U.S. patent application number 12/946485 was filed with the patent office on 2011-03-17 for methods and compositions for preparing ge/si semiconductor substrates.
This patent application is currently assigned to The Arizona Board of Regents, a body corporate of the State of Arizona acting for and on behalf of A. Invention is credited to Yan-Yan Fang, John Kouvetakis.
Application Number | 20110062496 12/946485 |
Document ID | / |
Family ID | 43729635 |
Filed Date | 2011-03-17 |
United States Patent
Application |
20110062496 |
Kind Code |
A1 |
Kouvetakis; John ; et
al. |
March 17, 2011 |
Methods and Compositions for Preparing Ge/Si Semiconductor
Substrates
Abstract
The present disclosure describes methods for preparing
semiconductor structures, comprising forming a Ge layer on a
semiconductor substrate using an admixture of (a)
(GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6; (b)
GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6; or (c)
(GeH.sub.3).sub.2CH.sub.2, GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6,
wherein in all cases, Ge.sub.2H.sub.6 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 (GeH.sub.3).sub.2CH.sub.2 and/or GeH.sub.3CH.sub.3 and
Ge.sub.2H.sub.6 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 10.sup.5/cm.sup.2 which can be useful in
semiconductor devices.
Inventors: |
Kouvetakis; John; (Mesa,
AZ) ; Fang; Yan-Yan; (Tempe, AZ) |
Assignee: |
The Arizona Board of Regents, a
body corporate of the State of Arizona acting for and on behalf of
A
Scottsdale
AZ
|
Family ID: |
43729635 |
Appl. No.: |
12/946485 |
Filed: |
November 15, 2010 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
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12133225 |
Jun 4, 2008 |
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12946485 |
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60933023 |
Jun 4, 2007 |
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Current U.S.
Class: |
257/192 ;
257/E29.255 |
Current CPC
Class: |
H01L 21/02381 20130101;
Y10S 438/933 20130101; H01L 21/02532 20130101; H01L 21/0262
20130101 |
Class at
Publication: |
257/192 ;
257/E29.255 |
International
Class: |
H01L 29/78 20060101
H01L029/78 |
Goverment Interests
STATEMENT OF GOVERNMENT INTEREST
[0002] The 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.
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 10.sup.5/cm.sup.2.
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 SiN.sub.x, Ta.sub.2O.sub.5,
Al.sub.2O.sub.3, HfSiON, HfO.sub.2, HfSiO, ZrO.sub.2, HfZrSiO
ZrSiO, La.sub.2O.sub.3, LaAlO.sub.3, PZT, or mixtures thereof.
26-41. (canceled)
Description
CROSS-REFERENCE TO RELATED APPLICATIONS
[0001] 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 .sctn.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.
BACKGROUND OF THE INVENTION
[0003] Elemental Ge exhibits many device advantages over pure Si.
Its smaller bandgap is attractive for photodetector and modulator
applications in the 1.3-1.6 .mu.m 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 .about.10.sup.8
cm.sup.-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.
[0004] Low temperature (T<375.degree. C.) chemical vapor
deposition (CVD) of GeH.sub.4 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.degree. 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
Si.sub.1-xGe.sub.x in which the Ge content is gradually increased
up to 100% to relieve the misfit strain with the substrate.
Typically 10 .mu.m is required to achieve acceptable levels of
threading defects (.about.10.sup.6 cm.sup.-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)).
[0005] 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 .about.10.sup.8 cm.sup.-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 Si.sub.1-xGe.sub.x 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 10.sup.7
cm.sup.-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 (.about.11 .mu.m)
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 .about.50 nm in thickness is first
grown at low temperatures of .about.350.degree. C. This layer is
intended to facilitate subsequent bulk growth at higher
temperatures .about.800-900.degree. 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.times.10.sup.7 cm.sup.-2 are purportedly
present even after thermal cycling of the samples between
780-900.degree. C. Further reductions in defect densities to levels
as low as .about.2.times.10.sup.6 cm.sup.-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 .about.1 .mu.m Ge thick films with defect densities of
.about.2.times.10.sup.7 have been demonstrated at 700.degree. 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.
[0006] 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 INVENTION
[0007] 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)
(GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6; (b)
GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6; or (c)
(GeH.sub.3).sub.2CH.sub.2, GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6,
wherein in all cases, Ge.sub.2H.sub.6 is in excess.
[0008] In a second aspect, the invention provides semiconductor
structures produced by the method of the first aspect.
[0009] 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 10.sup.5/cm.sup.2.
[0010] 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) (GeH.sub.3).sub.2CH.sub.2 and
Ge.sub.2H.sub.6; (b) GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6; or (c)
(GeH.sub.3).sub.2CH.sub.2, GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6,
wherein in all cases, Ge.sub.2H.sub.6 is in excess under conditions
whereby a layer comprising a Ge material is formed on the
substrate.
[0011] In a fifth aspect, the invention provides compositions
consisting essentially of (GeH.sub.3).sub.2CH.sub.2 and
Ge.sub.2H.sub.6 in a ratio of between 1:10 and 1:30.
[0012] In a sixth aspect, the invention provides compositions
consisting essentially of GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6 in
a ratio of between 1:5 and 1:30.
[0013] In a seventh aspect, the invention provides compositions
consisting essentially of GeH.sub.3CH.sub.3 and
(GeH.sub.3).sub.2CH.sub.2 in a ratio of 1:5 to 1:30 with
Ge.sub.2H.sub.6.
BRIEF DESCRIPTION OF THE DRAWINGS
[0014] FIG. 1 shows an AFM of Ge films grown using pure
Ge.sub.2H.sub.6 (top) and (GeH.sub.3).sub.2CH.sub.2/Ge.sub.2H.sub.6
in 1:15 ratio (bottom) at 420.degree. C. indicating rough and
smooth surface layer growth, respectively.
[0015] FIG. 2 shows an (a) and (b) Diffraction contrast micrographs
of Ge layers with 50 and 800 nm thickness, respectively; Note the
lack of penetrating defects and the very smooth surface
morphologies in both cases (c) HR XTEM image of the Ge/Si(100)
hetero-interface. where the arrows indicate the position of the
edge dislocations.
