U.S. patent application number 13/033361 was filed with the patent office on 2011-08-18 for methods and compositions for preparing tensile strained ge on ge1-ysny buffered 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. Invention is credited to Yan-Yan Fang, John Kouvetakis.
Application Number | 20110198729 13/033361 |
Document ID | / |
Family ID | 43769845 |
Filed Date | 2011-08-18 |
United States Patent
Application |
20110198729 |
Kind Code |
A1 |
Kouvetakis; John ; et
al. |
August 18, 2011 |
Methods and Compositions for Preparing Tensile Strained Ge on
Ge1-ySNy Buffered Semiconductor Substrates
Abstract
The present disclosure describes methods for preparing
semiconductor structures, comprising forming a Ge.sub.1-ySn.sub.y
buffer layer on a semiconductor substrate and forming a tensile
strained Ge layer on the Ge.sub.1-ySn.sub.y buffer layer using an
admixture of (GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6 in a
ratio of between 1:10 and 1:30. The disclosure further provides
semiconductor structures having highly strained Ge epilayers (e.g.,
between about 0.15% and 0.45%) as well as compositions comprising
an admixture of (GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6 in a
ratio of between about 1:10 and 1:30. The methods herein provide,
and the semiconductor structure provide, Ge epilayers having high
strain levels which can be useful in semiconductor devices for
example, in optical fiber communications 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
Scottsdale
AZ
of Arizona State University
|
Family ID: |
43769845 |
Appl. No.: |
13/033361 |
Filed: |
February 23, 2011 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
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12133221 |
Jun 4, 2008 |
7915104 |
|
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13033361 |
|
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60933013 |
Jun 4, 2007 |
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Current U.S.
Class: |
257/616 ;
257/E29.1 |
Current CPC
Class: |
H01L 21/0262 20130101;
H01L 21/0245 20130101; H01L 21/02631 20130101; H01L 21/0251
20130101; H01L 21/02381 20130101; H01L 21/02532 20130101; H01L
21/02452 20130101 |
Class at
Publication: |
257/616 ;
257/E29.1 |
International
Class: |
H01L 29/12 20060101
H01L029/12 |
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. A semiconductor structure produced by a method comprising: a)
forming a Ge.sub.1-ySn.sub.y buffer layer on a semiconductor
substrate; and b) forming a tensile strained Ge layer on the
Ge.sub.1-ySn.sub.y buffer layer using an admixture of
(GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6 in a ration of
between 1:10 and 1:30.
2. A semiconductor structure comprising: a semiconductor substrate,
a Ge.sub.1-ySn.sub.y buffer layer formed over the substrate, and a
tensile strained Ge layer formed over the Ge.sub.1-ySn.sub.y buffer
layer.
3. The semiconductor structure of claim 2, wherein the tensile
strained Ge layer has an essentially atomically flat surface.
4. The semiconductor structure of claim 2, wherein the substrate
comprises silicon.
5. The semiconductor structure of claim 2, wherein y is between
0.015-0.045.
6. The semiconductor structure of claim 2, wherein the
Ge.sub.1-ySn.sub.y buffer layer is at least 95% relaxed.
7. The semiconductor structure of claim 2, wherein the symmetry of
the tensile Ge layer is tetragonal.
8. The semiconductor structure of claim 2, wherein the Ge layer has
a tunable tensile strain of between 0.15% and 0.45%
9. The semiconductor structure of claim 2, wherein the Ge layer is
formed directly on the Ge.sub.1-ySn.sub.y buffer layer.
Description
CROSS-REFERENCE TO RELATED APPLICATIONS
[0001] This application is a Divisional of U.S. patent application
Ser. No. 12/133,221, filed Jun. 4, 2008 which claims the benefit
under 35 USC .sctn.119(e), of U.S. Provisional Application Ser. No.
60/933,013, filed 4 Jun. 2007, which is hereby incorporated by
reference in its entirety.
BACKGROUND OF THE INVENTION
[0003] Germanium has a direct band gap E.sub.0=0.81 eV at room
temperature (see, MacFarlane and Roberts, Phys. Rev. 97, 1714
(1955)), which corresponds to an optical wavelength of 1.54 .mu.m.
Although this is barely enough to reach the telecom C-band (1.53
.mu.m-1.56 .mu.m), the very strong wavelength dependence of the
absorption coefficient near the direct edge suggests that small
perturbations that shift E.sub.0 to lower energies should
dramatically improve the performance of this material for use in
optical fiber communications. The direct band gap of Ge can be
reduced by alloying with Sn and by applying stress. The dependence
of E.sub.0 on Sn-concentration has been measured and found to be
stronger than predicted, to the extent that the addition of only 2%
of Sn increases the absorption coefficient at 1.55 .mu.m by more
than one order of magnitude (see, D'Costa et al., Phys. Rev. B 73,
125207 (2006)). The use of stress as a perturbation is problematic
because the direct band gap can only be lowered with tensile
strain.
[0004] This strain cannot be obtained by growing epitaxial Ge on
Ge.sub.1-xSi.sub.x alloys because the latter have a smaller lattice
parameter. In spite of this inherent limitation, tensile-strained
Ge has been obtained by depositing the material directly on Si at
relatively high temperatures and by exploiting the smaller thermal
expansion of the substrate to induce stress in the Ge epilayer when
the sample is quenched from .about.800.degree. C. (see, Ishikawa et
al., Appl. Phys. Lett. 82, 2044 (2003); Cannon, et al., Appl. Phys.
Lett. 84, 906 (2004); Liu et al., Appl. Phys. Lett. 87, 103501
(2005); Wietler et al., Thin Solid Films 508, 6 (2006); and Liu et
al., Phys. Rev. B 70, 155309 (2004)). This process leads to biaxial
tensile strains as high as 0.25% in films as thick as 1 .mu.m.
