Next generation mid-wave infrared cascaded light emitting diodes: growth of broadband, multispectral, and single color devices on GaAs and integrated circuitsTheses and Dissertations
Next generation mid-wave infrared cascaded light emitting Next generation mid-wave infrared cascaded light emitting
diodes: growth of broadband, multispectral, and single color diodes: growth of broadband, multispectral, and single color
devices on GaAs and integrated circuits devices on GaAs and integrated circuits
Sydney R. Provence University of Iowa
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Copyright © 2016 Sydney R, Provence
This dissertation is available at Iowa Research Online: https://ir.uiowa.edu/etd/5827
Recommended Citation Recommended Citation Provence, Sydney R.. "Next generation mid-wave infrared cascaded light emitting diodes: growth of broadband, multispectral, and single color devices on GaAs and integrated circuits." PhD (Doctor of Philosophy) thesis, University of Iowa, 2016. https://doi.org/10.17077/etd.xii1sjrk
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COLOR DEVICES ON GAAS AND INTEGRATED CIRCUITS
by
Sydney R. Provence
A thesis submitted in partial fulfillment of the requirements for the Doctor of Philosophy
degree in Physics in the Graduate College of The
University of Iowa
Thesis Supervisors: Professor John P. Prineas Professor Thomas F. Boggess
Approved for Public Release - Distribution is unlimited.
Graduate College The University of Iowa
Iowa City, Iowa
CERTIFICATE OF APPROVAL
Sydney R. Provence
has been approved by the Examining Committee for the thesis requirement for the Doctor of Philosophy degree in Physics at the August 2016 graduation.
Thesis Committee: John P. Prineas, Thesis Supervisor
Thomas F. Boggess, Thesis Supervisor
Markus Wohlgenannt
ACKNOWLEDGMENTS
I would like to thank my advisors, Tom Boggess and John Prineas, for their
guidance and support during my graduate career. I very much appreciate Tom’s
encouragement at trying my hand at many different aspects of research and John’s
guidance in the art of molecular beam epitaxy. Their collaborations have allowed me
to pursue a vast number of skills and topics, for which I am truly grateful.
I have had the pleasure of working with quite a few graduate students who have
influenced this thesis. My initial training in molecular beam epitaxy was largely done
by Lee Murray and Asli Yildirim. Dennis Norton initially trained me in processing
and electrical characterization, the bulk of which was largely undertaken by Russell
Ricker for most of this thesis. I would also like to thank Kailing Zhang and Aaron
Muhowski for their help in the molecular beam epitaxy lab, and Yigit Aytac for his
collaborations. The training of Holly Morris in the Chemistry Department on the
atomic force microscope was also extremely helpful.
I have gained both practical knowledge and essential technical support in keeping
the molecular beam epitaxy lab and characterization equipment running from the
Mikes: Mike Thornburg, Mike Miller, and Mike Fountain. Without their help,
absolutely nothing in this department would work.
Brennan Van Alderwerelt and our cats, Squash and Zucchini, have been an
endless source of love and support. Gratitude is also due to Steve & Lisa, who have
actively discouraged very few of my pursuits.
Lastly, I would like to thank John’s Grocery, late night purveyor of decent-enough
wines and beer, and Wake Up Iowa (City), an outfit with admirable dedication to
fantastic light roasts. The contributions of both companies to this dissertation have
been nothing short of heroic.
ii
ABSTRACT
InAs/GaSb superlattices are an attractive material system for infrared light
emitting diodes, due to the ability to tune the band gap throughout most of the
infrared regime. A key consideration in the epitaxial growth of these heterostructures
is crystalline material quality. In developing thick layers of epitaxially grown mate-
rial, there are moderate amounts of elastic strain that can be incorporated into a
heterostructure, beyond which deformations will form that will alleviate the lattice
mismatch. These deformations have the potential to damage or ruin the optical and
electrical quality of an optoelectronic device. This thesis investigates the optical
and electronic properties of lattice-mismatched and strained materials through the
study of thick dual-color light emitting diodes, broadband light emitting diodes, and
InAs/GaSb superlattice devices developed on GaAs substrates and GaAs integrated
circuits.
A dual-color infrared light emitting diode is demonstrated emitting in two
separate midwave infrared bands. The design of the device stacks two independently
operable InAs/GaSb superlattices structures on top of one another, so that 10 µm of
material is grown with molecular beam epitaxy. Each layer is lattice-matched to a
GaSb substrate. At quasi-continuous operation, radiances of 5.48 W/cm2-sr and 2.67
W/cm2-sr are obtained.
A broadband light emitting diode spanning the mid-wave infrared is demonstrated
with eight stages of InAs/GaSb superlattices individually tuned to a different color.
The performance of the device is compared with an identical eight stage device emitting
in the middle of the mid-wave infrared. The emission of the fabricated broadband
device spans from 3.2 µm to 6 µm with peak radiance of 137.1 mW/cm2-sr.
Growth of antimonide-based devices on GaAs is desirable to the relative trans-
parency of semi-insulating substrates throughout the infrared, and as semi-insulating
GaSb substrates are not available. The growth of bulk GaSb on GaAs is explored
iii
through different techniques in order to confine relaxation due to lattice mismatch
strain to the GaSb/GaAs interface. A low temperature nucleation technique with a
thin GaSb wetting layer is found to have the best overall surface morphology, although
screw dislocations are a prominent feature on all samples. The dislocations and overall
surface roughness are not found to destructively impact the overall device quality,
as four stage InAs/GaSb superlattice devices grown on GaAs substrates are found
to have superior radiance emission and external quantum efficiency compared to an
identical device grown on a GaSb substrate due to the higher substrate transparency
and superior thermal properties.
Epitaxy on electronics growth techniques on GaAs integrated circuits are devel-
oped to bypass the hybridization process in light emitting diode development. Chips
obtained from Quorvo, Inc. are found to endure ultra-high vacuum molecular beam
epitaxy environment at higher temperatures with silicon nitride encapsulation, and a
low temperature oxide removal technique is developed using an atomic hydrogen source.
Chemical-mechanical polishing techniques are developed to create an epiready sub-
strate surface. Ultimately, no photoluminescent emission is observed from InAs/GaSb
superlattices grown on these GaAs integrated circuits, although electroluminescent
emission is still possible.
InAs/GaSb superlattices are an attractive material system for infrared light
emitting diodes, due to the ability to tune the band gap throughout most of the
infrared regime. A key consideration in the epitaxial growth of these heterostructures
is crystalline material quality. In developing thick layers of epitaxially grown mate-
rial, there are moderate amounts of elastic strain that can be incorporated into a
heterostructure, beyond which deformations will form that will alleviate the lattice
mismatch. These deformations have the potential to damage or ruin the optical and
electrical quality of an optoelectronic device. This thesis investigates the optical
and electronic properties of lattice-mismatched and strained materials through the
study of thick dual-color light emitting diodes, broadband light emitting diodes, and
InAs/GaSb superlattice devices developed on GaAs substrates and GaAs integrated
circuits.