[0016] FIG. 3 shows (Top left) SIMS of a Ge film deposited using
(GeH.sub.3).sub.2CH.sub.2/Ge.sub.2H.sub.6 in 1:15 ratio (region
Ge1) followed by in situ growth of pure Ge.sub.2H.sub.6 (region
Ge2). (right) Corresponding XTEM bright field image showing the
complete Ge layer which is free of penetrating defects and displays
a smooth surface. Arrow and dotted line indicate the switchover
point in the growth. (Bottom) SIMS profile of an 800 nm Ge film
showing no measurable C contamination across the layer. Impurity
level C and O peaks are observed at the sharp Ge/Si interface.
[0017] FIG. 4 is an optical image of an etched Ge surface showing
the location of penetrating defects (marked by oval shape ring).
Inset illustrates an enlarged site of a typical defect, seen as a
dark spot inside the ring.
[0018] FIG. 5 shows an XTEM of a Ge film grown on Si(100) at
420.degree. C. (a) Phase contrast micrograph showing a 2.5 .mu.m
film thickness and a perfectly flat surface. (b) high-resolution
image of the heteroepitaxial interface showing the location of
Lomer defects (circles).
[0019] FIG. 6 shows (224) RSM plots of the Ge film shown in FIG. 4
and Si substrate. The relaxation line passes near the center of the
Ge peak indicating that the material is almost strain-free. This is
consistent with the nearly equal values of the (a) plane and (c)
lattice constants (shown inset) derived from (224) and (004) Bragg
reflections.
[0020] FIG. 7 is a Raman spectrum of a 650 nm Ge film grown on Si
(solid line) compared with that of a bulk Ge substrate (dotted
line). The inset shows room temperature photoreflectance spectra
for both samples, using the same convention for the lines. The
measured energy range corresponds to the direct gap of Ge.
[0021] FIG. 8 is a Raman spectrum of a 500 nm Ge film on Si (solid
line) compared to that of bulk Ge (dotted line). The overlap
between the two spectra is perfect.
[0022] FIG. 9 shows AFM and XTEM images of a Ge sample containing
surface depressions. (a) AFM image (1 .mu.m.times.0.8 .mu.m) in the
vicinity of two representative "match box" depressions. (b)
Diffraction contrast micrograph showing the cross-sectional
morphology for one of these features, including one vertical edge
and a section of the horizontal bottom surface (c) XTEM image of
the sample showing an atomically smooth layer devoid of threading
defects penetrating to the surface within the field of view.
[0023] FIG. 10 is a SIMS profile of an .about.700 nm Ge film
deposited using GeH.sub.3CH.sub.3/Ge.sub.2H.sub.6 in 1:10 ratio at
360.degree. C. The data show no measurable C contamination across
the layer. Impurity level C and O peaks are observed at the sharp
Ge/Si interface.
[0024] FIG. 11 is an XTEM micrograph of a 5 nm Si film grown on a
Ge buffer. The interface is fully epitaxial and the Si epilayer is
partially strained (1.5%) because the film has exceeded the
critical thickness.
[0025] FIG. 12 shows (Left panel) a Raman spectra showing the
Ge--Si, and Si--Si peaks of a fully coherent
Si/Si.sub.0.50Ge.sub.0.50 structure. The tensile strained Si
overlayer with thickness .about.2.0 nm is grown on a Ge buffer via
a self-assembled sub-nanometer thick .about.Si.sub.0.50Ge.sub.0.50
interface layer. (Right panel) Schematic showing a fully coherent
heterostructure comprised of the cubic Ge buffer (light gray
squares), the tetragonally distorted SiGe interface (light and dark
gray) and the tensile strained Si overlayer (dark gray).
[0026] FIG. 13 shows (a) Ge trenches grown selectively in the
"source" and "drain" areas of a device. Note the absence of any
deposition on nitride spacers or on the polysilicon gate hardmask.
(b) Micrograph of the Ge/Si interface showing the location of an
edge dislocation (circle). (c) XTEM image illustrating the
formation of smooth layers devoid of threading defects on Si
exposed areas within the complex wafer structure.
[0027] FIG. 14 shows surface intermediates formed by the reaction
of the GeH.sub.3CH.sub.3 and (GeH.sub.3).sub.2CH.sub.2 with the
hydrogenated film surface. Dark grey, grey and white spheres
represent Ge, C and H atoms, respectively.
[0028] FIG. 15 shows proposed two-step reaction process involving
the evolution of methane and the adsorption of germyl (third
scenario described in the text).
[0029] FIG. 16 shows a proposed two-step reaction process involving
adsorption of GeH.sub.3--CH.sub.2--GeH.sub.3 compound via evolution
of methane.
[0030] FIG. 17 shows supercell geometries used for the calculation
of adsorption energies of GeH.sub.3CH.sub.3 on a hydrogenated
(2.times.1) Ge(001) surface. Reactions Rx1, Rx2 and Rx3 are
depicted in panels (a), (b) and (c), respectively. The dotted line
represents the center of the slab, where bulk conditions
prevail.
DETAILED DESCRIPTION OF THE INVENTION
[0031] 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)
(GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6; (b)
GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6; or (c)
(GeH.sub.3).sub.2CH.sub.2, GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6,
wherein in all cases, Ge.sub.2H.sub.6 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.
[0032] 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.
[0033] 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.
[0034] 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
(GeH.sub.3).sub.2CH.sub.2, GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6
reactants in which the Ge.sub.2H.sub.6 concentration in mixtures
involving either one or both of the metal-organic compounds remains
in excess.
[0035] 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, SiO.sub.2, Si.sub.3N.sub.4,
SiON, Ge, GeO.sub.2, GeON, or mixtures thereof; and wherein the Ge
layer deposits selectively over the first surface material.
[0036] In further embodiments, the Ge layer has a density of
threading defects of 10.sup.5/cm.sup.2 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.
[0037] 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 Ge.sub.2H.sub.6 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
10.sup.5/cm.sup.2 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.
[0038] 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.
[0039] The semiconductor substrate can be any substrate suitable
for semiconductor use, including but not limited to silicon,
silicon on insulator, SiO.sub.2, 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.
