (e.g. Thomas,
1965) quoted a value as high as 25-27. Although the basal plane can also
be involved in the reaction (Yang and Wong, 1981) and its reactivity may
101
be enhanced by the presence of lattice vacancies (Hennig, 1964),
the
majority of the actions still take place at the prism plane.
Similarly to the case of oxygen attack, the basal plane surface of
graphite is less affected by reaction with atomic hydrogen compared to the
prism plane surface (Balooch and Olander, 1975). An apparent reaction
probability of graphite in producing CH 4 or C 2 H 2 products by atomic
hydrogen was found to be approximately one order magnitude higher on
the prism plane than on the basal plane at a given temperature (see Figure
3.1.22). A sticking coefficient of atomic hydrogen was found to be 0.02 on
the prism plane and only 0.006 on the basal plane of graphite. The study
(Balooch and Olander, 1975) also found that the rate constants of the
reactions for the two planes are comparable in activation energy but much
higher in the preexponential factor for the prism plane.
One interesting
thing to mention here is that the graphite surface is inert at temperature
between 800 and
1000
K and it acts simply as a recombination site for
atomic H to form H 2 molecule as implied by the reaction probability gap
shown in Figure 3.1.22.
The sticking coefficient of CHn radicals to an ideal graphite surface
is believed to be very insigficant as well. The value had been estimated to
be close to 0.03 or smaller (Vandentop et al., 1990 ; Kline et al., 1989 ;
Toyoda et al., 1989).
Therefore, reactions involving the attachment of
CHn radicals to the graphite basal plane surface is extremely unlikely. This
102
Target te m p e ra tu re , Ts (*K)
2800
1800
1000
I------i
i
1 0 - ac—
800
400
l
Basal plane
- — — a o Prism plane
. I0 * 8»10,# a t o m s / c m 2-sec
\ f * 2 0 Hz
a*
—
Apparent reaction probability,
700
r \
\
Theory
-
10
-3
r
-
ch4
\
T h eo r y
10 - 4
10
12
14
16
18
20
22
24
26
104/ T 8
Figure 3.1.22 Probabilities of producing methane and acetylene by
reaction of graphite with atomic hydrogen. (Balooch and Olander, 1975).
103
is further supported by Vandentop et al. work (1991) by the extremely
slow initial deposition rate of a-C:H onto the graphite surface. They
suggested that a disruption of the graphite surface by high-energy ions is
indispensable for the radicals to be attached chemically to the graphite
surface.
A study of graphite etching by atomic hydrogen was therefore
carried out in time series. The results, as examined using STM and AFM
techniques, are given in Figure 3.1.23. A perfect hexagonal image of the
HOPG graphite lattice is clearly resolved prior to the atomic hydrogen
etching experiment.
Deterioration of the lattice by atomic hydrogen starts
at a very early stage (i.e. after approximately
1
minute etching) by the
creation of many lattice vacancies as predicted earlier (Feates, 1968). After
three-minutes of plasma exposure, the surface of graphite has been severely
roughened that the atomic imaging becomes an impossible task. Hexagonal
etch pits are seen and becoming clearer after five minutes of the etching
experim ent,as given by the AFM image in Figure 3.1.23(d).
The
formation of the etch pits with hexagonal symmetry are formed because of
the anisotropic reactivity of carbon along the principal crystallographic
orientations on the basal planes (< 1 0 1 Q> and < 1 1 2 0 >).
In the previous oxidation studies of graphite (Thomas, 1965), the
orientation of the hexagonal pits on the basal plane was put in perspective
with the twin bands which are formed either along the
<10lo> or <112Q>
104
Figure 3.1.23 Results of atomic H etching of HOPG crystals in
microwave plasma (H 2 : 20 seem ; He : 20 seem ; p o w er: 100 W ; 10 Torr
; 100 - 400 ° C ) : (a) STM image before etching ; (b) STM image after 1
min. etching ; (c) STM image after 3 min. etching ; (d) AFM image after 5
min. etching.
105
direction. This is schematically illustrated in Figure 3.1.24. The formation
o f either perpendicular or parallel pit depends on the oxidation rates in
those principal crystallographic orientation which in turns depend on the
oxidation temperature. In the case of hydrogenation reaction of graphite, a
reaction involving the hydrogen insertion at the {112L} lattice ("arm ­
chair") planes, but not the {101L} planes ("zig-zag"), of graphite
is
considered to be the most likely reaction path (Zielke and Gorin, 1955 ;
Tom ita and Tamai, 1974) in the typical temperature range for CVD
diamond growth. Schematic diagrams of both elementary mechanisms are
given in Figure 3.1.25 and 3.1.26.
The more likelihood of the occurence of the "arm -chair"reaction
can be explained by two simple arguments. The first argument is based on
the electronic structure of the carbon edge atoms on graphite. Coulson's
model (Coulson, 1960), which is based on the consideration of the
unsatisfied nature of the extra electrons at the graphite edge atoms, pointed
out the fact that the carbon atoms at the "arm-chair" configuration attain a
partial triple bonding character upon pairing the electrons between the two
neighboring carbon atoms.
This can be formed without significantly
affecting the behavior o f the resonating electrons in the n bond system.
The atoms at the "zig-zag" form however tend to be in divalent states based
on s2p2 hybridization. This type of configuration which allows the two
neighboring carbon atoms to form an extra o bond is energetically more
favorable with respect to the original sp2;i state. It is more stable and thus
106
<1010>
(•)
<
1120>
(10?L) FACE
TZIO-ZAO)
(I iI l ) pace
CAKM-CHAOT)
<b)
Figure 3.1.24 (a) A schematic diagram of the crystallographic
orientation of hexagonal lattice in graphite ; (b) The relation between the
hexagonal etch pits and twin planes in graphite lattice (Thomas, 1965).
107
H
H
H
H
\
H
H
Figure 3.1.25 The elementary mechanism for the hydrogenation of
graphite via the attack of the atoms on the { 1 12L} plane ("arm-chair")
(after Tom ita and Tamai, 1974).
106
Figure 3.1.26 The elementary mechanism for the hydrogenation of
graphite via the attack of the atoms on the {101L} plane ("zig-zag")
(after Tom ita and Tamai, 1974).
109
less reactive than the triple bond.
steric argument.
The second explanation is based on the
The hydrogenation reaction involving the atoms in the
"zig-zag" planes entails a complicated rearrangement o f ^-electron systems
and the steric repulsions, while no rearrangement in the ^-electron systems
is necessary to provoke the "armchair" reactions (Tomita and Tam ai,
1974).
3 .1 .4 The Epitaxial Relationship between Diamond and
C arbon-im planted Copper
From the arguments presented in the previous section, it is very
reasonable to say that a simpler and energetically more favorable way of
nucleating diamond on graphite is by first hydrogenating the carbon atoms
at the graphite edge to form sp 3 hybridized CH 3 bonds . This is done
through the disruption of carbon atoms in the {112L} planes (Zielke and
Gorin, 1955 ; Tomita and Tamai, 1974). It is then followed by reaction
with hydrocarbon radical precursors from the gas phase to form diamond.