v
1 INTRODUCTION . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1
1.1 Semiconductor Background . . . . . . . . . . . . . . . . . . . . . 2 1.1.1 Band Structure . . . . . . . . . . . . . . . . . . . . . . . . 2 1.1.2 Band Alignment . . . . . . . . . . . . . . . . . . . . . . . 4 1.1.3 Electrons and Holes . . . . . . . . . . . . . . . . . . . . . 5 1.1.4 Recombination Processes . . . . . . . . . . . . . . . . . . 7
1.2 Light Emitting Diodes . . . . . . . . . . . . . . . . . . . . . . . . 10 1.2.1 Efficiency . . . . . . . . . . . . . . . . . . . . . . . . . . . 11
1.3 III-V Semiconductors for Infrared LEDs . . . . . . . . . . . . . . 13 1.3.1 III-V Compound Semiconductors . . . . . . . . . . . . . . 13 1.3.2 InAs/GaSb Superlattices . . . . . . . . . . . . . . . . . . 14 1.3.3 Cascading InAs/GaSb Superlattices . . . . . . . . . . . . 17
1.4 Applications for InAs/GaSb Superlattices . . . . . . . . . . . . . 18
2 MOLECULAR BEAM EPITAXY . . . . . . . . . . . . . . . . . . . . 20
2.1 Epitaxial Growth Modes . . . . . . . . . . . . . . . . . . . . . . 21 2.2 Strained-Layer Epitaxy . . . . . . . . . . . . . . . . . . . . . . . 25 2.3 Crystalline Defects . . . . . . . . . . . . . . . . . . . . . . . . . . 27 2.4 Molecular Beam Epitaxy Growth Equipment . . . . . . . . . . . 29
2.4.1 Vacuum and Ion Gauges . . . . . . . . . . . . . . . . . . 31 2.4.2 Effusion Cells . . . . . . . . . . . . . . . . . . . . . . . . 34 2.4.3 Dopants . . . . . . . . . . . . . . . . . . . . . . . . . . . 36 2.4.4 Desorption . . . . . . . . . . . . . . . . . . . . . . . . . . 38
2.5 In Situ Growth Monitoring Techniques . . . . . . . . . . . . . . . 41 2.5.1 Reflection High-Energy Electron Diffraction . . . . . . . . 42 2.5.2 Temperature . . . . . . . . . . . . . . . . . . . . . . . . . 45
3 TWO-COLOR SUPERLATTICE LEDS . . . . . . . . . . . . . . . . . 51 3.1 Background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 51
3.1.1 Dual-Color Systems . . . . . . . . . . . . . . . . . . . . . 53 3.2 TCSA Design and Molecular Beam Epitaxial Growth . . . . . . 54 3.3 Processing and Results . . . . . . . . . . . . . . . . . . . . . . . 62 3.4 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 64
4 BROADBAND INFRARED LIGHT EMITTING DIODES . . . . . . 67 4.1 Motivation and Methods of Creating Broad Spectrum LEDs . . . 67 4.2 Molecular Beam Epitaxial Growth . . . . . . . . . . . . . . . . . 68 4.3 Processing Results and Discussion . . . . . . . . . . . . . . . . . 71
vi
5 INAS/GASB SUPERLATTICE LIGHT EMITTING DIODES GROWN ON GAAS SUBSTRATES . . . . . . . . . . . . . . . . . . . . . . . . 77
5.1 Advantages of Growth of Antimonide-Based Devices on GaAs Substrates . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77
5.2 Methods Developed for Growth of GaSb on GaAs . . . . . . . . 79 5.2.1 Interfacial Misfit Technique . . . . . . . . . . . . . . . . . 80 5.2.2 Low Temperature Nucleation . . . . . . . . . . . . . . . . 81
5.3 Antimonide-Based Devices Realized on GaAs and GaAs Integrated Circuits . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 82
5.4 GaSb Buffer Optimization on GaAs . . . . . . . . . . . . . . . . 84 5.5 InAs/GaSb Superlattice LED Molecular Beam Epitaxial Growth
and Fabrication on GaAs and GaSb . . . . . . . . . . . . . . . . 88 5.6 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 90 5.7 Superlattice LED Growth on GaAs Integrated Circuits . . . . . . 99
5.7.1 GaAs Integrated Circuit Substrate Characterization . . . 100 5.7.2 Heat Testing of GaAs Chip Circuitry . . . . . . . . . . . 101 5.7.3 GIC Polishing . . . . . . . . . . . . . . . . . . . . . . . . 105 5.7.4 Results of Growth on GaAs Integrated Circuits . . . . . . 106 5.7.5 Discussion and Future Work . . . . . . . . . . . . . . . . 110
REFERENCES . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 113
1.1 Common compound semiconductor elements in the Periodic Table. . 13
2.1 Summary of InSb melting point on various substrates. . . . . . . . . . 49
4.1 Broad spectrum MWIR LED active region summary, with a fixed GaSb thickness of 16 ML. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69
5.1 Thermal properties of GaSb, GaAs, and Si at 300 K. . . . . . . . . . 78
5.2 Summary of GaSb buffer growths on GaAs. . . . . . . . . . . . . . . 85
5.3 ABC Coefficients determined from fit to recombination rate equation for InAs/GaSb superlattices grown on GaSb and GaAs. . . . . . . . . 93
5.4 RMS roughness of epiready GaAs and GICs, as measured by AFM. . 102
5.5 Percentage current output drop on transistor compared to pristine thinned GaAs chip after 5 hour heat test on thinned GaAs chips. . . 103
5.6 RMS roughness of polishing abrasives on thinned GICs, as measured by AFM over a 5×5 µm2 area. . . . . . . . . . . . . . . . . . . . . . . 107
5.7 Summary of RMS roughness of photoluminescent emission samples grown on unpolished, thinned GICs. . . . . . . . . . . . . . . . . . . . 108
viii
1.2 Simple model of a parabolic band structure. . . . . . . . . . . . . . . 5
1.3 Band alignment of semiconductor heterostructures. . . . . . . . . . . 6
1.4 Radiative and non-radiative recombination mechanisms. . . . . . . . . 8
1.5 Lattice constants and energy gaps of select III-V compound semicon- ductors at 300 K. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 14
1.6 A representation of energy levels in a quantum well, multiple quantum well, and superlattice. . . . . . . . . . . . . . . . . . . . . . . . . . . . 15
1.7 A schematic of a cascaded superlattice. . . . . . . . . . . . . . . . . . 17
2.1 Schematic diagrams of Frank-van der Merwe, Stranski-Krastanov, and Volmer-Weber growth modes. . . . . . . . . . . . . . . . . . . . . . . 22
2.2 Schematic diagrams of pseudomorphic and metamorphic epitaxial growth. 26
2.3 Schematic diagram of a Burgers circuit. . . . . . . . . . . . . . . . . . 28
2.4 Schematic diagram of an MBE growth chamber. . . . . . . . . . . . . 30
2.5 The mean free path of a Ga atom as a function of pressure at 300 K . 32
2.6 Schematic diagram of a Hall measurement. . . . . . . . . . . . . . . . 38
2.7 Schematic diagram of a RHEED setup in an MBE growth chamber. . 42
2.8 Diagram of RHEED in reciprocal space. . . . . . . . . . . . . . . . . 43
2.9 RHEED patterns from deoxidized GaAs and GaSb substrates . . . . 45
2.10 RHEED intensity oscillations taken during growth of GaSb. . . . . . 46
2.11 BandiT blackbody measurement of temperature uniformity across the surface of a 3-inch GaSb substrate. . . . . . . . . . . . . . . . . . . . 47
2.12 Sample temperature reconstruction plot for GaSb, GaAs, and GaAs integrated circuit. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50
3.1 Stack diagram of a dual-color LED. . . . . . . . . . . . . . . . . . . . 55
ix
3.2 Band diagram of a four stage dual-color LED . . . . . . . . . . . . . 57
3.3 Calculated SL band offsets for a fixed GaSb thickness of 16 ML and measured photoluminescence spectra peak wavelengths at 77 K. . . . 58
3.4 Normalized electroluminescence and photoluminescence for longer and shorter wavelength InAs/GaSb superlattices at 77K. . . . . . . . . . . 59
3.5 High-resolution x-ray diffraction scan of a dual-color LED. . . . . . . 61
3.6 Interferometric micrographs of the surface of a dual-color LED. . . . 62
3.7 A schematic diagram of a TCSA pixel. . . . . . . . . . . . . . . . . . 63
3.8 Voltage vs. current taken for a two-color LED at 77 K with a quasi- continuous input. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 64
3.9 Radiance vs. current graph of a two-color LED at 77 K. . . . . . . . 65
4.1 Stack diagrams of a broad spectrum LED and a single-color LED. . . 70
4.2 High-resolution x-ray diffraction scan of a broad spectrum LED. . . . 71
4.3 Electroluminescence spectra for broadband and single color devices at 77 K. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72
4.4 Normalized photon flux versus photon energy for broadband and single color devices at 77 K. . . . . . . . . . . . . . . . . . . . . . . . . . . . 73
4.5 Radiance vs. current density curves for broadband and single-color devices. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 74
4.6 Current vs. voltage curves for broadband and single-colored devices at 77 K. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 75
4.7 Calculated dynamic resistances for broadband and single-color LEDs. 75
4.8 Wallplug efficiency vs. current for broadband and single-color LEDs at 77 K. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 76
5.1 Schematic diagrams of spiral growth pattern formation around a screw dislocation . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 80
5.2 Stack diagrams of 3 µm GaSb buffer samples grown on GaAs substrates using an interfacial misfit technique, low temperature nucleation, and AlSb and GaSb wetting layers. . . . . . . . . . . . . . . . . . . . . . . 84
5.3 Interferometric micrographs of GaSb buffers grown on GaAs substrates 86
x
5.4 Atomic force microscopy scans of 3 µm GaSb buffer samples grown on GaAs substrates using low temperature nucleation . . . . . . . . . . . 87
5.5 Stack diagrams of four-stage SLEDs and GSLEDs devices . . . . . . . 89
5.6 Photoluminescence spectra of 8ML/16ML InAs/GaSb superlattice sam- ples grown on GaAs and GaSb substrates at 77 K . . . . . . . . . . . 91
5.7 Atomic force microscopy scans of SLEDS and GSLEDs devices . . . . 91
5.8 Carrier recombination rates as a function of excess carrier density measured from time-resolved differential transmission measurements at 77 K. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 92
5.9 Electroluminescence spectra for SLEDs devices grown on GaAs and GaSb substrates at 77 K for 100 × 100 µm2 mesas. . . . . . . . . . . 94
5.10 Radiance vs. current density for SLEDs and GSLEDs devices at 77 K for 100 × 100 µm2 mesas. . . . . . . . . . . . . . . . . . . . . . . . . 94
5.11 Measured current vs. voltage curves for SLEDS devices grown on GaAs and GaSb substrates. . . . . . . . . . . . . . . . . . . . . . . . . . . . 95
5.12 Calculated dynamic resistance for SLEDS and GSLEDs devices. . . . 95
5.13 Absorption coefficient measured through a light n-doped GaSb substrate and an unintentionally doped GaAs substrate. . . . . . . . . . . . . . 96
5.14 Per stage external and internal quantum efficiencies vs. current for GSLEDs and SLEDs devices. . . . . . . . . . . . . . . . . . . . . . . 98
5.15 Radiance vs. current density measurements for SLEDs and GSLEDs devices with 100 × 100 µm2 mesas at various duty cycles at 83 K. . . 99
5.16 AFM images of an epi-ready GaAs substrate, an unthinned GIC, and a thinned and mechanically polished GIC. . . . . . . . . . . . . . . . . 101
5.17 Microscope images of the surface of encapsulated GICs after heat testing.104
5.18 Illustration of an oxide desorption recipe using atomic hydrogen for GICs.105
5.19 AFM images of a 5×5 µm2 area of GICs polished using colloidal silica and alumina-silica slurry. . . . . . . . . . . . . . . . . . . . . . . . . . 107
5.20 Images of InAs/GaSb superlattices grown on unpolished GaAs inte- grated circuits at different scales, taken on both an interferometric microscope and with AFM. . . . . . . . . . . . . . . . . . . . . . . . . 109
xi
5.21 Interferometric micrographs of photoluminescence samples grown on GICs polished using alumina-silica and colloidal silica as grits. . . . . 110
5.22 Schematic diagram of a four-stage SLEDs device grown on a GIC. . . 111
xii
1
INTRODUCTION
One can generalize most interesting semiconductor devices as having two or
more different semiconducting materials arranged in a heterostructure, or an arrange-
ment in which the material composition changes with respect to position. An ideal
heterojunction would layer two materials together with identical crystalline structure,
including identical lattice spacing, thermal properties, optical transparency, etc. This
particular configuration, however, rarely, if ever, occurs. A more realistic scenario
requires a delicate balancing act at the interface of even similar materials that takes
into account an accumulation of strain. Without this maneuvering, a deformation of
layers due to strain energy may result.