[0040] 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) (GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6; (b)
GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6; or (c)
(GeH.sub.3).sub.2CH.sub.2, GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6
wherein in all cases, Ge.sub.2H.sub.6 is in excess. In one
embodiment, the admixture can be an admixture of
(GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6 in a ratio of between
1:10 and 1:20. In another embodiment, the admixture can be an
admixture of GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6 in a ratio of
between 1:5 and 1:30. In another embodiment, the admixture can be
an admixture of GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6 in a ratio of
between 1:5 and 1:20. In yet another embodiment, the admixture can
be an admixture of GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6 in a ratio
of between 1:21 and 1:30. In yet another embodiment, the admixture
can be an admixture of GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6 in a
ratio of between 1:15 and 1:25.
[0041] In a further embodiment, the admixture can be an admixture
of a combination of (GeH.sub.3).sub.2CH.sub.2 and GeH.sub.3CH.sub.3
at a 1:5 to 1:30 ratio with Ge.sub.2H.sub.6. In another further
embodiment, the admixture can be an admixture of a combination of
(GeH.sub.3).sub.2CH.sub.2 and GeH.sub.3CH.sub.3 at a 1:5 to 1:20
ratio with Ge.sub.2H.sub.6. In another further embodiment, the
admixture can be an admixture of a combination of
(GeH.sub.3).sub.2CH.sub.2 and GeH.sub.3CH.sub.3 at a 1:21 to 1:30
ratio with Ge.sub.2H.sub.6. In another further embodiment, the
admixture can be an admixture of a combination of
(GeH.sub.3).sub.2CH.sub.2 and GeH.sub.3CH.sub.3 at a 1:15 to 1:25
ratio with Ge.sub.2H.sub.6. 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, H.sub.2 or N.sub.2 or other carrier gases that are
sufficiently inert under the deposition conditions and process
application.
[0042] 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.degree. C. and
450.degree. C., more preferably between 350.degree. C. and
430.degree. C., and even more preferably between 350.degree. to
420.degree., 360.degree. to 430.degree., 360 to 420.degree.,
360.degree. to 400.degree., or 370.degree. to 380.degree..
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.
[0043] In various further embodiments, the gaseous precursor is
introduced at a partial pressure between 10.sup.-8 Torr and 1000
Torr. In one embodiment, the gaseous precursor is introduced at
between 10.sup.-7 Torr and 10.sup.-4 Torr gas source molecular beam
epitaxy or low pressure CVD. In another embodiment, the gaseous
precursor is introduced at between 10.sup.-7 Torr and 10.sup.-4
Torr for gas source molecular beam epitaxy. In yet another
embodiment, the gaseous precursor is introduced at between
10.sup.-6 Torr and 10.sup.-5 Torr for gas source molecular beam
epitaxy.
[0044] 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.degree. 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.
[0045] 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 SiO.sub.2. For example the high-k
dielectric layer can comprise SiN.sub.x, Ta.sub.2O.sub.5,
Al.sub.2O.sub.3, HfSiON, HfO.sub.2, HfSiO, ZrO.sub.2, HfZrSiO
ZrSiO, La.sub.2O.sub.3, LaAlO.sub.3, PZT (lead zirconium titanate),
and mixtures thereof.
[0046] In a further aspect, the present invention provides
semiconductor structures made by the any one of the preceding
methods of the invention.
[0047] In another aspect, the present invention provides
semiconductor structures, comprising [0048] a silicon-based
semiconductor substrate, and [0049] a Ge layer formed directly on
the silicon-based semiconductor substrate, wherein the Ge layer has
a threading dislocation density below 10.sup.5/cm.sup.2.
[0050] All definitions and embodiments described above for the
methods of the invention apply to the semiconductor structure
aspects of the invention.
[0051] 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.
[0052] 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.
[0053] In another aspect, the present invention provides a
composition, comprising or consisting of (GeH.sub.3).sub.2CH.sub.2
and Ge.sub.2H.sub.6 in a ratio of between 1:10 and 1:30. In one
embodiment, the present invention provides a composition,
comprising or consisting of (GeH.sub.3).sub.2CH.sub.2 and
Ge.sub.2H.sub.6 in a ratio of between 1:10 and 1:20. In another
embodiment, the present invention provides a composition,
comprising or consisting of (GeH.sub.3).sub.2CH.sub.2 and
Ge.sub.2H.sub.6 in a ratio of between 1:21 and 1:30. In another
embodiment, the present invention provides a composition,
comprising or consisting of (GeH.sub.3).sub.2CH.sub.2 and
Ge.sub.2H.sub.6 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.
[0054] In another aspect, the present invention provides a
composition, comprising or consisting of GeH.sub.3CH.sub.3 and
Ge.sub.2H.sub.6 in a ratio of between 1:5 and 1:30. In one
embodiment, the present invention provides a composition,
comprising or consisting of GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6
in a ratio of between 1:5 and 1:20. In one embodiment, the present
invention provides a composition, comprising or consisting of
GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6 in a ratio of between 1:21
and 1:30. In another embodiment, the present invention provides a
composition, comprising or consisting of GeH.sub.3CH.sub.3 and
Ge.sub.2H.sub.6 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.
[0055] In another aspect, the present invention provides a
composition, comprising or consisting of GeH.sub.3CH.sub.3 and
(GeH.sub.3).sub.2CH.sub.2 in a ratio of 1:5 to 1:30 with
Ge.sub.2H.sub.6. In one embodiment, the present invention provides
a composition, comprising or consisting of GeH.sub.3CH.sub.3 and
(GeH.sub.3).sub.2CH.sub.2 in a ratio of 1:5 to 1:20 with
Ge.sub.2H.sub.6. In one embodiment, the present invention provides
a composition, comprising or consisting of GeH.sub.3CH.sub.3 and
(GeH.sub.3).sub.2CH.sub.2 in a ratio of 1:21 to 1:30 with
Ge.sub.2H.sub.6. In another embodiment, the present invention
provides a composition, comprising or consisting of
GeH.sub.3CH.sub.3 and (GeH.sub.3).sub.2CH.sub.2 in a ratio of 1:15
to 1:25 with Ge.sub.2H.sub.6. 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 Procedures
[0056] Ge films were grown directly on Si by gas source molecular
beam epitaxy (GS-MBE) at 350-420.degree. C. and 5.times.10.sup.-5
Torr using admixtures of either (GeH.sub.3).sub.2CH.sub.2 or
GeH.sub.3CH.sub.3 and Ge.sub.2H.sub.6 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
(GeH.sub.3).sub.2CH.sub.2 (248 Torr at 25.degree. C.) or
GeH.sub.3CH.sub.3 which is a gas at room temperature. The flask was
connected to a gas injection manifold which was pumped to
.about.10.sup.-8 Torr on the gas source MBE chamber. A boron doped
(1-10 .OMEGA.-cm), Si (100) wafer was RCA cleaned and then cleaved
to a 1 cm.sup.2 size that fits the dimensions of the sample stage.
