Note: Descriptions are shown in the official language in which they were submitted.
CA 02489199 2004-12-03
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CERAMIC THIN FILM ON VARIOUS SUBSTRATES, AND PROCESS FOR
PRODUCING SAME
CROSS REFERENCE TO RELATED APPLICATIONS
This patent application is a continuation-in-part
application of co-pending US application serial No. [not yet
assigned], filed November 22, 2004, entitled "Ceramic Thin
Film On Various Substrates, And Process For Producing Same",
which is the 371 National Phase of International Application
No. PCT/CA03/000763, filed May 23, 2003, which was published
in English under PCT Article 21(2) as WO 03/100123 A1. All
of these applications are incorporated herein in their
entirety.
FIELD OF THE INVENTION
The present invention relates to a thin film of an
amorphous silicon-based material on a substrate.
BACKGROUND OF THE INVENTION
Kitabatake et al., disclose in US Patent
6,270,573, CVD and CVD-related methods of producing silicon
carbide substrates, including the growing of silicon carbide
film by supplying separate silicon atoms and carbon atoms on
a surface. The silicon-carbon bond formation occurs mainly
on the surface of the substrate, a step that usually
requires high temperatures, in this particular case the
required temperature being 1300°C. MBE and MO-CVD may use
species that contain a limited number of pre-existing Si-C
bonds in the precursor, this number being usually related to
precursor synthesis requirements. Further, MBE and MO-CVD
rely mainly on the creation of Si-C bonds on the surface of
the substrate and prevent rearrangement of Si-Si/C-C bonds
adventitiously produced in the process to the desirable Si-C
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bonds. The rate of SiC production is limited to the number
of Si-C bonds produced on the surface of the substrate. The
synthesis of the material is based on Si-H and C-H bond-
dissociation and Si-C bond production, inducing a
significant concentration of residual chemical defects,
including residual dangling bonds.
Kong et al. (European Patent No. EP 0,970,267)
describe a susceptor design for silicon carbide resulting in
minimizing or eliminating thermal gradients between the two
surfaces of a substrate wafer. The CVD and CVD-related
deposition procedures of Kong et al., require strict control
of the temperature field and the gas flow at the surface of
the substrate, where the Si-C bond formation is occurring.
Grigoriev et al. (Grigoriev, D. A., Edirisinghe,
M. J., Bao, X., Evans, J. R. G. and Luklinska, Z. B.(2001)
"Preparation of silicon carbide by electrospraying of a
polymeric precursor," Philosophical Magazine Letters (UK),
81, 4, 2001 by Dept. of Mater., Queen Mary Univ. of London,
UK) present silicon carbide coatings and films prepared for
the first time by electrostatic atomization of a solution of
a polymeric precursor and deposition onto alumina and
zirconia substrates. In the method of Grigoriev et al., the
polymeric source already contains most of the Si-C bonds
required for the formation of the SiC film; however, the
molecular source is carried to the surface inside cages of
solvent molecules, implicitly leading to contamination of
the film, shrinking and outgassing phenomena, due to solvent
evaporation and polymer cracking. These effects will be
present in any polymer-source method (for example, spin-
coating, spraying, laser ablation, etc.). The synthesis of
material directly from a condensed source (the liquid
solution of the polymer in aerosol, or "spray"-form) as
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opposed to the synthesis from a gaseous precursor, prevents
a submicron-level roughness required for the efficient
exploitation of semiconductor/electronic properties of the
SiC-material.
Lau et al. (Lau, S. P., Xu, X. L., Shi, J. R.,
Ding, X. Z., Sun, Z: and Tay, B. K. (2001) "Dependences of
amorphous structure on bias voltage and annealing in
silicon-carbon alloys," Materials Science & Engineering, B85
(16), Sch. of Electr. & Electron. Eng., Nanyang Technol.
Inst., Singapore) report amorphous silicon-carbon alloy
films that have been obtained by a filtered cathodic vacuum
arc (FCVA) technique. They have observed that the disorder
of the Si-C network increased with using the high bias
voltages during the deposition. This high disorder in the
film with high bias voltages induces the smaller nanometer
crystallites after annealing at 1000°C rather than low bias.
The Raman peaks shift to the high frequency with increasing
the annealing temperature up to 750°C due to the increase of
nanometer grain size at the same bias. A sharp transition
from nanocrystalline to polycrystalline can be observed when
the films are annealed under 1000°C. The nanometer
crystallites of Si-C alloy used to obtain an increase of
grain-size via annealing induced new Si-C bond formation in
the bulk of the material. However, there was no
redistribution of pre-existent Si-C bonds in a polymeric
residue.
Jana of al. (Dana, T., Dasgupta, A. and Ray, S.
(2001) "Doping of p-type microcrystalline silicon carbon
alloy films by the very high frequency plasma-enhanced
chemical vapor deposition technique" Journal of Materials
Research, 16(7) 2001, 2130-5, Energy Res. Unit, Indian
Assoc. for the Cultivation of Sci., Calcutta, India) present
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the synthesis of p-type silicon-carbon alloy thin films by
very high frequency plasma-enhanced chemical vapor
deposition technique using a SiH4, H2, CH4, and BZH6 gas
mixture at low power (55 mW/cm2) and low substrate
temperatures (150-250°C). Effects of substrate temperature
and plasma excitation frequency on the optoelectronic and
structural properties of the films were studied. A film
with conductivity 5.75 Scm 1 and 1.93 eV optical gap Eo4 was
obtained at a low substrate temperature of 200°C using 63.75
MHz plasma frequency. The crystalline volume fractions of
the films were estimated from the Raman spectra. They
observed that crystallinity in silicon carbon alloy films
depends critically on plasma excitation frequency. When
higher power (117 mW/cm2) at 180°C with 66 MHz frequency was
applied, the deposition rate of the film increased to 5.07
nm/min without any significant change in optoelectronic
properties. The introduction of the dopant is based on bond-
dissociation and bond reconstruction, increasing the
concentration of residual chemical defects, such as residual
dangling bonds.
Yamamoto et al. (Yamamoto et al., Diam. Relat.
Mater., vol. 10 (no. 9-10), 2001, pp. 1921-6) present a
doping procedure whereby amorphous SiCN films were
prepared on Si (100) substrates by nitrogen ion-assisted
pulsed-laser ablation of a SiC target. The dependence of
the formed chemical bonds in the films on nitrogen ion
energy and the substrate temperature was investigated by
X-ray photoelectron spectroscopy (XPS). The fractions of
sp2 C-C, spa C-C and sp2 C-N bonds decreased, and that of
N-Si bonds increased when the nitrogen ion energy was
increased without heating during the film preparation. The
fraction of sp C-N bonds was not changed by the nitrogen
ion irradiation below 200 eV. Si atoms displaced carbon
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atoms in the films and the spa bonding network was made
between carbon and silicon through nitrogen. This tendency
was remarkable in the films prepared under substrate
heating, and the fraction of spa C-N bonds also decreased
when the nitrogen ion energy was increased. Under the
impact of high-energy ions or substrate heating the films
consisted of sp2 C-C bonds and Si-N bonds, and the
formation of spa C-N bonds was difficult. The Yamamoto
procedure proposes a doping step separate from the
synthesis step.
Budaguan et al. (Budaguan, B. G.; Sherchenkov,
A. A.; Gorbulin, G. L.; Chernomordic, V. D. (2001) "The
development of a high rate technology for wide-bandgap
photosensitive a-SiC:H alloys," Journal of Alloys and
Compounds, 327(30) Aug., 146-50, Inst. of Electron
Technol., Moscow, Russia) discuss in their paper the
deposition process and the properties of a-SiC:H alloy
fabricated for the first time by 55 kHz PECVD. It was found
that 55 kHz PECVD allows an increase in the deposition rate
of a-SiC:H films.
Modiano et al. (Japanese patent No. 145138/95)
present a process for producing silicon carbide fibers
having a C/Si molar ratio from 0.85 to 1.39, comprising the
steps of rendering infusible the precursory fibers made from
an organosilicon polymer compound, then primarily baking the
infusible fibers in a hydrogen gas-containing atmosphere.
This process for producing silicon carbide thin films
comprises the steps of imparting semiconductor properties to
passivating or dielectric thin films from volatile
precursory species produced from organosilicon polymer
compounds.
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Yang et al. (Yang, Lixin; Chen, Changqing; Ren,
Congxin; Yan, Jinlong; Chen, Xueliang, "Synthesis of SiC
Using Ion Beam and PECVD", International Conference on
Solid-State and Integrated Circuit Technology Proceedings,
pp. 811-814) present a process for producing silicon carbide
thin films comprising the steps of conferring semiconductor
properties to passivating or dielectric thin films from
volatile precursory species produced from organosilicon
polymeric compounds.