Figure 3.1.27 presents a simple lattice model of a diamond nucleus
on the edge o f the graphite lattice.
%
According to the proposed model, the
{ 1 1 1 } plane of diamond is perpendicular to the basal plane o f graphite and
the
<1
10>diamond is parallel to the < 1 120>graphite-
Since the < 1 120> has
three-fold symmetry, there are therefore three possible orientations of
diamond crystal about < 1 1 1> axis. The proposed model is reasonable
110
<
1010>
Figure 3.1.27 A simple model of diamond nucleus on the edge of
graphite basal plane. { I l l }diamond II {OOOtygnphite; < 1 10> diamond // the
< 1 120>graphite-
Ill
because the carbon-carbon length in the < 1 10> direction of diamond (2.507
A ) matches that in the <1120> direction of graphite (2.459A) within 2 %.
In fact, in the high pressure synthesis of diamond from graphite using the
explosion shock quenching method, the row of carbon atoms along the
<1120> of graphite is believed to transform into the <110> of diamond
(W heeler and Lewis, 1975). The model is also in agreement with the
previous study of graphite formation on CVD diamond films in which
graphite was found to form with its basal plane matches on or close to the
{111} planes o f diamond (Zhu et al., 1989).
112
3 .2
Growth o f Diamond Composite Films on Iron Surfaces
Diamond crystals with sizes ranging from 2000-5000 A have been
successfully deposited on carbon steel and 304 stainless steel substrates.
The crystals, with the typical cubical facets, do not show any preferred
orientation. The scanning electron micrograph (SEM) shown in Figure
3.2.1 also reveals the five fold twinnings in the crystals. Nucleation density
in terms o f surface area is designed to be approximately 50-75 % o f the
total substrate surface. Hydrogenated a-C and fluorocarbon films, each
approximately 4000 A thick, are then deposited on the diamond coated
carbon steel and 304 stainless steel substrates. The resulting composite
films are relatively smooth and uniform in thickness. A cross sectional
view of the diamond and a-C:H composite film on steel is shown in Figure
3.2.2. Figure 3.2.3 shows the case when the substrates are coated with
continuous diam ond film s. T he film s blister o ff the
su b stra tes
instantaneously during cooling due to the difference in thermal expansion
properties. This illustrates that continuous diamond films cannot be used
for protective coating purposes directly on these substrates .
3 .2 .1
Three-Point Bending Test
Three-point bending tests had been carried out to the multilayer
diamond composite films schematically shown in Figure 3.2.4. The results
are shown in Figure 3.2.5 where the load (which has been normalized with
113
Figure 3.2.1
SEM micrographs of diamond crystals grown on carbon
steel substrate.
Substrate
Substrate
Figure 3.2.2 Cross sectional view of diamond composite film on carbon
steel substrate.
Figure 3.2.3 SEM micrographs of continuous diamond film grown
carbon steel substrate. The white markers indicate the spontaneous
cracking o f the film upon sample coding.
116
<»)
[
r r y y y i
Substrate
Substrate
(c)
n ry in ri
Substrate
O
Diamond particles
a-C:H film
EZ2 Fluorocarbon film
Figure 3.2.4 Schematic diagrams o f the multilayer structure o f diamond
composite film s :
(a) a-C:H, fluorocarbon, and diamond on 304 stainless steel
(b) a-C:H and diamond on 304 stainless steel
(c) fluorocarbon and diamond on 304 stainless steel.
Bending angle * 2 9
* 2 arctan{28A*}
0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
D is p la c e m e n t (m m )
Figure 3.2.5 Plot of bending load versus sample displacement
(a) bare 304 stainless steel
(b) a-C:H, fluorocarbon, and diamond on 304 stainless steel
(c) a-C:H and diamond on 304 stainless steel
(d) fluorocarbon and diamond on 304 stainless steel.
118
respect to the width and the thickness of the specimen) is plotted versus
vertical displacement (d) for samples : (a) bare 304 stainless steel ; (b) aC.H film, fluorocarbon film, and diamond particles on 304 stainless steel ;
(c) a-C:H film and diamond particles on 304 stainless steels ; (d )
fluorocarbon films and diamond particles on 304 stainless steels.
Unlike sample (d) of Figure 3.2.5, sample (b) and sample (c)
experience film at the very early stage of plastic deformation, as indicated
by the black arrow markers which depict a sudden stress relaxation. These
crackings, as shown in Figure 3.2.6 (b), 3.2.7 (b), and 3.2.8 (b), are
generated by the loading forces after the threshold deform ation in the
composite coatings is exceeded. In contrary to the behavior of fracture in
brittle bulk solids in which the failure mechanism originates from the
weakest links in the materials and results in freely spread fracture without
any crack arrest, the fracture mechanism of thin brittle coatings on metallic
substrates is strongly affected by the interactions between the coating and
the substrate materials (Gille, 1984). The propagation o f the cracks
produced by the bending are restricted to one dimension and are stopped
by the steel substrate due to the large differences in fracture toughness
between the coating and the substrate The cracks are generated parallel
and with nearly equal distances to each other and perpendicular to the
loading force.
Arrest of cracks are also observed, as shown in Figure
3.2.7 (b), possibly due to inhomogeneity of the composite coatings and
the interactions o f cracks with each other (Gille, 1984 ; Hu and
119
Figure 3.2.6 Optical micrographs of a-C:H, fluorocarbon, and diamond
coated 304 stainless s te e l: (a) before and (b) after the bending test.
120
2 0 0 um
4
1
Figure 3.2.7 Optical micrographs of a-C:H and diamond coated 304
stainless s te e l: (a) before and (b) after the bending test.
121
Figure 3.2.8 Optical micrographs of fluorocarbon and diamond coated
304 stainless ste e l: (a) before and (b) after the bending test.
122
Evans, 1969).
The initial stages of coating fracture, as indicated by the arrows in
Figure 3.2.5, correspond to the critical sample displacements (6c) equal
to 0.42 mm for sample (b) and sample (c), and 0.56 mm for sample (d).
From the results of sam ple (d), it is expected that the fluorocarbon
interlayer in sample (b) would experience cracking at 6c som ewhere
intermediate between sample (c) and sample (d), and provides the extra
protection in case o f cracking failure of the top layer (a-C:H film).
However, the fluorocarbon and the top a-C:H layer in sample (b) crack
simultaneously. Hence, no improvement in the bendability of the film can
be obtained. The reason is that there is good adhesion o f the fluorocarbon
film to the substrate and the top a-C:H film.
As a result, the fluorocarbon
interlayer is experiencing the strong adhesive forces upon bending the
sample not only at the interface with the substrate, but also at the interface
with the a-C:H layer.
The critical bending angles corresponding to the precracking state
have been evaluated to be ten degrees for sample (b) and (c), and 13
degrees for sample (d). These values are obtained using the following
eq u atio n . :
Bending angle = 26 = 2 arc tan (26c/L)
(3.1).
123
As a comparison, Gille and Wetzig (1963) reported the 6c values for the
commonly used TiC and TiN coatings on C100W1 and X82WMo6.5 steels
based on four-point bending tests. The equivalent bending angle values
range from two to four degrees.