Each substrate was sonicated in methanol for 5 minutes, dried under
a stream of purified N.sub.2, and inserted through a load lock into
the MBE chamber at a base pressure of 8.times.10.sup.-10 Torr. The
sample was then heated at 600.degree. 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.degree. C. for 1 second to remove any remaining
contaminants, and then flashed again at 1150.degree. C. for 5 times
to remove the native oxide from the silicon. To commence growth,
the wafer was heated to 350-420.degree. 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.times.10.sup.-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 .mu.m at rates approaching 10 nm
per minute.
Example 2
Structural and Optical Characterization Methodology
[0057] The 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).
[0058] Raman studies were carried out at room temperature using
several laser lines. The laser beam was focused on the sample using
a 100.times. 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 Discussion
[0059] The 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 (GeH.sub.3).sub.2CH.sub.2
digermylmethane or CH.sub.3GeH.sub.3 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
(GeH.sub.3).sub.2CH.sub.2, the driving force for its deposition
reactions is the facile elimination of robust CH.sub.4 byproducts.
Growth of Ge films using the CH.sub.3GeH.sub.3 compound proceeds
via a similar mechanism and by the reaction decomposition pathway
of Eq. (1):
GeH.sub.3CH.sub.3.fwdarw.Ge+H.sub.2+CH.sub.4 (1)
[0060] 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 (GeH.sub.3).sub.2CH.sub.2, GeH.sub.3CH.sub.3 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 GeH.sub.3CH.sub.3 and
1:15 for (GeH.sub.3).sub.2CH.sub.2 in digermane.
Example 3a
Strain-Free Ge Growth Via (GeH.sub.3).sub.2CH.sub.2
[0061] At 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.
[0062] We found that the ratio of the
(GeH.sub.3).sub.2CH.sub.2/Ge.sub.2H.sub.6 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 (GeH.sub.3).sub.2CH.sub.2 compound, which is
designed to proceed with complete elimination of the CH.sub.2
fragment as CH.sub.4, 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 .about.3 .mu.m.
Under these conditions the growth rate and film thickness were
found to depend sensitively on the growth temperature. More
specifically, between 430 and 350.degree. C. the nominal rates and
corresponding thicknesses were 10-5 nm/min and 3-0.8 .mu.m,
respectively, and no discernible growth was observed below
350.degree. C. It is interesting to note that deposition reactions
in the range of 450-470.degree. 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 GeH.sub.3CH.sub.3 compound, which yielded similar carbon
incorporations.
[0063] Finally we also conducted control experiments using pure
Ge.sub.2H.sub.6 (in complete absence of (GeH.sub.3).sub.2CH.sub.2
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.degree. C. via pure Ge.sub.2H.sub.6, exhibited an RMS roughness
of .about.30 nm for 5.times.5 .mu.m.sup.2 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 (GeH.sub.3).sub.2CH.sub.2 are enhanced, and the film
quality is optimized, for reactant ratios close to 1:15 and growth
temperatures in the 350-430.degree. C. range.
[0064] AFM studies indicate that of all Ge samples grown under
these conditions possess atomically flat surfaces. FIG. 1 (top)
shows a tapping mode image of a 400 nm Ge film grown at 420.degree.
C. using the optimal 1:15 ratio exhibiting an RMS of .about.0.21 nm
for 5.times.5 .mu.m.sup.2 areas. Note that a Ge film of comparable
thickness grown at the same conditions using pure Ge.sub.2H.sub.6
(FIG. 1 (bottom)) shows a high roughness of .about.16 nm. Random
and channeling RBS including carbon resonance indicate pure Ge
materials perfectly aligned with the substrate. In most samples the
ratio of the aligned versus the random peak heights
(.chi..sub.min), which measures the degree of crystallinity,
decreases from 10% at the interface to 5% across the layer
indicating a dramatic reduction in dislocation density. HR-XRD
scans including (224) and (004) reciprocal space maps were used to
determine the average unit cell dimensions parallel (a) and
perpendicular (c) to the interface plane. The results showed that
all Ge films are virtually strain free which is consistent with our
observation of atomically smooth surfaces. We note that several
thin samples (<80 nm) revealed a slight tensile strain which is
as high as -0.04% at room temperature (a topic of ongoing studies).
XRD also indicated nominal mosaics spreads of .about.0.14 degrees
and lateral correlations (grain size) of .about.0.25 .mu.m.
[0065] 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 .mu.m,
independently of the layer thickness (50-850 nm) (FIG. 2). The film
surface in all cases was atomically flat. High resolution images in
[110] projection near the Ge/Si heterojunction indicated sharp,
commensurate interfaces and revealed the presence of periodic edge
dislocations confined to the plane of the interface. The arrows in
FIG. 2(c) show the location of well-separated stress-relieving
dislocation cores. The distance between cores corresponds to an
alignment of twenty nine Si rows with every twenty eight Ge rows,
which closely matches the ratio of the lattice constants between Si
and Ge [5.657 .ANG.(Ge)/5.431 .ANG.(Si)=1.04 and 29/28=1.035]. This
alignment perfectly accommodates the lattice mismatch between
Si(100) and Ge(100) leading to layer-by-layer growth of stress free
films with .about.0.2 nm RMS roughness. The edge dislocations can
also be seen in the Z-contrast high resolution images (not shown)
which further confirm the presence of a well defined interface
between Si and Ge. Electron energy loss nanoanalysis across the
interface corroborates this interpretation of the Z-contrast
results and confirms a chemically abrupt transition between Si and
Ge, indicating no elemental intermixing and no measurable levels of
impurities above the usual background.