Molecular beam epitaxy (MBE) is uniquely suited for developing semiconductor
heterostructures due to its ability to deposit extremely fine, high-quality layers with
abrupt interfaces. While metal-organic vapor deposition (MOCVD) can deposit
extremely thin atomic layers, the growth of antimonides is much more difficult than
that of other III-V materials such as arsenides, nitrides, or phosphides.
InAs/GaSb superlattices (SLs) are a particularly good example of a heterostruc-
ture that has been well-developed using MBE. A superlattice is defined as a periodic
arrangement of thin layers of two semiconducting materials in which the periodicity
occurs along one dimension. As described by Chang and Esaki, “the superlattice...can
be considered a new synthesized semiconductor not present in nature, which is ex-
pected to exhibit unusual electronic and optical properties [1].” The primary interest
in InAs/GaSb stems from the ability to tune the bandgap throughout most of the
infrared spectrum (3-30 µm) by varying the relative thicknesses of each material in
the system. The potential for band-engineered suppression of Auger recombination
rates is also attractive.
All electromagnetic signals experience some degree of attenuation when traveling
2
through the atmosphere, due to absorption, scattering, reflection, or diffusion. In
the infrared, absorption by atmospheric gases is largely predictable as a function of
wavelength. The infrared spectrum is often subclassified into separate bands, which
are largely dictated by the atmospheric absorption of light traveling through the
Earths atmosphere. Most of the the divisions in infrared bands can be attributed to
absorption by either water vapor or CO2 molecules [2]. These bands are commonly
referred to as the near (750 nm-1.4 µm), short-wavelength (SWIR) (1.4-2.7 µm),
mid-wavelength (MWIR) (3-5 µm), long-wavelength (LWIR) (8-12 µm), and very-long
wavelength (>12 µm) infrared bands (see Fig. 1.1).
Wavelength (m)
Wavelength ( m)
SW IR
MW IR
N E A R
Figure 1.1: The electromagnetic spectrum. The infrared band is further subdivided into near IR, SWIR, MWIR, LWIR, and very-long wavelength IR.
1.1 Semiconductor Background
1.1.1 Band Structure
The electronic properties of a semiconductor are best described through quantum
mechanics through solving the Schrodinger equation in its time-independent form:
[− h
2m0
3
In Equation 1.1, h is the Planck constant divided by 2π, m0 is the mass of the electron,
V (~r) is the potential energy, ψ(~r) is the wavefunction, and E represents the energy
levels of the system. In basic quantum mechanical systems, a classic example of which
is the hydrogen atom, the energy levels are discrete. In a semiconductor, the addition
of a relatively infinite number of atoms bound together in a lattice creates a drastically
more complicated picture. Here, atoms form a periodic potential dependent on the
lattice spacing, and discrete energy states for electrons are blurred into a range, or
band, of energy values that they may occupy. The states between these energy values
form the band gap of the structure.
For an electron in a crystalline structure with periodic potential, the potential
can be described by
V (~r) = V (~r + ~R), (1.2)
where ~R is a lattice translation vector. ~R describes any discrete translation operation.
In this hypothetical lattice, for any choice of ~R, the lattice will be identical in the
translation site. ~R can be defined as
~R = n1 ~a1 + n2 ~a2 + n3 ~a3, (1.3)
where n1, n2, and n3 are integers and ~a1, ~a2, and ~a3 are vectors that can be directed
to any other position in the lattice. This system also describes a Bravais lattice. III-V
semiconductors all have a zincblende crystalline structure, which forms a face-centered
cubic (fcc) Bravais lattice. The symmetric set of primitive vectors for the fcc lattice
are
4
The Bloch theorem states that any solution to the Schrodinger equation (Equa-
tion 1.1) in a periodic potential (Equation 1.2) can be described as a set of plane
waves modulated by a periodic function in accordance with the lattice periodicity. In
this case, all wavefunctions can be written as Bloch functions:
ψk(~r) = ei ~k·~ruk(~r). (1.5)
Here, uk(~r) must be a periodic function such that uk(~r) = uk(~r + ~R). Using uk(~r) in
Equation 1.1 can then used to obtain a dispersion relation, providing the bandstructure
of the system. Near k = 0, the E vs. k relation can be approximated as
E(k) = h2k2
2m∗ (1.6)
Typically, band structure references four primary bands, demonstrated in Fig. 1.2.
The conduction band is the “upper” band above the bandgap Eg, and represents
the lowest range of vacant electron energy states at 0 K, whereas the valence band
represents the highest range of filled electron energy states. Significant bands in the
valence band include the heavy hole, light hole, and split off bands. The parabolic
bandstructure schematic in Fig. 1.2 shows a direct bandgap semiconductor, or a
semiconductor in which the minimum point in the conduction band as well as the
maximum point of the valence band occur at k = 0. For an indirect bandgap
semiconductor, the bands are offset in k-space, and a momentum must be transferred
to support a carrier transfer across the bandgap [3,4].
1.1.2 Band Alignment
alignment. There are three basic heterojunction subsets (Fig. 1.3). Type-I alignment
5
k
E
EG
Figure 1.2: A simple model of a parabolic band structure demonstrating the conduction, heavy hole, light hole, and split off bands as a function of the electron wavevector.
has a heterojunction in which a material has a conduction band minimum lower than
the adjacent materials conduction band minimum, and a valence band maximum (hole)
energy that is higher than that of the adjacent material, such that the wider bandgap
material serves as a barrier to electrons and holes, confining them to the same layer in
the material with the smaller bandgap. In type-II alignments, both the conduction
band minima and the valence band maxima are lower than their counterparts for
the abutting material. A subset of the type-II alignment is the type-II broken-gap
configuration, in which the conduction band edge is below the valence band-edge of
the adjacent material. InAs/GaSb superlattices are an example of a type-II broken
gap configuration.
1.1.3 Electrons and Holes
Electrons and holes are the charge carriers in a semiconductor lattice. At 0
K, all electrons will occupy the valence band. It is important, however, to be able
to model the carrier concentration at realistic temperatures. The concentration of
electrons, n, and the concentration of holes, p, are given by
6
EC1
EV1
(a)
EC2
EV2
EC1
EV1
(b)
EC2
EV2
EC1
EV1
(c)
EC2
EV2
Figure 1.3: Band alignment of semiconductor heterostructures. For a (a) type-I band alignment, EC1 > EC2 and EV 1 < EV 2. For (b) a type-II staggered band alignment, EC1 > EC2 and EV 1 > EV 2, and EV 1 < EC2. A (c) type-II broken-gap alignment has EC1 > EC2 and EV 1 > EV 2, but EV 1 > EC2.
n =
∞∫ Ec
gh(E)fh(E)dE (1.8)
In Equations 1.7 and 1.8, Ec and Ev are the conduction band minimum energy and
valence band maximum energy, respectively, gc and gv are the density of states, and
fe and fh are carrier distribution functions.
For a non-degenerate bulk semiconductor, the density of states for the holes and
electrons are given by
)3/2√ Ev − E (1.10)
where me is the effective electron mass and mh is the effective hole mass. The
occupation probability at energy E for a given temperature T is governed by the
7
1 + e(E−EF )/kBT . (1.11)
EF is the Fermi energy. At E = EF , the occupation probability will always be 0.5.
Typically, the Fermi energy is centered in the bandgap for intrinsic semiconductors
(which have an ideal crystalline lattice) at 0 K, but can shift toward the conduction
or valence band with material doping or temperature increases [3, 4].
1.1.4 Recombination Processes
semiconductor processes that can be harnessed for optoelectronic devices occur under
non-equilibrium conditions (i.e. carrier injection or light is applied to the semicon-
ductor). After any event that excites carrier populations, the carriers will eventually
assemble towards their equilibrium concentrations through a carrier recombination
process. In an ideal light-emitting diode (LED), all carrier recombination would be
radiative, or all electrons would enter the conduction band, recombine with a hole in
the valence band, and emit a photon. Practically, other types of nonradiative recombi-
nation are present in optoelectronic devices, the effects of which can be diminished with
growth optimization and band structure engineering. Fig. 1.4 illustrates the primary
recombination processes of interest here, which are radiative, Shockley-Read-Hall
(SRH), and Auger recombination.