SUMMARY OF THE INVENTION
According to one aspect of the present invention,
there is provided a film of an amorphous silicon-based
material on a substrate, the film having a carrier
concentration of 1.013 to lOlg cm 3 in a depletion zone next to
the substrate.
According to another aspect of the present
invention, there is provided a film of an amorphous silicon-
based material on a substrate, the film having an electron
mobility of 5 to 30 cm2V-ls-1.
According to still another aspect of the present
invention, there is provided a film of an amorphous silicon-
based material on a substrate, the film having a dangling
bond concentration of 1012 to 1O19 Cm 3.
According to yet another aspect of the present
invention, there is provided a film of an amorphous silicon-
based material on a substrate, the film having no solvent-
related defects.
According to a further aspect of the present
invention, there is provided a film of an amorphous silicon-
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based material on a substrate, the film having a residual
hydrogen concentration of 0 to 25 atomic %.
According to yet a further aspect of the present
invention, there is provided a semiconductor device
comprising a film of an amorphous silicon-based material on
a substrate, the film having a carrier concentration of 1013
to 101$ cm 3 in a depletion zone next to the substrate.
According to still a further aspect of the present
invention, there is provided a semiconductor device
comprising a film of an amorphous silicon-based material on
a substrate, the film having an electron mobility of 5 to 30
cm2V-ls-1.
According to another aspect of the present
invention, there is provided a semiconductor device
comprising a film of an amorphous silicon-based material on
a substrate, the film having a dangling bond concentration
of 1012 to 1019 Cm 3.
According to yet another aspect of the present
invention, there is provided a semiconductor device
comprising a film of an amorphous silicon-based material on
a substrate, the film having no solvent-related defects.
According to another aspect of the present
invention, there is provided a semiconductor device
comprising a film of an amorphous silicon-based material on
a substrate, the film having a residual hydrogen
concentration of 0 to 25 atomic o.
In an embodiment, the semiconductor device is a
solar cell, light-emitting diode, Schottky diode, a
transistor, a photothyristor or an integrated monolithic
device on a single chip.
CA 02489199 2004-12-03
BRIEF DESCRIPTION OF THE DRAWINGS
In the accompanying drawings, which illustrate an
exemplary embodiment of the present invention:
Figure la is a schematic diagram of a Polymer
Source Chemical Vapor Deposition (PS-CVD) reactor;
Figure 1b is a diagram of the temperature-profile
in the reactor;
Figure 2 is a schematic diagram of the FT-IR cell
for monitoring the production of thin film by the PS-CVD
process;
Figure 3 is an infrared spectrum of a gaseous
precursor formed during deposition of an amorphous silicon
carbide thin film prepared by the PS-CVD process;
Figure 4 is a representation of the set of
chemical reactions leading to an n-type amorphous silicon
carbide thin film from a generic polysilane precursor;
Figure 5 is a series of substrates that may
support a thin film produced by the PS-CVD process;
Figure 6 is a thermogram from a thermogravimetric
analysis of a polymeric source subjected to the PS-CVD
process;
Figure 7 illustrates a simple Schottky solar cell;
Figure 8 illustrates a p-n junction barrier
photovoltaic cell based on amorphous silicon carbide thin
film;
Figure 9 illustrates a stacked p-n junction solar
cell based on amorphous silicon carbide thin film;
CA 02489199 2004-12-03
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Figure 10a illustrates an energy band diagram of
Schottky structure before intimate contact between metal and
semiconductor;
Figure lOb illustrates an energy band diagram of
Schottky structure after intimate contact between metal and
semiconductor;
Figure lla illustrates the energy band of a p-n
junction solar cell;
Figure llb illustrates the energy band of a p-n
junction solar cell;
Figure 12 is an ERD graph illustrating the
relationship between concentration and depth for an n-type
amorphous silicon carbide thin film;
Figure 13 is a graph illustrating the relationship
between carrier concentration and width of the depletion
zone for an n-type amorphous silicon carbide thin film
produced by the PS-CVD process;
Figure 14 is a graph illustrating the relationship
between current and voltage for an n-type amorphous silicon
carbide thin film produced by the PS-CVD process;
Figure 15 is an image from a scanning electron
microscope of an n-type amorphous silicon carbide thin film;
Figure 16 is a spectrum of the photoluminescence
of an n-type amorphous silicon carbide thin;
Figure 17 is an EPR spectrum of an n-type
amorphous silicon carbide thin;
CA 02489199 2004-12-03
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Figure 18 is an infrared spectrum of a p-type
amorphous silicon carbide thin film on a silicon single
crystal wafer;
Figure 19a is a graph illustrating the
relationship between the capacitance and voltage of a p-
type amorphous silicon carbide thin film on a silicon
single crystal wafer;
Figure 19b is a graph illustrating the
relationship between the carrier concentration and depth of
a p-type amorphous silicon carbide thin film on a silicon
single crystal wafer;
Figure 20 is a UV spectrum of a p-type amorphous
silicon carbide thin film on a silicon single crystal
wafer;
Figure 21 are ZgSi NMR spectra recorded at
reaction temperatures of 300°C, 400°C, 500°C,
600°C, 700°C,
800°C, and 1100°C for an amorphous silicon carbonitride thin
film;
Figure 22 is a contour plot illustrating the
relationship between the 29Si NMR chemical shifts and
reaction temperatures of an amorphous silicon carbonitride
thin film;
Figure 23 is an infrared spectrum of an amorphous
silicon carbonitride thin film;
Figure 24 is an infrared spectrum of an amorphous
silicon nitride thin film;
Figure 25a are infrared spectra of the oxidation
of SiONl in air for 10 minutes at 300, 550, 600, 700°C;
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Figure 25b are infrared spectra of the oxidation
of Si0N2 in air at 500°C for 10, 20, 30, 42 and 52 minutes;
and
Figure 26 are infrared spectra of amorphous
silicon oxycarbide thin films (a), (b), and (c) annealed at
1100°C for 8, 16, and 24 hours, respectively.
DETAILED DESCRIPTION OF THE INVENTION
The present invention is concerned with an
innovative method of forming volatile pyrolysis products
from silicon-based polymeric compounds or sources that are
to be used as gaseous precursors to thin films. This method
is herein called Polymer Source Chemical Vapor Deposition
(PS-CVD). The thin films produced by the method are
durable, may be flexible due to their thinness, and may be
used for high-performance semiconductors. Consequently,
the present invention provides a thin film of an amorphous
silicon-based material on a substrate with improved physical
characteristics and properties.
The term "an amorphous silicon-based material"
refers to both an amorphous and a nanocrystalline silicon-
based material. An amorphous silicon-based material is the
non-crystalline form of a silicon-based material. A silicon-
based material normally has a silicon atom tetrahedrally
bonded to four neighbouring atoms. This is also the case in
amorphous silicon-based materials, however, they do not show
long-range order of the tetrahedra in the crystalline
lattice as in crystalline silicon-based materials. In
addition, some silicon and/or carbon atoms may have dangling
bonds, which occur when they do not bond to four neighboring
atoms. These dangling bonds are defects in the continuous
random network. A nanocrystalline silicon-based material is
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similar to an amorphous silicon-based material, in that it
lacks long-range order (starting even in the second
coordination sphere), as is the case of an amorphous phase.
Where they differ, however, is that nanocrystalline silicon-
based materials have small grains of crystallites of
silicon-based materials ordered in the first coordination
sphere within the totally disordered amorphous phase. This
is in contrast to polycrystalline silicon-based material
which consists solely of crystalline silicon grains
separated by grain boundaries.
The film comprises a silicon-based material that
may be, but is not limited to, a silicon carbide, a silicon
carbonitride, a silicon nitride, a silicon oxycarbide, a
silicon oxynitride, or a silicon oxycarbonitride, in pure or
doped forms. In an embodiment, the film comprises a
dopant/impurity such as N, B, O, H, Cl, Al, Ga, In, Tl, P,
As, Sb, O, S, Se, Te, or Bi .
The film may be deposited on commonly available
substrates as a completely to partially amorphous
semiconductor, at a desired micron-range thickness, onto a
variety of commonly available and resistant substrates of
varying composition, degree of stiffness, shape, density, or
color over a small to an exceptionally large surface. For
example, the substrate may be regular ceramic material of a
complicated shape, quartz, electronic-grade sintered
alumina, polished alumina, silicon single crystal wafer,
graphite, stainless steel, polycrystalline semiconductor and
other commonly available and relatively inexpensive
materials, as well as any of the above materials coated with
metals or alloys.