While such tensile strain values may be sufficient for some
photodetector designs, higher strain values are necessary for most
optoelectronic applications which require tunable direct gaps (see,
Liu et al., 2004, supra).
[0005] Further limitations of the thermal expansion process include
lack of precise strain control and a maximum predicted strain value
of 0.3% for growth at 900.degree. C. (see, Cannon, et al., supra).
Moreover, the use of high temperatures (800-900.degree. C.)
typically induces inter-diffusion of the elements across the Si--Ge
heterojunction, resulting in non-uniform and potentially defective
interfaces. In the context of laser applications, spatial
confinement requires abrupt interfaces, which are precluded using
this high temperature process due to the inherent elemental
intermixing at the interface. In addition, precise and systematic
control of the final strain state has not been demonstrated using
this method, and this hampers the design of devices.
[0006] Thus, there exists a need in the art for improved methods of
preparing tensile strained Ge on semiconductor substrates.
SUMMARY OF THE INVENTION
[0007] In one aspect, the present invention provides methods for
preparing a semiconductor structure comprising: [0008] a) forming a
Ge.sub.1-ySn.sub.y buffer layer on a semiconductor substrate; and
[0009] b) forming a tensile strained Ge layer on the
Ge.sub.1-ySn.sub.y buffer layer using an admixture of
(GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6 in a ratio of between
1:10 and 1:30.
[0010] In a second aspect, the present invention provides
semiconductor structures made by the methods of the first aspect of
the invention.
[0011] In a third aspect, the present invention provides
semiconductor structures comprising a semiconductor substrate; a
Ge.sub.1-ySn.sub.y buffer layer formed over the substrate, and a
tensile strained Ge layer formed over the Ge.sub.1-ySn.sub.y buffer
layer.
[0012] In a fourth aspect, the present invention provides
compositions comprising (GeH.sub.3).sub.2CH.sub.2 and
Ge.sub.2H.sub.6 in a ratio of between 1:10 and 1:30.
[0013] Herein is provided an approach based for the generation of
tensile-strained Ge epilayers based on the creation of
Ge.sub.1-ySn.sub.y alloys with tunable lattice dimensions above
that of Ge which serve as the critical facilitating platforms for
the subsequent tensile-Ge growth. In addition this approach
subsumes the following distinctive features: (i) low growth
temperature that promotes the assembly of highly-strained
tetragonally-distorted Ge structures that remain robust despite the
inherent metastability. (ii) layer-by-layer growth mechanisms
leading to flat surfaces, chemically abrupt interfaces devoid of
chemical intermixing and relatively defect free layer
microstructures. Both features are enabled by exploiting the high
reactivity and the pseudo-surfactant behavior of the
(GeH.sub.3).sub.2CH.sub.2 species. Collectively this methodology
has allowed the systematic production of Ge layers with very high
tensile strains and all of the desired morphological and structural
properties as discussed below. Thereby, the utility of Ge can be
extended into the wider infrared optoelectronic domain by tuning
its fundamental optical properties using strain as a main
parameter
BRIEF DESCRIPTION OF THE DRAWINGS
[0014] FIG. 1 shows the electronic band structure of Ge near the
lowest direct band gap E.sub.0. Here, k is the wave vector along
the <100> direction, and a the cubic lattice constant. The
bands are indicated as cb (conduction band), hh (heavy hole band),
lh (light hole), and so (split off band). The valence band
spin-orbit splitting at k=0 is .DELTA..sub.0. The right panel shows
the bands calculated for a tensile strain of 1% using absolute
deformation potentials as in Ref. 15. The strain dependence of the
effective masses was ignored. The zero of energy is chosen at the
top of the valence band in the unstrained case (left panel).
[0015] FIG. 2 is a high resolution image of the corresponding an
Ge/GeSn interface (arrow) for an embodiment of the invention
showing perfect commensuration of the (111) lattice planes.
[0016] FIG. 3 is a diffraction contrast XTEM image of an entire
Ge/GeSn/Si structure according to an embodiment of the
invention.
[0017] FIG. 4 is a room temperature Raman spectrum comparing a
tensile-strained Ge film (solid line) and a bulk Ge crystal (dotted
line). The spectra were obtained with 514.5 nm
[0018] FIG. 5 is a room temperature Raman spectrum comparing a
tensile-strained Ge film (solid line) and a bulk Ge crystal (dotted
line). The spectra were obtained with 514.5 nm
[0019] FIG. 6 is a graph illustrating the plane strain induced
Raman shifts in tensile-strained Ge/GeSn films (solid line) as a
function of the in-plane strain measured via XRD.
[0020] FIG. 7 is an XRD reciprocal space map of the (224)
reflections, showing a Ge epilayer and a Ge.sub.0.975Sn.sub.0.025
buffer relative to the Si substrate peak for an embodiment of the
invention. The Ge peak falls above the relaxation line connecting
the Si peaks and the origin, indicating tensile strain.
[0021] FIG. 8 is an XRD reciprocal space map of the (135)
reflections, showing a Ge epilayer and a Ge.sub.0.975Sn.sub.0.025
buffer relative to the Si substrate peak for an embodiment of the
invention. As for the (224) peak, the (135) Ge peak lies above the
relaxation line connecting the Si peaks and the origin, indicating
tensile strain.
[0022] FIG. 9 is an XRD reciprocal space map of the (224)
reflections, showing the Ge epilayer and the
Ge.sub.0.965Sn.sub.0.035 buffer relative to the Si substrate peak
for an embodiment of the invention. The Ge peak falls above the
relaxation line connecting the Si peaks and the origin, indicating
tensile strain of .about.0.45%.
