Ceramic thin film on various substrates, and process for producing same

A thin film of an amorphous silicon-based material on a substrate. The thin film has the property of any one of a carrier concentration of 1013 to 1018 cm−3 in a depletion zone next to the substrate, an electron mobility of 5 to 30 cm2V−1s−1, a dangling bond concentration of 1012 to 1019 cm−3, no solvent-related defects, or a residual hydrogen concentration of 0 to 25 atomic %. The thin film may be used to fabricate many devices such as solar cells, light-emitting diodes, transistors, photothyristors, and integrated monolithic devices on a single chip.

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Description
CROSS REFERENCE TO RELATED APPLICATIONS

This patent application is a continuation-in-part application of co-pending U.S. application Ser. No. ______ [not yet assigned], filed Nov. 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 U.S. Pat. No. 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 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 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. (Jana, 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 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 B2H6 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 E04 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, sp3 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 atoms in the films and the sp3 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 sp3 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 sp3 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.

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 1013 to 1018 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−1s−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 1019 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-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 1018 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−1s−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 %.

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.

BRIEF DESCRIPTION OF THE DRAWINGS

In the accompanying drawings, which illustrate an exemplary embodiment of the present invention:

FIG. 1a is a schematic diagram of a Polymer Source Chemical Vapor Deposition (PS-CVD) reactor;

FIG. 1b is a diagram of the temperature-profile in the reactor;

FIG. 2 is a schematic diagram of the FT-IR cell for monitoring the production of thin film by the PS-CVD process;

FIG. 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;

FIG. 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;

FIG. 5 is a series of substrates that may support a thin film produced by the PS-CVD process;

FIG. 6 is a thermogram from a thermogravimetric analysis of a polymeric source subjected to the PS-CVD process;

FIG. 7 illustrates a simple Schottky solar cell;

FIG. 8 illustrates a p-n junction barrier photovoltaic cell based on amorphous silicon carbide thin film;

FIG. 9 illustrates a stacked p-n junction solar cell based on amorphous silicon carbide thin film;

FIG. 10a illustrates an energy band diagram of Schottky structure before intimate contact between metal and semiconductor;

FIG. 10b illustrates an energy band diagram of Schottky structure after intimate contact between metal and semiconductor;

FIG. 11a illustrates the energy band of a p-n junction solar cell;

FIG. 11b illustrates the energy band of a p-n junction solar cell;

FIG. 12 is an ERD graph illustrating the relationship between concentration and depth for an n-type amorphous silicon carbide thin film;

FIG. 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;

FIG. 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;

FIG. 15 is an image from a scanning electron microscope of an n-type amorphous silicon carbide thin film;

FIG. 16 is a spectrum of the photoluminescence of an n-type amorphous silicon carbide thin;

FIG. 17 is an EPR spectrum of an n-type amorphous silicon carbide thin;

FIG. 18 is an infrared spectrum of a p-type amorphous silicon carbide thin film on a silicon single crystal wafer;

FIG. 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;

FIG. 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;

FIG. 20 is a UV spectrum of a p-type amorphous silicon carbide thin film on a silicon single crystal wafer;

FIG. 21 are 29Si 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;

FIG. 22 is a contour plot illustrating the relationship between the 29Si NMR chemical shifts and reaction temperatures of an amorphous silicon carbonitride thin film;

FIG. 23 is an infrared spectrum of an amorphous silicon carbonitride thin film;

FIG. 24 is an infrared spectrum of an amorphous silicon nitride thin film;

FIG. 25a are infrared spectra of the oxidation of SiON1 in air for 10 minutes at 300, 550, 600, 700° C.;

FIG. 25b are infrared spectra of the oxidation of SiON2 in air at 500° C. for 10, 20, 30, 42 and 52 minutes; and

FIG. 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 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.

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 1017 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−1s−1, in a preferred embodiment the electron mobility is 10 to 25 cm2V−1s−1, and in a more preferred embodiment the electron mobility is 15 to 20 cm2V−1s−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).