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In an embodiment, the thin film of an amorphous
silicon-based material on a substrate has a carrier
concentration of 1013 to 1018 cm-3 in a depletion zone next to
the substrate, in a preferred embodiment the carrier
concentration is 1013 to 101' cm 3, and in a more preferred
embodiment the carrier concentration is 1013 to 1016 Cm 3. The
depletion zone is the zone where an electrical field exists,
sweeping the mobile charge carriers. In a further embodiment
the thin film of an amorphous silicon-based material on a
substrate has an electron mobility of 5 to 30 cm2V-ls-1, in a
preferred embodiment the electron mobility is 10 to 25 cm2V-
ls-1, and in a more preferred embodiment the electron
mobility is 15 to 20 cm2V-ls-1. Carrier concentration and
electron mobility may be determined by Hall measurements as
originally described by van der Pauw (L. J. van der Pauw,
Philips Res. Repts, 13, pp. 1-9, 1958 and L.J. van der Pauw,
Philips Tech. Rev, 20, pp. 220-224, 1958).
In an embodiment, the thin film of an amorphous
silicon-based material on a substrate has a dangling bond
concentration of 1012 to 1019 cm-3. By contrast, prior art
polycrystalline silicon-based thin films have a dangling
bond concentration of at least 1019 cm 3 (T. Christidis et
al., 2002, Applied Surface Science, 184, p. 268). A
dangling bond is a defect involving usually an unpaired
electron, or an "unsatisfied" bond. A dangling bond occurs
when an atom is missing a neighbour to which it would be
able to bind. Such dangling bonds are defects that disrupt
the flow of electrons as the dangling bonds may capture
electrons. These defects usually segregate to grain
boundaries, crystalline defects, and couple to chemical and
structural impurities. The dangling bonds in the film may
be assigned and their density evaluated via electron
paramagnetic resonance (EPR) (T. Christidis et al., 2002).
CA 02489199 2004-12-03
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In an embodiment, the thin film of an amorphous
silicon-based material on a substrate has no solvent-related
defects, for example voids in the thin film due to solvent
evaporation or alternatively solvent trapped in the thin
film during the deposition. Solvent-related defects do not
occur in the thin films produced by PS-CVD since no solvent
is used as in traditional deposition methods. Although a
unique defect that may arise in a thin film produced by the
PS-CVD process is the accidental incorporation of a polymer
chain.
In an embodiment the thin film of an amorphous
silicon-based material on a substrate has a residual
hydrogen concentration of 0 to 25 atomic o, in a preferred
embodiment the residual hydrogen concentration is 0 to 200,
and in a more preferred embodiment the residual hydrogen
concentration is 0 to 15 atomic %. "Residual hydrogen"
arises due to the nature of the starting materials used for
preparation of the thin films. For example, if
polyethylsilane, (SiC2H6)n were used as a polymeric source to
produce a thin film of a-SiC, 500 of the carbon and 1000 of
the hydrogen must be removed from the polymeric source to
obtain a stoichiometric product. However, formation of the
thin film is not always stoichiometric, typically resulting
in removal of between 50-1000 of the carbon and 60-1000 of
the hydrogen. Anything other than stoichiometric removal of
hydrogen will result in "residual hydrogen".
The grain size of the thin films described herein
may be 2 to 10 nm. In an embodiment, the grain size of
"green" or as deposited thin film is 2-3 nm, and may be
increased to 4-6 nm after annealing.
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The thin films described herein were prepared by
the PS-CVD process. Figure 1A illustrates an apparatus used
for the PS-CVD process while Figure 1B depicts the
temperature profile used within the apparatus. Referring to
Figure 1A, there is a quartz reactor (2), into which one or
several polymer derived precursors (19), enter through a gas
inlet (1). The quartz reactor (2) is also referred to as
the furnace or the PS-CVD reactor. Furthermore, a gaseous
atmosphere (60), either inert or active, also enters through
the gas inlet (1). The inert atmosphere may include argon,
nitrogen or other inert gases while the active atmosphere
may include gases such as ammonia, carbon monoxide or
similar gases. Before operation the reactor (2) is purged
with a selected atmosphere (60).
The gas inlet (1) has a high-vacuum seal to
minimize the ingress of oxygen impurities from the
surrounding air drawn into the reactor (2). The total
pressure in the reactor is measured with a pressure
controller (3) that also controls the flowrate into the
reactor (2). The outside of the reactor is heated with
electric heating elements (4), which surround the reactor
(2) to produce a temperature gradient as illustrated in
Figure 1B. There are additional heating elements (4) near
the inlet of the reactor (2), while there are fewer
surrounding the deposition area of the reactor. The PID
(Proportional-Integral-Derivative) temperature controller
(5) ensures that the temperature within the reactor (2) is
in the appropriate range for the polymer-derived precursor
(19) used, the type of gaseous atmosphere (60) and substrate
(6) to be coated. The substrate (6) is placed in the
deposition area of the reactor (2). Typically the substrate
(6), is a piece or part made of silicon, quartz, metal,
ceramics, or other materials described herein. The gas
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phase near the gas outlet (9), of the reactor (2), is
analyzed by a FT-IR spectrometer (7). The FT-IR
spectrometer (7) allows the in situ verification of the
deposition process and the presence of oxygen impurities. A
silicon-based film (8) is deposited on the substrate (6)
through the pyrolysis of the silicon based polymeric sources
and their chemical rearrangement. The deposited film can be
a single or multiple layered film.
Referring to Figure 2, there is represented a
cross-sectional view of a FT-IR cell based on a silicon
single crystal wafer (68), which is designed for the
monitoring of a film (8) deposited on the substrate (6) by
the PS-CVD process. The substrate (6) with a coating (8) is
held in place, the thickness of the deposited layer has been
exaggerated so that the upper half and lower half of the FT-
IR cell actually sit one on top of the other and are sealed
by the represented O-ring (67). There is a protective gas
swept through the FT-IR cell from an inlet (64) to an outlet
(65), which maintains the appropriate inert atmosphere. The
IR beam (61) is projected onto the substrate and it is bent
and reflected through the deposited film (8), and the
collected through a microscope objective (62) and detected
by a MCT (mercury-cadmium-tellurium, Hg-Cd-Te) detector
(63). The FT-IR cell is mounted on an adjustable 2D
micrometric stand (66) which allows the FT-IR to be adjusted
appropriately with respect to the IR beam (61).
Referring to Figure 3, the in situ FT-IR spectrum
analysis of a gaseous precursor at the outlet (9) of the
reactor (2) shows the numerous peaks that correspond to the
SiH bonds formed when chemical change in the structure of
the solid polymeric source produces a polymer derived
precursor (19). The increasing temperature near the inlet
CA 02489199 2004-12-03
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of the reactor (2) breaks down the polymeric source into
various subunits to produce the gaseous precursor (19).
The PS-CVD process distinguishes itself from other
forms of chemical vapor deposition because the synthesis is
based on high-density gaseous precursors (for example, MW
higher than 298 amu, from TGA-MS experiments) allowing high
mass transport to the substrate.
The PS-CVD process distinguishes itself from other
forms of chemical vapor deposition since the synthesis is
based on polymeric gaseous precursors that produce in situ
reactive networking functionalities (Si-H bonds, as captured
in the FT-IR analysis of the trapped gas phase precursor).
The PS-CVD process distinguishes itself from other
forms of chemical vapor deposition since the Si-H networking
functionalities required for the SiC-synthesis are 3-6 times
more productive than in regular CVD procedures due to the
high-mass of the individual networking precursors thereby
resulting in high deposition rates.
The PS-CVD process distinguishes itself from other
forms of chemical vapor deposition since the networking
functionalities do not require the intermediate production
of dangling bonds, a process replaced by a
thermodynamically-driven redistribution of pre-existent Si-C
bonds in the polymeric-source and in the in situ formed
gaseous precursor. This process may result in a lower
concentration of residual dangling bonds in the material.