[0023] FIG. 10 is a strain equilibration in the
Ge/Ge.sub.0.965Sn.sub.0.035 bilayer. The relaxed Ge and GeSn
in-plane lattice parameters (dotted lines) are shifted (arrows)
towards a common minimum (open circle) to minimize the combined
strain energy. Note that the buffer layer is only slightly
compressed.
[0024] FIG. 11 is a diffraction contrast XTEM image of the
Ge/Ge.sub.0.965Sn.sub.0.035/Si (100) heterostructure:
[0025] FIG. 12 is a high resolution XTEM micrograph of the
corresponding interface of FIG. 11.
[0026] FIG. 13 is a 2.times.2 .mu.m.sup.2 AFM scan showing a smooth
surface created by the coalescence of atomically flat terraces.
Inset is a 0.5.times.0.5 .mu.m.sup.2 AFM scan with RMS roughness of
0.2 nm.
DETAILED DESCRIPTION OF THE INVENTION
[0027] As demonstrated herein, the tensile-strained Ge layers of
the invention display homogeneous compositional and strain
profiles, low threading dislocation densities and atomically planar
surfaces, and the resulting semiconductor structure and methods for
making it are compatible with, for example, selective growth,
optical fiber communication applications, and back-end CMOS
telecommunication applications.
[0028] In the case of pure Ge, increasing the tensile strain beyond
the values allowed by the thermal expansion method is highly
desirable. The so-called telecom U-band can be covered with a
biaxial tensile strain of 0.4%. Our approach has provided strains
as high as 0.45% which is suitable for generating U-bands and
represents the highest value ever observed. As described below, the
availability of Sn-based buffers allows the growth of strained Ge
films to be conducted at unprecedented low temperatures
(350-380.degree. C.) which are compatible with selective growth and
back-end CMOS telecommunication applications.
[0029] The Ge.sub.1-ySn.sub.y buffer layer can be a single layer,
or can be a plurality of layers, such as a graded layer, and can be
formed with a thickness ranging between 50 nm and several microns.
In various embodiments, the semiconductor buffer layer has a
thickness in a range from 50 nm to 10 microns. In various other
embodiments, the semiconductor buffer layer has a thickness in a
range from about 50 nm to 1 micron. In various embodiments, the
semiconductor buffer layer has a thickness in a range from 50 nm to
500 nm or in a range from about 20 nm to 300 nm.
[0030] In various embodiments, the Sn content varies between 1.5%
and 3.5% (i.e., y is between 0.015 and 0.035). In a further
embodiment, the method further comprises in situ thermal cycling
(i) at a temperature of between 500.degree. C. and 600.degree. C.
for a period of time ranging from about 1 sec. to 120 minutes; or
(ii) by rapid thermal annealing at temperatures ranging up to about
850-900.degree. C. for 1-100 seconds, following forming of the
Ge.sub.1-ySn.sub.y buffer layer on a semiconductor substrate, to
reduce residual strain. In a further embodiment, the
Ge.sub.1-ySn.sub.y buffer layer is at least 90% relaxed; in various
further embodiments, at least 91%, 92%, 93%, 94%, 95%, 96%, 97%, or
more relaxed.
[0031] The Ge layer can be formed as a single layer, and can be
formed over or formed directly on the one or more
Ge.sub.1-ySn.sub.y buffer layers (as noted above) with a thickness
ranging between 30-200 nm or thicker depending on the critical
thickness of the sample. In certain embodiments, the Ge layer can
be formed as a single layer, and can be formed over or formed
directly on the one or more Ge.sub.1-ySn.sub.y buffer layers (as
noted above) with a thickness ranging between about 50-200 nm or
thicker depending on the critical thickness of the sample.
[0032] In certain embodiments, the Ge layer can be formed as a
single layer directly on the one or more Ge.sub.1-ySn.sub.y buffer
layers. In certain other embodiments, the Ge layer can be formed as
a single layer formed directly on the one or more
Ge.sub.1-ySn.sub.y buffer layers (as noted above) with a thickness
ranging between about 30-200 nm or about 50-200 nm or thicker
depending on the critical thickness of the sample. 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] In various embodiments, the Ge layer has a tunable tensile
strain of between 0.15% and 0.45% In a further embodiment, the Ge
layer is tetragonal; tetragonal distortion leads to the split of
the heavy/light hole bands which ultimately leads to novel
optoelectronic properties such as direct gaps and high mobilities
in doped systems. As used herein, "tensile strain" refers to
deformation along a layer segment that increases in length and
width when a load/perturbation is applied that stretches the layer
within the horizontal pane and reduces its vertical dimension.
[0034] The semiconductor substrate can be any substrate suitable
for semiconductor use, including but not limited to silicon,
silicon on insulator, SiO.sub.2, Si:Ge alloys, and Si:C alloys. In
a preferred embodiment, the substrate comprises silicon, including
but not limited to Si(100). 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).
[0035] In various embodiments, the semiconductor buffer layer has a
thickness in a range from 50 nm to several microns. In various
embodiments, the semiconductor buffer layer has a thickness in a
range from 50 nm to 10 microns. In various other embodiments, the
semiconductor buffer layer has a thickness in a range from about 50
nm to 1 micron. In various embodiments, the semiconductor buffer
layer has a thickness in a range from 50 nm to 500 nm or in a range
from about 20 nm to 300 nm. In further embodiments, the
semiconductor buffer layers have a density of threading defects of
10.sup.6/cm.sup.2 or less.