The continuity equations for electrons and holes in a semiconductor are given by
∂n
and
∂p
∂t =
1
8
(a)
EC
EV
ET
(b)
EC
EV
Figure 1.4: Radiative and non-radiative recombination mechanisms. Non-radiative recombination mechanisms including (a) SRH and (c) Auger recombination, recombine into the lattice emitting phonons, while (b) radiative recombination emits a photon.
where the variation of the carrier concentration n or p over time is dependent on the
fundamental charge q, the charge currents Jn,p, the net generation rates Gn,p, and the
net recombination rates Un,p. Here, the carrier concentrations are denoted by the sum
of the equilibrium carrier concentration n0 and the excess carrier concentration n,
n = n0 + n p = p0 + p. (1.14)
The net recombination rate in one dimension can be related to the electron and hole
carrier lifetimes τn and by τn
Un = n
τn Up =
τp (1.15)
It is generally useful to describe the total carrier lifetime τ , or the average amount
of time it takes for a minority carrier to recombine into the lattice, by using both
radiative and non-radiative components:
τ−1 = τ−1SRH + τ−1Rad + τ−1Auger. (1.16)
In Equation 1.16, τSRH is the SRH lifetime, is τRad the radiative lifetime, and τAuger
is the Auger lifetime. Combining the contributions from radiative and non-radiative
recombination into the total carrier recombination R, the carrier recombination can
9
R = An+Bn2 + Cn3, (1.17)
where A is the SRH coefficient, B is the radiative recombination coefficient, and C is
the Auger coefficient [5, 6].
SRH recombination occurs as a result of defects in the crystalline lattice, creating
a trap level illustrated in Fig. 1.4(a). Rather than recombining with a hole and emitting
a photon, a carrier will recombine in the trap level and the released energy will be
converted into a phonon. In a heavily p-type semiconductor, the SRH carrier lifetime
for electrons is given by
1
τSRH =
1
τn0
= nTvnσn. (1.18)
For a heavily n-type semiconductor, the SRH carrier lifetime for holes would be
1
τSRH =
1
τp0 = nTvpσp. (1.19)
In Equations 1.18 and 1.19, nT is the defect density, σn and σp are the capture
cross-sections of the traps, and vn and vp are the thermal velocities of the electrons
and holes, given by
vn =
√ 3kBT
me
vp =
√ 3kBT
mh
. (1.20)
As deduced by Shockley, Read, and Hall, the SRH recombination rate RSRH is then
given as
τp0(n+ n1) + τn0(p+ p1) (1.21)
for a trap level with energy ET [7, 8]. The quantities n1 and p1 are the electron and
10
(EF−ET )/kBT . (1.22)
The SRH recombination rate can be reduced by epitaxial material quality improve-
ments, such as reducing crystal relaxation and defect concentration in the material.
Auger recombination occurs when an electron and a hole recombine, and the
energy is absorbed by another carrier (Fig. 1.4(c)). The excited carrier will lose
energy by phonon emission until relaxation occurs. Auger scattering affects the
performance of optoelectronic devices at higher carrier densities due to the cubic
carrier dependence, and can place large limitations on performance in narrow bandgap
semiconductors [9, 10]. The primary Auger processes are band-to-band, phonon-
assisted, or trap-assisted Auger recombination [11].
1.2 Light Emitting Diodes
The development of light emitting diodes (LEDs) forms the core of this work. An
LED is formed by a p-n or p-i-n junction, and emits light over a range of wavelengths.
The spectral bandwidth is temperature dependent governed by the Fermi distribution
(Equation 1.11). A p-n junction may be a homojunction of the same material at
different doping concentrations or a heterojunction of two different materials (that
may also be at different doping concentrations). Regardless of the junction structure,
under equilibrium conditions the Fermi level of the system will align so that it is
a constant across the junction. This creates a process under which electrons will
diffuse to the p-type material and holes will diffuse to the n-type material, creating a
depletion region, the width of which is given by
WD =
√ 2ε
1
NA
+ 1
ND
). (1.23)
In Equation 1.23, ε is the dielectric permittivity of the semiconductor, e is the
11
fundamental charge of an electron, V is the bias voltage, and VD is the diffusion
voltage. NA and ND are the acceptor and donor concentrations, respectively, which
can be considered equal to the electron and hole concentrations if the assumption is
made that all dopants are ionized. The diffusion voltage is given by
VD = kBT
e ln(
The Shockley equation gives the current-voltage (IV) characteristics of the diode,
and can be written as
I = I0e eV/nkBT (1.25)
where, under forward bias,
) . (1.26)
and n is the ideality factor. In Equation 1.26, A is the diode cross section and Dp and
Dn are hole and electron diffusion constants, respectively [12].
The spectral peak of an LED emission under a drive current is governed by its
electroluminescence. When a voltage above the threshold voltage VTH (VTH ≈ VD)
is applied to the p-n or p-i-n junction of an LED, the injected carriers recombine to
emit photons with frequency ν roughly proportional to the bandgap energy Eg, or
hν ≈ Eg ≈ eV. (1.27)
1.2.1 Efficiency
In describing and comparing the performance of LEDs, efficiency is used as a
good standard benchmark to quantify electroluminescence available from a device.
12
However, there are a number of different metrics for efficiency. Quantum efficiency is
one such metric, and can be subdivided into internal and external quantum efficiencies.
Internal quantum efficiency (IQE) provides a metric for the performance of the
active region of an LED, and is a good general characterization of material quality.
IQE measures the rate of photons generated compared to the rate of electrons injected
into the active region. For an ideal LED, every electron would recombine with a hole
and this value would be unity. The IQE can be written as
IQE = Pint/(hν)
I/e , (1.28)
where Pint is the the optical power emitted from the material active region and I is
the injected current. Alternatively, the IQE can be calculated by using the radiative
recombination rate and dividing it by the overall recombination rate, which is used in
Section 5.6.
External quantum efficiency (EQE) is determined by the rate at which photons
are able to escape the LED compared to the rate of electrons injected into the LED.
EQE is a better metric for measuring losses associated with the design of the LED,
which can include light reflected or absorbed by the substrate or metal contacts of the
LED before it escapes. Indeed, EQE can also be written as the product of the IQE
and extraction efficiency ηext,
I/e = ηext × IQE, (1.29)
where Pout is the total output power from the LED and the extraction efficiency can
be written as
ηext = Pout Pint
. (1.30)
Wallplug efficiency ηW , which describes the efficiency of converting energy to
optical power, can also be used. Wallplug efficiency compares the electrical power
13
input into the LED (P = IV ) to the total output power, or
ηW = Pout IV
1.3.1 III-V Compound Semiconductors
Although historically silicon and SiC were the dominant semiconducting material,
III-V compound semiconductors were not very well studied until well after the discovery
that GaAs and related III-V compounds were, in fact, semiconductors in 1952 by
Heinrich Welker [13]. A III-V compound semiconductor can consist of any of the III
or V element semiconductors listed in Table 1.1 in binary, terniary, or higher order
configurations. The term III-V, despite drawing from elements in groups 13 and 15 in
the periodic table, remains due to two archaic periodic group nomenclature schemes,
the Chemical Abstracts Service and old International Union of Pure and Applied
Chemistry numbering, which utilized roman numerals signifying the number of valence
electrons of the group and a system of lettering that made identifying which element
was being discussed impossible without explicit mention of the numbering scheme
being used. For solid-state physicists and chemists, these schemes endured [14].
Period Group II Group III Group IV Group V Group VI
2 B C N O
3 Mg Al Si P S
4 Zn Ga Ge As Se
5 Cd In Sn Sb Te
6 Hg Pb
Table 1.1: Common compound semiconductor elements in the Periodic Table.
14
Many III-V compound semiconductors are suitable for infrared or visible ma-
terials. III-V compounds are often grown in alloy or multiple layer materials due to
having similar lattice constants, minimizing strain in the device (see Fig. 1.5). Of
particular interest in this thesis are AlSb, InAs, and GaSb, which are approximately
lattice matched around 6.1 A, with room temperature energy gaps of 1.61 eV, 0.36 eV,
and 0.73 eV [15]. With lattice constants of 6.0959 A and 6.0583 A, respectively, GaSb
and InAs have a 0.62% lattice mismatch. This disparity can be compensated for at
the heterointerfaces is to grow devices with minimal strain, and the small bandgaps of
InAs and GaSb make it favorable for MWIR and LWIR devices.
B a n d g
a p (
e V
0.5
1
1.5
2
2.5
3
Figure 1.5: Lattice constants and energy gaps of select III-V compound semiconductors at 300 K.
1.3.2 InAs/GaSb Superlattices
A superlattice is a periodic semiconductor heterostructure consisting of rapidly
alternating layers in one direction. First proposed by Esaki and Tsu in 1970, it is
possible to modify the band structure of a superlattice by varying the thickness of
15
its components [16]. One can compare a superlattice heterostructure with that of
a multiple quantum well (MQW). A single quantum well is formed by a well and
barrier material on either side, such that the well material has a bandgap lower
than the barrier material and electrons and holes are confined within the well. The
wavefunction within these barriers can then be considered as standing waves, with
discrete (or quantized) energy levels. A MQW is constructed by separating two or
more wells by multiple barriers. The primary difference between a superlattice and a
MQW is the barrier thickness, so that wavefunctions in adjacent quantum wells do
not overlap in MQWs. Once the barrier thickness is thinned down to the level of the
de Broglie wavelength, however, the band structure of the material begins to change.
The wavefunctions of adjacent wells overlap, and the discrete energy levels form bands
characteristic of a superlattice. A schematic representation of this description is shown
in Fig. 1.6.
(a) (b) (c)
Conduction Band
Valence Band
Figure 1.6: A representation of energy levels in a (a) quantum well, (b) multiple quantum well, and (c) a superlattice heterostructure.