The PS-CVD process further distinguishes itself
from other forms of chemical vapor deposition because it
does not require the following more sophiscated driving
forces:
CA 02489199 2004-12-03
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a) ion implantation, ion beam enhanced deposition,
reactive ion beam sputtering and plasma enhanced chemical
vapor deposition (PECVD);
b) RF power as described in "High temperature annealing
of hydrogenated amorphous silicon carbide thin films" INS
Ol-17 6910830 A2001-11-6855-060 (PHA) NDN- 174-0691-0829-
5 Yihua l~Tang; Jianyi Lin; Cheng Hon Alfred Huan; Zhe
Chuan Feng; Soo Jin Chua;
c) IR and UV laser photolysis as described in "Laser gas-
phase photolysis of organosilicon compounds: approach to
formation of hydrogenated Si/C, Si/C/F, Si/C/0 and Si/O
phases" INS 00-50 6791677 A2001-03-8250F-001 (PHA) NDN-
174-0679-1676-1 Pola, J., Proceedings of the Indian
National Science Academy, Part A (Physical Sciences); and
d) electron cyclotron resonance as described in
"Application of electron cyclotron resonance chemical
vapor deposition in the preparation of hydrogenated SiC
films. A comparison of phosphorus and boron doping" INS
98-04 5814339 A98058115H-021 (PHA); B9803-0520F-017 (EEA)
NDN- 174-0581-4338-2, S.F. Yoon and R. Ji.
Referring to Figure 1B, the temperature profile
within and along the length of the reactor (2) is
represented. The input zone (10) shows a constant lower
temperature associated with the gas inlet (1). In the
heating zone (11), there is an increase in the temperature
due to the large amount of heating elements (4) at the
inlet. In the heating zone (11), the rearrangement of the
silicon-based polymeric sources occurs which leads to the
formation of the poly(carbosilane) species. The next
temperature is that of the pyrolysis zone (12), where there
is direct precursor formation and doping occurring. This is
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followed by a reduction in the number of heating elements
(4) in the deposition zone (13) which consequently cools the
particularly zone of the reactor (2) thus lowering the
temperature. The deposition zone (13) is represented by a
constant temperature wherein the silicon-based film is being
deposited on the substrate (6). Finally, there is the gas
exit zone (14), where the temperature falls at the gas
outlet (9) of the reactor (2) and approaches that of the
ambient temperature outside the reactor (2). The
temperature in the reactor varies between 100°C and 1000°C
depending on the stage and the specific local requirements
of the process steps in the aforementioned reactions. The
gaseous species are monitored by FT-IR spectroscopy (7) of
samples extracted near the reactor outlet (9). The
amorphous silicon-based thin films may also be characterized
by IR spectroscopy while the concentration of adventitious
oxygen in the thin film can be measured by using a
Czochralski silicon window as a standard (Scarlete, M., J.
Electrochem. Soc., 1992, 139(4), p. 1207).
The PS-CVD design allows the use of a broad series
of polymeric precursors for use in the synthesis of silicon-
based thin films including, but not limited to, oxides,
nitrides, carbides and variously weighted combinations in
homogeneous phases or multi-layered structures. The
resulting films are of considerable interest as electronic
and optoelectronic materials as well as for protective
coatings. A large variety of appropriate silicon-based
polymeric sources that do not present Kumada-type
rearrangements, but can be cracked, vaporized directly or
via an intermediate liquid phase, or chemically transformed
in gaseous species in the reactor, may be used, such as
polysilanes, polycarbosilanes, polycarbosilazanes,
polysiloxanes and polysiloxazanes. Other appropriate
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polymeric sources may be used, such as carbon nitride
polymeric sources and boron nitride polymeric sources.
Possible carbon nitride polymeric sources (CxNy) may be, for
example, polyepoxy-, polyamides, polyamines, polyimides,
polyureas and polyurethanes. It is noted that polyamides,
polyamines, polyimides and polyureas may be used in mixtures
with other polymeric sources and as a possible source of
nitrogen in the reaction. Polymers sources with other
backbones are envisaged, comprising: A1, B, Ge, Ga, P, As,
N, In, Sb, S, Se, Te, In and Sb.
In the PS-CVD process, gaseous precursors from
polymeric sources are to be produced first directly from the
solid phase via sublimation, or via an intermediate liquid
phase subsequently subjected to vaporization, contrary to
the classical polymeric route. A definite advantage of this
process is a purification that involves the polymer source
during the sublimation process. The purification reduces
the effect of adventitious oxidation on the initial solid
polymeric source by the decreased capacity of oxidized
backbones to produce volatile material (e. g., at the limit,
a high degree of oxidation produces SiOz with negligible
volatility). The oxidized material is therefore
concentrated in the solid residue, while the precursors
reaching the substrate are purified this way. This
purification helps to produce films that have very few
chemical impurities and consequently fewer surface and bulk
defects.
The tolerance of the polymeric source to cracking,
thermal pyrolysis, and/or depolymerization processes is
related directly to silicon-carbon, silicon-nitrogen and
silicon-oxygen relative bond stability under given thermal
and pyrolysis conditions. The PS-CVD process was tested by
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subjecting the above-mentioned silicon-based polymers to
various thermal budgets controlling the depolymerization
conditions, for example, thermal cracking, chemical
decomposition and polymeric disproportionation).
PS-CVD further allows the polymeric source to
self-adjust to the temperature field because each polymeric
source will develop a set of gaseous precursors adapted to
the particular thermal conditions (various gradients and
temperature-cycles produce different gaseous precursors,
thermodynamically stable under the specific conditions).
This self adjustment is different for different polymeric
sources. The objective of the PS-CVD process is to create
at the outset, in the gas phase, the majority of the
required bonds that will constitute the solid silicon-based
film. Consequently, the role of the chemical reactions
occurring on the substrate is limited to the completion of
the remaining small number of bonds required for the
silicon-based structure. This technique facilitates a high
rate of mass transfer during desublimation of the large
precursor molecules, thereby increasing the growth rate of
silicon-based material on the substrate. This technique
permits much lower operating temperature during growth of
the silicon-based thin film than standard industrial
practices. For this reason the nature of the substrate is
less important, in terms of thermal stability, size or
shape. A large number of the SiC, SiO, or SiN bonds pre-
exist in the polymeric precursor. Furthermore, the gaseous
precursor is deposited in chemical chains, similar to
physical chains, that then rearrange and bind to one
another. The bond formed between the precursor and the
substrate is a van der Waals physical bond. Consequently,
the lower operating temperature provides an environment for
lowering the amount of unintentional impurities in the
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deposited film. Desublimation is herein defined as a change
of phase from a gaseous species directly to a solid species.
An example of a chemical reaction in the PS-CVD
process is represented in Figure 4. The first reaction of
the deposition involves significantly higher a-SiC
deposition rates compared to traditional CVD methods because
of chemically-synchronized Si-C bond redistribution in
organo-polysilanes. Still referring to Figure 4, in the
first chemical reaction (15), a thermally activated
methylene insertion into silicon-silicon bonds takes place
to produce poly(carbosilane) precursor. This intramolecular
reaction, known as the Kumada rearrangement, (Shiina, K.;
Kumada, M., 1958, J. Org. Chern., 23, p. 139), provides the
structural framework of silicon carbide, (Scarlete, M.;
Brienne, S.; Butter, S.S. and Harrod, J.F., 1994, Chem.
Mater., 6, p. 977). By simply heating the polymer
precursor, a very large number of Si-C bonds are
appropriately redistributed at a very fast rate.
The second reaction (16) also represented in
Figure 4 leads to the introduction of nitrogen atoms as
donor impurities into the silicon carbon precursors. The
formation of the aminocarbosilane precursor in reaction
(16), occurs via a reaction with ammonia, found either in
the atmosphere or in the polymer derived source.
While still referring to Figure 4, the third
reaction (17) results in high molecular weight species
through the formation of secondary amine species, leading to
increased desublimation capacity. The formation of the
secondary amine species is via the Si-H/N-H dehydrogenation.
The fourth reaction (18) governs the formation of
the film derived from the third reaction (17) onto the
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substrate (6). This example illustrates the formation of
high molecular mass ternary amine species by transamination,
the amine species being direct precursors are deposited via
desublimation to yield an n-type a-SiC film on the
substrate.
Figure 5 shows various types of substrates that
may be coated, with a-SiC film deposited by the PS-CVD
process, as well as, their nature and complexity. The a-SiC
thin film can be deposited on a regular ceramic material of
a complicated shape, quartz, electronic-grade sintered
alumina, polished alumina, silicon single crystal wafer,
graphite, and other commonly available and relatively
inexpensive materials. As seen in Figure 5, several of the
materials have been coated on one side (the dark surface) by
the PS-CVD process while the other side was masked during
deposition. A pale surface remained after mask removal.
Therefore, the PS-CVD process is also compatible with
conventional techniques such as masking understood by those
skilled in the art.