[0036] In a further embodiment, the semiconductor substrates of the
invention comprise a Ge.sub.1-ySn.sub.y buffer layer and/or a Ge
layer having a substantially atomically planar surface morphology
(i.e., essentially atomically flat). 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.
[0037] In other embodiments, the Ge layers formed according to the
present methods of the invention are Ge layer 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.
[0038] The methods comprise depositing the Ge layer on the buffer
layer, which may involve introducing into a reaction chamber a
gaseous precursor comprising or consisting of an admixture of
(GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6 in a ratio of between
1:10 and 1:30 (or, in other embodiments, 1:10 to 1:25; 1:11 to
1:25, 1:12 to 1:25; 1:13 to 1:25; 1:14 to 1:25; 1:15 to 1:25; 1:10
to 1:20; 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), under conditions whereby the Ge layer material is
formed on the buffer layer. In a particular embodiment, the
admixture comprises (GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6
in a ratio of between 1:15 and 1:25.
[0039] In another embodiment, the methods comprise depositing the
Ge layer on the buffer layer, which may involve introducing into a
reaction chamber a gaseous precursor comprising or consisting of an
admixture of (GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6 in a
ratio of between about 1:20 and 1:30 (or, in other embodiments,
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; and 1:29 to
1:30.
[0040] In various embodiments, the step of introducing the gaseous
precursor comprises introducing the gaseous precursor in
substantially pure form. 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.
[0041] In the methods of the invention, 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
300.degree. C. and 420.degree. C., more preferably between
350.degree. C. and 400.degree. C., and even more preferably between
350.degree. C. to 380.degree. C. 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.
[0042] 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.-8 Torr and 10.sup.-3 Torr for gas source molecular
beam epitaxy. 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.
[0043] All definitions and embodiments described above for the
methods of the invention apply to the semiconductor structure
aspects of the invention.
[0044] The semiconductor structures of the invention 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.
[0045] In another aspect, the present invention provides
composition, comprising (GeH.sub.3).sub.2CH.sub.2 and
Ge.sub.2H.sub.6 in a ratio of between 1:10 and 1:30. In various
preferred embodiments, the composition comprises the individual
components in a ratio of 1:10 to 1:25; 1:11 to 1:25, 1:12 to 1:25;
1:13 to 1:25; 1:14 to 1:25; 1:15 to 1:25; 1:10 to 1:20; 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. In a
further embodiment of each of these embodiments, the composition is
in a gaseous form. In particular embodiment, the composition
comprises (GeH.sub.3).sub.2CH.sub.2 and Ge.sub.2H.sub.6 in a ratio
of between 1:15 and 1:25.
[0046] In another embodiment, the present invention provides
composition, comprising (GeH.sub.3).sub.2CH.sub.2 and
Ge.sub.2H.sub.6 in a ratio of between about 1:20 and 1:30 (or, in
other embodiments, 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;
and 1:29 to 1:30.
EXAMPLES
Example 1
General Deposition Procedures
Example 1a
GeSn Buffer Deposition
[0047] Ge.sub.1-ySn.sub.y buffer layers (y=0.02-0.04) were
deposited on hydrogen-passivated Si(100) wafers at 330-350.degree.
C., as described previously (see, Bauer et al., Appl. Phys. Lett.
81, 2992 (2002), which is hereby incorporated by reference in its
entirety). The as-grown Ge.sub.1-ySn.sub.y films were .about.93-95%
relaxed (even for thicknesses less than 100 nm) and achieve strain
relief from the substrate by generating Lomer dislocations that run
parallel to the film/substrate interface. The residual strain was
relieved by in situ thermal cycling for 30 minutes at
500-600.degree. C. or by rapid thermal annealing up to 850.degree.
C. for several seconds, depending on composition. These steps also
reduces the density of threading defects penetrating to the surface
to levels below 10.sup.6/cm.sup.2. Atomic force microscopy (AFM)
shows planar surfaces for both the as-grown and annealed
Ge.sub.1-ySn.sub.y buffers that provide an ideal platform for
subsequent growth. The typical RMS roughness was in the range of
0.5-0.8 nm for 10.times.10 .mu.m.sup.2 areas.
Example 1b
Tensile Strained Ge Deposition
[0048] Growth of Ge epilayers was conducted ex situ on the relaxed
Ge.sub.1-ySn.sub.y buffers (y=0.02-0.04) via gas-source molecular
beam epitaxy (MBE) at 340-380.degree. C. and 5.times.10.sup.-5 Torr
using 1:15 admixtures of (GeH.sub.3).sub.2CH.sub.2 and
Ge.sub.2H.sub.6. This combination of compounds was designed to
provide built-in pseudo surfactant growth behavior enabling the
fabrication of dislocation free, and atomically flat Ge films with
no measurable carbon incorporation. SIMS measurements indicate C
content at the detection limit (<3.times.10.sup.17 cm.sup.-3)
(see, Wistey et al., Appl. Phys. Lett., 90, 082108 (2007)).
[0049] The reaction mixture of (GeH.sub.3).sub.2CH.sub.2 in
Ge.sub.2H.sub.6 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.). The flask
was connected to a gas injection manifold which was pumped to
.about.10.sup.-8 Torr on the gas source MBE chamber.
[0050] Prior to Ge growth, the Ge.sub.1-ySn.sub.y/Si(100)
substrates were sonicated for 5 minutes in methanol, dried by
flowing N.sub.2 over their surface, inserted through a load lock
into an ultra high vacuum (UHV) chamber at a base pressure of
5.times.10.sup.-10 Torr, and then heated on the sample holder for
one hour at 250.degree. C. to desorb any volatile surface
contaminants until the chamber pressure was restored to the base
value of 10.sup.-10 Torr. Under these conditions we find that the
Ge.sub.1-ySn.sub.y surfaces exhibit the typical (2.times.1) to
(1.times.2) reconstruction, indicating a well-ordered
crystallographic state suitable for subsequent heteroepitaxial
growth.