[0066] SIMS was routinely used to determine the Ge, C and O atomic
profiles of representative samples with variable thicknesses grown
at 350-420.degree. 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.times.10.sup.17
cm.sup.-3 (bottom of FIG. 3). In view of the relatively high
concentration level of C in the reaction mixture (.about.3 at %
with respect to Ge) this finding confirms the complete elimination
of CH.sub.4 from the (GeH.sub.3).sub.2CH.sub.2 molecules, as
expected according to the reaction: (GeH.sub.3).sub.2CH.sub.2
(gas).fwdarw.CH.sub.4 (gas)+2Ge(solid)+2H.sub.2 (gas). We note that
the O content through the film thickness was also at background
levels. Minor C and O contamination peaks, commonly observed in CVD
grown materials can be also seen at the Ge/Si interface in our
samples. These C and O signatures vary from sample to sample
indicating that C (in particular) may not be related to any
measurable decomposition of (GeH.sub.3).sub.2CH.sub.2 in the
vicinity of the interface. However, some C detected at impurity
level concentrations could be due to side reactions of the
precursor with the reactive Si surface. Nevertheless, the fact that
oxygen levels track those of carbon in all of our samples cannot be
explained in terms of the precursor decomposition on the
surface.
[0067] An etch pit density technique was used to estimate the
concentration of threading dislocations in selected samples with
thickness close to 0.8 .mu.m. These were etched for typically 200
sec at an average rate of .about.2 nm/sec using a mixture of 700 mL
CH.sub.3COOH, 70 mL HNO.sub.3, 4 mL HF and 270 mg I.sub.2. The
resulting pits were counted from images obtained by an optical
microscope (FIG. 4). The images showed no dislocations within the
first 300-400 nm over areas of 85.times.64 .mu.m. At 500 nm the
detectable dislocation densities were found to be below
.about.10.sup.5 cm.sup.-2, and this value increased abruptly
towards the interface region. This is consistent with XTEM images
that show no threading defects throughout the upper portion of the
film but a substantial "pile up" near the interface region. These
images also showed atomically flat film surfaces and confirmed the
RBS measured thicknesses which ranged from 0.50-2.5 .mu.m in most
samples (FIG. 5a). The high resolution micrographs revealed a
periodic array of Lomer-type edge dislocations confined to the
interface plane (FIG. 5b). These likely form in the early stages of
growth and provide the necessary strain relief between the highly
mismatched substrate and epilayer. Atomic resolution Z-contrast
images (not shown) confirmed the presence of edge dislocations and
revealed an abrupt transition between the substrate and the
epilayer. The interface roughness is less than 1 nm, which appears
to be similar to that of the underlying substrate. Electron energy
loss nanoanalyis of the interface region indicated no elemental
intermixing on the subnanometer scale. The intediffusion at these
low temperatures is therefore suppressed, leading to atomically
sharp interfaces. No measurable levels of impurities above the
usual background levels were detected by EELS.
[0068] 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 .about.2.5 .mu.m
thickness (FIG. 6) shows that the layer possesses a residual
tensile strain. The full width at half maximum of the (004) Bragg
reflection for this sample is approximately 400 arcseconds and the
mosaic spread derived from the (224) peak is 250 arcseconds,
indicating slight tilt between grains. Secondary ion mass
spectroscopy (SIMS) was also performed on selected samples produced
across the entire 350-430.degree. C. temperature regime. Samples
grown at 420.degree. C. show a carbon content of
.about.3.times.10.sup.16 atoms per cm.sup.3 across the layer, which
is below the noise level of the SIMS measurement (see, Wistey, et
al., J. Appl. Phys. Lett. 2007, 90(8), 082108). The oxygen content
in these films was similarly low and was dominated by the
background level in the SIMS apparatus. We note, however that at
the Ge/Si interface there is a sharp increase in the C and O
concentrations to levels as high as 10.sup.19 cm.sup.-3, as
indicated by the presence of two distinct peaks corresponding to
these elements. Similar features are commonly observed in CVD-grown
materials and they are typically associated with surface
contamination. FIG. 10 at a later section shows similar interface
signatures for the C, and O impurities.
[0069] Raman, Photoreflectance and Photoluminescence of
Ge/Si(100)
[0070] FIG. 7 compares the room temperature Raman spectrum of a 650
nm thick Ge film with that of a reference Ge bulk substrate,
obtained with 532 nm excitation. For this wavelength, the
penetration depth of light in Ge is approximately 18 nm (so that
the Raman signal originates mainly from the first 9 nm). According
to the X-ray data, the film strain is relaxed. Since, as mentioned
above, the only allowed first-order Raman peak corresponds to the
highest frequency vibration in the perfect crystal, the presence of
defects should shift and broaden the line toward lower Raman
shifts. In our case, the film Raman spectrum is very similar to
that of the Ge bulk substrate, confirming the high quality of our
film. A very small frequency upshift is observed, which we assign
to a residual amount of compressive strain. Using Eq. (1), we
estimate that this compressive strain is 0.04%. It is apparent from
FIG. 7 that this is close to the detection limit of the Raman
technique and within the error of the x-ray measurements. The inset
in FIG. 7 shows room temperature photoreflectance results in the
spectral region corresponding to the lowest direct gap E.sub.0. We
fit the spectra with standard expressions for critical point
features, and for the reference Ge substrate we find
E.sub.0=0.803(1) eV. For the Ge film, the transition is slightly
shifted to 0.806(1) eV, and the broadening parameter is virtually
identical to that of bulk Ge. Interestingly, if we use standard
deformation potential theory to calculate the strain shift of the
band gap for 0.04% compressive strain, we find that the heavy-hole
band gap increases by 3 meV, in agreement with the observed shift.
Thus the photoreflectance measurement is consistent with the Raman
result and both confirm the high quality of the Ge-film and the
very small magnitude of the residual strain. Moreover, the example
discussed in FIG. 7 represents in a certain sense a "worst-case"
scenario: in other samples we find absolutely no discernible
difference between the Raman spectrum of the Ge-film and that of
the Ge-substrate. In FIG. 8 we show data for one such sample. Its
Raman spectrum essentially overlaps the Raman spectrum of bulk Ge.
This sample also shows a weak but clearly observable
photoluminescence peak upon excitation with 250 mW of 488 nm light.
The peak maximum at 0.745 eV is very close to the position of the
indirect edge in pure Ge.