InAs/GaSb are an attractive infrared superlattice material system due to the
broken-gap band alignment (Fig. 1.3), which allows the bandgap of the composite
system to be less than the individual bandgaps of either InAs or GaSb. Under
equilibrium conditions, electrons are primarily localized in InAs, and holes are primarily
localized in GaSb, leading to indirect transitions.
16
Other materials systems are more popular in various aspects of infrared opto-
electronics. HgCdTe (MCT) is an alloy with a direct energy bandgap that can be
tuned throughout the infrared, making it a popular material for infrared detectors.
InSb has a relatively small bandgap, and is common in MWIR detectors. InAs/GaSb
superlattices have been proposed as a comparable material system for infrared de-
tectors [17], although the SRH-limited lifetime has been an issue for most MWIR
and LWIR detector fabrication [15]. Rather, InAs/GaSb superlattices have several
advantages that make them ideal for the development of LEDs.
InAs/GaSb is attractive due to its tunability throughout most of the infrared by
varying the thickness of either component layer. Decreasing the InAs layer thickness for
a fixed GaSb thickness typically causes a blue shift, while increasing it causes a red shift.
Increasing the periodicity beyond a critical thickness can change the material from
a semiconductor to a semimetal [18]. The potential for band structure engineering
for the reduction of Auger recombination is also critical for use of InAs/GaSb in
LEDs [9, 10,19].
Epitaxially grown InAs/GaSb generally has a relatively low defect density and
high material uniformity. There are trade-offs in optimizing the growth process,
however. For example, higher superlattice growth temperatures can be associated
with poorer material quality due to increased interdiffusion between layers and V
adatom desorption during shutter transitions [20, 21]. Lower growth temperatures are
associated with poorer crystalline material quality, with the downside of decreased
photoluminescent emissions from higher SRH recombination rates [22,23]. The largest
format GaSb substrates available are 4-inch, which necessitates using other substrates
available in larger diameters for large scale focal plane arrays or thermal pixel arrays,
although these substrates are associated with rougher material growth and increased
threading dislocations for growth of InAs/GaSb.
17
1.3.3 Cascading InAs/GaSb Superlattices
In developing infrared LEDs, improving emission power and device efficiency are
desirable outcomes. One such method for improving device performance is cascading
of InAs/GaSb active regions. Active region cascading is the growth of successive
superlattices alternated with p-n tunnel junctions. In a cascaded device, an electron
can recombine with a hole in one active region, tunnel through the tunnel junction
into the next active region and recombine with a hole there as well [24]. This process
can be repeated more or less ad infinitum, or until other complications due to device
thickness occur. A schematic of the process is provided in Fig. 1.7.
E n e rg y
Position
Active
Region
n-doped
Tunnel
Junction
p-doped
Tunnel
Junction
Active
Region
Conduction Band
Valence Band
Figure 1.7: A schematic of a cascaded superlattice. An electron recombines in the active region, and is able to tunnel through the tunnel junction to recombine with another hole in the adjacent active region.
If adequate confinement barriers arent imposed on the active region, carriers can
traverse the tunnel junction without radiatively recombining in the active region [25].
Tunnel junctions must be designed with high barriers for electrons in the conduction
band on the p side of the tunnel junction, and high barriers for holes in the valence
18
band on the n side. A tunnel junction is essentially a reverse-biased junction in
comparison to the LED junction. As the LED is increasingly forward biased, the
tunnel junction is increasingly reversed biased, thinning the tunneling region.
The internal and external quantum efficiencies for N cascaded active regions
can scale by a factor of N . As operating current is reduced by a factor of N , effects
due to Joule heating from series resistance are also scaled down by a factor of N2 [26].
Murray et al. studied various tunnel junction designs and found that a p-GaSb/n-
AlInAsSb tunnel junction lattice matched to GaSb provided minimal resistance to
carriers and displayed good overall performance and surface morphology [27]. This
design is utilized in all devices described in this work.
1.4 Applications for InAs/GaSb Superlattices
Optoelectronic devices in the infrared are useful in a variety of academic, in-
dustrial, and military applications. InAs/Ga(In)Sb superlattices have been used to
develop infrared single element detectors [28], focal plane arrays [29–31], and dual-band
detectors [32], as well as laser diodes [19, 33,34]. Single element infrared InAs/GaSb
LEDs have also been developed [?], with applications in remote sensing, gas sensors,
process controls, spectroscopy, optical alignment, and infrared countermeasures [35].
This dissertation focuses primarily on potential applications in thermal scene
generation. As infrared focal plane array detectors become more advanced, more
sophisticated methods of testing these devices, or infrared scene projectors, are required.
Infrared scene projection aims to simulate accurate real-world phenomena for the
sensor being tested. Previous efforts at developing thermal pixel arrays involved the
use of resistive array devices [36–39], although they are limited in frame rate due to
large rise and fall times, as well as lower maximum apparent temperatures. InAs/GaSb
LEDs are well-suited to this application, due to their rapid switching speed, cost,
and reliability [40], and large-format InAs/GaSb MWIR arrays have already been
19
MOLECULAR BEAM EPITAXY
Epitaxy is a form of crystal growth, in which a crystalline layer (or epilayer) is
deposited onto a seed crystal (or substrate) in such a way that the arrangement of
atoms in the epilayer are either lattice-matched or strained by the substrate. Epilayers
may be either the same as the substrate material, which describes homoepitaxy, or
one or more different materials, which describes heteroepitaxy. The development
of the superlattice can be largely tied to advances in epitaxial techniques. Prior
to the development of molecular beam epitaxy (MBE) and metalorganic chemical
vapor deposition (MOCVD), creating the ultra-thin layers required was a virtually
impossible task [42].
MOCVD is a vapor phase epitaxial technique that takes place at atmospheric
pressure and uses metalorganic precursors, typically with hydride sources [43]. Al-
though MOCVD can grow thin, precise layers, the higher pressure makes it difficult
to use in situ growth characterization methods, such as reflection high-energy electron
diffraction (RHEED). Typically, MOCVD lends itself better to higher bandgap mate-
rials grown at hotter temperatures, as well as commercial applications due to higher
growth rates and faster production. It is difficult to grow InAs/GaSb superlattices
with MOCVD, however, as the growth of antimonides is much more complex using
MOCVD. The vapor pressure of elemental Sb is very low, requiring keeping the V/III
ratio near unity during growth [44]. Deviations from unity can lead to antimonide
material forming Ga droplets (lower than unity) or Sb crystals (higher than unity).
The low melting point of Sb also places a limitation on MOCVD growth of antimonides,
as growth temperatures must be kept low (< 520°C), at which point most precursors
for Sb do not fully dissociate [45–47]. Antimony also has no stable hydrides at room
temperature, and other precursors can lead to contamination issues in MOCVD-grown
material [48].
21
Molecular beam epitaxy can be thought of as a highly evolved form of sputtering.
MBE takes place in an ultra-high vacuum chamber, so that the epilayer growth is
primarily governed by the kinetics of the component molecular beams on the substrate.
The growth in ultra-high vacuum produces high quality material with few impurities.
The beam fluxes are controlled by shutters, and growth rates are relatively low (< 1
ML/s), which lends itself to forming the abrupt interfaces characteristic of superlattices
or quantum wells. MBE is used exclusively in this dissertation.
2.1 Epitaxial Growth Modes
The surface morphology of epitaxially developed crystalline structures is, perhaps
obviously, directly driven by the physical processes occurring at the surface of a
structure during growth. A number of different processes can occur during growth,
driven by thermodynamics, strain, substrate orientation, etc. The results of these
processes are generally classified into epitaxial growth modes. Three classical growth
modes are generally identified (and are pictured in Fig. 2.1): Frank-van der Merwe
(FM) layer-by-layer growth mode, Volmer-Weber (VM) island growth mode, and
Stranski-Krastanov (SK) layer plus island growth mode [49,50]. Additionally, other
distinct growth modes have been identified, which include columnar, step-flow, and
screw-island growth modes.
Of the classical growth modes, FM-mode growth occurs when adatoms approach-
ing the substrate are more strongly attracted to the substrate than each other, forming
a complete monolayer before the second monolayer begins [51]. FM growth requires
a flat substrate surface and almost no lattice mismatch strain. Step-flow growth
is generally considered a distinct two-dimensional growth mode from FM growth
and results from a slight misorientation in the cut of a substrate wafer, creating
monatomic steps along the terraces from which growth proceeds [52]. Occasionally, a
slight misorientation can be used as a strategy to avoid island coalescence by inducing
step-flow growth mode. Both FM and step-flow are two-dimensional growth modes
22
FM
VM
SK
Figure 2.1: Schematic diagrams of Frank-van der Merwe (FM) layer-by layer growth, Stranski-Krastanov (SK) layer and island growth, and Volmer-Weber (VW) island growth modes.
and generally considered desirable for high-quality epitaxially grown material, as they
produce low surface roughness and sharp interfaces.
VW growth mode involves clusters of adatoms that are more strongly bound
to one another than the substrate. Adatom clusters will nucleate on the substrate
surface, forming distinct islands. As growth proceeds, the islands will coalesce when
they merge, forming a rough surface. Columnar growth is a similar growth mode,
although the islands do not merge, and instead form an unconnected array of columns.
SK growth is characterized by an initial two-dimensional layer-by-layer growth that
eventually gives way to a three-dimensional island growth mode after a critical layer
thickness is attained. SK growth typically occurs in cases in which there is a large
lattice mismatch between the substrate and epilayers, creating strain conditions that
can manifest in island growth after a critical thickness is reached [50].