[0051] The temperature was then increased to 360-380.degree. C. and
the reactant gases were admitted at a final pressure of
5.times.10.sup.-5 Torr to commence film growth. The pressure was
maintained constant (5.times.10.sup.-5 Torr) during growth via
dynamic pumping using a corrosion resistant turbomolecular pump.
Typical growth times from 30-60 minutes yielded films of thickness
in the range 30-60 nm, respectively, at a fixed growth temperature
of 360.degree. C. The Ge films were deposited at a rate of 2
nm/min. Under these conditions Ge growth was observed to proceed
via nucleation of nanoscale atomically-flat mesas which gradually
coalesce to produce continuous films with planar surfaces as
evidence by AFM characterizations.
[0052] At growth temperatures of 380.degree. C. and higher, a
significant increase in growth rate and a surface morphology
consisting of two-dimensional tiling formations based on
rectangular mesas with variable shape and size were observed. On
the length scale of the mesas (.about.1 .mu.m) the film surface is
atomically flat; however, on the scale of 2-5 .mu.m the surface
roughness is higher due the presence of vertical steps between the
mesas.
Example 2
Structural and Optical Characterization
[0053] The samples prepared according to Example 1 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).
[0054] As detailed below, the precise strain state of the Ge
epilayers case can be systematically manipulated by varying the
thickness and composition of the underlying template, for example,
via tuning of the Sn content in the Ge.sub.1-ySn.sub.y buffer.
Growth of Ge layers on buffers with smaller/larger lattice
constants, such as Ge.sub.1-ySn.sub.y with y=0.015-0.035,)
systematically produced larger strains in the Ge overlayers with
increasing y, within the 1.5-3.5% range.
[0055] We have used the Ge.sub.1-ySn.sub.y buffer layers
(y=0.015-0.035) to demonstrate that the highest reported tensile
strains on Ge can be easily reproduced within this composition
range, although record strains as high as 0.45% have already been
achieved by our method. Advantageously, these buffer compositions
have been shown to possess high thermal stability (up to
800.degree. C.) and to be compatible with conventional CMOS
processing.
[0056] Combined Raman analysis and high-resolution x-ray
diffraction using multiple off-axis reflections revealed
unequivocally that the symmetry of tensile Ge was perfectly
tetragonal. A downshift of the direct gap consistent with tensile
strain has been observed.
[0057] The degree of tensile strain alters the Ge band structure
and induces a tunable redshift of the direct gap E.sub.0 absorption
edge FIG. 1 shows the quantitative changes in the band structure
for 1% tensile strain in Ge. Small reductions in E.sub.0 can
dramatically improve the detection performance of this material at
1.55 .mu.m and beyond to cover the technologically important
near-IR telecom bands for use in optical fiber communications. For
sufficiently large tensile strain it has been predicted that Ge may
become a direct gap system, with potential applications as an
active lasing material (see, Menendez and Kouvetakis, J. Appl.
Phys. Lett. 2004, 85, 1175).
Example 2a
Surface Characterization
[0058] In XTEM studies of the materials prepared according to
Example 1 revealed monocrystalline Ge epilayer films possessing
commensurate interfaces and atomically smooth surfaces regardless
of the epilayer thickness. Electron micrographs demonstrated
heteroepitaxial growth of a 60 nm thick Ge film on a 200 nm thick
Ge.sub.1-ySn.sub.y buffer with .epsilon..sub..parallel.=0.25% (see,
FIGS. 2 & 3). The diffraction contrast image of the entire film
thickness in (110) projection (FIG. 3) shows an atomically flat
surface and a film devoid of threading dislocations within the
field of view.
[0059] Surface flatness was confirmed on a larger 10.times.10
.mu.m.sup.2 scale by AFM scans, which show RMS roughness values of
0.5 nm. A corresponding high resolution image showed a sharp
interface between the two materials that is fully commensurate and
defect-free. Occasional threading defects are observed to
propagate, however, from the Si interface through the buffer and
terminate at the Ge/GeSn interface, suggesting that deflection and
bending of dislocations into the plane of this interface may be
taking place.
[0060] The typical RBS spectra (in random and channeled modes)
showed consistently monocrystalline and perfectly aligned
materials. The ratio of the aligned and random spectra
(.chi..sub.min), which measures the degree of epitaxial registry
was extremely low, an .chi..sub.min value of 10% within the Ge
layer, close to the 3% limit for Si (100) wafers.
Example 2b
Raman Analysis
[0061] Combined Raman analysis and high-resolution x-ray
diffraction using multiple off-axis reflections reveal
unequivocally that the symmetry of tensile Ge epilayers, with
varying thickness, prepared according to Example 1 are perfectly
tetragonal, while the strain state of the buffer remains
essentially unchanged. A downshift of the direct gap consistent
with tensile strain has been observed.
[0062] The growth strategies herein provide controlled and
reproducible Ge strain values exceeding those obtained to date,
leading to values of tensile strain as high as 0.8%. The strain in
this case can be systematically manipulated by varying the
thickness and composition of the underlying template.