Example 3b
Strain-Free Depositions Via GeH.sub.3CH.sub.3
[0071] The GeH.sub.3CH.sub.3, analog of (GeH.sub.3).sub.2CH.sub.2,
was also explored as a potentially practical source to grow Ge
films on Si at low temperatures between 360 and 420.degree. C. The
experiments were performed at constant deposition pressure of
5.times.10.sup.-5 Torr using suitable mixtures of the compound with
Ge.sub.2H.sub.6 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 (GeH.sub.3).sub.2CH.sub.2.
[0072] Preliminary depositions using a dilute reactant ratio of
1:20, GeH.sub.3CH.sub.3/Ge.sub.2H.sub.6 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.degree. 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.
[0073] 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 (.chi..sub.min), 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.degree. C. with a concomitant increase in
growth rate from 5 to 10 nm/min., respectively.
[0074] Among these samples the AFM examinations of films deposited
at 400 and 420.degree. 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 (FIG. 9). XTEM
micrographs indicate that their lateral and vertical dimensions are
100-150 and 10-15 nm, respectively, corroborating the surface
patterns and morphologies observed by AFM. In addition the
horizontal surface and vertical edges were perfectly planar and
approximately orthogonal, indicating that these depressions have a
parallelepiped "match-box"-like shape. The low magnification XTEM
images also revealed an average layer thickness of 400 nm and an
overall flat surface morphology interrupted by the occasional
presence of the depressions in these samples. Finally, we note that
by lowering the sample growth temperature into the 360-400.degree.
C. range produced samples in which the density of these depressions
could be reduced to an arbitrarily low value, but not entirely
eliminated.
[0075] Accordingly, we further increased the concentration of
GeH.sub.3CH.sub.3 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.degree. C. produced thick
films exhibiting substantial concentrations of the "match-box"
depressions. However, for this reactant stoichiometry lowering the
temperature to 360.degree. C. yielded perfectly flat films
completely devoid of these features. The growth rate in this case
(GeH.sub.3CH.sub.3 at 360.degree. C.) is 5 nm/minute, which is
essentially identical to that obtained from the deposition of
(GeH.sub.3).sub.2CH.sub.2 at 420.degree. C. Thus the
GeH.sub.3CH.sub.3 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 GeH.sub.3CH.sub.3 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
(GeH.sub.3).sub.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 (GeH.sub.3).sub.2CH.sub.2
compound.
[0076] FIG. 10 shows a typical profile for a sample deposited at
360.degree. C. using a 1:10 ratio of GeH.sub.3CH.sub.3 in
Ge.sub.2H.sub.6. Note that the carbon content remains below the
detection limit (.about.10.sup.16 at/cm.sup.3) throughout the
entire thickness of the sample, with a sharp rise to a value
<10.sup.20 at/cm.sup.3 at the interface. The C and O profiles
found here are also essentially identical to those observed in our
previous growth studies using the (GeH.sub.3).sub.2CH.sub.2
compound (see, Wistey et al, supra). These levels are in fact quite
typical for CVD materials grown under these conditions. Thus, the
carbon enhancement near the interface is not likely associated with
the reactivity of the particular precursor employed.
[0077] It is interesting to note that Ge:C films 30 nm thick have
been demonstrated in high temperature (<450.degree. C.)
depositions of GeH.sub.3CH.sub.3 by UHV-CVD. In this case the
intentional incorporation of carbon at the 1% level purportedly
produces surface roughness of .about.3 .ANG. and low defect
densities of 3.times.10.sup.5 cm.sup.-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 .about.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 GeH.sub.3CH.sub.3 were conducted on a
single-stage wafer configuration at significantly lower
temperatures (360.degree. C.) which precluded any the side
reactions which might lead to carbon contamination.
Example 3c
Ge Depositions Via Initiation Layers
[0078] To confirm that (GeH.sub.3).sub.2CH.sub.2 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
.about.250 nm was grown directly on the Si substrate using the 1:15
reactant ratio, as described above.
[0079] 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.
[0080] 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 Ge.sub.2H.sub.6, 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 Ge.sub.2H.sub.6 is 4-5 times higher that
that using the 1:15 mixture.
[0081] XTEM and AFM examinations of the full sample thickness
revealed a complete, continuous and monocrystalline layer with an
atomically flat surface (AFM RMS .about.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
(GeH.sub.3).sub.2CH.sub.2. SIMS profiles showed a constant Ge
content throughout the entire thickness and the typical C and O
impurity peaks located at the interface (FIG. 3). In addition, a
small and distinct C and O signal can be seen slightly above the
noise at the switchover point, 250 nm above the Si/Ge interface.
These minor impurities likely originate from exposure of the
surface to the reactor ambient during the .about.1 hour
switch-over. These results indicate that the "CH.sub.2" fragment of
the precursor is continuously eliminated as methane at the growth
front and does not accumulate on the Ge surface or become
incorporated into the film. Furthermore this two step process
yields high growth rates >20 nm/min. making the method viable
for large scale fabrication.
[0082] 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 Si
[0083] A pure silicon film was grown on top of a Ge initiation
layer (prepared according to Example 1) via decomposition of
SiH.sub.3SiH.sub.2SiH.sub.3. This compound incorporates highly
reactive SiH.sub.2 groups, allowing the formation of
monocrystalline Si at unprecedented low temperatures
(400-420.degree. C.) despite of the lower surface energy of the Ge
template, which is known to severely hinder such growth in
conventional MBE-based processes.
[0084] 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. FIG. 11 shows the entire Si film thickness (.about.5 nm)
including the interface region with the Ge buffer in high
resolution. The microstructure appears to be fully commensurate and
the layer surface atomically smooth. FIG. 12 shows the Raman
spectra of a sample with a 2.5 nm Si film on Ge. The spectrum shows
features similar to the Si--Ge and Si--Si Raman peaks in a
Si.sub.0.5Ge.sub.0.5 alloy and a high-energy peak that can be
assigned to a pure Si layer. This peak is down shifted by about 30
cm.sup.-1 with respect to the Raman frequency in bulk Si, which
according to Eq. (1) is consistent with a strain level of 4%. A
full Raman study of the growth of Si on Ge buffers will be
presented elsewhere. This study indicates that the
Si.sub.0.5Ge.sub.0.5 alloy is confined to the immediate interface
between Ge and Si and that highly strained Si films grow upon
subsequent Si deposition (see FIG. 12 for a schematic
interpretation of the growth).
