Methods and Compositions for Preparing Ge/Si Semiconductor Substrates
The present disclosure describes methods for preparing semiconductor structures, comprising forming a Ge layer on a semiconductor substrate using an admixture of (a) (GeH3)2CH2 and Ge2H6; (b) GeH3CH3 and Ge2H6; or (c) (GeH3)2CH2, GeH3CH3 and Ge2H6, wherein in all cases, Ge2H6 is in excess. The disclosure further provides semiconductor structures formed according to the methods of the invention as well as compositions comprising an admixture of (GeH3)2CH2 and/or GeH3CH3 and Ge2H6 in a ratio of between about 1:5 and 1:30. The methods herein provide, and the semiconductor structures provide, Ge layers formed on semiconductor substrates having threading dislocation density below 105/cm2 which can be useful in semiconductor devices.
Latest The Arizona Board of Regents, a body corporate of the State of Arizona acting for and on behalf of A Patents:
This application is a divisional application of U.S. non-provisional application Ser. No. 12/133,225, filed Jun. 4, 2008 which claims the benefit under 35 USC §119(e), of U.S. Provisional Application Ser. No. 60/933,023, filed 4 Jun. 2007, which is hereby incorporated by reference in its entirety.
STATEMENT OF GOVERNMENT INTERESTThe invention described herein was made in part with government support under grant number FA9550-06-01-0442, awarded by AFOSR under the MURI; and under grant number DMR-0526734, awarded by the National Science Foundation. The United States Government has certain rights in the invention.
BACKGROUND OF THE INVENTIONElemental Ge exhibits many device advantages over pure Si. Its smaller bandgap is attractive for photodetector and modulator applications in the 1.3-1.6 μm wavelength range (see, Oh et al., IEEE J. Quantum Electron. 38, 1238 (2002); Liu et al., Appl. Phys. Lett. 87, 103501 (2005); and Kuo et al., Nature 437, 1334 (2005)). Transistors based on Ge should also provide greater speed performance due to higher carrier mobilities of this material. Since the manufacturing infrastructure for Si—Ge technologies is well established, the direct growth of Ge on Si could produce new classes of opto- and microelectronic systems, but such growth has been problematic. The conventional formation of mismatched (4%) Ge on Si typically proceeds via the Stranski-Krastanov mechanism yielding islands (after deposition of 3-4 monolayers) rather than relaxed, continuous layers. For thick films a high roughness is obtained and threading dislocation densities of ˜108 cm−2 are commonly observed (see, Kroemer et al., J. Cryst. Growth 95, 96 (1989)). Carrier scattering and traps at defect sites reduce mobility in electronic devices and also increase dark current in photodetectors.
Low temperature (T<375° C.) chemical vapor deposition (CVD) of GeH4 has produced Ge layers directly on Si(100) possessing fairly smooth surfaces with occasional pits, and threading dislocation densities that appear to be too high for certain applications (see, Cunningham et al., Appl. Phys. Lett. 59, 3574 (1991)). Higher temperature growth (T>400° C.) invariably produces rougher layers that display the classic cross-hatched patterns created by strain relaxation and defect formation. The higher temperatures also increase the propensity of microcrack formation upon cooling of the samples (see, for example, Currie et al., Appl. Phys. Lett. 72, 1718 (1998); and Fitzgerald et al., J. Vac. Sci. Technol. B 10, 1807 (1992)) making such approaches incompatible with back end (post-metallization) CMOS processes. A more recent method utilizes thick graded buffers of Si1-xGex in which the Ge content is gradually increased up to 100% to relieve the misfit strain with the substrate. Typically 10 μm is required to achieve acceptable levels of threading defects (˜106 cm−2) and a complicated chemical mechanical polishing step is necessary to produce a smooth surface, making device processing expensive (see, Luan et al., Appl. Phys. Lett. 75, 2909 (1999)).
Growth of Ge on Si typically proceeds via the Stranski-Krastanov mechanism, yielding islands (after deposition of 3-4 monolayers) rather than relaxed, continuous layers. For thick films a high roughness is obtained and threading dislocation densities of ˜108 cm−2 are commonly observed, eventually producing the classic crosshatched surface morphologies (see, Fitzgerald and Samavedam, Thin Solid Films, 1997, 294, 3). Scattering and traps at defect sites reduce carrier mobility in electronic devices and also increase dark current in photodetectors. A variety of growth schemes have been developed in an attempt to circumvent some of these problems, including (i) the use of a graded Si—Ge buffer layer (see, Fitzgerald and Samavedam, supra), (ii) a two-step growth in which an initial thin buffer layer is deposited at low temperature, followed by the high temperature growth of the bulk material (see, Luan et al., supra), and (iii) surfactant-mediated epitaxy using As and Sb atomic beams (see, Wietler et al., Appl. Phys. Lett. 2005, 87, 182102). The compositionally graded Si1-xGex buffer layer approach has been demonstrated via UHV-CVD. The Ge concentration is gradually increased as a function of layer thickness, and the terminal Ge portion of the stack exhibits a defect density of 107 cm−2 and a high AFM RMS roughness of 50 nm. A post growth chemical mechanical polishing step is then conducted to reduce the surface roughness to a level that allows subsequent growth of lower defect density overlayers of the Ge material. The drawbacks of this method include excessive final film thicknesses (˜11 μm) and a relatively large residual surface roughness, both of which are problematic for device fabrication. An alternative two-step UHV-CVD process has also been developed (current state-of-the-art) to produce relaxed Ge on Si films with relatively flat surfaces. Here an initiation layer of ˜50 nm in thickness is first grown at low temperatures of ˜350° C. This layer is intended to facilitate subsequent bulk growth at higher temperatures ˜800-900° C. and significantly enhanced rates. In this process the excess hydrogen on the growth surface is believed to act as a surfactant, thereby promoting the formation of misfit dislocations parallel to the Ge/Si interface which relieve the misfit strain. The surface morphology of the resultant films reveals an AFM RMS roughness value of 0.5 nm with no sign of the crosshatch pattern attributed to strain relaxations. However, defect densities of 2.3×107 cm−2 are purportedly present even after thermal cycling of the samples between 780-900° C. Further reductions in defect densities to levels as low as ˜2×106 cm−2 can be obtained using this method via selective growth on oxide patterned Si wafers (see, Fitzgerald and Samavedam, supra). In recent years more conventional surfactant-based approaches have been implemented via MBE to grow Ge layers with suitable morphologies using solid sources of As or Sb. This is typically achieved by first depositing a completed surfactant monolayer on clean Si prior to growth of pure Ge. Using this method ˜1 μm Ge thick films with defect densities of ˜2×107 have been demonstrated at 700° C. (the surface roughness was not reported in this case; see, Wietler et al., supra). The resultant films have been found to exhibit tensile strains as high as 0.2% due to the thermal mismatch with the Si substrate. They are also unintentionally doped by the As/Sb surfactant.