[0063] FIG. 4 shows a Raman spectrum from a tensile-strained
Ge-film in a Ge/Ge.sub.0.975Sn.sub.0.025 structure. The Raman peak
is clearly downshifted relative to the corresponding Raman peak in
bulk Ge. Tetragonal strain splits the three-fold degenerate optical
phonons at the center of the Ge Brillouin zone into a doublet and a
singlet. For the (001) backscattering configuration used here, only
the singlet is observable, and its shift relative to bulk Ge is
given by .DELTA..omega.=b.epsilon..sub.p, with
b=[q-p(C.sub.12/C.sub.11)]/.omega..sub.0. Here p and q are the
optical phonon anharmonic parameters, as defined, for example, in
Anastassakis and Cardona in High Pressure in Semiconductor Physics
II, ed. by T. Suski and W. Paul (Academic Press, New York, 1998)),
and .omega..sub.0 is the Raman frequency in bulk Ge. Using values
from Anastassakis (supra), we obtain b=-415.+-.40 cm.sup.-1. Thus
from the observed Raman shift .DELTA..omega.=-1.0.+-.0.1 cm.sup.-1,
we deduce .epsilon..sub..parallel.=-0.24.+-.0.04%, in excellent
agreement with the X-ray analysis. Moreover, initial
photoreflectance experiments on this sample confirm the downshift
of the direct gap induced by the tensile strain.
[0064] FIG. 5 shows a typical Raman spectrum of a
Ge/Ge.sub.0.965Sn.sub.0.035 sample compared to that of a bulk Ge
reference. Again, the peak of the Ge film is shifted to lower
energies, as expected for tensile strain. The results for this and
similar samples are summarized in FIG. 6, where we plot the Raman
shifts of the Ge-films with strains intermediate to 0.15 and 0.43%
(with respect to the Raman shift in pure Ge) as a function of the
measured in-plane strain .epsilon..sub..parallel. obtained from XRD
(infra). A monotonic and approximately linear dependence was
observed, which confirms the XRD results. A fit with Eq. (1) yields
b.sub.G e=436.+-.18 cm.sup.-1. This overlaps considerably with the
earlier results of Cerdeira et al. (see, Phys. Rev. B 1972, 5,
580). Thus we conclude that, as in the case of Si, the original
strain shift coefficients determined in the 70's are sufficiently
accurate to characterize strain using Raman spectroscopy. Finally
we note that preliminary photoreflectance studies of selected
samples have shown a downshift of the direct gap consistent with
tensile strain.
Example 2c
XRD Analysis
[0065] The strain state of the Ge films was further investigated by
recording line scans and reciprocal space maps for the symmetric
(004), and asymmetric (224) Bragg reflections. In each case the
data were referenced to the corresponding reflections of the Si
wafer.
[0066] The data were acquired with a PANalytical.TM. X'Pert MRD
system (PANalytical B.V., Almelo, Netherlands) and were referenced
for each sample to the corresponding reflections of the Si wafer.
FIG. 7 shows the (224) reciprocal space map for a
Ge/Ge.sub.0.975Sn.sub.0.025 structure with epilayer and buffer
thicknesses of 60 nm and 200 nm, respectively.
[0067] For the Ge.sub.0.975Sn.sub.0.025 layer, the diffracted
intensity peaks lie very close to the full relaxation line ((224)
line in FIG. 7). A numerical fit yields
.alpha..sub..parallel.=5.6742 .ANG. and
.alpha..sub..perp.,Ge=5.6794 .ANG., which corresponds to a
compressive strain of -0.05%. The vertical alignment of the Ge and
GeSn (224) peaks indicates a coherent heterostructure, while the
considerable offset of the Ge peak with respect to the relaxation
line implies a significant tensile strain. The measured in-plane
(.alpha..sub..parallel.,Ge=5.6719 .ANG.) and vertical
(.alpha..sub..perp.,Ge=5.6476 .ANG.) lattice parameters are related
by the tetragonal strain relation
.epsilon..sub..perp.=2C.sub.12/C.sub.11.epsilon..sub..parallel.,
where C.sub.11 and C.sub.12 are bulk Ge elastic constants, and the
strain is defined as
.epsilon..sub..parallel.=(.pi..sub..parallel.-.alpha..sub.0)/.alpha..sub.-
0 and
.epsilon..sub..perp.=(.alpha..sub..perp.-.alpha..sub.0)/.alpha..sub.-
0. By averaging published experimental values for the elastic
constants (see, O. Madelung, in Semiconductors, Landolt Borstein
New Series III (Springer-Verlag, Berlin, N.Y., 2001), Vol. 41-A),
we obtain 2C.sub.12/C.sub.11=0.72.+-.0.04, from which we calculate
.alpha..sub.0=5.6578 .ANG.. This is in excellent agreement with
measurements in relaxed Ge films and with published values for the
lattice constant of bulk Ge (see, Madelung, supra). Using the
calculated relaxed lattice constant, we estimate for our sample a
strain value .epsilon..sub..parallel.=-0.25.+-.0.02%, where the
error arises from the scattering of the elastic constant data and
the uncertainty in the determination of the maxima in the
reciprocal space maps.
[0068] In previous work we showed that Ge.sub.1-ySn.sub.y alloys
display a compliant behavior when used as buffer layers for the
growth of lattice-mismatched Ge.sub.1-x-ySi.sub.xSn.sub.y alloys
(see, Tolle, et al., Appl. Phys. Lett. 88, 252112 (2006)). The
lattice constant .alpha..sub..parallel. for the buffer and epilayer
was found to approach the value predicted from a simple strain
energy minimization expression that neglects the interaction with
the Si substrate. For the sample in FIG. 7 such a model predicts a
compressive buffer strain .epsilon..sub..parallel.=-0.08, in
reasonably agreement with the observed value. To exp incorporate
the same buffer layer but a much thicker Ge film (140 nm) than the
previously discussed sample (60 nm). Then the compressive strain on
the buffer layer should increase to
.epsilon..sub..parallel.=-0.14.