The PS-CVD process of the present invention
further distinguishes itself from other forms of CVD because
it does not require solvent to dissolve the precursor, the
evaporation of these solvents in other CVD methods produces
defects on the surface of the coating they produce, this is
one reason why the PS-CVD process produces films with fewer
and essentially no surface defects.
This invention incorporates the theoretical
concept of "anticeramic yield" of the gaseous precursors.
The traditional method for producing silicon-based materials
from polymeric sources is through rearrangement of the solid
residue left after pyrolysis of the precursor sources, with
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a typical ceramic yield amounting to 80% solid (F. Cao et
al., 2001, Korean Journal of Chemical Engineering, 18, p.
761). Recent efforts relating to traditional methods are
directed towards maximization of the amount of polymer
remaining as a solid, thereby increasing the yield to higher
values. The theory applying to the PS-CVD process generally
involves the opposite: to maximize the fraction of polymer
that is vaporized for the formation of the desired new
gaseous precursor leading to the deposit of an amorphous
silicon-based material. As such, almost the entire polymer
source is vaporized with the net result that the ceramic
yield is almost nil while the polymer transforms itself into
a new gaseous source resulting in almost 1000 anticeramic
yield. For example, this phenomenon is illustrated by the
TGA of a polymeric source subjected to the PS-CVD process,
shown in Figure 6.
Generally, the PS-CVD process according to the
present invention:
1) allows for the hydrogenation, heat treatment and makes
full use of many different types of gaseous reactants
such as, B2H6, NH3, PH3, AsH3, BC13, BZC16, NC13, PC13,
AsCl3, CO, O2, 03, C0, C02, as well as H2 or Dz, pure or
inert carried gases such as Ar or N2, or similar
mixtures, with the inert gases varying from 0.1 to 99.0 0
in volume;
2) may accommodate a wide variety of heat sources and
treatment lengths for the polymeric precursors or for the
deposited film under the gaseous atmospheres, for as
little as 1 second and up to, but not limited to, tens of
hours, and leading to a wide variety of passivating,
semiconductor, and dielectric thin film materials;
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3 ) allows secondary annealing under BH3, B2H6, NH3, PH3,
AsH3, BC13, B2C16, NC13, PC13, AsCl3, C0, O2, 03, C0, C02 or
H2 gases for several seconds, thereby increasing the
crystallinity and/or the degree of reticulation of the
deposited film;
4) may make use of standard heating methods as well
electronic beams, X-rays, UV and IR radiation microwave
power, laser beams and other energy transfer mechanisms
to produce objects and substrates in flat, or tubular or
complex shapes including, but not restricted to, rods,
cylinders, spheres, and ceramic boats; and
5) produces substrates with varying conductor,
semiconductor or dielectric properties, including, but
not restricted to polycrystalline or amorphous silicon;
quartz; graphite; metals; electronic-grade or refractive
ceramic materials, such as alumina or sintered oxides,
nitrides, phosphides; as well as AZB6, A3B5, ternary and
quaternary compounds in this class.
The thin films described herein may be used as
active, passivation, dielectric or protective coatings for
semiconductor discrete or integrated devices, or implantable
materials.
The thin film possesses highly desirable
electronic, optoelectronic and photonic properties that make
it highly suitable for standard, cost-effective fabrication
of a variety of electronic and optoelectronic devices,
including photovoltaic cells. The amorphous silicon-based
thin film may an n-type and/or p-type mono or heterojunction
with a donor concentration of 1013 to 101$ cm-3.
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By means of p-n homo- or heterojunctions and using
a variety of flexible or rigid substrates, it is possible to
fabricate solar cells, light-emitting diodes, transistors,
photothyristors, and similar devices. Using high breakdown
electrical field and high electron saturation velocity, it
is further possible to produce high frequency, high power
and high temperature electronic and optoelectronic devices.
By combining optical and electronic properties, the
materials may also serve to fabricate integrated monolithic
devices on a single chip.
Referring to Figure 7, there is represented a
simple Schottky solar cell. The cell comprises one layer
only of an amorphous silicon-based material (22), a metallic
substrate (20) acting as anode (when layer (22) is n-type)
or cathode (when layer (22) is p-type). The suitable metal
is an inexpensive conductive material and its thickness or
uniformity of thickness is not critical (viz. aluminum
foil). The ohmic contact layer (21) deposited by physical
evaporation or other physico-chemical means provides
effective contact with the overlying semiconductor layer as
well as the underlying metallic substrate. The substrate
may consist of aluminum or similar conductor (200 nm) if n-
type, or aluminum/nickel (100 nm/100 nm) if p-type; the
surface of which must be cleaned by chemical etching or
mechanical means to avoid oxidation with respect to layer
(22). Alternatively, layers (20) and (21) could be made or
fabricated as one composite layer over which layer (22)
could be deposited. The semiconductor layer of an amorphous
silicon-based material of n- or p-type (22) with free
carrier density between 1013 and lOla cm 3 and 0.2 to 1 um
thickness, produced by the PS-CVD process, acts as the heart
of the cell. The top layer may be gold (Au) layer (23) of 5
to 10 nm thickness acting as cathode if the semiconductor is
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n-type or as anode if it is p-type. The gold deposited
mechanically or chemically onto n-type semiconductor. The
gold layer is sufficiently thin as to allow light to reach
the semiconductor.
Figure 8 is a p-n junction barrier photovoltaic
cell based on an amorphous silicon-based thin film with
multiple layers produced by the process of the invention.
In this photovoltaic cell, a metallic substrate is acting as
an anode (24). Layer (25) is a metallized (aluminum) ohmic
contact layer 0100 nm). These layers are followed by an
n-type amorphous silicon-based thin film 0750 nm) layer
(26), p-type amorphous silicon-based thin film 0250 nm)
layer (27), a nickel ohmic layer (28) and an aluminum top
contact layer (29) serving as cathode which covers
approximately 10 percent of illuminated surface.
Figure 9 represents a stacked p-n junction solar
cell, the multiple layers produced by the method of the
invention. The layers of Figure 9 (with reference numbers
followed by the layer thicknesses, listed from bottom to
top) are: (30) metallic substrate acting as cathode; (3I)
aluminum-nickel ohmic contact layer 0100 nm/~100 nm); (32)
p-type a-Ge layer 0200 nm); (33) n-type a-Ge layer 0200
nm); (34) p-type a-Si 0200 nm); (35) n-type a-Si 0200 nm);
(36) p-type amorphous silicon-based thin film 0200 nm);
(37) n-type amorphous silicon-based thin film 0200 nm); and
(38) top aluminum anode contact, covering about 10 percent
of the surface.
In any semiconductor junction, such as in a
Schottky junction shown Figure lOB or a p-n junction such as
shown in Figure 118, there is an internal electrical field,
Eb;,, called "built-in electrical field", that prevents the
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charge carriers (electrons and holes) to stay in the a
region of the material called the "depleted zone." If the
depletion region of thickness W, is illuminated by photons
with energies greater than (E~ - Ev), the region develops
pairs of "electron-hole" which are separated by the internal
electrical field. The electrons are attracted towards the
semiconductor while holes are directed towards the metal,
creating a photocurrent when the device is connected to an
external load. The structure is called a photovoltaic cell
or solar cell.
The current generated in an amorphous
semiconductor is mainly due to a drift component because the
diffusion component is not significant due to low mobilities
of the charge carriers. In order to collect efficiently the
photon energies in an amorphous semiconductor junction, the
depleted region width must be as large as possible. The
depleted zone width, W, is given by:
W= ( ~ Vb~ / q ND ) 1/a
where ~ is the dielectric constant of the semiconductor, q
is the electron charge, ND is the electron concentration,
and Vbi is the built-in voltage given by:
_ (~a - (EC - EF) )
The width of the depletion region can be increased by
lowering the free carrier concentration of the material.
If a p-n junction (Figure 11) is used instead of a
Schottky one, the depletion region width can be increased
(in this case, each type of material has its own depleted
zone), therefore increasing the efficiency of the
photovoltaic cell.
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In the Schottky structure (Figure 8), the energy
band diagram is shown in Figure lOB before the intimate
contact between the metal and the semiconductor. The work
function, ~S (41m) is the energy difference between the
vacuum level (39) and the Fermi level, EF (40). The vacuum
level (39) is the zone where the electron is free from the
semiconductor atoms and has no kinetic energy. In elemental
solids such as a metal, represented in Figure 10A, the
values of the work function ~m (4Im) are well established,
(see for example lnTeast, R.C. (1990), CRC Handbook of
Chemistry and Physics, 70th Edition, CRC Press, E-93).