These issues have prompted us to consider an alternative, more straightforward approach which obviates the need for thick buffers and associated processing problems.
SUMMARY OF THE INVENTIONIn one aspect, the present invention provides methods for preparing a semiconductor structure, comprising forming a Ge layer on a semiconductor substrate using an admixture of (a) (GeH3)2CH2 and Ge2H6; (b) GeH3CH3 and Ge2H6; or (c) (GeH3)2CH2, GeH3CH3 and Ge2H6, wherein in all cases, Ge2H6 is in excess.
In a second aspect, the invention provides semiconductor structures produced by the method of the first aspect.
In a third aspect, the invention provides semiconductor structures comprising a silicon-based semiconductor substrate, and a Ge layer formed directly over the silicon-based semiconductor substrate, wherein the Ge layer has a threading dislocation density below 105/cm2.
In a fourth aspect, the invention provides methods for depositing a Ge layer on a substrate in a reaction chamber, comprising introducing into the chamber a gaseous precursor comprising an admixture of (a) (GeH3)2CH2 and Ge2H6; (b) GeH3CH3 and Ge2H6; or (c) (GeH3)2CH2, GeH3CH3 and Ge2H6, wherein in all cases, Ge2H6 is in excess under conditions whereby a layer comprising a Ge material is formed on the substrate.
In a fifth aspect, the invention provides compositions consisting essentially of (GeH3)2CH2 and Ge2H6 in a ratio of between 1:10 and 1:30.
In a sixth aspect, the invention provides compositions consisting essentially of GeH3CH3 and Ge2H6 in a ratio of between 1:5 and 1:30.
In a seventh aspect, the invention provides compositions consisting essentially of GeH3CH3 and (GeH3)2CH2 in a ratio of 1:5 to 1:30 with Ge2H6.
In one aspect, the present invention provides methods for preparing a semiconductor structure, comprising forming a Ge layer on a semiconductor substrate using an admixture of (a) (GeH3)2CH2 and Ge2H6; (b) GeH3CH3 and Ge2H6; or (c) (GeH3)2CH2, GeH3CH3 and Ge2H6, wherein in all cases, Ge2H6 is in excess. The methods of the present invention utilize a range of ratios of the recited components that are compatible for growth conditions that produce selective growth or blanket growth of pure Ge material with an essentially atomically flat surface morphology and device quality defect levels.
It should be understood that when a layer is referred to as being “on” or “over” another layer or substrate, it can be directly on the layer or substrate, or an intervening layer may also be present. It should also be understood that when a layer is referred to as being “on” or “over” another layer or substrate, it may cover the entire layer or substrate, or a portion of the layer or substrate. It should be further understood that when a layer is referred to as being “directly on” another layer or substrate, the two layers are in direct contact with one another with no intervening layer. It should also be understood that when a layer is referred to as being “directly on” another layer or substrate, it may cover the entire layer or substrate, or a portion of the layer or substrate.
As used herein, the terms “substantially atomically planar” and “essentially atomically flat” means that the referenced surface has an RMS roughness value of less than about 1.0 nm as measured by atomic force microscopy according to methods familiar to one skilled in the art. Preferably, that the referenced surface has an RMS roughness value of less than about 0.75 nm or an RMS roughness value ranging from about 0.2 to 1.0 nm or about 0.3 to about 0.75 nm.
The methods of the invention can be used, for example, to produce semiconductor structures for applications in MOSFETs, HBTs, optoelectronic devices and III/V integration on Si. Under conditions disclosed herein we have deposited highly conformal Ge features with atomically flat surfaces and monocrystalline microstructures devoid of penetrating dislocations and threading defects. These are deposited on patterned wafers containing arrays of transistors. The Ge films were found to form readily on the Si exposed regions of the wafer such as the “source” and “drain” component of the device, while no deposition whatsoever was observed on insulator covered poly silicon areas or on device sections masked by nitride-based thin films. The low growth temperatures and the facile, epitaxy driven mechanism afforded by the methods disclosed herein promote strain free Ge to be grown directly on the mismatched Si substrate via compensating edge type dislocation confined to the plane of the heterojunction. The morphology and crystallinity of the films produced in experiments disclosed herein were characterized by high resolution electron microscopy, x-ray diffraction, atomic force microscopy and optical microscopy. Other deposition parameters that may be compatible with selective deposition include widely variable ratios of the (GeH3)2CH2, GeH3CH3 and Ge2H6 reactants in which the Ge2H6 concentration in mixtures involving either one or both of the metal-organic compounds remains in excess.
The Ge layer can be formed as a single layer or a plurality of layers, including selective growth on patterned substrates and can be formed with a total thickness ranging between 40 nm-3 microns or more. For example, the Ge layer can be selectively deposited, utilizing the methods described previously, on a substrate having two or more surface materials, wherein the first surface material comprises an elemental semiconductor material or alloy, such as Si, Ge, or SiGe; the second surface material comprises an oxide, nitride, or oxynitride of an elemental semiconductor material, for example, SiO2, Si3N4, SiON, Ge, GeO2, GeON, or mixtures thereof; and wherein the Ge layer deposits selectively over the first surface material.
In further embodiments, the Ge layer has a density of threading defects of 105/cm2 or less, is virtually strain free, and/or has a substantially atomically planar surface morphology. The Ge layer can be deposited directly on the semiconductor substrate, obviating the need for thick buffer layers.
In certain embodiments, a Ge initiation layer can be formed on a substrate according to the methods described in the first aspect followed by changing the gas source to a second gas source essentially consisting of Ge2H6 to continue to deposit Ge on the Ge initiation layer to form the Ge layer on the substrate. Generally, the Ge initiation layer can have a thickness ranging from about 50 nm to about 1000 nm; preferably, the Ge initiation layer has a thickness ranging from about 50 nm to about 500 nm. In other embodiments, the Ge initiation layer has a thickness ranging from about 50 nm to about 250 nm. The Ge layers formed according to the preceding method can have a density of threading defects of 105/cm2 or less, can be virtually strain free, and/or can have a substantially atomically planar surface morphology. The Ge layer can have a thickness ranging from 50 nm to several microns, for example, up to 1-10 microns.
In other embodiments, the Ge layers formed according to the present methods of the invention are epitaxial. The term “epitaxial” as used herein, means that a material is crystalline and fully commensurate with the substrate. Preferably, epitaxial means that the material is monocrystalline, as defined herein. The term “monocrystalline” as used herein, means a solid in which the crystal lattice of the entire sample is continuous with no grain boundaries or very few grain boundaries, as is familiar to those skilled in the art.