Spiral growth is a distinct growth mode associated with strained growth, in
which screw dislocations can trigger step-flow growth around a central point, creating
23
pyramids with a spiral pattern. Spiral growth mode is discussed more extensively in
Section 5.2.
A classical method of determining epitaxial growth modes relies on the free
energies of the substrate (γs) and epilayer (γe), and their interplay with the interfacial
free energy between the substrate and interface, γi. One can then define a parameter
for the change in free energy associated with epitaxial growth, γ, where
γ = γe + γi − γs. (2.1)
Under these conditions, two-dimensional (FM) growth will occur when the
adatoms are more strongly attracted to the surface than each other, or when γ < 0.
Three-dimensional growth modes, corresponding to the VM growth mode, will occur
when γ > 0 and it’s more energetically favorable for adatoms to nucleate onto the
substrate. The presence of strain due to lattice mismatch between the substrate and
epilayer can manifest itself so that once the epilayer reaches some critical thickness, it
becomes energetically favorable to relieve mismatch strain through relaxation, creating
islands on top of full epilayers [43].
Molecular beam epitaxy, however, is a non-equilibrium process. In epitaxy,
non-equilibrium can be defined as growth occurring with a large supersaturation factor
S, such that
Pinf
, (2.2)
where P0 is the partial pressure of the nucleating species and Pinf is equilibrium vapor
pressure of the substrate. One model that uses an atomistic approach for adatoms
by Tersoff, Denier van der Gon, and Tromp, assumes circular monolayer islands with
uniform radius growing under constant adatom flux F [53]. The nucleation rate ω of
islands on the substrate is assumed to be
ω = DN2 0 η
24
in which D is the diffusion constant for surface atoms, N0 is the atomic density on
the surface, η is the normalized adatom density, and ν is the number of atoms in the
smallest possible stable island. The adatom density is subject to the diffusion equation
dη
F
N0
(2.4)
Using boundary conditions at the edge of an island of radius R
dη
√ N0 = 0 (2.5)
and assuming a steady-state solution gives the following solution for η as a function
of r, the distance from the center of the island:
η = η0 − R
η0 = F
(R2 +RLα). (2.7)
In Equations 2.5 and 2.7, α is the probability per unit time that an adatom on an
island will jump off the terrace upon reaching the island edge, divided by the rate
of hops onto the terrace. If there is assumed to be a barrier energy Es for hopping
over the edge of the island, α also varies with temperature such that α ∼ e−Es/kBT ,
where T is the temperature and kB is the Boltzmann constant. Lα is the characteristic
length of a diffusion barrier at an island edge, and Lα = 2
α √ N0
.
The total rate of nucleation on top of the island is obtained by integrating
the island nucleation rate over the area of the island:
=
] . (2.8)
The fraction of islands f that will nucleate a second layer upon them will vary in time
25
as
df
The radius of islands with time is assumed to be
R2 = FL2
ordinary differential equation yields
f = 1− e−(R/Rc)m (2.11)
in which Rc is the critical island radius at which point growth transitions from the
FM growth mode to SK growth, and m is a unitless parameter that depends on the
critical cluster size ν.
Equation 2.11 demonstrates that for islands with size R > Rc, the probability
that a second layer will nucleate on top of an island approaches 0, indicating FM
growth. For islands with size R < Rc, the probability f approaches unity, indicating
SK growth. The most interesting aspect of this nucleation model, however, is its
assessment of the probability of step-flow growth. If a characteristic length Ls is
introduced as the separation between steps on a vicinal substrate, in which Ls = h tan θ,
where h is the step height and θ is the substrate miscut angle, at high temperatures
or Ln > Ls, step flow growth will occur. At low temperatures or Ln < Ls, the FM
growth mode will occur. For Lα < Ln, island growth can be reasonably expected.
2.2 Strained-Layer Epitaxy
Strain and its influence on growth is a concern in all heteroepitaxial structures.
The lattice mismatch strain between two different materials f is defined as
f = as − ae ae
, (2.12)
26
where as is the lattice constant of the substrate and ae is the lattice constant of the
epilayer. In epilayers thinner than the critical thickness, the in-plane strain can adapt
to the crystalline structure by means of an elastic deformation of the crystal lattice.
This process is also called pseudomorphic growth, pictured in Fig. 2.2. The nature of
pseudomorphic growth restricts thicker growths, unless the strain is suitably minimized
through lattice matching. At larger strain, once the epilayer reaches some critical
thickness, strain relaxation will occur through the development of misfit dislocations.
Once strain relaxation occurs, thicker growths can proceed as metamorphic layers.
(a) (b)
Figure 2.2: Schematic diagrams of (a) pseudomorphic and (b) metamorphic epitaxial growth.
The critical thickness hc can be loosely defined as the thickness of an epilayer
at which point it becomes energetically favorable to develop misfit dislocations to
accommodate the strain. The Matthews and Blakeslee model is a model that was
developed for heterostructures that are not egregiously lattice mismatched. By
balancing the strain force acting on an existing threading dislocation and the tension
in the dislocation, the critical thickness for an isotropic single layer can be calculated
27
as
(1 + ν) cosλ
) + 1
) , (2.13)
where ν is the Poisson ratio, b is the magnitude of the Burgers vector, α is the angle
between the Burgers vector and the dislocation line, and λ is the angle between
the Burgers vector and the line in the interface plane that is perpendicular to the
intersection of the glide plane with the interface [54]. For III-V materials, which have
anisotropic bulk elastic constants, the anisotropic model of critical thickness for a
single layer is more applicable:
hc = 2D
Y f
) + 1
) . (2.14)
In Equation 2.14, D represents the average shear modulus at the interface and Y is
Young’s modulus under biaxial stress. Detailed derivations of both critical thickness
models are available in Ref. [55]. Experimental work has shown agreement between
these models and experimental values for epilayers with low levels (<1%) of lattice
mismatch [56].
2.3 Crystalline Defects
One quirk of epitaxial growth is the improbability of obtaining defect-free
epilayers. Although defects break the translational symmetry of the crystalline lattice
and can become sites of non-radiative recombination, they can often be optimized or
even advantageous in growth. Technically, doping is the intentional introduction of
extrinsic point defects. Most defects that are relevant to this thesis can be categorized
as either point defects or dislocations.
A point defect is a single blemish or imperfection in a crystal lattice, although
the effects of point defects can manifest throughout epitaxial growth by propagating
through the crystal. The simplest point defects are intrinsic, or defects native to
28
the material. Vacancies occur when an atom is missing from the lattice, and self-
interstitials occur when an additional atom is wedged into the lattice. Extrinsic point
defects involve impurities, or atoms different from the host lattice, that are introduced
into the material. This type of defect is also referred to as substitutional.
Dislocations are linear defects that occur in a crystal, characterized by the
interruption of the crystal lattice at a strain site that propagates through the crystal.
A misfit dislocation generally refers to any dislocation that is generated at the interface
to relieve strain due to lattice mismatch. Threading dislocations are associated with
misfit dislocations and run throughout the epilayer from the interface, terminating at
the surface [55]. Dislocations can be represented by a line vector, ~l, which describes
the direction of the dislocation, and the Burgers vector, ~b. The Burgers vector is an
invariant representation of the displacement of a dislocation from an ideal crystalline
lattice, which can be determined by forming a closed loop around the dislocation. An
example of a Burgers vector is shown in Fig. 2.3.
Burgers vector
(a) (b)
Figure 2.3: Schematic diagrams of a Burgers circuit in (a) a crystal with a screw dislocation and (b) an idealized version of the same crystal. The Burgers vector, in blue, quantizes the amount that the crystal has distorted to accommodate the dislocation.
The angle between the Burgers vectors and line vector can characterize the
type of dislocation. A screw dislocation is the result of a shear stress across the
29
crystal, causing the atoms on one side of the shear place to be displaced by an
atomic spacing (see Fig. 2.3). In this case, the Burgers vector is parallel to the
direction of the dislocation, such that it can be referred to as a 0° dislocation. An edge
dislocation involves an extra half-plane of atoms inserted into the crystal, similar to
the representation in Fig. 2.2(b). The Burgers vector is perpendicular to the direction
of the dislocation, and these dislocations are generally referred to as 90° dislocations.
In practice, dislocations can manifest both edge and screw characters, and tend to be
some combination of the two.
The glide plane of a crystal refers to the plane that contains both the line and
Burgers vector. Zinc blende semiconductors, which are exclusively used in this thesis,
typically manifest dislocations along the {111} glide plane. 60° dislocations, defined
as any dislocation in which there is a 60° angle between the Burgers vector and line
vector, are common in zinc blende crystals along the {111} glide plane. In some cases,
notably with GaSb heterostructures on GaAs, arrays of 60° dislocations have been
predicted to completely relieve strain in the heterostructure by canceling out at their
intersections and completely relaxing the first few monolayers [56, 57]. These types of
dislocations are confined to the interface and are known as Lomer dislocations.
2.4 Molecular Beam Epitaxy Growth Equip- ment
Solid source molecular beam epitaxy (MBE) is an extremely precise epitaxial
technique that is characterized by atomic layer-by-atomic layer deposition of materials
onto a seed substrate in ultrahigh vacuum. Component sources are controlled by
individual effusion cells, which heat materials to the point of sublimation into a
molecular beam. These beams then coalesce and crystallize on the substrate material
at extremely low deposition rates, typically around 1 monolayer/second. The growth
rate, as well as relative doping levels, depends on the evaporation rate of material in
the effusion cells. The extreme vacuum levels allow for the development of relatively
30
high purity crystals, as well as the use of in situ diagnostics to control growths, such
as reflection high energy electron diffraction (RHEED).
heating block rotating substrateRHEED gun
heating block
Figure 2.4: A schematic diagram of an MBE growth chamber.