Their results are clearly inferior to the
current findings obtained by three-point bending which is known to be a
more severe test.
The maximum deflection of a beam subjected to four-
point bending is larger than that subjected to three point bending at a given
load (Suhir, 1990).
No quantitative analysis, such as Young's modulus and flexural
strength of the films, will be performed here. According to the A STM
standard (E855-84) for three-point bending (ASTM), the span length and
the total specimen length shall be 100 times and 165 times the nominal
thickness in the range exceeding 0.51 mm respectively.
The sample
dimensions used here (thickness = 0.6-0.66 mm ; length = 1 0 mm ; width =
6-6.6 mm) are clearly not within the standard due to the limitation o f
diamond CVD reactor for growing a larger sample.
3 .2 .2
S c ra tc h T est
Adhesion tests had been performed on carbon steel coated with aC:H and diamond. The results are given in Figure 3.2.9. A slight wear of
coating on the diamond stylus track upon imposing a load of 8.8 newtons is
clearly observed (Figure 3.2.9 (a)). However, upon increasing the load.
124
*
Figure 3.2.9 SEM micrographs of scratch tracks of a~C:H and diamond
composite film coated 304 stainless steel with the stylus load :
(a) 5 N, (b) 20 N, (c) 29 N, (d) 39 N, (e) 49 N, (f) 68 N.
The arrow markers indicate the direction of diamond stylus motion.
125
the film start to experience some crackings with semicircular trajectories
parallel to the leading edge of the diamond stylus (Figure 3.2.9 (b)-(e». In
addition, the width of the scratch track gets wider as well due to the
increasing applied loads . This type of film failure is commonly called the
conformal cracking which occurs without any adhesive failure on the
coating and substrate interface (Burnett and Rickerby, 1987). This failure
mode is schematically represented in Fig. 3.2.10 (a). It essentially consists
o f cracking only within the scratch track. It is due to tensile bending
moments within the coating as the diamond stylus is deforming the coating
and the substrate. As the load is increased to 68 newtons ( ~7kg), the
scratch track width increases to as wide as 200 pm . Significant pile up of
removed materials on the edges of the track is also observed (Fig. 3.2.9
(0)-
However, the cracks are now running almost perpendicular to the
direction of the stylus track.
This type o f cracking failure mode is
somewhere in between the conformal and the tensile crackings (Figure
3.2.10 (b)).
The interesting thing to note here is that the film still adheres to the
steel substrate at this such high applied load. This implies the existing of
strong interfacial chemical bonding. In contrary, a-C:H or fluorocarbon
film alone does not adhere quite as well to the steel substrate.
The films
are easily scratched off by a Ti tweezer even at a very small load.
Interface modification by ion implantation for example is necessary to
improve the adhesion property of these carbon films on steel (see e.g.
126
(ft) Conform*! Cracking
(Top Vitw)
S ib ]
Crockl
(b ) Tensile Cracking
Figure 3.2.10 Schematic diagrams of (a) conformal and (b) tensile
cracking failure modes in thin film after scratch adhesion tests (after
Burnett and Rickerby, 1987).
127
Ferber et al, 1967).
The diamond particles presumably play an important
role in obtaining this extraordinary adherent films due to the similar nature
of carbon-carbon bonding between the diamond and the a-C:H films. No
higher scratch load has been tested due to the limitation o f our equipment.
Consequently, the value o f the critical load (Lc) cannot be determined other
than stating that Lc > 68 newtons. As a comparison, the typical values for
Lc for few microns thick of TiC or TiN coatings on steel range from 10-50
new tons, depending on the film thickness and substrate hardness
(Steinmann and Hintermann, 1985 ; Valli et al, 1985). It had also been
reported (M cCune et al., 1989) that the Lc values of only less than 10
newtons was obtained in the case of diamond films on SiAlON tool inserts.
3 .2 .3
C hem ical T est
The composite coatings (samples (b), (c) and (d) of Figure 3.2.5)
had been subjected to a variety of chemical tests (i.e. HNO 3 : H C 1: H 2 O
(3:7:30), 48% HF, 96 % H 2S04, 37 % HC1, 71% HNQ 3 , 5 g NaCl inlOO
cc H 2 O, CCI4 , and CHCI3 solutions) to evaluate their chemical resistance
and permeability. The tests were performed on samples before and after
bending, but prior to film cracking. The salt solution test was perform ed
on the coated and uncoated carbon steel substrate, while the rest of the
chemical tests were carried out on the coated and uncoated 304 stainless
steels. In the case of corrosion test, the samples, which had been protected
128
on the edges and the backside by epoxy, were immersed into the solutions
for ~ 4 hours. For other tests, a drop of the chemical solutions was placed
on top o f the coatings for several minutes. The results, examined using
optical microscope, show that the films are resistant to all o f the chemicals,
and impermeable to the salt solutions (see Figure 3.2.11). However, upon
film cracking, all of the samples, including sample (b) of Figure 3.2.5, fail
the tests. Delaminations of the films due to attack at the film and substrate
interface by the chemicals seeping through the cracks are always observed
in such a case.
(b)
Before Ihe test
A f t er the lest
Figure 3.2.11 The results of four-hour-corrosion tests on : (a) a-C:H and
diamond coated carbon steel and (b) bare carbon steel in a solution of
5 g NaCl and 100 cc water.
130
3 .3
Growth of smooth diamond films
3 .3 .1
Effect o f cycle duration and temperature on diamond
grain size
The time duration of each cycle ( A t ) and the final temperature
(Tf) can be varied to change the diamond grain size (see Figure 2.13). It is
found that the grain size increases with At and Tf, as shown in Figure
3.3.1. For a given At, the total number of cycles used to form an uniform
and continuous film is different. More process cycles are needed for the
case of
smaller At to cover the substrate surface with a continuous
diamond film.
In addition, the average temperature (T avg) over one
process cycle is lower as At is reduced. The Tavg is obtained by using :
(3.2)
w here T av g is the average temperature, T(t) is the instantaneous
temperature, t is the time , and tf is the time to reach the Tf. The variable
t does not include the time when the plasma is off for cooling the substrate.
Since the nucleation process is constantly occurring during each process
cycle, diamond films with higher nucleation density and sm aller final
average grain size are obtained as we reduce At (Figure 3.3.1 (a) and (b)).
By reducing T f and At further, the average tem perature can
be
substantially lowered. This results in films with much finer grains, as
131
800
(b)
600
g <00
H
200
100 nm
0
200
100
300
r.m.s. * 150 A
t (s e c )
(C)
600
g
H
400
200
100 nm
0
0
10
20
30
l (s e c )
40
50
r.m.s. = 100 A
Figure 3.3.1 Dependence of average diamond grain size on the
process cycle time interval At and the final temperature Tf when the
plasma is turned o f f :
(a) At = 60 min.,Tf = 800 °C , Tavg = 800 °C, one cycle
(b) At = 4 min.,Tf = 800 °C , Tavg = 600 °C, four cycles
(c) At = 45 sec., T f = 600 °C, Tavg = 415 °C, 16 cycles.
132
shown in Figure 3.3.1 (c) with Tf and At equal to 600 °C and 45 sec
respectively.