Figure 10A illustrates the work function of a
semiconductor (41s) normally denoted by ~S. The energy
difference between the vacuum level (39) and the bottom of
the conduction band (42), denoting electron affinity (~, is
used as reference since the Fermi level depends on the
carrier concentration in the semiconductor. However,
still represents the energy required to remove an electron
from the semiconductor. Referring to Figure 11 and 10B, the
conduction level (42) E~, the valence level (43) E~; the
affinity (44) x; the work function (45) ~S; and the energy
gap (46) E~ are of the semiconductor layer (22).
The Fermi level represents the energy for which
the probability to find a free electron in equilibrium and
near zero Kelvin equals 0.5. The probability of finding an
electron at a given energy level is obtained according to
the Fermi-Dirac function:
F(E) - 1 - [1 + exp (E-EF) / kT ]
The Fermi level in a semiconductor depends on the
free carrier concentration, and it is closer to the
conduction band than the valence band in n-type
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semiconductor. Assuming that (gym > ~s, and that the metal-
semiconductor system at equilibrium of Figure 11A, the
Fermi level is at the same both in the metal and the
semiconductor. Therefore, an internally built-in electric
field (Ebi) develops between the metal and the
semiconductor. The field is oriented from the positive
charges to the negative charges, that is, towards the
metal. The resulting built-in voltage is equal to [(gym -
~s) / q]. A depletion layer, of thickness W (53b), is
formed where there are no free charges. The potential
energy barrier for electrons moving from the metal to the
semi-conduction is known as the Schottky barrier height,
~B, and is given by: (~B - (gym -x). Under reverse bias or
zero bias electrical conditions, there is no net current
I5 flowing through the metal-semiconductor junction.
For photon energies greater than E~, the
electron-hole pairs are generated in the depleted zone.
In the p-n structure of Figure 11B, where the p
type amorphous silicon-based thin film (47) and the n-type
amorphous silicon-based thin film (48) are represented.
The principal electronic phenomena takes place in the
depleted zones (52) and (53). In this case, the built-in
voltage Vbi depends on the carrier concentration in the
semiconductor:
Vbi= ( kT/q) Ln (NAND / nit)
where k is the Boltzmann constant, T the temperature, NA
the hole concentration and ni the intrinsic carrier
concentration.
The depletion region widths are given by:
W"= ( a Vbi / q ND ) 1/2
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Wp = ( a Vbi / q NA ) 1/2
With the conduction level (49) E~, the Fermi Level (50) EF
and the valence level (51) also represented in Figure 11B.
Examples
The following examples are provided to illustrate
the invention. It will be understood, however, that the
specific details given in each example have been selected
for the purpose of illustration and are not to be construed
as limiting in scope of the invention.
Example 1. Synthesis of an n-type amorphous silicon carbide
(N-doped) thin film via the PS-CVD process.
Synthesis - Charges of polydimethylsilane and in-
house prepared polymethylsilane (M. Scarlete et al., 1995,
Chem. Mater., 6, p.1214) have been pyrolyzed in a single-
zone Lindberg ceramic furnace. The synthesis was performed
in a 2" quartz tube attached to a silicone-based hydraulic
lock bubbler and to a vacuum line capable of providing a
reduced pressure of 5*10-2 torr. The synthesis was
performed in a gaseous atmosphere of UHP-Ar or in home
purified NH3. The purification of NH3 was obtained via
passing the gas through a 1000 mm column of KOH and a 250 mm
column of a mixture of 3 and 4 A molecular sieves. The
furnace was operated via a PID Eurotherm temperature-
controller proving ~0.5°C in the range of 110 to 1100°C at
10 torr above the atmospheric pressure. The atmosphere of
the pyrolysis was carefully purified from oxygen and water
vapours via a series of evacuations/Ar-fillings. The
temperature cycle during the pyrolysis was the following:
a) a temperature slope of 4°C/min in the 110 to 450°C, in
order to allow silane-carbosilane transformation via
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Kumada rearrangement (Shiina, K. and Kumada, M.,
J.Org.Chem., 1958, 23, p.139);
b) a 20 min carbosilane annealing at 450 °C;
c) a temperature slope of ZO°C/min in the range of 450 to
1050°C to allow nitridation and cracking of the polymer
via reaction with the NH3-enriched inert gas; and
d) a 30 min annealing at 1050°C.
The process can be configured in a single-zone furnace,
using the temperature cycle above, or in a three-zone
furnace.
Various substrates have been placed downstream of
the gas flow in a cooler region of the furnace, in order to
collect the desublimation products of the gaseous species
produced during the pyrolysis. Among the substrates used
were 1-10 S2cm, type P(B) [100] silicon single crystal
wafers, electronic-grade alumina ceramic, quartz and steel
plates. The native oxide from the silicon substrate was
removed by etching in an acidic 10:3:1 DI-H20:CH3COOH:HF
solution for 15-30 sec, then dried in acetone in an
ultrasound bath for 0.5-2 minutes prior to introduction in
the pyrolysis furnace. The acetone sonication step was used
to prepare all substrates. All manipulations prior to
pyrolysis were conducted under continuous inert gas flow.
Characterization - The a-SiC film has been
characterized by FT-IR spectroscopy having a peak at
approximately 800 cm 1 with the FWHH around 150 cm-1 in the
green or as deposited material. Subsequent to annealing,
the FWHH decreases to 100 crri 1.
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An elastic recoil detection (ERD) study was
performed to reveal that the thin film had a C/Si ratio of
1-1.4, a residual hydrogen concentration of 0 to 15%, and a
residual oxygen concentration of 0 to 15% (Figure 12).
Film Properties - The carrier concentration in an
n-type semiconductor of an amorphous silicon carbide thin
film was measured by the capacitance-voltage (CV) method
(Schroder. D.K., 1990, Semiconductor Materials and Device
Characterization, Gdiley Interscience, p. 41) using a
Schumberger* impedance analyzer Solartron* 3200. The voltage
was varied between -6 and 0 V and the resulting capacitance
was observed to increase with increasing voltages. On a
sample film, six Schottky diodes were fabricated using
mercury as anode metal. The mercury probe used provided a
diode area of 0.453 mm2. The capacitance is given by:
C= ~A/w
where A is the diode area. In the presence of an applied
voltage V, the depletion region width is given by:
W = ( ( ~ (~lb~ -V) ) / q ND ) 1/2
The derived value for ND (electron concentration) in the
diode was 9x101' ~ 0.2x101' cm-3.
The electron mobility was measured using the Van
der Pauw method (Van der Pauw, L.J., 1958, Phil. Tech.
Rev., 20, p. 220). The measured mobility was defined as a
Hall mobility since the technique is based on the Hall
effect. The measurements were carried out with a current
of 1 mA and a magnetic field of 5 kG. The correction
factor f derived was 0.67, the derived resistivity was
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22.32 S2cm'1, the derived mobility was 4.48 cmzV-ls-1, and the
carrier concentration was 6.25x1016 cm'3.
Referring to Figure 13, the carrier properties of
an n-type amorphous silicon carbide thin film produced by
the PS-CVD process are represented. The graph represents
the donor concentration, n (cm-3) versus the width of the
depletion zone, W (um). We observed that the sample tested
has a low donor concentration, which can be below 1013 but
range to 1015cm'3. A preferred range is that less than 1014
cm 3. These low donor level values were before any doping.
The width of the depletion zone (W) which is measured in
(um) is a function of the material connected at the
junction, in the case of Figure 11, that of an n-type film
with the substrate. W must not be confused with the film
thickness. Thicker films (above 20 um) were required for the
method of detection used to quantify the carrier
concentration in the depletion zone and the thin films
obtained by this method (100 to 0.1 um) have the same type
of curve as found in Figure 6 at the far lower thicknesses.
Figure 14 represents the semiconductor properties
of an n-SiC PS-CVD film which is a qualitative indication of
the quality of the film, indicated by the curve of current
versus voltage.
An image from scanning electron microscopy (Figure
15) indicates that the thin film had a thickness of 73 nm,
and a deposition rate of 25 nm/min.
The excitation of the n-type a-SiC with photons at
3.25 eV induced photoluminescence of the material at room
temperature (Figure 16). The maximum is located in the blue
region, with a tail extending to the red region practically
covering the whole visible spectrum.
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An EPR spectrum of dangling bonds and other
paramagnetic centers of the thin film is presented in Figure
17. The noisy spectrum qualitatively indicated a somewhat
low dangling bond concentration.
Example 2. Synthesis of a p-type amorphous silicon carbide
(B-doped) thin film via the PS-CVD process.
Synthesis - Charges of poly(dimethylsilane) (PDMS)
were pyrolyzed in a Lindberg* Bluets single-zone furnace.