The semiconductor substrate can be any substrate suitable for semiconductor use, including but not limited to silicon, silicon on insulator, SiO2, and Si:C alloys. The semiconductor substrates can be n- or p-doped as is familiar to those skilled in the art; for example, n- or p-doped Si(100). In a preferred embodiment, the substrate comprises silicon, including but not limited to Si(100) and various buffer layers grown on Si. While the methods of the invention do not require a buffer layer, such buffer layers can be prepared with defect densities at least one order of magnitude lower than those possible in the prior art.
The methods comprise depositing the Ge layer on the semiconductor substrate, which may involve introducing into a reaction chamber a gaseous precursor comprising or consisting of an admixture of (a) (GeH3)2CH2 and Ge2H6; (b) GeH3CH3 and Ge2H6; or (c) (GeH3)2CH2, GeH3CH3 and Ge2H6 wherein in all cases, Ge2H6 is in excess. In one embodiment, the admixture can be an admixture of (GeH3)2CH2 and Ge2H6 in a ratio of between 1:10 and 1:20. In another embodiment, the admixture can be an admixture of GeH3CH3 and Ge2H6 in a ratio of between 1:5 and 1:30. In another embodiment, the admixture can be an admixture of GeH3CH3 and Ge2H6 in a ratio of between 1:5 and 1:20. In yet another embodiment, the admixture can be an admixture of GeH3CH3 and Ge2H6 in a ratio of between 1:21 and 1:30. In yet another embodiment, the admixture can be an admixture of GeH3CH3 and Ge2H6 in a ratio of between 1:15 and 1:25.
In a further embodiment, the admixture can be an admixture of a combination of (GeH3)2CH2 and GeH3CH3 at a 1:5 to 1:30 ratio with Ge2H6. In another further embodiment, the admixture can be an admixture of a combination of (GeH3)2CH2 and GeH3CH3 at a 1:5 to 1:20 ratio with Ge2H6. In another further embodiment, the admixture can be an admixture of a combination of (GeH3)2CH2 and GeH3CH3 at a 1:21 to 1:30 ratio with Ge2H6. In another further embodiment, the admixture can be an admixture of a combination of (GeH3)2CH2 and GeH3CH3 at a 1:15 to 1:25 ratio with Ge2H6. In various non-limiting embodiments, the admixtures can be in ratios between 1:5 and 1:15, between 1:5 and 1:10, between 1:10 and 1:20, between 1:0 and 1:15, between 1:21 and 1:30, between 1:22 and 1:30, between 1:23 and 1:30, between 1:24 and 1:30, between 1:25 and 1:30, between 1:26 and 1:30, between 1:27 and 1:30, between 1:28 and 1:30, or between 1:29 and 1:30; or admixtures in ratios of 1:5, 1:6, 1:7, 1:8, 1:9; 1:10; 1:11, 1:12; 1:13; 1:14; 1:15.1:16, 1:17, 1:18, 1:19, 1:20, 1:21, 1:22, 1:23, 1:24, 1:25, 1:26, 1:27, 1:28, 1:29, or 1:30. In various embodiments, the step of introducing the gaseous precursor comprises introducing the gaseous precursors in substantially pure form in the absence of dilutants. In a further preferred embodiment, the step of introducing the gaseous precursor comprises introducing the gaseous precursors as a single gas mixture. In another embodiment, the step of introducing the gaseous precursor comprises introducing the gaseous precursor intermixed with an inert carrier gas. In this embodiment, the inert gas can be, for example, H2 or N2 or other carrier gases that are sufficiently inert under the deposition conditions and process application.
In these aspects, the gaseous precursor can be deposited by any suitable technique, including but not limited to gas source molecular beam epitaxy, chemical vapor deposition, plasma enhanced chemical vapor deposition, laser assisted chemical vapor deposition, and atomic layer deposition. In a further embodiment, the gaseous precursor is introduced by gas source molecular beam epitaxy at between at a temperature of between 350° C. and 450° C., more preferably between 350° C. and 430° C., and even more preferably between 350° to 420°, 360° to 430°, 360 to 420°, 360° to 400°, or 370° to 380°. Practical advantages associated with this low temperature/rapid growth process include (i) short deposition times compatible with preprocessed Si wafers, (ii) selective growth for application in high frequency devices, and (iii) negligible mass segregation of dopants, which is particularly critical for thin layers.
In various further embodiments, the gaseous precursor is introduced at a partial pressure between 10−8 Torr and 1000 Torr. In one embodiment, the gaseous precursor is introduced at between 10−7 Torr and 10−4 Torr gas source molecular beam epitaxy or low pressure CVD. In another embodiment, the gaseous precursor is introduced at between 10−7 Torr and 10−4 Torr for gas source molecular beam epitaxy. In yet another embodiment, the gaseous precursor is introduced at between 10−6 Torr and 10−5 Torr for gas source molecular beam epitaxy.
Further, in any of the preceding aspects and embodiments thereof, a second silicon-based layer can be deposited over the Ge layer. For example, a second silicon-based layer can be deposited over the Ge layer by contacting the same with a silane, such as trisilane, under conditions whereby the second silicon-based layer is deposited. In certain embodiments, the silicon-based layer comprises elemental Si; preferably, the silicon-based layer comprises monocrystalline Si. Such silicon-based layers can be deposited at unprecedented low temperatures, for example, about 400 to about 420° C., despite of the lower surface energy of the Ge template (i.e., the Ge layer), which is known to severely hinder such growth in conventional MBE-based processes. The silicon-based layers can have a thickness ranging from about 2 nm to about 1000 nm; preferably, the thickness ranges from about 2 nm to about 100 nm; more preferably, the thickness ranges from about 2 nm to about 50 nm; or from about 2 nm to about 25 nm.
In other embodiments, a high-k dielectric layer can be formed over the second Si-based layer according to methods familiar to those skilled in the art for depositing high-k dielectric layers over a Si-based layer (e.g., elemental Si or Si(100)). “High-k dielectrics” as used herein means a material having a dielectric constant greater than that of SiO2. For example the high-k dielectric layer can comprise SiNx, Ta2O5, Al2O3, HfSiON, HfO2, HfSiO, ZrO2, HfZrSiO ZrSiO, La2O3, LaAlO3, PZT (lead zirconium titanate), and mixtures thereof.
In a further aspect, the present invention provides semiconductor structures made by the any one of the preceding methods of the invention.
In another aspect, the present invention provides semiconductor structures, comprising
-
- a silicon-based semiconductor substrate, and
- a Ge layer formed directly on the silicon-based semiconductor substrate, wherein the Ge layer has a threading dislocation density below 105/cm2.
All definitions and embodiments described above for the methods of the invention apply to the semiconductor structure aspects of the invention.
The semiconductor structure may further comprise other features as desired, including but not limited to the inclusion of dopants, such as boron, phosphorous, arsenic, and antimony. These embodiments are especially preferred for semiconductor substrates used as active devices. Inclusion of such dopants into the semiconductor substrates can be carried out by standard methods in the art or by use of specialty chemicals.