The surface roughness of the films is measured by a Tencor
Alpha Stepper profilometer, and taken as the r.m.s. of the peak to valley
profile o f the film surfaces. The film with the finest grain size (sample
(c» has the smoothest texture (r.m.s. = 100 A). No preferred orientation
o f crystals is observed in the obtained films for all cases. The lowest
average temperature is about 400 °C in the present experiments.
Figure 3.3.2 shows the SEM picture of diamond film grown using
the following conditions : At = 20 sec ; Tf = 500 °C ; 700 W ; 20 cycles.
Notice that the surface texture is extremely smooth as confirmed by the
step profilometry spectra given in Figure 3.3.3. Again, it demonstrates the
significant effect of the At and T f on the microstructure and su rface
smoothness o f diamond films.
3 .3 .2
R a m an Spectroscopy
Figure 3.3.4 shows the Raman spectra of two diamond films grown
at two different pulsed conditions.
The spectra of the high pressure
synthetic diamond powders is also given for comparison. Figure 3.3.4 (a)
reveals a peak at 1332 cm-1 indicating the presence of crystalline diamond
bonding in the film grown using four-cycle pulse.
The broad peak
centered near 1500 cm*l is also observed. The origin of this peak is still
uncertain. Nemanich et al. (1968) postulated that the feature might be due
133
Figure 3.3.2 Scanning electron micrographs of very smooth diamond
films : (a) low magnification (1000 x ) ; (b) high magnification (40,000 x)
grown at At = 20 sec., Tf = 300 °C, 20 cycles.
0 2 H " |
|'i~“ 1
Figure 3 3 .3
lo p
Hm
**
Surface profile o f diamond film shown in Figure 3.3.2.
Intensity (arb. unit)
134
1800 1700 1600 1500 1400 1300 1200 1100 1000
Wavenumber (cm-1)
Figure 3.3.4 Raman spectroscopy of diamond thin films grown at Tf =
800 °C : (a) 4 cycles, (b) 1 cycle, (c) diamond powder
135
to an amorphous sp3 and sp2 coordinated carbon or im purity carbon
network. Figure 3.3.4 (b) however shows a stronger diamond peak at
1332 c n r 1 and a broad graphitic peak at ~1580 cm-1.
Although the
intensity of the 1332 cm-1 peak is comparable to the 1580 cm-1 peak, the
absolute scattering cross sections of graphite has been measured to be
greater than that of diamond by a factor of ~ 50 (W ada et al., 1980).
Therefore, the ratio of sp^/sp2 bonding content in the film is much larger
than the observed Raman scattering intensities show. The results indicate
the presence of larger amorphous carbon content in the film grown using
more process cycles.
This is understood knowing the fact that the
am orphous phase is known to form preferentially at low tem perature
(<400 °C).
3 .3 .3
Transm ission Electron Microscopy
The crystalline structure of diamond films is examined using the
standard transmission (200 keV Hitachi H-700) and high resolution (300
keV Hitachi H-9000) electron microscopes.
The bright-field (BF) image
of the film, as shown in Figure 3.3.5, indicates the distinct diamond crystal
habit.
The sharp continuous rings in the diffraction pattern confirm the
polycrystalline nature of the film. Table IV shows the crystallographic
plane spacings (d) measured from the diffraction rings. The tabulated
values are in excellent agreement with the corresponding values for natural
diamond (ASTM 6-675) and not graphite-2H (ASTM 23-64), which has
136
Figure 3.3.5 (a) Transmission electron micrograph of the diamond film
(bright field : 200 kV), (b) The corresponding transmission electron
diffraction, (taken by Hitachi H-700 TEM).
Table IV. Transmission election diffraction data for CVD diamond film, the reported values of natural
cubic diamond (ASTM 6-675) and hexagonal graphite-2H (ASTM 23-64).
(Hitachi H -700,200 kV)
Observed
= = ====
Ring no. d(A °)
Reported (ASTM 6-675)
= = = =======
=
hid
d ^ j (A0)
Graphite-2H (ASTM 23-64)
========== =
hkl
dh k |(A°)
% Deviation of the
observed spadngs from
ASTM 6-675
ASTM 23-64
1
2.06
111
2.059
101
2.033
+0.048
+ 1.32
2
1.266
220
1.261
110
1.232
+0.39
+2.76
3
1.082
311
1.075
200
1.067
+0.67
+ 1.40
4
1.032
222(0
1.030*
201
1.053
+0.19
-1.99
5
0.891
400
0.8917
204
0.9002
-0.078
-1.02
6
0.816
331
0.8183
210
0.8062
-0.28
+1.21
7
0.723
422
0.7281*
300
0.7110
-0.70
+ 1.69
8
0.6806
511/333
0.6865*
118
0.6935
-0.86
-1.86
(Q : due to second order diffraction
* Calculated values based on lattice constant = 3.5667 A°
138
values for d spacings closests to that of diamond. This comparison is very
important, since recently Kitahama etal. (1986) mistakenly identified the
carbon films prepared by ArF excimer laser-induced CVD as diamond.
The films were in fact later re-identified as graphite-2H (Kitahama, 1988).
For the case o f film s in this present study,
the previous Raman
spectroscopy measurements confirm that they are truly diamond particles.
Figure 3.3.6 presents a high resolution electron micrograph from a
single diamond crystal. The crystal is oriented in a < 011> orientation, and
therefore the lattice images of the {111} planes are clearly visible, as
shown by the cross fringes with 2.06 A spacing (Figure 3.3.6 and Figure
3.3.7).
A closer view of the picture reveals that the crystals contain
defects in the form of twins and stacking faults as shown in Figure 3.3.7
(b) and Figure 3.3.8.
The 2= 3 primary twins have been c o m m o n ly
observed in CVD diamond crystals (Williams et al., 1989 ; Matsumoto and
Matsui, 1983, Narayan et al., 1988). Notice from Figure 3.3.7 (b) that
there are three twin planes already observed within only 100 A region.
The 2 = 3 type of grain boundaries in CVD diamond films have been
studied in details by Narayan (Narayan, 1990). He simulated the atomic
structure at the boundary based on energy minimization techniques, with a
preliminary structure modeled using coincidence site lattice (CSL) theory.
The calculated atomic structure in [001] projection is shown in Figure 3.3.8
(b), with {111} plane as the twin boundary. The sim ulated structure
matches reasonably well with the present experimental observations. He
139
Figure 3.3.6 High resolution TEM micrograph of a single crystal
diamond (taken using Hitachi H-9000 HREM, Ong et al., 1990).
140
<*>
o
19 A
(b )
Figure 3.3.7 High resolution TEM micrograph of a single crystal
diamond oriented in a [110] direction. The lattice image of {111} plane is
clearly observed with lattice spacing of 2.06 A : (a) image of the perfect
lattice , (b) image of the defected lattice (twins and stacking faults)
(Ong et. al, 1990)
141
Figure 3.3.8 A detailed HREM view of a primary twin plane (2 = 3) in
diamond crystal twin plane along [011] projection (the twin plane is
{111}): (a) experimental results (Ong et al., 1990), (b) calculated
structure (Narayan, 1990).