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 %, in a preferred embodiment the residual hydrogen concentration is 0 to 20%, 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, 50% of the carbon and 100% 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 50-100% of the carbon and 60-100% 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.

The thin films described herein were prepared by the PS-CVD process. FIG. 1A illustrates an apparatus used for the PS-CVD process while FIG. 1B depicts the temperature profile used within the apparatus. Referring to FIG. 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 FIG. 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 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 FIG. 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 FIG. 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 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:

    • 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 01-17 6910830 A2001-11-6855-060 (PHA) NDN-174-0691-0829-5 Yihua Wang; 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/O 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 FIG. 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 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 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: Al, 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 SiO2 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 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 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 FIG. 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 FIG. 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. Chem., 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 FIG. 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 FIG. 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 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.

FIG. 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 FIG. 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 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 100% anticeramic yield. For example, this phenomenon is illustrated by the TGA of a polymeric source subjected to the PS-CVD process, shown in FIG. 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, BCl3, B2Cl6, NCl3, PCl3, AsCl3, CO, O2, O3, CO, CO2, as well as H2 or D2, pure or inert carried gases such as Ar or N2, or similar mixtures, with the inert gases varying from 0.1 to 99.0% 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;
    • 3) allows secondary annealing under BH3, B2H6, NH3, PH3, AsH3, BCl3, B2Cl6, NCl3, PCl3, AsCl3, CO, O2, O3, CO, CO2 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 A2B6, 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 1018 cm−3.

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 FIG. 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 1018 cm−3 and 0.2 to 1 μm 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 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.

FIG. 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 (˜100 nm). These layers are followed by an n-type amorphous silicon-based thin film (˜750 nm) layer (26), p-type amorphous silicon-based thin film (˜250 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.

FIG. 9 represents a stacked p-n junction solar cell, the multiple layers produced by the method of the invention. The layers of FIG. 9 (with reference numbers followed by the layer thicknesses, listed from bottom to top) are: (30) metallic substrate acting as cathode; (31) aluminum-nickel ohmic contact layer (˜100 nm/˜100 nm); (32) p-type a-Ge layer (˜200 nm); (33) n-type a-Ge layer (˜200 nm); (34) p-type a-Si (˜200 nm); (35) n-type a-Si (˜200 nm); (36) p-type amorphous silicon-based thin film (˜200 nm); (37) n-type amorphous silicon-based thin film (˜200 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 FIG. 10B or a p-n junction such as shown in FIG. 11B, there is an internal electrical field, Ebi, called “built-in electrical field”, that prevents the 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 (Ec-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=(εVbi/qND)1/2
where ε 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:
Vbi=(ΦB−(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 (FIG. 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.

In the Schottky structure (FIG. 8), the energy band diagram is shown in FIG. 10B before the intimate contact between the metal and the semiconductor. The work function, Φs (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 FIG. 10A, the values of the work function Φm (41m) are well established, (see for example Weast, R. C. (1990), CRC Handbook of Chemistry and Physics, 70th Edition, CRC Press, E-93).

FIG. 10A illustrates the work function of a semiconductor (41s) normally denoted by ΦS. The energy difference between the vacuum level (39) and the bottom of the conduction band (42), denoting electron affinity (χ), is used as reference since the Fermi level depends on the carrier concentration in the semiconductor. However, Φs still represents the energy required to remove an electron from the semiconductor. Referring to FIG. 11 and 10B, the conduction level (42) Ec, the valence level (43) Ev; the affinity (44) χ; the work function (45) Φs; and the energy gap (46) EG 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 semiconductor. Assuming that (Φm>Φs, and that the metal-semiconductor system at equilibrium of FIG. 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 [(Φm−Φs)/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, ΦB, and is given by: (ΦB=(Φm−χ). Under reverse bias or zero bias electrical conditions, there is no net current flowing through the metal-semiconductor junction.

For photon energies greater than EG, the electron-hole pairs are generated in the depleted zone.