[0069] XRD data including the (224) and (004) reflections indicated
that the buffers remain virtually unchanged with
.alpha..sub..parallel.,GeSn=5.6731 .ANG. and
.alpha..sub..perp.,GeSn=5.6794 .ANG., which correspond to
.epsilon..sub..parallel.=-0.06. The lattice constants
.alpha..sub..parallel.,GeSn=5.6704 .ANG. and
.alpha..sub..parallel.,GeSn=5.6731 .ANG. are again very close,
indicating the same degree of coherency as in the previous
structure. The calculated unstrained Ge lattice constant
(.alpha..sub.0=5.6576 .ANG.) is virtually identical to the one
obtained for the previous sample. The strain analysis yields
.epsilon..sub..parallel..about.0.23%, which is nearly the same as
that of the 60 nm thick Ge sample. This value is within the 0.02%
error shown above. Therefore, these experiments do not show
evidence for a compliant buffer behavior. It is interesting to
point out, however, that in the present case the strain energy
densities are smaller than in Tolle et al. (supra) by up to one
order of magnitude, so that neglecting the buffer-substrate
interaction may no longer be a good approximation.
[0070] The larger scattering volume of the 140 nm thick sample
enabled us to measure typically weak off-axis peaks such as the
(135) and (135) reflections. These have h<k and accordingly they
are sensitive to deviations from a tetragonal distortion. If the Ge
film distortion is of lower symmetry than tetragonal, the in-plane
value derived from these reflections will deviate significantly
from their (224) counterpart. Table 1 compares both the in-plane
and vertical lattice constants (.alpha..sub..parallel. and
.alpha..sub..perp. respectively), for the Si substrate, the GeSn
buffer layer and the strained Ge epilayer. In all three cases the
in-plane lattice constant values obtained independently from the
(135), (135) and (224) analysis are identical, within instrumental
resolution.
TABLE-US-00001 TABLE 1 In-plane and perpendicular lattice constants
for the Si substrate, Ge.sub.0.975Sn.sub.0.025 buffer layer and
strained Ge epilayer as determined from independent x-ray
reflections (004), (224), (135) and (135) Reflection a (.ANG.) c
(.ANG.) Si (004) 5.4306 (224) 5.4305 5.4305 (135) 5.4306 5.4307 (
135) 5.4305 5.4305 5.4305 5.4306 GeSn (004) 5.6801 (224) 5.6731
5.6794 (135) 5.6735 5.6798 ( 135) 5.6730 5.6800 5.6702 5.6798 Ge
(004) 5.6484 (224) 5.6703 5.6482 (135) 5.6704 5.6485 ( 135) 5.6702
5.6484 5.6703 5.6484
A similar independent analysis of the (004), (135), (135) and (224)
peaks also yields vertical lattice constants that are virtually
identical within experimental error. Collectively these data
demonstrate that the Ge epilayer displays a perfect tetragonal
distortion, verifying the underlying assumption in our strain
analysis above. In addition, Table 1 shows that Si is perfectly
cubic as expected, and that the buffer is only slightly compressed,
as shown in FIG. 8, where the diffraction maps for the (135) peaks
of the Ge epilayer and GeSn buffer are shown. It should be noted
that for samples grown on unrelaxed GeSn buffers we consistently
observe discrepancies in the .alpha..sub..parallel. parameter
derived from the (224), (135) and (135) reflections, demonstrating
a non-tetragonal deformation. This indicates that residual
compressive strains in the as grown buffers lead to tensile
strained Ge layers possessing complicated lower symmetries
associated with monoclinic or orthorhombic distortions. This
emphasizes the crucial need to relax the buffers prior to any
subsequent depositions. Interestingly, we find that a combination
of X-ray diffraction and Raman spectroscopy provided (see, Example
2b) a quick way to identify non-tetragonal distortions. If we use
the (224) reciprocal space maps and assume a tetragonal distortion
to deduce the strain, the value obtained from the X-ray analysis
and from Raman spectroscopy agree within experimental error for
perfectly tetragonal samples, but differ by factors as high as two
if the sample distortion is of lower symmetry than tetragonal.
[0071] FIG. 9 shows the (224) spectra of a representative sample
with composition Ge/Ge.sub.0.965Sn.sub.0.035, and epilayer/buffer
thicknesses of .about.60/230 nm. The thicker
Ge.sub.0.965Sn.sub.0.035 layer was also found to be essentially
relaxed and nearly cubic (.alpha..sub..parallel.,GeSn=5.6872 .ANG.,
.alpha..sub..perp.,GeSn=5.6955 .ANG.). This was further confirmed
by the close proximity of the GeSn peak maximum to the relaxation
line connecting the corresponding Si (224) peak to the origin of
the plot. The plots also show a near perfect vertical alignment of
the Ge and GeSn (224) peaks indicating a fully coherent stack as
shown by the dotted line in FIG. 9. Note the considerable vertical
offset of the Ge peak with respect to the relaxation line,
indicating a significant tensile strain which is manifested as a
tetragonal distortion that accommodates the large lattice mismatch
with the buffer. The ideal strain equilibration behavior of the
Ge/Ge.sub.0.965Sn.sub.0.035 sample is plotted in FIG. 10 which
shows that most of the strain energy minimization is accounted for
by the expansion of the Ge epilayer in the a-b plane. This
indicates that for thick (>200 nm) buffers very little change in
the dimensions of the buffer is expected.