142
also calculated the energy for the £= 3 boundary using Tersoff potentials,
and the results are very close to zero indicating the preferential formation
of this type of twin in CVD diamond crystals.
Figure 3.3.9 shows the image of fivefold symmetry in diamond
crystals.
This formation results from twinning in {111} planes.
The
angles between the {111} planes which enclose the fivefold symmetry are
measured within 0.5° to be falling between 70.5° (ideal case) and 74°, and
the sum of all the angles is 360°. The plane boundaries are of coherent
type.
Since the diamond lattice is very rigid and with the high energy
constraints in forming defects such as dislocations, it is reasonable to say
that the misfit of the plane boundaries is mainly accommodated by elastic
strains (Matsumoto and Matsui, 1983 ; Narayan, 1990).
T he m echanism
fo r
the
fo rm atio n
of
th ese
fiv efo ld
microcrystallites was proposed by Matsumoto and Matsui (1983) based on
the hydrocarbon cage compounds. These compounds are schematically
shown in Figure 3.3.10. They suggested these em bryos (hexacyclo
pentadecane and dodecahedrane) as the basic skeleton for the fivefold
twinned diamond crystals.
Devries (1967) had constructed a model of
fivefold twin diamond crystal based on the nucleus structure shown in
Figure 3.3.11.
Narayan et al. (1988) contended that the fivefold
microcrystallites are formed by nucleating at the core of a/2 <110>{001}
edge dislocations which are almost normal to the substrate. The presence
143
1 nm
9, = 71°
Oj = 70.5°
04 = 70.5°
03= 74°
03= 74°
Figure 3.3.9
An HREM micrograph of fivefold twin in diamond crystal
taken along [011] projection (the twin plane is {111})
(Ong et al., 1990).
144
b
Figure 3.3.10 The schematic diagram of five fold twinned diamond
crystal and the hydrocarbon cage compounds proposed as the precursors
(a) decahedral-Wulff-polyhedron
(b) hexacyclo [5.5.1.1 * A 1 8,12.o3,11.05,9] pentadecane
(c) icosahedron
(d) dodecahedrane
(M atsumoto and Matsui, 1983).
Figure 3.3.11 A proposed model of a fivefold twin diamond shown along
[110] projection based on 4(111) twins (Devries, 1967).
146
of 10 dangling bonds, a pair of penta and septa rings make up the five twin
boundaries in the microcrystallites. At this point a valid argument is still
unclear.
3 .3 .4
Effect o f film surface roughness on optical transm ission
spectra
In examining the effect o f film surface roughness on o p tic al
transm ission properties, transmission measurements from 0.2 to 2 /<m
wavelength region were carried out on two diamond film coated quartz
substrates which differ in surface roughness due to different average grain
size. The optical transmission study was carried out using a double-beam
Perkin Elm er 330 U V /V isible/near IR spectrom eter. T ungsten and
deuterium light sources were used, from which the light wave was sent
from the quartz substrate side at normal incidence angle.
The film
transmittance versus wavelength in p m is plotted in Figure 3.3.12.
The transmission of the rough film (sample d; r.m.s. = 2000A) is
lower than that of the smooth film (sample b; r.m.s. = 200 A) of equivalent
thickness by as much as 20 %. The effect of surface irregularities on the
transparency of the films is clearly significant. The optical absorption edge
for both film s is at 0.225 /<m, similar to the type Ila diamond.
The
observed tailing near the absorption edge for both type of films is believed
to be from the structural imperfection and impurities in the film (such as
147
100
C
O
0)
(0
e^ a
E
(0
c
(Q
(Quartz)
(Diamond)
1
0 .0
0.5
i
I
1.0
i
I
1.5
i
I
i
2.0
Wavelength (nm)
Figure 3.3.12 Optical transmission spectra o f :
(a) 1-mm-thick quartz
(b) 0.92 n m diamond coated quartz (roughness = 200 A)
(c) 1-mm-thick type Ila natural diamond (*)
(d) 1.25 n m diamond coated quartz (roughness = 2000 A)
(* Klocek et al., 1968).
L_
2 .5
148
a-C, graphite, and sputtered Si from fused silica reaction tube), as well as
the internal light scattering at the grain boundaries. The degradation o f the
film transparency near the absorption edge can also be accounted by the
enhanced incoherent light scattering from the rough surfaces as o is
approaching the wavelength of incident light.
However, the 60+% optical
transparency from 0.6 - 2 p m wavelength is sufficiently high for most
practical applications. The oscillations in the transmission spectra are due
to Fresnel interference phenomena of the thin film . Figure 3.3.13 depicts
the microstructures of the two samples, as examined by SEM.
Optical transmission measurements were also taken on samples with
different diamond film thicknesses. The results are shown in Figure 3.3.14
- 3.3.17 for film thicknesses ranging from 3000 A - 8000 A. Again, the
oscillations in the spectra are observed. Although the sample in Figure
3.3.17 (t = 8000 A) is thicker than the sample shown in Figure 3.3.14 (t=
3000A), there is no discernible grain size as presented by the SEM
micrographs in Figure 3.3.18. The film thickness versus deposition time
(taken as the number of pulses multiplied by At) is given in Figure 3.3.19.
From this figure, the growth rate is calculated to be approximately 0.5
^m /hr, based on the slope of the straight line.
149
Figure 3.3.13 SEM micrographs showing the crystal morphology and
grain size of samples shown in : (a) Figure 3.3.12 ( b ) . and (b) Figure
3.3.12(d).
150
100
X Transmission
3000 A
80
60
40
20
0
0
400
800
1200
1600
2000
2400
Wavelength (nm)
Figure 3.3.14 Optical transmission spectra of 0.3 ftm thick diamond film
on 1-mm-thick quartz.
151
100
X Transmission
4000
A
80
60
40
20
0
0
400
800
1200
1600
2000
2400
Wavelength (nm)
Figure 3.3.15 Optical transmission spectra of 0.4 p m thick diamond film
on 1-mm-thick quartz.
152
100
X Transmission
5000
A
80
60
40
20
0
0
400
800
1200
1600
2000
2400
Wavelength (nm)
Figure 3.3.16 Optical transmission spectra of 0.5 p m thick diamond film
on 1-mm-thick quartz.
153
100
X Transmission
8000
A
80
60
40
20
0
0
400
800
1200
1600
2000
2400
Wavelength
Figure 3.3.17 Optical transmission spectra of 0.8 /<m thick diamond film
on 1-mm-thick quartz.
154
Figure 3.3.18 SEM micrographs of samples shown in Figure 3.3.14
(a and a ’) and Figure 3.3.17 (b and b’).
155
Film Thickness (pm)
1
CH4 :H2 : 0 2 - 0.3 MOO: 0.10
Total flowrata - 100 aeem
0.8
At - 4 min
Mlerowava powor - 310 W
Praaaura - 40 mbor
0.6
0 .4
0.2
0
0
20
40
60
Deposition time (min.)
Figure 3.3.19 Plot of Him thickness versus deposition time. Note that the
deposition time is taken as the number of cycles times the cycle period (At).