In a further embodiment, the semiconductor structure may further comprise varying quantities of carbon or tin, as desired for a given application. Inclusion of carbon or tin into the semiconductor substrates can be carried out by standard methods in the art. The carbon can be used to reduce the mobility of the dopants in the structure and more specifically boron. Incorporation of Sn can yield materials with novel optical properties such as direct emission and absorption leading to the formation of Si-based lasers and high sensitivity infrared photodetectors.
In another aspect, the present invention provides a composition, comprising or consisting of (GeH3)2CH2 and Ge2H6 in a ratio of between 1:10 and 1:30. In one embodiment, the present invention provides a composition, comprising or consisting of (GeH3)2CH2 and Ge2H6 in a ratio of between 1:10 and 1:20. In another embodiment, the present invention provides a composition, comprising or consisting of (GeH3)2CH2 and Ge2H6 in a ratio of between 1:21 and 1:30. In another embodiment, the present invention provides a composition, comprising or consisting of (GeH3)2CH2 and Ge2H6 in a ratio of between 1:15 and 1:25. In various non-limiting embodiments, the composition is present in ratios of 1:10 to 1:19; 1:10 to 1:18; 1:10 to 1:17; 1:10 to 1:16; 1:10 to 1:15; 1:21 to 1:30; 1:22 to 1:30; 1:23 to 1:30; 1:24 to 1:30; 1:25 to 1:30; 1:26 to 1:30, 1:27 to 1:30; 1:28 to 1:30; or 1:29 to 1:30; or in ratios of 1:10, 1:11, 1:12, 1:13, 1:14, 1:15, 1:16, 1:17, 1:18, 1:19, or 1:20. In a further embodiment of all of these embodiments, the composition is present in a gaseous form.
In another aspect, the present invention provides a composition, comprising or consisting of GeH3CH3 and Ge2H6 in a ratio of between 1:5 and 1:30. In one embodiment, the present invention provides a composition, comprising or consisting of GeH3CH3 and Ge2H6 in a ratio of between 1:5 and 1:20. In one embodiment, the present invention provides a composition, comprising or consisting of GeH3CH3 and Ge2H6 in a ratio of between 1:21 and 1:30. In another embodiment, the present invention provides a composition, comprising or consisting of GeH3CH3 and Ge2H6 in a ratio of between 1:15 and 1:25. In various non-limiting embodiments, the composition is present in ratios of 1:5 to 1:19; 1:5 to 1:18; 1:5 to 1:17; 1:5 to 1:16; 1:5 to 1:15; 1:5 to 1:14; 1:5 to 1:13; 1:5 to 1:12; 1:5 to 1:11; 1:5 to 1:10; 1:21 to 1:30; 1:22 to 1:30; 1:23 to 1:30; 1:24 to 1:30; 1:25 to 1:30, 1:26 to 1:30, 1:27 to 1:30; 1:28 to 1:30; or 1:29 to 1:30; or in ratios of 1:5, 1:6, 1:7. 1:8, 1:9, 1:10, 1:11, 1:12, 1:13, 1:14, 1:15, 1:16, 1:17, 1:18, 1:19, 1:20, 1:21, 1:22, 1:23, 1:24, 1:25, 1:26, 1:27, 1:28, 1:29, or 1:30. In a further embodiment of all of these embodiments, the composition is present in a gaseous form.
In another aspect, the present invention provides a composition, comprising or consisting of GeH3CH3 and (GeH3)2CH2 in a ratio of 1:5 to 1:30 with Ge2H6. In one embodiment, the present invention provides a composition, comprising or consisting of GeH3CH3 and (GeH3)2CH2 in a ratio of 1:5 to 1:20 with Ge2H6. In one embodiment, the present invention provides a composition, comprising or consisting of GeH3CH3 and (GeH3)2CH2 in a ratio of 1:21 to 1:30 with Ge2H6. In another embodiment, the present invention provides a composition, comprising or consisting of GeH3CH3 and (GeH3)2CH2 in a ratio of 1:15 to 1:25 with Ge2H6. In various non-limiting embodiments, the composition is present in ratios of 1:5 to 1:19; 1:5 to 1:18; 1:5 to 1:17; 1:5 to 1:16; 1:5 to 1:15; 1:5 to 1:14; 1:5 to 1:13; 1:5 to 1:12; 1:5 to 1:11; 1:5 to 1:10; 1:21 to 1:30; 1:22 to 1:30; 1:23 to 1:30; 1:24 to 1:30; 1:25 to 1:30; 1:26 to 1:30, 1:27 to 1:30; 1:28 to 1:30; or 1:29 to 1:30; or in ratios of 1:5, 1:6, 1:7. 1:8, 1:9, 1:10, 1:11, 1:12, 1:13, 1:14, 1:15, 1:16, 1:17, 1:18, 1:19, 1:20, 1:21, 1:22, 1:23, 1:24, 1:25, 1:26, 1:27, 1:28, 1:29, or 1:30. In a further embodiment of all of these embodiments, the composition is present in a gaseous form.
EXAMPLES Example 1 General Ge Deposition ProceduresGe films were grown directly on Si by gas source molecular beam epitaxy (GS-MBE) at 350-420° C. and 5×10−5 Torr using admixtures of either (GeH3)2CH2 or GeH3CH3 and Ge2H6 at optimized molar ratios. The reaction mixture was prepared prior to each deposition by combining the pure compounds in a 100 mL vacuum flask. The total pressure was 115 Torr, which is well below the vapor pressure of (GeH3)2CH2 (248 Torr at 25° C.) or GeH3CH3 which is a gas at room temperature. The flask was connected to a gas injection manifold which was pumped to ˜10−8 Torr on the gas source MBE chamber. A boron doped (1-10 Ω-cm), Si (100) wafer was RCA cleaned and then cleaved to a 1 cm2 size that fits the dimensions of the sample stage. Each substrate was sonicated in methanol for 5 minutes, dried under a stream of purified N2, and inserted through a load lock into the MBE chamber at a base pressure of 8×10−10 Torr. The sample was then heated at 600° C. under UHV to remove surface contaminants until the pressure of the chamber was restored to background levels. The sample was subsequently flashed 10 times to 1000° C. for 1 second to remove any remaining contaminants, and then flashed again at 1150° C. for 5 times to remove the native oxide from the silicon. To commence growth, the wafer was heated to 350-420° C. as measured by single-color pyrometer and the temperature was then allowed to stabilize for 5 minutes. The heating was performed by passing direct current though the samples. The precursor mixture was introduced into the chamber at a flow rate of approximately 0.08 sccm through a manual leak valve. The pressure was maintained constant (5×10−5 Torr) during growth via dynamic pumping using a corrosion resistant turbomolecular pump. The typical deposition times were 0.5-5 hours depending on the desired film thickness. Under these conditions Ge on Si films were produce with thicknesses in the range of 0.30-3 μm at rates approaching 10 nm per minute.