In the p-n structure of FIG. 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/ni2)
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:
Wn=(εVbi/qND)1/2
Wp=(εVbi/qNA)1/2
With the conduction level (49) Ec, the Fermi Level (50) EF and the valence level (51) also represented in FIG. 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 Å 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 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 10° 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 Ωcm, 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-H2O: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 cm−1.

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% (FIG. 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, Wiley 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=εA/w
where A is the diode area. In the presence of an applied voltage V, the depletion region width is given by:
W=((ε(Vbi−V))/qND)1/2
The derived value for ND (electron concentration) in the diode was 9×1017±0.2×1017 cm3.
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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 22.32 Ωcm−1, the derived mobility was 4.48 cm2V−1s−1, and the carrier concentration was 6.25×1016 cm−3.

Referring to FIG. 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 (μm). We observed that the sample tested has a low donor concentration, which can be below 1013 but range to 1015 cm−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 (μm) is a function of the material connected at the junction, in the case of FIG. 11, that of an n-type film with the substrate. W must not be confused with the film thickness. Thicker films (above 20 μm) 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 Å to 0.1 μm) have the same type of curve as found in FIG. 6 at the far lower thicknesses.

FIG. 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 (FIG. 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 (FIG. 16). The maximum is located in the blue region, with a tail extending to the red region practically covering the whole visible spectrum.

An EPR spectrum of dangling bonds and other paramagnetic centers of the thin film is presented in FIG. 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* BlueM 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 BCl3 was created in the reactor.
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The substrates used were p-type (B doped) (100) oriented single crystal silicon substrates, electronic grade quartz substrates, and Ni coated (200 nm) 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 10:3:1H2O/CH3COOH/HF solution for 15-30 sec, then dried in acetone, in a ultrasound bath for 15 minutes prior to introduction into 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 FIG. 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 νas(Si—O). The oxygen content of the film was calculated as being in the range of 4×1017-1018 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 (FIGS. 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* S1260 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*1017 cm−3.
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The UV spectrum (FIG. 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 λmax 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 displacement level. CP-MAS NMR analysis (FIG. 21) indicated a large, continuous transformation of the polymer into silicon carbonitride and eventually into silicon nitride. Table 1 and FIG. 22 quantify the results of the NMR studies.

TABLE 1 SiN2C2 SiN2CH SiN4 High desity T[° C.] (−5 ppm) (−22 ppm) (−43 ppm) (−50 < −0 ppm) 300  −5 ppm −20 ppm −38 ppm — 24.4% 56.5 19.0 — 400 −10 ppm −6 −23 ppm −43 ppm — 6.9% 8.9 59.4 24.8 — 500 −11 ppm 29 ppm −25 −22 −45 −43 −41 −50 ppm  2.3% 4.8.0% 13.8 9.2 36.6 12.1 5.9 15.2 600 — −30 ppm −45 ppm −42 −60 ppm −52 — 4.2% 54.9% 29.0 2.2% 9.7 700 — — −47 −44 −40 −58 ppm −51 — — 11.4 42.4 21.6 7.8% 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 FIG. 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 Ωcm.

No residual Si—H and C—H were observed in the FT-IR spectrum of the material, although the presence of 0 to 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 Å-1 μm via single/multiple layered deposition.

Characterization—A representative IR spectra is presented is FIG. 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 850 cm−1 and 940 cm−1 peaks associated with the Si—N stretching mode, and the Si—O symmetric stretching mode at 1040 cm−1. 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 478 cm−1 was associated with the Si—N breathing mode, 680 cm−1 with the Si—H wag, and 1159 cm−1 with the N—H2 bend. The large peak at 2300 cm−1 was due to an excess of CO2 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*10−2 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 KOH and a 250 mm column of a mixture of 3 and 4 Å 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, 2.5×2.0 cm and 10×2.5 cm. 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 FIGS. 25a and 25b. The samples will be referred to as SiON1, which was oxidized at constant temperature (FIG. 25b) and SiON2, which was oxidized at various temperatures (FIG. 25a). In sample SiON2, the main peak shifted from 958 to 1082 cm−1 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 517 cm−1 indicates increased oxygen incorporation. The single sharp absorption band at 1100 cm−1 as in a SiO2 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 SiON1, the shift in wavenumbers of the main peak from 896 to 1050 cm−1 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 4.92 eV compared to the literature value of 5.0 eV (Klaus Mogenson, Peter Friis, Jörg Hübner, Nickolaj Peterson, Anders Jargenson, Pieter Telleman, Jorg P. Kutter, Optics Letters, 26(10), 2001).