156
3 .3 .5
T rib o test
Optical micrographs shown in Figure 3.3.20 depicts the results of
the tribotest. Significant amount of material wear for the bare quartz is
clearly observed, and this is further substantiated by the measured surface
profile across the
grinding track (Figure 3.3.20 (b)) using the Tencor
Alpha Stepper piofilometer. A U-shape groove as deep as three microns
due to removal of the materials is detected.
The diamond film coated
quartz, however, is completely intact after being tested under the same
condition (Figure 3.3.20 (c) and (d)).
As a matter of fact,
m aterials are transferred onto the film surface.
bearing
This is shown in the
optical micrograph (Figure 3.3.20 (c)) by the white striations along the
direction o f the grinding. The more detailed features are depicted clearly
in the SEM micrograph of Figure 3.3.21.
This finding implies that the
diamond film adheres extremely well to the quartz substrate. The results
are, of course, not too surprising given the fact that the thermal expansion
coefficient of diamond matches reasonably well to that of quartz over a
wide range of temperature.
157
Figure 3.3.20 Results o f block-on-ring tribotest on bare and diamondcoated quartz : (a) optical micrograph of the grinding track on bare quartz,
(b) surface profile across the grinding track of sample (a), (c) optical
micrograph o f the grinding track on 0.35 p m diamond-coated quartz, and
(d) surface profile across the grinding track of sample (c).
The small arrow indicates the materials transferred from the steel bearing.
158
'i
Figure 3.3.21
SEM micrograph showing the bearing materials which are
left on the diamond surface after 30 min tribotest.
4.
CONCLUSIONS
Three surface issues related to diamond nucleation and growth by
microwave plasma enhanced CVD had been studied. First, initial study of
diamond nucleation and growth on single crystal copper was initiated. It
involved attempts to nucleate diamond on the modified surface o f copper
crystal and hoped to provide some clues about how diamond thin films can
be grown heteroepitaxially on the substrate. Second, the various difficulties
of dealing with CVD diamond nucleation and growth on metal substrates
were confronted. These included the surface catalytic effects o f transition
elements, the large thermal expansion coefficient, and the rapid carbon
diffusion into the metals.
Third, it entailed the issue of film surface
roughness.
First of all, a study of nucleation and growth of diamond on single
crystal Cu had been performed. The interest on this study stems from the
fact that Cu has a lattice constant close to diamond (acu - 3.6148 A and
^diamond = 3.5671 A). Currently, the heteroepitaxial diamond growth is
being actively pursued due to the unusual electrical properties of diamond
which are highly desirable for high power electronic device applications.
The present technology of high quality single crystal diamond film growth
thus far is only limited to homoepitaxy (diamond on diamond). Diamond
nucleating on non-diamond substrate is very difficult due to the extremely
high surface energy of diamond ( y ( l l l ) ~ 5.3 J/m 2 ; y(lOO) ~ 9.2 J/m 2)
(Field, 1979). An attempt to grow epitaxial diamond film on single crystal
159
160
nickel substrate (aNi = 3.5239 A) had been carried out (Belton and
Schmieg, 1969) with no success. The reason was that nickel preferentially
turns the carbon gas species into soot or graphitic-like carbon structure
(Banaijee et al., 1961 ; Trimm, 1983). Koizumi et. al. (1990) reported
that epitaxial diamond films can be grown on single crystal c-BN substrate
(a = 3.616 A). However, a large single crystal of this material is even
harder to come by than diamond itself.
One of the unique properties of Cu is that it does not form carbide
when it interacts with carbon. This is considered desirable for diamond
heteroepitaxial growth. Nonetheless, the surface energy o f Cu, as of most
other materials, is very small compared to that of diamond (ycu(lOO) = 2.08
J/m 2). It is therefore necessary to modify the surface of Cu in order to
alter its surface energy. In this thesis, the surface of single crystal Cu was
implanted with carbon at the following conditions :
energy = 60 - 125 keV
ion dose =
1
x
1 0 18
cm ' 2
temperature = room temperature - 860 °C
Rutherford backscattering/channeling analysis revealed that the carbon
layer implanted at elevated temperature was mostly on the surface of the
Cu. The thickness of the layer was approximately 900 A as estimated from
the ion dose and density of graphite. In the case of Cu crystal implanted at
161
room temperature, the carbon layer was embedded approximately 700 A
deep inside the Cu.
X-ray and TEM analysis showed that the carbon layer implanted at
elevated temperature (>820 °C) was turbostratic graphite, with the c-axis
of the hexagonal lattice perpendicular to the Cu surface. The films,
however, were nearly randomly oriented about azimuthal direction. The
(0002) reflection of the graphite was found to be misoriented within 0.5°
with respect to the C u ( lll ) or Cu(200) reflections. The d<0 0 0 2 ) spacing
was calculated to be 3.356 A, in good agreement with the theoretical value
of 3.354 A. The graphite layers were in the form of 20-30 /<m islands and
were form ed by the out-diffusion process during cooling through the
carbon dissolution-precipitation mechanism. Interestingly, because of the
significant difference in coefficient of thermal expansion between graphite
and Cu ( a a axis,graphite = 0.9 x 10*6 /K at 800 °C , c.f. a cu = 24 x 10*6
/K), the graphite films experienced a large com pressive stress during
cooling, resulting in crackings of approximately
1000
A wide.
Diamond crystals had been grown on virgin Cu (111) and Cu (111)
implanted with carbon at 820 °C. The diamond crystals obtained were
polycrystalline and no indication o f epitaxial growth or growth with
preferred orientation had been observed in any of these samples.
The
graphite layer formed on top of Cu substrate at high tem perature was
found to enhance diamond nucleation quite significantly. However,
162
diamond crystals grown on Cu were found to be smaller in the case of
implantation carried out at room temperature than that at 820 °C under
exactly the same diamond deposition conditions. This was believed to be
caused by the induction period for the carbon species embedded inside Cu
substrate to come up to the surface to provide nucleation sites for diamond.
Interestingly, diamond crystals were found to nucleate preferentially on the
edges of the graphite layers.
This result was further supported by the
experiments conducted on HOPG crystals, in which diamond crystals were
found primarily at the edges of the graphite sheet. A striking difference in
the diamond nucleation density on the basal and prism plane o f graphite
was also observed.
T he observed phenomena is not surprising at all when one
considers the fact that the chemistry of the basal and prism (edge) planes of
graphite is very different. Knowing that the bonding on the basal plane of
graphite lattice is saturated with 7t (sp2) bonding, the carbon-carbon bond
on this plane is very strong, and in fact it is stronger than carbon-carbon
bond in diamond.
Each carbon at the prism plane, however, has one
dangling bond resulting in a lower bond strength.
T his is strongly
supported by the many previous studies o f catalyzed and uncatalyzed
reactions of graphite with oxygen and hydrogen, which indicated a strong
anisotropy behavior between the prism and basal plane of graphite.