The furnace was operated via a PID Eurotherm temperature-
controller providing an accuracy of ~0.5°C in the ranger of
110 to 1100°C. The furnace was equipped with a 2.5" quartz
tube attached to a silicone-based hydraulic lock bubbler and
a vacuum line capable of providing 5*10-2 torr vacuum as
measured with a Welch Pirani-gauge. The atmosphere prior to
pyrolysis was carefully purged of moisture and oxygen via a
series of evacuations and UHP-Ar fillings. The residual
oxygen level obtained after the purging procedure was
measured using an Innovative Technology gauge and was found
to be below 1 ppm. Once purged, a partial pressure of BC13
was created in the reactor.
The substrates used were p-type (B doped) (100)
oriented single crystal silicon substrates, electronic grade
quartz substrates, and Ni coated (200nm) quartz substrates.
The latter permits bulk resistivity measurements since Ni
forms an ohmic contact on an amorphous silicon carbide
semiconductor layer. Native oxide on the silicon substrate
was removed by etching in an acidic 10:3:1 HZO/CH3COOH/HF
solution for 15-30 sec, then dried in acetone, in a
ultrasound bath for 15 minutes prior to introduction into
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the pyrolysis furnace. The quartz substrates were cleaned in
the acetone sonication bath for 15 min.
Characterization - Vibrational spectroscopy has
been used to analyze the layer deposited on Czochralski-
silicon single crystal wafer. The IR absorption of the
deposited film at 803 cm 1 in Figure 18 is characteristic of
silicon carbide. XRD showed no signal associated to
crystalline material, therefore the obtained film contained
amorphous silicon carbide. Adventitious oxygen
incorporation in the film was readily observed in the IR
spectrum of the film due to the strong absorptivity of the
vas(Si-0). The oxygen content of the film was calculated as
being in the range of 4x101'-1018 Cm 3 by the external
reference method (ASTM F-118 procedure).
Film properties - Hall measurements were performed
on quartz substrates that were previously covered with 200
nm of high purity Ni via vapor deposition. The type of
conductivity and the concentration of charge carriers was
determined by measurement of the CV profiles (Figures 19a
and 19b) of the layers. Quartz substrates were used for
metal-deposition via vaporization, then the metalized
substrate was used for the PS-CVD deposition of a-SiC(B). A
Schottky contact was formed on the a-SiC(B) layer by a
mercury probe having an area of 0.453 mm2. A Schlumberger*
51260 impedance analyzer was used to measured the
capacitance with an applied sinusoidal voltage having an
amplitude of 15 mV and a frequency of 1 MHz. The
concentration of charge carriers was determined to be 2*101'
Cm 3.
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The UV spectrum (Figure 20) was measured on a
Hewlet Packard HP8452 UV/Vis spectrophotometer and the
optical band gap was calculated from the obtained spectra
assuming that strong optical absorption marks the energy
level of band to band transitions. The ?~maX is located
around 270 nm and the optical bandgap of the sample was
calculated to be 2.82 eV using the intercept of the line
corresponding to high absorption rate with the x axis.
Other generic properties of the p-type amorphous
silicon thin film have not yet been measured, but they are
expected to be similar to those of the n-type amorphous
silicon thin film since the boron concentration is low.
Example 3. Synthesis of an amorphous silicon carbonitride
thin film via the PS-CVD process.
Synthesis - In-house prepared polymethylsilane and
poly(dimethyl)silane were pyrolyzed under NH3. The
nitrogen-containing amorphous layers were deposited on
electronic-grade substrates such as silicon single crystal
wafers, quartz, and alumina. The pyrolyses were conducted
at 5-20 torr above the atmospheric pressure and were
undertaken in a Lindberg single-zone, programmable furnace
equipped with a Eurotherm PID temperature controller with a
maximum temperature of 1100 ~ 0.5°C. The temperature cycle
during the PS-CVD process was the following: 4°/min 110-
450°C, 30 min at 450°C, 4-8 °C/min up to 1050°C,
and 30 min
at 1050°C (batch process).
Characterization - Prior to analysis of the film,
the products of the chemical reactions involving the carbon-
for-nitrogen exchange were followed in the ceramic residue
of the polymer. There was continuous C/N replacement
activated by a critical temperature necessary for the
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displacement level. CP-MAS NMR analysis (Figure 21)
indicated a large, continuous transformation of the polymer
into silicon carbonitride and eventually into silicon
nitride. Table 1 and Figure 22 quantify the results of the
NMR studies.
Table 1.
T SiN2C2 SiN2CH SiN4 High
[ desity
C]
(-5ppm) (-22ppm) (-43ppm) (-50
<
-OpPm)
300 -5ppm -20ppm -38ppm -
24.40 56.5 19.0 -
400 -lOppm -6 -23ppm -43ppm -
6.9~ 8.9 59.4 24.8 -
500 -llppm 29ppm -25 -22 -45 -43 -41 -50ppm
2.3~ 4.8.0$13.8 9.2 36.6 12.1 5.9 15.2
600 - -30ppm -45ppm -42 -60ppm -52
- 4.20 54.9$ 29.0 2.2$ 9.7
700 - - -47 -44 -40 -58ppm -51
- - 11.4 42.4 21.6 7.80 16.8
800 - - -47 -43 -40 -61 -56 -51
- - 41.0 14.5 6.0 4.8 10.7 23.
1100 - - -45 -42 -39 -63 -56 -50
- - 25.8 15.8 13.8 7.51 15.2 16.0
Once the possibility of the N-for-C displacement
was proven as a chemical possibility, the polymer was
subjected to PS-CVD under similar conditions. The FT-IR
spectrum of the thin film is shown in Figure 23, where the
broad absorption at around 850 cm 1 is assigned to the
formation of SiCXNY silicon carbonitride species.
Film properties - The silicon carbonitride thin
film deposited on silicon, quartz, and alumina substrates
were characterized electrically. The carrier concentrations
of the films were found to be in the range of 1014 - 1016 cm-3
in the bulk. The corresponding calculated resistivities of
the films deposited on quartz, alumina and electronic-grade
ceramics vary between 3-50 S2cm.
No residual Si-H and C-H were observed in the FT-
IR spectrum of the material, although the presence of 0 to
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15 atomic ~ of molecular hydrogen trapped in the film cannot
be excluded, based on similar characteristics found in other
silicon-based thin films produced by the PS-CVD method.
Example 4: Synthesis of an amorphous silicon nitride via the
PS-CVD process.
Synthesis - An appropriate time-dependent
temperature gradient was programmed in the furnace, so that
quantitative polysilazane formation was promoted. Possible
ranges for the gradients are 1-5 °K min-1 and 3-50 °K cm 1
where the temperature increases in a 2 inch / 150 cm
horizontal quartz reactor. A second step involves pyrolysis
of the polysilazane in the reaction zone where the
temperature is relatively constant. This step may be
optionally followed by a transamination processes induced
directly in the deposition zone, via a carefully monitored
(flow, pressure - parameters and PID temperature parameters
where P=1-25, I=10-250, and D=0.1-10) reaction under pure
electronic-grade gaseous ammonia introduced in the
temperature zone at a pressure level of 1 to 50 torr over
atmospheric pressure. The resultant precursor gaseous
species were transported in the deposition zone, where they
are desublimed onto the substrate that can be placed in a
horizontal, vertical or a tilted position, can be mobile or
immobile during the deposition. The resultant material is
a-SiXNY, with a x/y ratio in the range of 0.75 to 1. The
thickness of the resulted film may be adjusted in the 100
- 1 um via single/multiple layered deposition.
Characterization - A representative IR spectra is
presented is Figure 24. The major peak was observed at 896
cm 1. This peak was believed to be the product of overlap
of several other peaks. These included the 850cm-1 and
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940cm-1 peaks associated with the Si-N stretching mode, and
the Si-O symmetric stretching mode at 1040crri1. The
presence of a large shoulder on the main absorption peak
verifies the presence of more than one absorbance and is in
the range of the Si-O symmetric stretching mode. The peak
found at 478ciri1 was associated with the Si-N breathing
mode, 680cm 1 with the Si-H wag, and 1159ciri 1 with the N-H2
bend. The large peak at 2300cn11 was due to an excess of C02
in the machine during the sample scan compared to the
background scan. The presence of oxygen observed in the IR
may be due to oxygen contamination during the synthesis
process or surface oxide formation upon handling the sample.