Example 2 Structural and Optical Characterization MethodologyThe samples were extensively characterized for morphology, microstructure, purity and crystallographic properties by atomic force microscopy (AFM), Rutherford backscattering (RBS), secondary ion mass spectrometry (SIMS), cross sectional transmission electron microscopy (XTEM) and high resolution x-ray diffraction (XRD). The threading defects densities were estimated using an etch pit technique (EPD) (see, Luan et al, supra).
Raman studies were carried out at room temperature using several laser lines. The laser beam was focused on the sample using a 100× objective. Typical incident powers were in the 1-2 mW range. The scattered light was dispersed with either an Acton 27.5 cm or an Acton 50.0 cm spectrometer equipped with 1800 and 2400 grooves/mm gratings and Charge Coupled Device detectors. Photoreflectance experiments were performed at room and cryogenic temperatures using a laser wavelength of 514.5 nm. The modulating laser beam was chopped at 1 kHz. Light from a quartz tungsten halogen source was reflected off the sample, dispersed through the Acton spectrometer and focused on an InGaAs detector. The AC component of the signal was extracted using standard lock-in techniques. The same setup minus the halogen lamp was used for photoluminescence experiments.
Example 3 Results and DiscussionThe present invention involves the design and application of purpose built molecular precursors targeted to tailor the surface reactions at the growth front. This is achieved by the deposition of highly reactive compounds such as (GeH3)2CH2 digermylmethane or CH3GeH3 with built-in “pseudosurfactant” capabilities intended to promote layer-by-layer growth of smooth, continuous and relaxed Ge films devoid of deleterious threading dislocations. In the initial stages of growth, the deposition of the molecules proceeds via formation of edge dislocations, which heal the interfacial strain associated with the pseudomorphic growth and suppress the propagation of dislocation cores throughout the layer. During subsequent growth, the compound facilitates a self organized assembly of thick films with atomically flat and defect free surfaces. In the case of (GeH3)2CH2, the driving force for its deposition reactions is the facile elimination of robust CH4 byproducts. Growth of Ge films using the CH3GeH3 compound proceeds via a similar mechanism and by the reaction decomposition pathway of Eq. (1):
GeH3CH3→Ge+H2+CH4 (1)
It is well-known that under the temperature/pressure parameters described above, Ge film growth using pure digermane proceeds via the classic Stranski-Krastanov process by formation and coalescence of three dimensional islands, ultimately leading to the creation of rough films dominated by surface undulations. Collectively, our experiments demonstrate that small concentrations of the (GeH3)2CH2, GeH3CH3 organic additives can profoundly alter this conventional film growth mechanism leading to the assembly of flat, strain-free films with record low defect densities. The highest quality films are obtained for optimum concentration ratios of 1:10 for GeH3CH3 and 1:15 for (GeH3)2CH2 in digermane.
Example 3a Strain-Free Ge Growth Via (GeH3)2CH2At the onset of the deposition experiments we systematically varied the substrate temperature and the reactant concentration in the gaseous mixture to determine the optimum growth parameters and the most favorable surface reaction conditions that would yield the best possible film quality in terms of purity, morphology and microstructure.
We found that the ratio of the (GeH3)2CH2/Ge2H6 precursor flux had a profound effect on all of these material properties. For example, growth using molar ratios in the range form 1:2 to 1:10 produced discontinuous layers with rough surfaces dominated by islands of variable size and shape. These samples were analyzed by RBS carbon resonance profiles, which revealed a significant carbon contamination throughout the layers, primarily in samples grown using molecular ratio of 1:2 and 1:5. This indicates that in such highly concentrated mixtures the “pseudosurfactant” reaction mechanism of the (GeH3)2CH2 compound, which is designed to proceed with complete elimination of the CH2 fragment as CH4, is compromised, resulting in the incorporation of residual C—H impurities. Further reduction of the compound molar ratio in the range of 1:10-1:15 yielded higher-purity, carbon-free films, with atomically smooth surfaces (AFM RMS values of 0.3 nm) and thickness up to ˜3 μm. Under these conditions the growth rate and film thickness were found to depend sensitively on the growth temperature. More specifically, between 430 and 350° C. the nominal rates and corresponding thicknesses were 10-5 nm/min and 3-0.8 μm, respectively, and no discernible growth was observed below 350° C. It is interesting to note that deposition reactions in the range of 450-470° C. produced films that were contaminated with carbon at levels as high as several atomic percent, and exhibited relatively rough surfaces with RMS values or 3-5 nm. This is consistent with previous UHV-CVD experiments of the related GeH3CH3 compound, which yielded similar carbon incorporations.
Finally we also conducted control experiments using pure Ge2H6 (in complete absence of (GeH3)2CH2 from the mixture, i.e, 0:1 ratio), which produced films dominated by surface undulations exhibiting extremely high RMS roughness, as expected. Variation in deposition conditions such as temperature, pressure and compound flux rate did not yield any improvement in the film quality and all layers produced via this method were defective and rough as determined by XTEM and AFM studies. For example AFM scans of Ge films with 450 nm thickness grown at 420° C. via pure Ge2H6, exhibited an RMS roughness of ˜30 nm for 5×5 μm2 areas. In contrast, the roughness for films grown using the 1:15 precursor ratio was found to be 0.2-0.3 nm, which is comparable with atomic step heights on Si. These results collectively indicate that the “pseudosurfactant” properties of (GeH3)2CH2 are enhanced, and the film quality is optimized, for reactant ratios close to 1:15 and growth temperatures in the 350-430° C. range.
AFM studies indicate that of all Ge samples grown under these conditions possess atomically flat surfaces.