163
A study of graphite etching by atomic hydrogen had been carried
out in time series. The results showed that deterioration of the lattice by
atom ic hydrogen started at a very early stage (i.e. after approximately
1
minute etching) by the creation of many lattice vacancies. A fter threeminutes of plasma exposure, the surface of graphite had been severely
roughened. Hexagonal etch pits were seen and becoming clearer after five
minutes o f the etching experiment. The formation o f the etch pits with
hexagonal symmetry stems from the anisotropic reactivity of carbon along
the principal crystallographic orientations on the basal planes of graphite
(< 1 0 l 0 > and < 1 1 2 0 >).
It was concluded from this study that a simpler and energetically
fav o rab le way o f nucleating diam ond on graphite was by first
hydrogenating the carbon atoms at the edges o f graphite to form sp 3
hybridized CH 3 bonds . This was done through the disruption o f carbon
atom s in the {112L} planes which was known to be the most likely
hydrogenation mechanism (Zielke and Gorin, 1955 ; Tom ita and Tamai,
1974).
It was then follow ed by reaction with hydrocarbon radical
precursors from the gas phase to form diamond. A simple lattice model of
a diamond nucleus on the edge of the graphite lattice was then constructed.
A ccording to the proposed model, the {111} plane o f diam ond was
perpendicular to the basal plane of graphite and the < 1 1 0 >diamond is
parallel to the <112Q>graphite- The proposed model was reasonable because
the carbon-carbon length in the < 1 1 0 > direction of diamond (2.507 A)
164
matched that in the < 1 1 2 0 > direction of graphite (2.459A) within
2
%. In
fact, in the high pressure synthesis of diamond from graphite using the
explosion shock quenching method, the row of carbon atoms along the
m
< 1 1 2 0 > o f graphite was believed to transform into the < 1 1 0 > of diamond
(W heeler and Lewis, 1975).
The model was also in agreement with the
previous study o f graphite formation on CVD diamond films in which
graphite was found to form with its basal plane matches on or close to the
{111} planes of diamond (Zhu et al., 1969).
The results concluded that diamond heteroepitaxial growth on the
modified surface of single copper were not obtained eventhough surface
modification of Cu by carbon implantation had been carried out. There
are still many phenomena which need further detailed studies, such as a
thorough study o f the carbon precipitation after the ion implantation at
elevated temperature.
Confronting the issue of diamond nucleation and growth on metal
substrates, particularly the transition elements such as Fe, Co, Ni or Pt, is
nothing but simple.
One of the major problems is the widely known
catalytic effect of those elements in forming coke from carbon containing
gases at elevated temperature.
The rate of carbon diffusion into most
metals is known high (e.g. Dc-->Fe, 800 °C ~ lO-8 cm 2 /sec) in such a way
that the carbon species from the diamond forming plasma environment are
consumed significantly by the materials. It is commonly known also that
165
the thermal expansion coefficient o f most metals far exceeds that of
diamond. Differences in thermal properties between the metal substrates
and the deposited films can introduce severe stress at the film substrate
interface. It is also pivotal to consider the fact that diamond films alone
cannot provide suitable protection for sheet metal owing to the stiffness of
diamond.
A novel concept o f composite film s consisting o f d ia m o n d
crystallites, hydrogenated amorphous carbon (a-C:H) and/or fluorocarbon
polymer was invented for coatings on metal surfaces. A buffer amorphous
Si layer was deposited on transition metal elements prior to CVD diamond
growth. This was to prevent the catalytic effect o f those elements and to
isolate the carbon from diffusing into the metals. The main reason o f using
a-C:H and fluorocarbon films was that the films are relatively more
flexible com pared to diamond.
Moreover, a-C:H and flu o ro c a rb o n
polymer films are known to possess diamond-like and teflon-like properties
respectively, making the composite films suitable for sheet metal coating
applications.
In constructing the composite films, diamond crystals with sizes
ranging from 2000-5000 A were grown on carbon steel and 304 stainless
steel substrates using the previously developed microwave pulsed plasma
technique. The small diamond grain size was highly desirable in producing
smooth composite films. Diamond nucleation density in terms o f surface
166
area was designed to be approximately 50-75% of the total substrate
surface. The a-C:H and fluorocarbon polymer films, each approximately
4000 A thick, were then deposited on the diamond crystals covered
substrates. Three different types of multilayer diamond composite film s
were o b tain ed :
1.
a-C:H film and diamond on 304 stainless steel
2. fluorocarbon film and diamond on 304 stainless steel
3. a-C:H, fluorocarbon films, and diamond on 304 stainless steel.
Three-point bending tests were carried out to ev alu ate the
bendability of the composite films.
From the load versus displacement
data, a critical sample displacement corresponding to the first stage of film
cracking ( 6 C) was obtained. The values obtained were ten degrees for
samples (1) and (3) and thirteen degrees for sample (2). It was originally
expected that the fluorocarbon polymer film inserted between diamond and
a-C:H film in sample (3) would experience cracking at
6
C somewhere
intermediate between that corresponding to sample (1) and sample (2). In
contrary, the good adhesion of the polymer film to the substrate and the aC:H film resulted in polymer film fracture at approximately the same 6 c as
the a-C:H film.
The fracture behavior of the composite films was found to be
different from the case of fracture in brittle bulk solids, in which the
167
failure mechanism originates from the weakest links in the materials and
spreads freely without any crack arrest. The propagation of the cracks in
the com posite films produced by the bending was restricted to one
dimension and stopped by the substrate due to the large differences in
fracture toughness between the coating and the substrate materials. The
generated cracks were found to be parallel and with nearly equal distances
to each other and perpendicular to the loading force. Arrest of cracks was
also observed, possibly due to inhomogeneity of the composite films and
the interactions of cracks with each other.
The scratch adhesion tests using a 200 pm diamond stylus had been
carried out. The results showed that upon imposing a load as high as
68
newtons (~7 kg), the composite films still adhered to the steel substrate.
However, film cracks were observed even at load as low as 20 newtons.
The cracks showed semicircular trajectories parallel to the leading edge of
the diamond stylus.
This type of film failure is commonly called the
conformal cracking, and it was caused by the tensile bending moments
within the coating as the diamond stylus was deforming the coating and the
substrate.
As the scratch load reached
68
newtons, the cracks were
developing almost perpendicular to the direction o f the stylus track. This
cracking mode is intermediate between the conformal and the tensile
cracking. It is interesting to note that when a continuous diamond film was
grown on the steel substrate, the film spontaneously blistered off owing to
thermal expansion difference.
168
Subjecting the composite films with various chemical tests (i.e.
H N 0 3 : HC1 : H 2 O (3:7:30), 48% HF, 96 % H 2 SO 4 , 37 % HC1, 71%
HNQ 3 , 5 g NaCl in 100 cc H 2 O, CCI4 , and CHCI3 solutions) resulted in no
effects at all on the films.
This implies that the films were inert and
impermeable to the chemicals.
Finally, the effect of surface roughness o f diamond films on the
optical properties of the materials was examined in this thesis. The issue is
considered very important for the technological applications o f CVD
diamond thin films. As well known from the elements o f optical theory,
surface roughness (a ) can strongly affect the optical tranparency of a
sample. According Filinski's formulation (1972), the optical transmittance
o f a sample is exponentially dependent o f (on/X )2, where n and X are
respectively the refractive index of a sample and wavelength. For a 1 /<m
thick diamond sample where n = 2.4, there is a factor of ~2.75 difference
in the intensity of light transmission for o =X/100 and o = X/10. Surface
roughness of much less than 800 A is required for the diamond film
surface to be considered smooth in the visible spectrum according to the
Raleigh criterion.