A Scanning Electron Microscope (SEM) thickness profile of a representative sample nitride showed that the thickest region in the film is 600 nm 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 1000 mm column of KOH and a 250 mm column of a mixture of 3 and 4 Å 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 5 ppm. 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.
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The silicon carbonitride thin film was further thermally oxidized in a second Minibrute furnace in 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 Kcm−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, O2, O3, CO at a partial pressure level of 10−4-10−1 torr in a UHP-Ar (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—FIG. 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−1, 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 100 Å-1 μm 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”.

Claims

1. A film of an amorphous silicon-based material on a substrate, the film having a carrier concentration of 1013 to 1018 cm−3 in a depletion zone next to the substrate.

2. The film according to claim 1, wherein the carrier concentration is 1013 to 1017 cm−3 in a depletion zone next to the substrate.

3. The film according to claim 1, wherein the carrier concentration is 1013 to 1016 cm−3 in a depletion zone next to the substrate.

4. The film according to claim 1, wherein the amorphous silicon-based material is an amorphous silicon carbide.

5. The film according to claim 1, wherein the amorphous silicon-based material is an amorphous silicon carbonitride.

6. The film according to claim 1, wherein the amorphous silicon-based material is an amorphous silicon nitride.

7. The film according to claim 1, wherein the amorphous silicon-based material is an amorphous silicon oxynitride.

8. The film according to claim 1, wherein the amorphous silicon-based material is an amorphous silicon oxy(carbo)nitride.

9. The film according to claim 1, wherein the amorphous silicon-based material is an amorphous silicon oxycarbide.

10. The film according to claim 1 further comprising a dopant.

11. The film according to claim 10, wherein the dopant is selected from the group consisting N, B, O, H, Cl, Al, Ga, In, Tl, P, As, Sb, O, S, Se, Te, and Bi.

12. The film according to claim 1, wherein the substrate is selected from the group consisting of quartz, metal, ceramic, electronic-grade sintered alumina, polished alumina, silicon single crystal wafer, graphite, and stainless steel.

13. The film according to claim 1, wherein the film has a thickness of at least 100 Å.

14. The film according to claim 1, wherein the film has a grain size of 2-10 nm.

15. The film according to claim 1, wherein the film has a grain size of 2-5 nm.

16. The film according to claim 1, wherein the film has a grain size of 2-3 nm.

17. A film of an amorphous silicon-based material on a substrate, the film having an electron mobility of 5 to 30 cm2V−1s−1.

18. The film according to claim 17, wherein the electron mobility is 10 to 25 cm2V−1s−1.

19. The film according to claim 17, wherein the electron mobility is 15 to 20 cm2V−1s−1.

20. The film according to claim 17, wherein the amorphous silicon-based material is an amorphous silicon carbide.

21. The film according to claim 17, wherein the amorphous silicon-based material is an amorphous silicon carbonitride.

22. The film according to claim 17, wherein the amorphous silicon-based material is an amorphous silicon nitride.

23. The film according to claim 17, wherein the amorphous silicon-based material is an amorphous silicon oxynitride.

24. The film according to claim 17, wherein the amorphous silicon-based material is an amorphous silicon oxy(carbo)nitride.

25. The film according to claim 17, wherein the amorphous silicon-based material is an amorphous silicon oxycarbide.

26. The film according to claim 17 further comprising a dopant.

27. The film according to claim 26, wherein the dopant is selected from the group consisting N, B, O, H, Cl, Al, Ga, In, Tl, P, As, Sb, O, S, Se, Te, and Bi.