A schematic of a basic MBE chamber is displayed in Fig. 2.4. The beam flux
extending from the effusion cells is controlled with a system of mechanical shutters,
allowing for precision control of flux and well-defined interfaces. Control over material
deposition is exerted by both shutters and effusion cell temperatures. The substrate
temperature is controlled using a heating block behind the substrate, which is mounted
on a rotating holder to promote material uniformity. This thesis uses both Veeco
Epi930 and Gen20 rectors, each of which is equipped with gallium, indium, aluminum,
antimony, and arsenic effusion cells and beryllium and tellurium as p-type and n-type
dopants, respectively. The group V materials, antimony and arsenic, have valved
31
2.4.1 Vacuum and Ion Gauges
In order to maintain the stringent vacuum conditions for MBE growth, most
MBE systems are comprised of an entry or load lock chamber, a transitional buffer
chamber, and the actual growth chamber. Substrates are mounted onto holders and
transferred from chamber to chamber with an interstage transfer system. Thus, the
growth and buffer chambers are opened to the atmosphere only during maintenance
cycles. The exposure to air is then cleaned with a bake out period of the system to
remove adsorbed contaminating species. Typical vacuum conditions in the growth
chamber are ∼ 10−10 Torr or lower, and are dictated by the purity requirements of
the material [43].
The vacuum in the chamber can be characterized by the mean free path of gas
molecules in the vacuum, λ, and the concentration of gas molecules in the vacuum per
unit volume, n. Neglecting particle interactions in the gas and assuming an isotropic
gas with a Maxwellian velocity distribution, the mean free path can be derived as
λ = 1√
2nπσ2 , (2.15)
where σ is the average molecular diameter and πσ2 is the collisional cross section [49].
The concentration of gas molecules can be estimated from the ideal gas law, such that
n = P
kBT , (2.16)
where P is the pressure, kB is the Boltzmann constant and T is the temperature. Sub-
stituting Equation 2.16 into Equation 2.15 yields the mean free path of an evaporated
particle:
. (2.17)
Equation 2.17 shows how the mean free path of species in the system can be
32
employed to determine the necessary vacuum for MBE growth. The mean free path of
constituent species must be at least as long as an MBE chamber to form a “molecular
beam,” which is generally ∼1 m. The mean free path as a function of pressure is
plotted for Ga (which has a molecular diameter of ∼270 pm) at 300 K in Fig. 2.5.
As shown in the plot, an MBE system can be conceivably used under a pressure of
10−4 Torr to maintain a mean free path of ∼1 m. In practice, however, the pressure
of typical MBE systems leads to mean free paths with an order of magnitude of 1000
km. This pressure is utilized to maintain purity-it is ideal to keep the mean free path
of contaminant species as high as possible to avoid collisions within the molecular
beam and incorporation into the epilayer.
Pressure (Torr)
re e P
100
102
104
106
108
Figure 2.5: The mean free path of a Ga atom as a function of pressure at 300 K.
Multiple pumps are typical in an MBE system to maintain low pressures. Pumps
are primarily either capture pumps, which collect residual gases via freezing or gettering,
or throughput pumps, which compress gas to a level that can be sucked out by a
roughing pump. Some of the more common pumps used, all of which are installed
on the Gen20 reactor, are helium cryo-pumps, ionization pumps, and turbomolecular
33
pumps, as well as liquid nitrogen-cooled cryo-panels, which thermally isolate effusion
cells and condenses gases on its surface.
A cryopump is a capture pump that relies on freezing particles onto a cryopanel.
Liquid helium is a preferred coolant, although cooled gaseous helium can be used as
well. A closed cycle compressor pumps liquid helium so that it cools the cryopanel,
trapping heavier particles on the surface. Lighter species, such as helium or hydrogen,
are difficult to pump with a cryopump as they will not adsorb at liquid helium
temperature (4 K) [58]. Cryosorption techniques can help pump lighter species as
certain porous materials, such as charcoal, will adsorb more lightweight particles at
low temperatures [59]. The pumping speed of a cryopump is generally proportional to
the effective surface area of the cryopanel. Regeneration on occasion is necessary to
remove condensate from the cooled walls of the pump.
Ion pumps are capture pumps that rely on trapping ionized particles. A Penning
cell is used to create an electric field in which electrons move towards an anode,
ionizing incoming gas particles during collisions. The ionized gas particles are then
accelerated into the cathode, which, in a sputter ion pump, is generally titanium. Due
to the large amount of kinetic energy obtained from the electric field, some fraction of
molecules, particularly lighter ones, will embed themselves into the cathode. Other
ions will sputter the cathode material over the anode and pump walls. The sputtered
film acts as a getter for gas molecules, which can react with the sputtered material
and form stable compounds, adsorbing onto the film surface [60,61]. Ion pumps can
be useful for pumping lighter materials, including hydrogen.
Turbomolecular pumps are throughput pumps that operate by transferring
momentum to gas particles through rapidly moving blades. The particles are directed
towards the exhaust of the pump and compressed by the rotors, where they are
removed via backing pump. Pumping speed is unaffected by molecular weight, so
lighter species are pumped as well [62]. Turbo pumps can achieve pressures as low as
34
10−11 Torr [60].
Ion gauges are used to determine the pressure levels in an MBE system, and
function similarly to an ion pump. A heated filament is used to produce electrons,
which are accelerated through a grid to ionize gas molecules. A collector wire is used
to measure the current generated by the ions produced, which is proportional to the
gas density in the gauge. Thus, the pressure in the gauge P can be determined from
the relationship
Ic = S × P × Ie, (2.18)
where Ic is the current measured by the ion collector, Ie is the electron current directed
to the grid, and S is the sensitivity of the ion gauge.
2.4.2 Effusion Cells
The molecular beams in solid state MBE are produced by effusion cells, dia-
grammed in Fig. 2.4, in which the solid elemental materials are evaporated. A Knudsen
cell is an idealized version of a modern MBE effusion cell, in which material in an
isothermal cell is evaporated from a small orifice. In this idealized enclosure, the
diameter of the orifice must be less than or equal to one-tenth of the mean free path
of atoms in the enclosure at equilibrium pressure, and the wall surrounding the orifice
must be infinitesimally thin so that impinging atoms will not be adsorbed or scattered
by the wall. Under these conditions, Knudsen derived a total effusion rate Γe from an
effusion cell in ultra-high vacuum, representing the number of molecules exiting the
cell per unit time [49],
Γe = Aepeq
2πMkBT , (2.19)
where Ae is the orifice area, peq is the equilibrium pressure in the effusion cell, Na
is Avogadro’s number, M is the molecular weight of the material in the cell, kB is
Boltzmanns constant, and T is the cell temperature. The equilibrium pressure can be
35
peq = Ceα/T , (2.20)
where C is a constant determined by the system geometry, and α is a constant
determined by the evaporated species. Equation 2.20 is useful for monitoring the
beam flux as a function of temperature, as the factor of T 1/2 in the denominator of
Equation 2.19 can be neglected to produce
Γe = Ae√ T eα/T ' A′eα/T . (2.21)
In Equation 2.21, A and A′ are constants. In practice, Equation 2.21 is a good
approximation that can be related to the growth rate (GR) of a given material, as the
beam flux and GR are proportional quantities. Thus, for a given effusion cell, the GR
as a function of cell temperature can be empirically modeled as
GR = A′′eα/T , (2.22)
where A′′ and α can be determined through RHEED measurements (see Section 2.5.1).
Practically, one cannot satisfy the requirements for a Knudsen cell as the walls
surrounding the orifice of an effusion cell will always have some finite thickness.
Nevertheless, effusion cells in MBE are occasionally referred to as k-cells. Real effusion
cells are also not strictly isothermal, which can cause potential problems due to cooling
at the end of the cell. Cooling at the exit orifice from radiative heat loss can result in
a build-up of condensed material, reducing the exit orifice dimensions and lowering
the total beam flux hitting the substrate surface.
Modern MBE effusion cells typically use pyrolytic boron nitride (PBN) as
the insulating material for sources as well as the crucible material as PBN can be
obtained with low impurity levels. A water-cooled shroud is deployed to dissipate
heat. Temperature control of each cell is provided by proportional-integral-derivative
(PID) controllers with thermocouple feedback. Each source has a shutter than can
36
abruptly stop or start the beam once the material in the effusion cell is heated to the
temperature corresponding with the desired growth rate.
Antimony and arsenic produce tetramers when directly evaporated, which tend
to have lower surface mobility than dimers and monomers. Tetramers also assemble
on the growth surface in clusters, which can lead to vacant sites and associated point
defects. Cracker cells are generally used with V sources, which are additions to a
regular effusion cell that add a literally-named, high-temperature region that can
dissociate tetramers into dimers or monomers. For arsenic, cracking zones have been
shown to improve optical and electrical properties of GaAs, as well as lower the defect
density [63]. Similarly, with cracking zone temperatures above 700°C, a mixture of Sb2
and Sb1 can be obtained. Above 1100°C, mostly antimony monomers are produced [64],
the use of which has been found to increase GaSb photoluminescent emission 5-10
times as strong as Sb dimers [65].