It therefore requires diamond thin films to be very
sm ooth in order to m inimize the incoherent light scattering which
eventually leads to surface haziness and poor transparency.
D iam ond
possesses very good optical transparency. Diamond films with large grain
size and thus rough surfaces render them futile for optical coating
purposes.
169
A pulsed microwave plasma technique was invented to deposit thin
and smooth diamond films on fused quartz. It is known from classical
theory of nucleation, the temperature for optimal nucleation is known to be
lower than that for growth. Therefore, this pulsed technique m anipulated
the idea of depositing diamond at relatively low average temperature.
There were three process parameters which were found to affect the final
diamond grains size and thus the film surface roughness : the time duration
o f each cycle (At), the final temperature (Tf) and the number o f process
cycles (np). For a given At, the np used to form an uniform and continuous
film was different. As At decreased, the np needed to cover the substrate
surface with continuous diamond films increased. The following results
were ob tain ed :
- At=60 min., Tf=800 °C, one cycle, TaVg=800 °C, o=1600 A
- At=4 min., Tp=800 °C, four cycles, TaVg=600 °C, o=150 A
- At=45 sec., Tp=600 °C, 16 cycles, TaVg=415 °C, o = 1 0 0 A
Diamond films as smooth as 50 A had been deposited on fused quartz
substrates using At=20 sec., Tf=500 oC, and n p=20 cycles.
Raman
spectroscopy results showed that as np increased, the content of amorphous
carbon phase increased as well. However, the films obtained using smaller
np were found to contain more graphitic components.
The effect of diam ond film surface roughness on optical
transmission spectra had been examined. The transmission of the rough
170
film (o=2000 A) was lower than that o f the smooth film (o=200 A) of
equivalent thickness by at least 20 %, which is clearly significant. The
optical absorption edge of the films was at 0.225 pm , similar to the type Ila
diamond. Tailing near the absorption edge was observed and believed to
be from the structural imperfection and impurities in the film (such as a-C,
graphite, and sputtered Si from fused silica reaction tube), as well as the
internal light scattering at the grain boundaries.
Optical tran sm ission
measurements taken on samples with diamond film thicknesses ranging
from 3000-8000 A showed no significant differences in the average optical
transmittance.
No discernible grain size was observed.
Block-on-ring
tribotest results revealed that the diamond films adhered to the quartz
substrates very well.
Transmission electron microscopy results revealed the presence of
many stacking faults and 1= 3 primary twins, with {111} planes as the
boundaries.
The energy of the 1 = 3 boundary as calculated using the
Tersoff potentials (Narayan, 1990) is very close to zero indicating the
preferential formation o f this type of twin in CVD diamond crystals.
Diamond crystals with fivefold symmetry resulting from twinning in the
{111} planes were also observed. The angles between the {111} planes
were measured to be falling between 70.5° (ideal case) and 74°, and the
sum of all the angles was 360°. The plane boundaries are of coherent type.
Since the diamond lattice is very rigid and with the high energy constraints
in forming defects such as dislocations, it is reasonable to say that the misfit
171
o f the plane boundaries was mainly accommodated by elastic strains
(Matsumoto and Matsui, 1983).
5.
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VITA
Tiong P. Ong
DATE O F BIRTH : March 3, 1964
PLACE OF BIRTH : Palembang, Indonesia
CITIZEN SH IP: Indonesia
EDUCATIONS:
N o rth w estern U niversity, Evanston, Illinois
PhD., Materials Science and Engineering (1991)
Thesis Advisor : Professor Robert P.H. Chang
M.S., Chemical Engineering (1986)
Project Advisor : Professor Harold Kung
F lo rid a In stitu te o f Technology, Melbourne, Florida
B.S., Chemical Engineering (1985, with highest honor)
PUBLICATIONS :
T .P. Ong, F.L. Xiong, R.P.H. Chang, and C.W. White, "A Mechanism
for Diamond Nucleation and Growth on Single Crystal Copper Surfaces,"
submitted to Applied Physics Letters.
T .P . Ong, F.L. Xiong, R.P.H. Chang, and C.W. White, "Diamond
Nucleation and Growth Studies on Carbon Implanted Copper Surfaces,”
in preparation for publication in Journal Materials Research.
T .P . O ng and R.P.H. Chang, "Properties of Diamond Composite Films
Grown on Iron Surfaces," AppI. Phys. Lett. 58, 358 (1991).
T .P . O ng and R.P.H. Chang, "Flexible Diamond Composite Films for
Sheet Metal Coating Applications," Proc. of the 2nd. Int. Conf. on the New
Diamond Science and Technology, Washington D.C., September 1990.
184
185
T .P. O ng and R.P.H. Chang, "Low Temperature Deposition of Thin
Diamond Films for Optical Coating Applications," AppI. Phys. Lett. 55,
2063 (1989).
T .P. Ong, W.A. Chiou, F.R. Chen, and R.P.H. Chang, "Preparation of
Nanocrystalline Diamond Films for Optical Coating Applications Using A
Pulsed Microwave Plasma CVD Method," Extended Abstract, Workshop
on the Science and Technology of Diamond Thin Films, May 20-24, 1990,
Concord, Ohio.
M.S. Wong, R. Meilunas, T.P. Ong, and R.P.H. Chang, "Tribological
Properties of Diamond Thin Films Grown by Plasma Enhanced Chemical
Vapor Deposition," AppI. Phys. Lett. 54, 2006 (1989).
M.S. Wong, R. Meilunas, T .P. Ong, and R.P.H. Chang , "Thin Diamond
Films for Tribological Applications," in New Materials Approaches to
Tribology : Theory and Applications, edited by L.E. Pope, L.
Fehrenbacher, W.A. Winer (Materials Research Society : Pittsburgh,
1989), Proc. Vol. 140.
R. Meilunas, K. Sheng, M.S. Wong, T.P. Ong, and R.P.H. Chang, "The
Initial Stages of Plasma Synthesis of Diamond Films," in Laser and Particle
Beam Chemical Processes on Surfaces, edited by A. Wayne Johnson, Gary
L. Loper, and T.W. Sigmon (Materials Research Society : Pittsburgh,
1989), Proc. Vol. 129, pp. 533-538.
R.P.H. Chang, R. Meilunas, T.P. Ong, and M.S. Wong, "Application of
Thin Diamond Films for Protective Coatings," in Technology Update on
Diamond Films, edited by R.P.H. Chang, D. Nelson, and A. Hiraki
(Materials Research Society : Pittsburgh, 1988), Proc. EA-19,
pp. 171-175.
HONORS AND AW ARDS:
Member of Tau Beta Pi and Alpha Sigma Mu, Florida Institute of
Technology Distinguished Scholar, W.P. Murphy Fellowship (1985),
American Institute of Chemist Award (1985), SPIE Scholarship (1989),
Lee Scholarship (1988-1991).
fOMPi
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