28. The film according to claim 17, wherein the substrate is selected from the group consisting of quartz, metal, ceramic, electronic-grade sintered alumina, polished alumina, silicon single crystal wafer, graphite, and stainless steel.

29. The film according to claim 17, wherein the film has a thickness of at least 100 Å.

30. The film according to claim 17, wherein the film has a grain size of 2-10 nm.

31. The film according to claim 17, wherein the film has a grain size of 2-5 nm.

32. The film according to claim 17, wherein the film has a grain size of 2-3 nm.

33. A film of an amorphous silicon-based material on a substrate, the film having a dangling bond concentration of 1012 to 1019 cm−3.

34. The film according to claim 33, wherein the dangling bond concentration is 1013 to 1018 cm−3.

35. The film according to claim 33, wherein the dangling bond concentration is 1014 to 1017 cm−3.

36. The film according to claim 33, wherein the amorphous silicon-based material is an amorphous silicon carbide.

37. The film according to claim 33, wherein the amorphous silicon-based material is an amorphous silicon carbonitride.

38. The film according to claim 33, wherein the amorphous silicon-based material is an amorphous silicon nitride.

39. The film according to claim 33, wherein the amorphous silicon-based material is an amorphous silicon oxynitride.

40. The film according to claim 33, wherein the amorphous silicon-based material is an amorphous silicon oxy(carbo)nitride.

41. The film according to claim 33, wherein the amorphous silicon-based material is an amorphous silicon oxycarbide.

42. The film according to claim 33 further comprising a dopant.

43. The film according to claim 42, wherein the dopant is selected from the group consisting N, B, O, H, Cl, Al, Ga, In, Tl, P, As, Sb, O, S, Se, Te, and Bi.

44. The film according to claim 33, wherein the substrate is selected from the group consisting of quartz, metal, ceramic, electronic-grade sintered alumina, polished alumina, silicon single crystal wafer, graphite, and stainless steel.

45. The film according to claim 33, wherein the film has a thickness of at least 100 Å.

46. The film according to claim 33, wherein the film has a grain size of 2-10 nm.

47. The film according to claim 33, wherein the film has a grain size of 2-5 nm.

48. The film according to claim 33, wherein the film has a grain size of 2-3 nm.

49. A film of an amorphous silicon-based material on a substrate, the film having no solvent-related defects.

50. The film according to claim 49, wherein the solvent-related defects are voids in the film due solvent evaporation or solvent trapped in the film during deposition.

51. The film according to claim 49, wherein the amorphous silicon-based material is an amorphous silicon carbide.

52. The film according to claim 49, wherein the amorphous silicon-based material is an amorphous silicon carbonitride.

53. The film according to claim 49, wherein the amorphous silicon-based material is an amorphous silicon nitride.

54. The film according to claim 49, wherein the amorphous silicon-based material is an amorphous silicon oxynitride.

55. The film according to claim 49, wherein the amorphous silicon-based material is an amorphous silicon oxy(carbo)nitride.

56. The film according to claim 49, wherein the amorphous silicon-based material is an amorphous silicon oxycarbide.

57. The film according to claim 49 further comprising a dopant.

58. The film according to claim 57, wherein the dopant is selected from the group consisting of N, B, O, H, Cl, Al, Ga, In, Tl, P, As, Sb, O, S, Se, Te, and Bi.

59. The film according to claim 49, wherein the substrate is selected from the group consisting of quartz, metal, ceramic, electronic-grade sintered alumina, polished alumina, silicon single crystal wafer, graphite, and stainless steel.

60. The film according to claim 49, wherein the film has a thickness of at least 100 Å.

61. The film according to claim 49, wherein the film has a grain size of 2-10 nm.

62. The film according to claim 49, wherein the film has a grain size of 2-5 nm.

63. The film according to claim 49, wherein the film has a grain size of 2-3 nm.

64. A film of an amorphous silicon-based material on a substrate, the film having a residual hydrogen concentration of 0 to 25 atomic %.