[0085] 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 Si.sub.1-xGe.sub.x (x=.about.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 Structures
[0086] Selective 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.degree. C. to remove
any residual contamination from the surface. The Ge films were
grown for 30 minutes at 370.degree. C. and 5.times.10.sup.-5 Torr
using the 1:15 mixture of (GeH.sub.3).sub.2CH.sub.2/Ge.sub.2H.sub.6
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.
[0087] 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 .about.2 nm per minute. The channeled spectra indicated that the
material was highly aligned and commensurate with the underlying
substrate. The XTEM micrograph in FIG. 13 confirms that the
deposition of Ge occurred selectively in the "source" and "drain"
areas of the devices and not on the nitride spacers or on the
polysilicon gate hardmask. High resolution images of the interface
showed a perfect epitaxial registry between the Ge film and the
surrounding Si wafer structure with typical edge dislocations
accommodating the strain differential between the materials. The
micrographs clearly demonstrated that the Ge deposited conformably
on the sidewalls and bottom of the trench portion of the device
feature entirely filling the drain/source region. The perfectly
flat surface morphology of the Ge layer, as shown in the AFM and
XTEM images (FIG. 13c), is consistent with a layer-by-layer growth
mode. We note that the polygermanium nodule that is observed at the
corner of the spacer and the gate hardmask (FIG. 13a) is where the
wet etch broke through the hardmask enabling germanium nucleation
from the underlying polysilicon. The occurrence of these nodules
can be controlled by further refining the HF-based etching process.
Future experiments will focus on perfecting the sample preparation
and growth conditions to produced relaxed and strained materials on
large area substrates to enable measurements of the electrical and
optical properties on entire arrays of device structures. The
seamless and fully conformal growth of strained Ge in the recessed
source/drain regions might provide a route to extending the use of
group IV materials to induce compressive strain in the silicon
channel for increased hole mobility applications.
Example 6
Calculations of Surface Energetics
[0088] 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 (GeH.sub.3).sub.2CH.sub.2, GeH.sub.3CH.sub.3 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 GeH.sub.3CH.sub.3 and
1:15 for (GeH.sub.3).sub.2CH.sub.2 in digermane.
[0089] 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.sup.-5 Torr) preclude gas phase
reactions amongst the various deposition species (GeH.sub.3CH.sub.3
or (GeH.sub.3).sub.2CH.sub.2, and Ge.sub.2H.sub.6), (ii) the strong
C--H bonds prevent facile degradation of the GeH.sub.3--CH.sub.3 or
GeH.sub.3--CH.sub.2--GeH.sub.3 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 GeH.sub.3 groups.
[0090] 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 (2.pi.mk.sub.BT)}, where P=pressure,
T=temperature, m=molecular mass and k.sub.B is Boltzmann's
constant. At a typical temperature of 360.degree. C., and
P=5.times.10.sup.-5 Torr, this model gives arrival rates of 70, 60
and 50 molecules/nm.sup.2/s for the pure, undiluted
(GeH.sub.3).sub.2CH.sub.2, Ge.sub.2H.sub.6 and GeH.sub.3CH.sub.3
compounds, respectively. For the specific case of a 1:10 mixture of
GeH.sub.3CH.sub.3 to Ge.sub.2H.sub.6 the corresponding flux ratio
is also 1:10, which indicates that a significant number of
GeH.sub.3CH.sub.3 molecules arrive at any given time. Using a
sticking coefficient .sigma.=0.05, which is a typical value of
Ge.sub.2H.sub.6 on Si surfaces at these conditions (see,
Schwarz-Selinger et al., Phys. Rev. B 2002, 65, 125317) we obtain a
growth rate R=.OMEGA.J.sigma. equal to .about.4-5 nm/min (where
.OMEGA.=22.7 .ANG..sup.3 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.degree. C. This
clearly indicates that higher reaction rates at 420.degree. C. must
contribute to the rate of growth.
[0091] In the case of deposition of pure GeH.sub.3CH.sub.3 at
360.degree. 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
(GeH.sub.3).sub.2CH.sub.2 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.
[0092] Accordingly, we assume that in both cases the
CH.sub.3GeH.sub.3 and (GeH.sub.3).sub.2CH.sub.2 react and bond to
the surface via the Ge atoms of the GeH.sub.3 ligands. This likely
involves the formation of surface intermediate species comprised
--GeH.sub.2--CH.sub.3 and --GeH.sub.2--CH.sub.2--GeH.sub.2--
complexes, respectively which remain intact at the low growth
operating temperatures of 360-420.degree. C. due to the strong C--H
and Ge--C bonds (see FIG. 14). These intermediates are envisioned
to be analogous in both function and size to the traditional Sb or
As surfactant atoms used to produce flat Ge films grown on
mismatched substrates. They promote wetting and suppress diffusion
of the Ge adatoms, thereby enforcing the formation of a flat
surface morphology. In our experiments the surface bonded
--GeH.sub.2--CH.sub.3 or GeH.sub.2--CH.sub.2--GeH.sub.2--
intermediates also promote flat layer growth by effectively serving
as site holders which facilitate the organized assembly of smooth
films and are continuously replenished by the incident molecular
precursor flux. Unlike conventional surfactants, which are admitted
at the onset of growth and then recycled, these organic species
react and desorb CH.sub.4, which is generated via Ge--C cleavage
induced by incoming Ge.sub.2H.sub.6 reactant molecules or by
H.sub.2 reaction byproducts. Thus as the molar concentration of
organics is increased from an optimal value to the lowest 1:0
limit, the accumulation of unreacted intermediates saturates the
available reaction sites preventing further Ge growth.
[0093] A similar mechanism is also operative in CVD of SiGe films
on Si where adsorbed H is well-known to suppress diffusion of
GeH.sub.x 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 H.sub.2 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.degree. C., 10.sup.-5 Torr
pressure).
[0094] Ultimately the elimination of the CH.sub.2 and CH.sub.3
groups as CH.sub.4 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.