XTEM in phase-contrast and Z-contrast modes was extensively used to characterize the microstructure and morphology of the films. The bright field images revealed essentially no penetrating threading defects over an extended lateral range of 2 μm, independently of the layer thickness (50-850 nm) (
SIMS was routinely used to determine the Ge, C and O atomic profiles of representative samples with variable thicknesses grown at 350-420° C. In all cases we find that the layers consist of high purity Ge. More importantly, the SIMS profiles show that the C content is at the detection limit, below 3×1017 cm−3 (bottom of
An etch pit density technique was used to estimate the concentration of threading dislocations in selected samples with thickness close to 0.8 μm. These were etched for typically 200 sec at an average rate of ˜2 nm/sec using a mixture of 700 mL CH3COOH, 70 mL HNO3, 4 mL HF and 270 mg I2. The resulting pits were counted from images obtained by an optical microscope (
XRD reciprocal space maps (RSM) in the vicinity of the (004) and (224) reflections were obtained to precisely determine the in-plane (a) and out-of-plane (c) lattice constants for the films. The data showed that the Ge films grown by this technique were fully relaxed. This result is consistent with the atomically flat surface morphology and suggests that a layer-by-layer growth mechanism must be in operation in the absence of strain. The relaxation in the samples was so complete that several of the films showed a slight tensile strain upon cooling to room temperature, presumably due to the thermal mismatch of Ge with the underlying Si substrate. The (224) RMS plot of a sample with ˜2.5 μm thickness (
Raman, Photoreflectance and Photoluminescence of Ge/Si(100)
The GeH3CH3, analog of (GeH3)2CH2, was also explored as a potentially practical source to grow Ge films on Si at low temperatures between 360 and 420° C. The experiments were performed at constant deposition pressure of 5×10−5 Torr using suitable mixtures of the compound with Ge2H6 at molar concentration ranging mostly from 1:5 to 1:15. The substrate preparation and sample handling procedures were essentially identical to those described above for the experiments involving (GeH3)2CH2.
Preliminary depositions using a dilute reactant ratio of 1:20, GeH3CH3/Ge2H6 produced crystalline but visibly rough layers exhibiting a cloudy surface appearance. AFM indicated an RMS roughness of 8-9 nm and revealed a surface morphology dominated by a network of two-dimensional terraces with variable size, shape and orientation. This is in stark contrast to the usual three-dimensional islands or surface ripples typically observed in conventional Ge growth at this temperature. To further improve the film morphology we increased the reactant molar ratio to the 1:15 and performed a series of depositions at 360, 380, 400 and 420° C. using this mixture. The resultant films displayed a smooth appearance indistinguishable from that of the underlying substrate, indicating that the samples were flat and crystalline.
Characterizations by RBS including ion channeling indicated highly aligned layers in perfect commensuration with the substrate. In most samples the ratio of the channeled versus the random peak heights (χmin), which measures the degree of crystallinity, decreases from 10% at the interface to 4%, across the layer indicating a dramatic reduction in dislocation density (the 4% value represents the theoretical limit in pure Si). RBS carbon resonance experiments were used to confirm the absence of carbon impurities within the bulk of the material. The corresponding depth profiles using a series of ion beam energies with variable penetration depths did not reveal any measurable carbon contamination within the uncertainty of the measurement (less than 0.5 at. %). The RBS derived thicknesses for the Ge/Si(100) layers varied from 400 nm to 600 nm with increasing temperature from 360-420° C. with a concomitant increase in growth rate from 5 to 10 nm/min., respectively.
Among these samples the AFM examinations of films deposited at 400 and 420° C. revealed flat surface morphology with an RMS roughness of 0.2 nm. However, the AFM images also revealed that approximately 5% of the sample surface was covered by an array of perfectly rectangular nanoscale depressions (
Accordingly, we further increased the concentration of GeH3CH3 beyond the 1:15 ratio to determine its effect on the surface properties and film growth rates. Depositions using 1:10 mixtures at temperatures near 420° C. produced thick films exhibiting substantial concentrations of the “match-box” depressions. However, for this reactant stoichiometry lowering the temperature to 360° C. yielded perfectly flat films completely devoid of these features. The growth rate in this case (GeH3CH3 at 360° C.) is 5 nm/minute, which is essentially identical to that obtained from the deposition of (GeH3)2CH2 at 420° C. Thus the GeH3CH3 compound is the best candidate to date for viable low temperature Ge growth compatible with selective area applications. Depositions using 1:5 mixtures produced comparable surface morphologies; however the growth rates were significantly reduced to levels below 3 nm per minute. In the limiting case of using the pure GeH3CH3 compound as the sole reactant (1:0 molar ratio) we obtained negligible film growth. This indicates that the Ge films in these reactions must be generated by the facile surface dissociation of the highly reactive (GeH3)2 species and the role of the organic analog is to catalyze or facilitate layer-by-layer growth. To confirm the absence of carbon in the bulk of the film we analyzed selected samples by SIMS compositional profiles. The data showed that all materials were essentially free of carbon and displayed very similar C and O profiles to those observed previously in the related growth studies using the (GeH3)2CH2 compound.
It is interesting to note that Ge:C films 30 nm thick have been demonstrated in high temperature (<450° C.) depositions of GeH3CH3 by UHV-CVD. In this case the intentional incorporation of carbon at the 1% level purportedly produces surface roughness of ˜3 Å and low defect densities of 3×105 cm−2 (see, Kelly et al., Appl. Phys. Lett. 2006, 88, 152101). SIMS profiles showed that the carbon segregates towards upper/lower portion of the films near the surface/interface regions. The layers are slightly compressed with a net relaxation of ˜78% and are typically too thin for optical applications but could prove useful for Ge-channel MOSFETs (see, Kelly et al., Semicond. Sci. Technol. 2007, 22, S204). Our deposition experiments using GeH3CH3 were conducted on a single-stage wafer configuration at significantly lower temperatures (360° C.) which precluded any the side reactions which might lead to carbon contamination.
Example 3c Ge Depositions Via Initiation LayersTo confirm that (GeH3)2CH2 does not deposit C during growth we conducted two control experiments designed to verify that carbon does not migrate to the surface in a typical surfactant fashion. In both cases an initiation Ge layer of ˜250 nm was grown directly on the Si substrate using the 1:15 reactant ratio, as described above.
In a first experiment, a second growth step is then performed in situ immediately after completion of the first layer in order to maintain an uninterrupted growth environment and—more importantly—an unperturbed “as grown” Ge surface.
In the second experiment, an overlayer of pure Ge was deposited on top of the initiation layer of Ge on Si that had been prepared as described previous. After deposition of the initiation layer, the gas source was switched to pure Ge2H6, to eventually produce a thick composite film (800 nm) representative of the two growth modes. We note that the growth rate during the latter stage using pure Ge2H6 is 4-5 times higher that that using the 1:15 mixture.
XTEM and AFM examinations of the full sample thickness revealed a complete, continuous and monocrystalline layer with an atomically flat surface (AFM RMS ˜0.4 nm). XTEM also showed that the microstructure throughout the growth transition region is continuous and indistinguishable form the bulk material, indicating that the layer-by-layer growth is uninterrupted in the absence of (GeH3)2CH2. SIMS profiles showed a constant Ge content throughout the entire thickness and the typical C and O impurity peaks located at the interface (
In contrast, recent growth studies of Ge on Si, suggest that conventional Sb or As surfactants alter the free energy of the growth surface and promote layer-by-layer growth far beyond the critical thickness (see, Stirman et al., Appl. Phys. Lett. 84, 2530 (2004)). These surfactants remain on the growth surface where they effectively serve a catalytic role throughout the subsequent growth process by mediating chemisorption interactions between the reactants and the surface, reactant diffusion and reducing surface tension.