[0072] To derive the precise values of the in-plane
(.epsilon..sub..parallel.) and perpendicular (.epsilon..sub..perp.)
strain values we conducted a quantitative analysis of the XRD data
assuming that the epitaxial Ge films adopt a purely tetragonal
distortion, relative to their unstrained cubic form, in which
.epsilon..sub..parallel.=(a.sub..parallel.-a.sub.0)/a.sub.0 and
.epsilon..sub..perp.=(a.sub..perp.-a.sub.0)/a.sub.0 (where
a.sub..parallel., a.sub..perp. and a.sub.0 are the measured
in-plane, vertical and relaxed lattice constants). The two strains
are related by
.epsilon..sub..perp.=2C.sub.12/C.sub.11.epsilon..sub..parallel.,
which we can write
.epsilon..sub..perp.=-.xi..epsilon..sub..parallel., with
.xi.=2C.sub.12/C.sub.11 (where C.sub.12 and C.sub.11 are the
elastic constants). The strain relation can be inverted to obtain
the unstrained lattice constant
a.sub.0=(a.sub..perp.+.xi.a.sub..parallel.)/(1+.xi.), the
.epsilon..sub..parallel.=(a.sub..parallel.-a.sub..perp.)/(a.sub..perp-
.+.xi.a.sub..parallel.) and the
.epsilon..sub..perp.=.xi.(a.sub..perp.-a.sub..parallel.)/(a.sub..perp.+.x-
i.a.sub..parallel.) of the film. For the elastic constants we adopt
the values C.sub.11=128.5 GPa, C.sub.12=48.3 GPa from Madelung
(supra). while for strained Ge film we use
a.sub..parallel.,Ge=5.6802 .ANG. and a.sub..perp.,Ge=5.6416 .ANG.
obtained from XRD. These data yield parallel and perpendicular
strain values of .epsilon..sub..parallel..about.+0.40% and
.epsilon..sub..perp..about.-0.30%, and an unstrained Ge lattice
constant of a.sub.0=5.658 .ANG., in excellent agreement with the
known bulk Ge value. Taking differentials of the strain relations
above yields the following expression:
.delta. = ( 1 1 + a .perp. / .xi. a ) .delta..xi. .xi. .apprxeq. 1
2 .delta..xi. .xi. ( 2 ) ##EQU00001##
Thus a fractional error in the elastic ratio
.xi.=2C.sub.12/C.sub.11 produces half this error in the in-plane
strain. For instance, if C.sub.11=129.+-.3 GPa and C.sub.12=48.+-.3
GPa, (see, Madelung, supra), then .xi.=0.75.+-.0.07 (.+-.9.3%)
which implies a corresponding fractional error of only .about.5% in
the in-plane strain. So, for the Ge/Ge.sub.0.965Sn.sub.0.035 we
obtain .epsilon..sub..parallel..about.0.40.+-.0.03%. The
corresponding strain value obtained by Raman (see below) is
.epsilon..sub..parallel..about.0.43 well within the 0.03% error.
This value represents the highest tensile strain observed in
elemental Ge semiconductors to date.
[0073] Extensive XRD studies of a wide range of tensile strained Ge
films grown on Ge.sub.1-ySn.sub.y (y=0.015-0.035) indicated strain
states between 0.15-0.43%, respectively. The XTEM data of these
samples revealed monocrystalline films possessing commensurate
interfaces and atomically smooth surfaces regardless of the
epilayer thickness. FIGS. 11 and 12 shows electron micrographs of
the Ge/Ge.sub.0.965Sn.sub.0.035 sample indicating a 60 nm thick Ge
layer grown epitaxially on a 230 nm thick Ge.sub.1-ySn.sub.y
buffer. The bright field image shows essentially no penetrating
threading defects over an extended lateral range of 1 .mu.m and the
high resolution micrograph near the Ge/Si interface reveals a
sharp, commensurate transition between the epilayer and the buffer,
AFM analysis of the sample shows that the surface is predominantly
flat, with an RMS roughness of .about.1.5 nm over a typical
5.times.5 .mu.m.sup.2 area [FIG. 13]. On a smaller 0.5.times.0.5
.mu.m.sup.2 scale we observe atomically smooth surfaces with a
reduced RMS roughness of .about.0.2 nm. The difference in these
values is associated with the presence of tiny voids located at the
intersection of individual coalescing terraces as described
previously. In thicker films these terraces coalesce fully and such
features are not longer visible. In contrast, the growth of pure
Ge.sub.2H.sub.6 in the absence of the organic additive
(GeH.sub.3).sub.2CH.sub.2 produces defective and rough films with
classic island-like surface morphology consistent with a mismatched
heteroepitaxy growth mode. This observation confirms the notion
that the organic additives significantly alter the surface
energetics, promoting organized assembly of planar Ge films at
conditions that would normally preclude their growth directly on
GeSn on the basis of surface energy differences. Note that
analogous differences in surface energetics prevent the
heteroepitaxial growth of silicon films directly on a Ge surface.
However, as we have demonstrated above, the pseudosurfactant
properties of the organic additives enable the remarkable and
unconventional growth of fully strained and perfectly planar Si
layers on Ge buffers.
[0074] In one example of Ge epilayer strain tuning via Sn content
in the buffer, we conducted growth of Ge layers on buffers with
smaller/larger lattice constants such as Ge.sub.1-ySn.sub.y with
y=0.015-0.035. For samples with the similar Ge/GeSn film thickness
ratio we find .epsilon..sub..parallel.=0.17% for y=0.015 and
.epsilon..sub..parallel.=0.45% for y=0.035 (see, FIG. 3).
Increasing y (within the 1.5-3.5% range) produces systematically
larger strains in the Ge overlayers.
[0075] In summary, we have demonstrated the growth of tensile
strained Ge films on Ge.sub.1-ySn.sub.y/Si composite substrates.
Our approach is straightforward and suitable for large-scale
integration. Contrary to the thermal expansion processes, our
growth proceeds at low temperatures (.about.380.degree. C.)
compatible with selective growth.
* * * * *