TABLE-US-00001 TABLE 1 Calculated bond enthalpies for bonds between
H, C and Ge (available experimental values are in parentheses). H C
Ge H 110 105 73 (104) (99) (69) C 80 56 (58) Ge 38
[0095] We consider three mechanisms for the reaction of the
GeH.sub.3CH.sub.3 molecules with the Ge growth surface as shown
FIG. 15. In reaction 1 shown below (Rx1), the molecule decomposes
via breaking the Ge--C bond, to adsorb the methyl (CH.sub.3) and
the germyl (GeH.sub.3) fragments onto the Ge--H surface. The
corresponding reaction energy obtained from the bond enthalpies
equals .DELTA.H=-2 Kcal/mole, assuming that the entropy change for
the process is zero.
GeH.sub.3CH.sub.3(g)+2Ge.sub.(ads)--H.fwdarw.Ge.sub.(ads)--CH.sub.3+Ge.s-
ub.(ads)--GeH.sub.3+H.sub.2(g) .DELTA.H=-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
--CH.sub.2GeH.sub.3 ligand is bonded to the surface via the
CH.sub.2 functionality (Rx2a),
GeH.sub.3CH.sub.3(gas)+Ge.sub.(ads)--H.fwdarw.Ge.sub.(ads)--CH.sub.2GeH.-
sub.3+H.sub.2(gas) .DELTA.H=+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.sub.(ads)--CH.sub.2GeH.sub.3+H.sub.2(gas).fwdarw.Ge.sub.(ads)--CH.sub-
.3+GeH.sub.4(gas) .DELTA.H=-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):
GeH.sub.3CH.sub.3(gas)+Ge.sub.(surface)--H.fwdarw.Ge.sub.(surface)--CH.s-
ub.3+GeH.sub.4(gas) .DELTA.H=0 kcal/mol (Rx2a+Rx2b)
The third reaction involves a two step mechanism in which the
precursor is adsorbed via the --GeH.sub.2 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):
GeH.sub.3CH.sub.3(gas)+Ge.sub.(surface)--H.fwdarw.Ge.sub.(surface)--GeH.-
sub.3+CH.sub.4(gas) .DELTA.H=-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 CH.sub.3--GeH.sub.3 precursor.
[0096] For the reaction GeH.sub.3--CH.sub.2--GeH.sub.3 compound we
consider only two plausible reaction schemes for the incoming
molecules with the Ge growth surface (FIG. 16). The first proceeds
with the attachment of a single germyl group to the surface, while
the second involves simultaneous binding of both germyl end members
to the surface. The third possibility in which the molecules binds
to the surface via the central --CH.sub.2-- is highly unfavorable
on energetic grounds and is not considered.
[0097] In the first case the --GeH.sub.3 ligand of
GeH.sub.3--CH.sub.2--GeH.sub.3 attaches to the surface via a single
proton transfer and releases GeH.sub.3CH.sub.3. 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 H.sub.2 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 GeH.sub.3CH.sub.3 then reacts with the surface via methane
abstraction as described by (Rx3, -14 Kcal/Mole) the net energy for
the entire deposition reaction is:
GeH.sub.3--CH.sub.2--GeH.sub.3+Ge.sub.(surface)--H.fwdarw.Ge.sub.(ads)---
GeH.sub.3+GeH.sub.3CH.sub.3(gas).fwdarw.2Ge.sub.(ads)--GeH.sub.3+CH.sub.4(-
gaF) .DELTA.H=+10 kcal/mole.
On the basis of this simple analysis growth via this decomposition
route is thus unfavorable.
[0098] In the case where both GeH.sub.3 end members of
GeH.sub.3--CH.sub.2--GeH.sub.3 bind to the surface, as shown in
FIG. 16, we obtain a strained methylene-bridged
--GeH.sub.2--CH.sub.2--GeH.sub.2-- complex via the liberation of 2
H.sub.2 molecules and an energy change of -4 Kcal/mole. In the next
step methane is released with a corresponding reaction energy of
-24 Kcal/mole. The net reaction energy for this entire process is
therefore -28 Kcal/mole, or -14 Kcal/mole per --GeH.sub.3 adsorbed.
This is therefore a highly favorable route comparable to that
described by (Rx3). Note that the significant bond strain
associated with the methylene-bridge
--GeH.sub.2--CH.sub.2--GeH.sub.2-- complex is due to the small size
of the carbon atoms compared to Ge and likely promotes the facile
removal of the CH.sub.2 as methane.
[0099] To corroborate the surface reaction energies obtained from
bond enthalpies we carried out a series of large-scale control
calculations for the GeH.sub.3CH.sub.3 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 Borstein 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 .ANG. 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 FIG. 17. Core
electrons were replaced by ultra-soft pseudoptentials and the
remaining valence electrons were expanded in a plane-wave basis
using an energy cutoff of 500 eV, and single k-point sampling at
.GAMMA.. Unlike our simple bond enthalpy models above the
large-scale slab representations adopted here implicitly
incorporate surface relaxation effects induced by the adsorbing
molecules. This may be particularly important in the case of
(GeH.sub.3).sub.2CH.sub.2 adsorption since the methylene-bridge
--GeH.sub.2--CH.sub.2--GeH.sub.2-- complex may transfer strain from
the complex to the surface. Our detailed simulations confirm that
the adsorbed complex bears the majority of the distortion while the
surrounding Ge substrate remains essentially unperturbed.
Accordingly, in this case the simple addition of bond enthalpies is
fairly quantitative. In fact our calculated energies confirm the
general trends predicted by the bond enthalpy calculations for all
cases considered. For example, in Reaction Rx2, where the methyl is
adsorbed and GeH.sub.4 is released, we find a energy change of -1
kcal/mol for the overall process. This is close to zero, as
expected on the basis of bond formation/breakage discussed above.
Similarly, the most energetic reaction involving the release of
methane and the binding of a GeH.sub.3 occurs with an energy
difference of -17 kcal/mol, which is very similar to the value -14
kcal/mol predicted using bond enthalpies. The main origin of this
energy difference is the use of the LDA in the slab calculations
and B3LYP treatment in the bond enthalpies. A .about.20-30%
overbinding is typical of LDA with respect to B3LYP, suggesting
that bond enthalpies in Table 1 can be used as a reasonable
quantitative tool for characterizing the simple surface reactions
in our H--C--Ge system.
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