Example 4 Deposition of SiA pure silicon film was grown on top of a Ge initiation layer (prepared according to Example 1) via decomposition of SiH3SiH2SiH3. This compound incorporates highly reactive SiH2 groups, allowing the formation of monocrystalline Si at unprecedented low temperatures (400-420° C.) despite of the lower surface energy of the Ge template, which is known to severely hinder such growth in conventional MBE-based processes.
A series of Si films with thicknesses ranging from 2-16 nm were deposited using conditions similar to those described above for the Ge growth. High resolution XTEM was used to characterize the crystallinity, surface morphology and epitaxial registry of the films.
The trisilane deposition directly on our Ge buffers described here establishes proof-of-principle routes for producing continuous, fully-pseudomorphic, Si layers with tensile-strained structures and atomically flat surfaces. The successful growth of crystalline Si showing a fully commensurate hetero-interface unambiguously confirms that the original Ge buffer layer surface must be free of significant carbon containing impurities originating from decomposition of the organic source. These buffers should also provide an ideal platform for producing Si epilayers with record high strain states that cannot be accessed via the currently available Si1-xGex (x=˜0.20) counterparts, due to their smaller lattice dimensions. Such materials are desirable for high mobility electronic device applications. The growth of Si directly on Ge would also create new opportunities for the development of Ge-based Metal-Oxide-Semiconductor (MOS) devices (see, Shang et al, IBM Journal of Research and Development 2006, 50, 377). Additionally, thin Si films can act as passivation layers for the growth of high-k dielectrics on Ge (see, De Jaeger et al., Microelectronic Engineering 2005, 80, 26).
Example 5 Selective Growth in Semicondutor Device StructuresSelective growth was conducted using a wafer provided by ASM America (Phoenix Ariz.), incorporating various device architectures including simple transistor structures and various patterns masked by amorphous nitride and oxide thin layers. The wafer was cleaved to produce substrates with suitable dimensions to fit the sample stage of the growth chamber. These were rinsed with methanol and then dipped in HF to remove the thin oxide covering the bare Si surface on the wafer while preserving the nitride masked areas essentially intact. The resulting samples were inserted into the reactor and flashed briefly for 1 sec at 1000° C. to remove any residual contamination from the surface. The Ge films were grown for 30 minutes at 370° C. and 5×10−5 Torr using the 1:15 mixture of (GeH3)2CH2/Ge2H6 which was routinely employed for deposition directly on Si. The “as deposited” samples exhibited an appearance essentially identical to that of the underlying patterned wafer material. Optical microscopy revealed that the appearance of the nitride/oxide masked regions was unchanged while the coloration of the Si-based areas was changed from a metallic grey, typical of Si, to a light brownish hue indicating that selective deposition had occurred. A comprehensive characterization of the wafers by RBS, XRD, AFM, XTEM and EDX (energy dispersive x-ray) revealed the presence of atomically flat Ge films with single crystalline and fully relaxed microstructures throughout the sample.
The film nominal thickness was estimated by the random RBS spectra to be in the 45-50 nm range yielding an average growth rate of ˜2 nm per minute. The channeled spectra indicated that the material was highly aligned and commensurate with the underlying substrate. The XTEM micrograph in
It is well-known that under the temperature/pressure parameters described above, Ge film growth using pure digermane proceeds via the classic Stranski-Krastanov process by formation and coalescence of three dimensional islands, ultimately leading to the creation of rough films dominated by surface undulations. Collectively, our experiments demonstrate that small concentrations of the (GeH3)2CH2, GeH3CH3 organic additives can profoundly alter this conventional film growth mechanism, leading to the assembly of flat, strain-free films with record low defect densities. The highest quality films are obtained for optimum concentration ratios of 1:10 for GeH3CH3 and 1:15 for (GeH3)2CH2 in digermane.
To elucidate how the organic derivatives influence the classic digermane growth process so dramatically we use fundamental bond energies as a guide to develop plausible surface reaction mechanisms. We proceed with the following assumptions: (i) The low temperatures and pressures (10−5 Torr) preclude gas phase reactions amongst the various deposition species (GeH3CH3 or (GeH3)2CH2, and Ge2H6), (ii) the strong C—H bonds prevent facile degradation of the GeH3—CH3 or GeH3—CH2—GeH3 precursors in this temperature range (iii) the precursors react at the growth front by forming Ge—Ge bonds with the underlying surface. This implies that they do not simply physisorb to act as kinetic dilutants or non-bonded surface diffusion barriers but participate in the growth of the film via the deposition of their GeH3 groups.
Simple gas kinetics dictates that the arrival rate of gaseous species at a surface is given by the classic formula J=P/√{square root over (2πmkBT)}, where P=pressure, T=temperature, m=molecular mass and kB is Boltzmann's constant. At a typical temperature of 360° C., and P=5×10−5 Torr, this model gives arrival rates of 70, 60 and 50 molecules/nm2/s for the pure, undiluted (GeH3)2CH2, Ge2H6 and GeH3CH3 compounds, respectively. For the specific case of a 1:10 mixture of GeH3CH3 to Ge2H6 the corresponding flux ratio is also 1:10, which indicates that a significant number of GeH3CH3 molecules arrive at any given time. Using a sticking coefficient σ=0.05, which is a typical value of Ge2H6 on Si surfaces at these conditions (see, Schwarz-Selinger et al., Phys. Rev. B 2002, 65, 125317) we obtain a growth rate R=ΩJσ equal to ˜4-5 nm/min (where Ω=22.7 Å3 is the volume per Ge atom in the film). The R value in this case is remarkable close to that of the growth rates obtained in the experiments described previously. However, the predicted growth rate using the above formula decreases with increasing temperature, in contradiction with our observation that the growth rates increase slightly in going to 420° C. This clearly indicates that higher reaction rates at 420° C. must contribute to the rate of growth.
In the case of deposition of pure GeH3CH3 at 360° C. the lack of a measurable growth rate in spite of the significant impingement flux indicates that a certain minimum concentration of digermane is required at the growth front to activate the release of methane, otherwise the accumulation of organic derivatives might lead to surface passivation and termination of growth. In a related control experiment using (GeH3)2CH2 at a concentration ratio of 1:2-1:5 no measurable growth was observed, corroborating the notion that excess organic derivatives could lead to “surface poisoning”. This assumption precludes a mechanism based on a simple physisorption model and also suggests that the molecules react with the growth surface in the presence of digermane.