65. A film of an amorphous silicon-based material on a substrate, the thin film having a residual hydrogen concentration of 0 to 20 atomic %.

66. A film of an amorphous silicon-based material on a substrate, the thin film having a residual hydrogen concentration of 0 to 15 atomic %.

67. The film according to claim 64, wherein the amorphous silicon-based material is an amorphous silicon carbide.

68. The film according to claim 64, wherein the amorphous silicon-based material is an amorphous silicon carbonitride.

69. The film according to claim 64, wherein the amorphous silicon-based material is an amorphous silicon nitride.

70. The film according to claim 64, wherein the amorphous silicon-based material is an amorphous silicon oxynitride.

71. The film according to claim 64, wherein the amorphous silicon-based material is an amorphous silicon oxy(carbo)nitride.

72. The film according to claim 64, wherein the amorphous silicon-based material is an amorphous silicon oxycarbide.

73. The film according to claim 64 further comprising a dopant.

74. The film according to claim 73, wherein the dopant is selected from the group consisting of N, B, O, H, Cl, Al, Ga, In, Tl, P, As, Sb, O, S, Se, Te, and Bi.

75. The film according to claim 64, wherein the substrate is selected from the group consisting of quartz, metal, ceramic, electronic-grade sintered alumina, polished alumina, silicon single crystal wafer, graphite, and stainless steel.

76. The film according to claim 64, wherein the film has a thickness of at least 100 Å.

77. The film according to claim 64, wherein the film has a grain size of 2-10 nm.

78. The film according to claim 64, wherein the film has a grain size of 2-5 nm.

79. The film according to claim 64, wherein the film has a grain size of 2-3 nm.

80. A semiconductor device comprising a film of an amorphous silicon-based material on a substrate, the film having a carrier concentration of 1013 to 1018 cm−3 in a depletion zone next to the substrate.

81. The semiconductor device according to claim 80, wherein the device is a p-type or an n-type.

82. The semiconductor device according to claim 80, wherein the device is selected from the group consisting of a solar cell, a light-emitting diode, a Schottky diode, a transistor, a photothyristor and an integrated monolithic device on a single chip.

83. 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−1s−1.

84. The semiconductor device according to claim 83, wherein the device is a p-type or an n-type.

85. The semiconductor device according to claim 83, wherein the device is selected from the group consisting of a solar cell, a light-emitting diode, a Schottky diode, a transistor, a photothyristor and an integrated monolithic device on a single chip.

86. 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.

87. The semiconductor device according to claim 86, wherein the device is a p-type or an n-type.

88. The semiconductor device according to claim 86, wherein the device is selected from the group consisting of a solar cell, a light-emitting diode, a Schottky diode, a transistor, a photothyristor and an integrated monolithic device on a single chip.

89. A semiconductor device comprising a film of an amorphous silicon-based material on a substrate, the film having no solvent-related defects.

90. The semiconductor device according to claim 89, wherein the device is a p-type or an n-type.

91. The semiconductor device according to claim 89, wherein the device is selected from the group consisting of a solar cell, a light-emitting diode, a Schottky diode, a transistor, a photothyristor and an integrated monolithic device on a single chip.

92. 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 %.

93. The semiconductor device according to claim 92, wherein the device is a p-type or an n-type.

94. The semiconductor device according to claim 92, wherein the device is selected from the group consisting of a solar cell, a light-emitting diode, a Schottky diode, a transistor, a photothyristor and an integrated monolithic device on a single chip.

Patent History
Publication number: 20050139966
Type: Application
Filed: Dec 3, 2004
Publication Date: Jun 30, 2005
Inventors: Mihai Scarlete (Roxboro), Cetin Aktik (Sherbrooke)
Application Number: 11/003,249
Classifications
Current U.S. Class: 257/632.000; 438/62.000