2.4.3 Dopants
Extrinsic semiconductors are produced by intentionally introducing impurities
into the crystalline lattice in order to change the carrier concentration. Dopants can
be input into the lattice from heated effusion cells, and are simultaneously sublimated
onto the substrate with other constituent materials, albeit in much lower quantities.
Dopants for III-V semiconductors are usually group II (p-type) or groups VI (n-type),
as group II materials can replace a III-type element and add an additional hole to the
lattice, or VI materials can replace a V-type element and add an additional electron
to the lattice [66]. Both the Gen20 and Epi930 reactors use beryllium as a p-type
dopant and gallium telluride as an n-type dopant.
Dopants generally compose a relatively small elemental fraction of a doped
semiconductor, to the point that the flux is difficult to measure with a beam flux
monitor (BFM). A BFM is a moveable ion gauge that measures the beam equivalent
37
pressure (BEP) from a given molecular beam that be placed in front of the substrate
and swung out of the way during growth. Thus, dopant level calibration as a function
of dopant cell temperature is generally carried out by a series of test growths designed
to measure carrier concentration. A calibration growth is generally composed of 1-2 µm
of material on a semi-insulating substrate. Hall and secondary ion mass spectroscopy
(SIMS) measurements are common for determining the dopant concentration as a
function of dopant cell temperature.
SIMS measurements bombard a sample’s surface with a primary ion beam and
use a mass spectrometer to analyze sputtered secondary ions. Dynamic SIMS uses a
higher intensity ion beam that actually etches away the surface of the sample during the
measurement, allowing depth profiles for species of interest to be taken. The relatively
high sensitivity of SIMS is its primary advantage, as it can detect concentrations
of up to 1010 atoms/cm3. By varying the temperature of the dopant k-cell over a
controlled depth during growth, one can obtain a depth profile using dynamic SIMS
that displays the density of a dopant at each cell temperature. Although the values
are proportional and often similar, dopant density is not necessarily representative
of the majority carrier concentration. Thus, SIMS measurements must be made in
tandem with Hall measurements for an accurate dopant calibration.
In principle, Hall measurements are voltage measurements obtained under a
steady magnetic field. Under the magnetic field B, a perpendicular current I directed
through a semiconductor will have electrons start to drift in the direction of the cross
product of the magnetic field and current vectors (the y-direction, in Fig. 2.6) due to
Lorentz forces. The accumulation of surface charge on one side of the sample creates
the Hall voltage VH , given by
|VH | = IB
end , (2.23)
where e is the elementary charge, n is the carrier density, and d is the thickness of the
sample. The van der Pauw technique [67] utilizes contacts placed at the corners of a
38
sample, and is generally employed to measure the Hall voltage and thus obtain the
sheet carrier density through a four-point probe technique.
I
Figure 2.6: A schematic diagram of a Hall measurement.
One can model the dopant carrier concentration in a similar manner as Equations
2.21 and 2.22, in that it can be assumed that the majority carrier concentration will be
proportional the effusion rate from the dopant cell. Thus, one can make approximately
linear graphs of ln(n) vs. (1/T ) to calibrate each dopant cell, using the SIMS dopant
density profile as a slope and the Hall measurement to determine the y-intercept.
2.4.4 Desorption
Surface oxide removal (or in the parlance of epitaxy, desorption) is an important
consideration and first step in epitaxial growth. Most III-V substrate surfaces are
chemically active, and can quickly form a few nanometers of oxide with exposure to air.
Surface defects and roughness occur without complete oxide removal, affecting device
performance. A clean, oxide-free substrate is imperative for high quality MBE growths.
39
Oxide formation and removal on GaAs and GaSb are discussed in this section, as they
are used extensively in this thesis.
Surface oxides on GaAs and GaSb form from chemisorbed oxygen, which moves
around on the substrate surface until it finds a stable bonding site [68]. Oxides initially
nucleate as islands that grow parallel to the substrate surface until the substrate
is covered. One the surface is covered, oxidation occurs either through the inward
diffusion of oxygen or the outward diffusion of the substrate material. The oxidation
rate of GaSb is logarithmic over time, whereas oxide growth on GaAs proceeds more
slowly up to 10 A, at which point the oxide growth rate is also roughly logarithmic.
Native oxides on GaSb substrates are primarily Ga2O3 and Sb2O3. Sb2O3 is
metastable on GaSb, and the following reaction takes place over 200°C [69]:
Sb2O3 + 2GaSb→ Ga2O3 + 4Sb. (2.24)
During thermal desorption in the MBE growth chamber, the reaction described in
Equation 2.24 will generally take place around 200°C, forming Ga2O3 and elemental
Sb. Sb2O3 is found to desorb fully around 380°C. Full oxide removal is found to occur
between 480°C-520°C, depending on the chemical etch used to thin oxides before
growth [70].
Native oxides on GaAs are slightly more complex to desorb. GaSb has a weaker
crystalline bond than GaAs, and oxide layers grow in nearly monolayer steps [68].
On the surface of a GaAs substrate at room temperature, elemental As, Ga2O3, and
As2O3 are present, while closer to the oxide-substrate interface GaAsO3 and GaAsO4
can be found [71]. The arsenic oxides can generally be removed by heating to 400°C,
while Ga2O3 need temperatures in excess of 500°C to desorb. The values for full
oxide desorption on GaAs compose a wide range of values dependent on the type of
oxides formed as well as the total oxide thickness, with maxima as high as 660°C.
For thin oxide layers, however, full oxide removal has been found to occur as low as
40
580°C [72,73].
Chemical thinning of oxides has been found to improve surface smoothness
of substrates during thermal oxide removal, particularly for GaSb substrates. All
GaSb substrates in this thesis use the chemical thinning steps outlined. Substrates
are degreased in both acetone and isopropanol baths for 5 minutes, and blown dry
with nitrogen. Although the degreasing step is a remnant from an earlier era in
which substrates were packed in oil, it is still used to promote surface cleanliness.
Substrates are etched in HCl for 4 minutes, rinsed with isopropanol, and blown dry
with nitrogen before being loaded into the MBE reactor. Vineis et al. found that an
HCl etch mainly etches Sb2O3, such that most of the Sb2O3 reacts with the GaSb
substrate during heating, leaving no Sb2O3 to desorb at 380°C. In turn, this promotes
a smoother growth surface than un-etched epi-ready substrates [74, 75]. The oxide
thickness of vacuum-packed, epi-ready GaAs substrates obtained from Wafertech
has been measured to be 20-25 A using spectroscopic ellipsometry, which confirms
results obtained by Allwood et al on epi-ready GaAs oxide thickness [71]. As this is a
relatively thin oxide layer, no additional chemical etching of epi-ready GaAs substrates
has been pursued in this thesis.
In lieu of thermal oxide desorption at temperatures in excess of 500°C, atomic
hydrogen is a viable option for low temperature oxide removal. To this end, a Veeco
Atom-H source with a high temperature (1800°C-2200°C) H2 cracker has been installed
on the prep chamber of the Gen20 reactor. An atomic hydrogen source (AHS) uses the
cracker cell to dissociate hydrogen pumped into the chamber, which then chemically
reacts with the oxides on the substrate surface. Ultra-high purity hydrogen is pumped
into the chamber at an estimated flow rate of 0.2-0.3 cc/min, to a background pressure
of 10−6-10−5 Torr. The hydrogen and removed oxides are then pumped out of the
chamber with an Agilent V301 turbo pump.
For GaAs oxide removal with atomic hydrogen, the cleaning process is two steps.
41
As2Ox + 2xH∗ → xH2O + Asx/( 1
2 As4), (2.25)
Where x = 1, 3, or 5, depending on the oxide, and Ga oxides react with the hydrogen
via
Ga2O3 + 4H∗ → 2H2O +Ga2O. (2.26)
With atomic hydrogen, most oxides present on GaAs substrates are found to be com-
pletely removed at 400°C [77,78], although substrate anneals under atomic hydrogen
overpressure are found to improve the surface quality. Oxide removal with atomic
hydrogen on GaSb has been found to take place at temperatures as low as 300°C [79].
Anneals with an AHS over 500°C have been found to introduce pitting and surface
roughness to the substrate [78,80].
Recipes developed for use with the Veeco Atom-H source roughly mirror the
findings in the above paragraph. Temperature determination during AHS desorption is
somewhat imprecise, as there is no real-time substrate temperature monitoring in the
prep chamber, aside from the thermocouple. Issues with temperature determination
are dealt with in Section 2.5.2. Nevertheless, GaSb samples have been found to desorb
under a hydrogen background pressure of 1 × 10−6 Torr at ∼300°C, as confirmed by
RHEED analysis. A variety of samples are used for GaAs oxide removal with the
AHS, with variable recipes dependent on the needs of the growth at hand. It has been
found that oxides on GaAs can be removed from the substrate within 15 minutes at
∼400°C, although increases in substrate temperature generate a brighter and clearer
(2×4) RHEED pattern.
2.5 In Situ Growth Monitoring Techniques
One advantage of the ultra-high vacuum of the MBE growth environment is
the ability to monitor and control the growth in real time. In situ methods must not
42
block the molecular beams or interfere with the growth. Although many methods
have been used in situ in MBE growth chambers for real-time film analysis, such as
low-energy electron diffraction (LEED), ellipsometry, and Raman spectroscopy, the
Gen20 and EPI930 reactors are equipped with a RHEED gun and kSA BandiT. Once
an MBE growth has started, the primary tools available during growth are RHEED
and temperature monitoring, which are detailed in the following sections.
2.5.1 Reflection High-Energy Electron Diffraction
Reflection high-energy electron diffraction (RHEED) is an important quali