Accordingly, we assume that in both cases the CH3GeH3 and (GeH3)2CH2 react and bond to the surface via the Ge atoms of the GeH3 ligands. This likely involves the formation of surface intermediate species comprised —GeH2—CH3 and —GeH2—CH2—GeH2— complexes, respectively which remain intact at the low growth operating temperatures of 360-420° C. due to the strong C—H and Ge—C bonds (see
A similar mechanism is also operative in CVD of SiGe films on Si where adsorbed H is well-known to suppress diffusion of GeHx and thus promote layer-by-layer growth. In our system the heavier and chemically robust organic units provide an effective diffusion barrier on the time scale of H2 desorption. Furthermore, the low temperature reduces the diffusion rate of mobile surface species such as H and Ge adatoms. Although the solubility of C into Ge is negligible, its incorporation into the lattice as a metastable substitutional impurity is nevertheless possible at these low temperatures. However, this requires breaking of multiple C—H bonds, which is unlikely on thermodynamic grounds under these conditions (350-420° C., 10−5 Torr pressure).
Ultimately the elimination of the CH2 and CH3 groups as CH4 is the dominant driving force in this process. This can be quantified using bond enthalpies, which allow approximate calculation of the various plausible surface reactions involving these compounds on H-terminated Ge surfaces which mimic the local growth environment. Although some of the required bond enthalpies are known, the values involving Ge are not fully available (see, Bills and Cotton, J. Phys. Chem. 1964, 68, 806). For this purpose, and for internal consistency in our estimates, we have conducted ab initio calculations to determine the bond enthalpies involving H, C and Ge binary combinations. Using the B3LYP functional at the N311++(2d,2p) level of theory we obtain the values listed in Table 1.
We consider three mechanisms for the reaction of the GeH3CH3 molecules with the Ge growth surface as shown
GeH3CH3(g)+2Ge(ads)—H→Ge(ads)—CH3+Ge(ads)—GeH3+H2(g) ΔH=−2 kcal/mol (Rx1)
The second reaction can viewed as a two step process involving the formation of a surface intermediate complex in which the —CH2GeH3 ligand is bonded to the surface via the CH2 functionality (Rx2a),
GeH3CH3(gas)+Ge(ads)—H→Ge(ads)—CH2GeH3+H2(gas) ΔH=+12 kcal/mol (Rx2a)
The reaction energy of this step is +12 kcal/mol (as shown below) indicating that the formation of GeC bonds on the surface is unfavorable. In the second step the complex decomposes with the release of germane and the binding of a methyl group (Rx2b):
Ge(ads)—CH2GeH3+H2(gas)→Ge(ads)—CH3+GeH4(gas) ΔH=−12 kcal/mol (Rx2b)
The net reaction energy is thus zero because the same number and type of bonds are broken and formed (Rx2a+Rx2b):
GeH3CH3(gas)+Ge(surface)—H→Ge(surface)—CH3+GeH4(gas) ΔH=0 kcal/mol (Rx2a+Rx2b)
The third reaction involves a two step mechanism in which the precursor is adsorbed via the —GeH2 functionality, with a reaction energy of −2 kcal/mol (compared with +12 kcal/mol in the second reaction scenario above). Here, however, the subsequent decomposition step releasing the extremely robust methane molecule, and binding the germyl group onto the Ge surface, evolves an additional −12 kcal/mol. As expected, this is the most favorable reaction with a net overall energy of −14 kcal/mole (Rx3):
GeH3CH3(gas)+Ge(surface)—H→Ge(surface)—GeH3+CH4(gas) ΔH=−14 kcal/mol (Rx3)
From a chemical point of view the release of methane thus represents the dominant driving force in the growth reactions involving the CH3—GeH3 precursor.
For the reaction GeH3—CH2—GeH3 compound we consider only two plausible reaction schemes for the incoming molecules with the Ge growth surface (
In the first case the —GeH3 ligand of GeH3—CH2—GeH3 attaches to the surface via a single proton transfer and releases GeH3CH3. This process is dominated by the breaking of one of the Ge—C bonds, leading to a net energy change of +24 Kcal/mole (note that H2 is not evolved in this process). Since the associated entropy change is small (unimolecular reaction) the process is dominated by the enthalpy change, and is thus highly endothermic. If we consider that the GeH3CH3 then reacts with the surface via methane abstraction as described by (Rx3, −14 Kcal/Mole) the net energy for the entire deposition reaction is:
GeH3—CH2—GeH3+Ge(surface)—H→Ge(ads)—GeH3+GeH3CH3(gas)→2Ge(ads)—GeH3+CH4(gaF) ΔH=+10 kcal/mole.
On the basis of this simple analysis growth via this decomposition route is thus unfavorable.
In the case where both GeH3 end members of GeH3—CH2—GeH3 bind to the surface, as shown in
To corroborate the surface reaction energies obtained from bond enthalpies we carried out a series of large-scale control calculations for the GeH3CH3 on a proton terminated Ge(001) surface using state-of-the-art electronic structure simulations at the LDA level. A parallel implementation of the VASP code (see, Madelung, Semiconductors, Landolt Börstein New Series III; Springer-Verlag: Berlin, N.Y., 2001) was used to obtain all of the optimized structures and electronic properties. The hydrogen terminated Ge(001) substrate was represented by a 180 atom slab with a thickness sufficient to ensure complete bulk behavior in the interior and a large supercell dimension of 80 Å normal to the surface was used to minimize coupling between periodic slab replicas. To preclude the development of long range fields we adopted configurations with symmetrical molecular adsorption geometries on both sides of the slab, as shown in
Claims
1-17. (canceled)
18. A semiconductor structure comprising:
- a silicon-based semiconductor substrate, and
- a Ge layer formed directly over the silicon-based semiconductor substrate, wherein the Ge layer has a threading dislocation density below 105/cm2.
19. The semiconductor structure of claim 18, wherein the substrate comprises Si(100).
20. The semiconductor structure of claim 18, wherein the Ge layer is virtually strain free.
21. The semiconductor structure of claim 18, wherein the Ge layer is virtually atomically flat.
22. The semiconductor structure of claim 18, further comprising a second Si-based layer formed over the Ge layer.
23. The semiconductor structure of claim 22, wherein the second Si-based layer comprises elemental Si.
24. The semiconductor structure of claim 22, further comprising a high-k dielectric layer formed over the second Si-based layer.
25. The semiconductor structure of claim 24, wherein the high-k dielectric layer comprises SiNx, Ta2O5, Al2O3, HfSiON, HfO2, HfSiO, ZrO2, HfZrSiO ZrSiO, La2O3, LaAlO3, PZT, or mixtures thereof.
26-41. (canceled)
Type: Application
Filed: Nov 15, 2010
Publication Date: Mar 17, 2011
Applicant: The Arizona Board of Regents, a body corporate of the State of Arizona acting for and on behalf of A (Scottsdale, AZ)
Inventors: John Kouvetakis (Mesa, AZ), Yan-Yan Fang (Tempe, AZ)
Application Number: 12/946,485
International Classification: H01L 29/78 (20060101);