The amount of residual bonded-hydrogen, evidenced as Si-H
and C-H bonds is function of the PS-CVD parameters, and
appears to decrease with the increase temperatures of the
zones 2 and 3 of a three-zone reactor. Residual carbon
(Auger) is in the order of 0-5 atomic %.
Example 5. Synthesis of an amorphous silicon oxynitride thin
film via PS-CVD.
Synthesis - Silicon nitride thin films were
produced by the PS-CVD process in various reaction
conditions. An organosilicon polymer was used as the
silicon source and NH3 as the nitride source. All nitride
depositions were performed in a one zone furnace equipped
with a 2" quartz tube attached to a silicone-based hydraulic
lock bubbler and to a vacuum line capable of providing a
reduced pressure of 5*10-z torr. When not in use the tube
was maintained at a temperature of 110°C. The gaseous
atmosphere consisted of a diluted mixture of in house
purified NH3 in a UHP-Ar carrier. The purification of NH3
was attained via passing the gas through a 1000 mm column of
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KOH and a 250 mm column of a mixture of 3 and 4 ~1 molecular
sieves.
The atmosphere of the pyrolysis was carefully
purified from oxygen and water vapor via a series of
evacuations by vacuum pump and N2 flushings. In the first
series of experiments the system was evacuated and then
filled with nitrogen. An ammonia flow was introduced into
the system throughout the temperature cycle at an
approximate rate of 1 "bubble" per second. In these initial
experiments the ammonia flow was gauged by counting
bubbles/time in the silicone-based hydraulic lock bubbler.
The temperature of the furnace increased at a constant rate
until it reached a maximum temperature of 550°C.
The films were deposited on both n- and p-type
polished silicon (100) wafers and quartz substrates. Two
sample sizes were employed, 2.5 X 2.Ocm and 10 X 2.5cm.
Prior to deposition, substrates were immersed in acetone and
placed in a sonic bath for four minutes. They were then
rinsed with acetone and dried in a stream of nitrogen.
The successful production of silicon nitride was
followed by oxidation of the films. The nitride samples
were oxidized in open air. The thermal profile of the PS-
CVD process was varied in separate trials. In the first
trial, the IR spectra were taken after thermal annealing of
the sample for 10 minutes at temperatures ranging between
300 and 700°C. In the second trial, the sample was oxidized
at a constant temperature of 500°C and the progression of
oxidation was monitored by FTIR analyses.
Characterization - The silicon oxynitride films
were analyzed using IR spectroscopy as seen in Figures 25a
and 25b. The samples will be referred to as SiONl, which
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was oxidized at constant temperature (Figure 25b) and Si0N2,
which was oxidized at various temperatures (Figure 25a). In
sample Si0N2, the main peak shifted from 958 to 1082 cm I
upon exposure to temperatures up to 700°C with a total
oxidation time of 50 min. This main peak shift according to
W.L. Scopel et al., 2003, Physical Review - Section B -
Condensed Matter, 68, p. 155332 indicates the oxidation of
the nitride. Shoulders present in the as deposited nitride
due to the Si-O and Si-N stretching vibrations absolved into
the main peak as oxidation proceeded, which indicated the
formation of one phase oxynitride. Moreover, the appearance
and increase of the Si-O bending mode in both samples in the
range of 478 to 517cm-1 indicates increased oxygen
incorporation. The single sharp absorption band at 1100 cm-
1 as in a Si02 network was not observed, rather a broad band
was seen with the shoulder at lower wavenumbers, which was
explained by the presence of Si-N bonding.
In the case of SiONl, the shift in wavenumbers of
the main peak from 896 to 1050cn11 indicated oxidation of
the nitride. The oxynitride produced by annealing at 400°C
stopped incorporating oxygen after one hour but oxidation
resumed once the temperature was increased to 500°C. This
indicated that as temperature is increased the rate of
reaction also increases.
Film properties - The optical band gap of the film
calculated from the UV spectrum was found to be 4.92eV
compared to the literature value of 5.0 eV (Klaus Mogenson,
Peter Friis, Jorg Hubner, Nickolaj Peterson, Anders
Jargenson, Pieter Telleman, Jorg P. Kutter, Optics Letters,
26 (10) , 2001) .
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A Scanning Electron Microscope (SEM) thickness
profile of a representative sample nitride showed that the
thickest region in the film is 600nm and that the thickness
falls off gradually with few inconsistencies.
Example 6. Synthesis of a silicon oxy(carbo)nitride thin
film via the PS-CVD process.
Synthesis - For the initial nitridation procedure,
commercial poly(dimethylsilane) from Gelest was used. The
experiments were performed in a Minibrute parallel diffusion
three-zone furnace at 5-10 torr above atmospheric pressure
capable of operating up to 1500 °C. It was equipped with a
2" quartz tube attached to a silicone-based hydraulic lock
bubbler, and to a vacuum line capable of providing a reduced
pressure of 5*10-2 torr. The gaseous atmosphere consisted of
a diluted mixture of in house purified NH3 in a UHP-Ar
carrier. The purification of NH3 was attained via passing
the gas through a 1000mm column of KOH and a 250mm column of
a mixture of 3 and 4 A molecular sieves. The atmosphere of
the pyrolysis was carefully purified from oxygen and water
vapour via a series of consecutive evacuations/Ar-fillings.
The residual oxygen level obtained after the flushing
procedure was measured with an Innovative* Technology gauge
and found to be below 5ppm. Silicon and quartz substrates
were used for the PS-CVD deposition. A continuous flow of
anhydrous ammonia was introduced into the system before
heating and during the full temperature cycle. The
pyrolysis cycle performed under pure ammonia flow led to a
deposit consistent with a silicon carbonitride species.
The silicon carbonitride thin film was further
thermally oxidized in a second Minibrute furnace in
Trademark
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conjunction with a three-inch Quartz tube with ends open to
the atmosphere. The IR spectrum of the silicon carbonitride
thin film was taken after heating for 10 minutes at a
variety of temperatures between 200 and 720°C. The sample
was further annealed for a period of one to ten hours in the
furnace at 700°C and their respective IR spectra were taken.
Example 7. Synthesis of an amorphous silicon oxycarbide via
the PS-CVD process.
Synthesis - Silicon oxycarbide is an amorphous
metastable phase wherein the silicon atoms are bonded to
oxygen and carbon simultaneously. In silicon oxycarbides,
high temperature properties and chemical stability have been
reported, exceeding those of ordinary vitreous silica.
Silicon oxycarbide materials have also the potential for use
in a variety of protective applications within the
semiconductor industry. Using PS-CVD technique, silicon
oxycarbides of various compositions have been deposited on
highly resistive single crystal silicon wafers, using
different conditions to vary the oxygen content in the
films.
An appropriate time-dependent temperature gradient
was programmed in the furnace to enhance quantitative
polycarbosilane formation. Possible ranges for the gradients
were 1-10 Kmin-1 and 3-50 Kcrri 1 in a 2 inch / 150 cm
horizontal quartz reactor. The resultant polymeric gaseous
species were transported in the deposition zone.
The polycarbosilane was oxidized via a carefully
monitored (flow, pressure, and FT-IR) reaction with oxygen
carrying species including, but not limited to, 02, 03, CO
at a partial pressure level of 10-q - 10-1 torr in a UHP-Ar
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(or N2) carrier flow. The oxygen carrying species were
introduced directly in the deposition zone.
The controlled oxidation products were desublimed
onto the substrate that may be placed in a horizontal,
vertical, or a tilted position, and may be mobile or
immobile during the deposition. The resultant material was a
SiOXCy glass with an oxygen content in the range from x = 10-
3 to x = 1.3 measured using an external standard of Cz-
silicon single crystal via ASTM F-1188.
Characterization - Figure 26 represents the FT-IR
spectra of three samples, (a), (b) and (c), of a synthesized
silicon oxycarbide film annealed for increasing time periods
of 8, 16 and 24 hours, respectively. It must be noted, that
the interstitial oxygen peak found in sample (a) at
approximately 1100 cm-l, increased as the film is annealed
for longer periods. This indicated the conversion limited
resistance of the film to oxidation.
Film properties - The thickness of the resultant
film can be adjusted in the 1001 - 1 um range via
single/multiple layered deposition.
Although various embodiments of the invention are
disclosed herein, many adaptations and modifications may be
made within the scope of the invention in accordance with
the common general knowledge of those skilled in this art.
Such modifications include the substitution of known
equivalents for any aspect of the invention in order to
achieve the same result in substantially the same way.
Numeric ranges are inclusive of the numbers defining the
range. In the claims, the word "comprising" is used as an
open-ended term, substantially equivalent to the phrase
"including, but not limited to".
