STRAINED Ge AND GeSn BAND ENGINEERING
FOR Si PHOTONIC INTEGRATED CIRCUITS
A DISSERTATION
SUBMITTED TO THE DEPARTMENT OF ELECTRICAL
ENGINEERING
AND THE COMMITTEE ON GRADUATE STUDIES
OF STANFORD UNIVERSITY
IN PARTIAL FULFILLMENT OF THE REQUIREMENTS
FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY
Yijie Huo
December 2010
© 2011 by Yijie Huo. All Rights Reserved.
Re-distributed by Stanford University under license with the author.
This work is licensed under a Creative Commons Attribution3.0 United States License.
http://creativecommons.org/licenses/by/3.0/us/
This dissertation is online at: http://purl.stanford.edu/fm704sg2739
ii
I certify that I have read this dissertation and that, in my opinion, it is fully adequate
in scope and quality as a dissertation for the degree of Doctor of Philosophy.
James Harris, Primary Adviser
I certify that I have read this dissertation and that, in my opinion, it is fully adequate
in scope and quality as a dissertation for the degree of Doctor of Philosophy.
Krishna Saraswat
I certify that I have read this dissertation and that, in my opinion, it is fully adequate
in scope and quality as a dissertation for the degree of Doctor of Philosophy.
Kamins Ted
Approved for the Stanford University Committee on Graduate Studies.
Patricia J. Gumport, Vice Provost Graduate Education
This signature page was generated electronically upon submission of this dissertation in
electronic format. An original signed hard copy of the signature page is on file in
University Archives.
iii
Abstract
The on-chip interconnect bandwidth limitation is becoming an increasingly
critical challenge for integrated circuits (ICs) as device scaling continues to push the
speed and density of ICs. Silicon photonics has the ability to solve this emerging
problem due to its high speed, high bandwidth, low power consumption, and ability
to be monolithically integrated on silicon. Most of the key devices for Si photonic
ICs have already been demonstrated.
However, a practical CMOS compatible
coherent light source is still a major challenge.
Germanium (Ge) has already been demonstrated to be a promising material
for optoelectronic devices, such as photo-detectors and modulators. However, Ge is
an indirect band gap semiconductor, which makes Ge-based light sources very
inefficient and limits their practical use. Fortunately, the direct Γ valley of the Ge
conduction band is only 0.14 eV higher than the indirect L valley, suggesting that
with band-structure engineering, Ge has the potential to become a direct band gap
material and an efficient light emitter.
In this dissertation, we first discuss our work on highly biaxial tensile strained
Ge grown by molecular beam epitaxy (MBE). Relaxed step-graded InGaAs buffer
layers, which are prepared with low temperature growth and high temperature
annealing, are used to provide a larger lattice constant substrate to produce tensile
strain in Ge epitaxial layers. Up to 2.3% in-plane biaxial tensile strained thin Ge
iv
epitaxial layers were achieved with smooth surfaces and low threading dislocation
density. A strong increase of photoluminescence with highly tensile strained Ge
layers at low temperature suggests that a direct band gap semiconductor has been
achieved.
This dissertation also presents our work on more than 9% Sn incorporation in
epitaxial GeSn alloys using a low temperature MBE growth method. This amount of
Sn is 10 times greater than the solid-solubility of Sn in crystalline Ge. Material
characterization shows good crystalline quality without Sn precipitation or phase
segregation. With increasing Sn percentage, direct band gap narrowing is observed
by optical transmission measurements. The studies described in this dissertation will
help enable efficient germanium based CMOS compatible coherent light sources.
Other possible applications of this work are also discussed in the concluding chapter.
v
Acknowledgement
First and foremost, I would like to sincerely thank my advisor, Professor
James Harris (“Coach”) for all his encouragement, support, and guidance in many
aspects. Coach is the wise, fearless captain of the group. He is amazingly visionary
at identifying future projects, meanwhile he is also very open-minded to allow us to
try many new projects. He is very experienced and knowledgeable in all the material
growth, device design and fabrication and system measurement, but he still
encourages us to try totally different approaches.
He is very experienced at
managing the project progress and reports to contractor, while he blocks most of the
pressure and provides a comfortable environment for the students. Due to his great
personality and high academic achievements, many previous group members and
scholars come to our group as visitor scholars and they bring in many great
opportunities to new research as well. Coach is not only my academic advisor, but
also my life-long mentor.
I would like to thank Professor Ted Kamins. Ted is very knowledgeable
about SiGe material growth and Si devices technologies. He is one of the most
responsible people I have met. The meetings with Ted always gave me many new
ideas leading to great progress. I would like to thank Professor Saraswat for his
thought-provoking advising in Ge materials and devices; Professor Vuckovic for her
advising in photonic crystals cavity and optical characterizations; Professor Miller
vi
for his helpful advising in quantum confined stark effects and SiGe modulators;
Professor Brognersma for his advising in nano-scale surface plasmon polariton light
source.
This Si phonics project cannot achieve so many achievements without all the
efforts from the team members. I would especially like to thank Hai and Robert for
their great work on strained Ge and GeSn projects, Yiwen for the work in material
growth with MBE system, and Yu-Husan who is the pioneer in our group to start the
SiGe modulator work. I would like also to thank all the collaborators in this project:
Dr. Makarova, Dr. Fiorentino, Dr. Ochalski and Professor Kim. Thanks to the entire
MBE grower team, Angie, Donghun, Hopil, Tomas, Weisheng, Ed, and Mingyang. I
would like to give my special thanks to Angie who always helps to organize the daily
operation of MBE lab and is always a great source for all kinds of help.
I have involved in many other projects during the Ph.D. studies. The first
project is the on-chip nonlinear optics wave guide and later changed to stopping light
with photonic crystal structure with the same team members. I would like to thank
Professor Fejer. He is one of the most knowledgeable people I have met. For all
kinds of questions, he can directly give me the previous conclusion as well as who
had worked on it and when. Also, when I face a problem, he can always identify the
most critical factor with very convincing explanations.
I would like to thank
Professor Fan. He is one of the smartest people I have met. His profound and
unique understanding of electromagnetic dynamics and related fields brings us to
many great new discoveries. Also, he can always break traditional opinions and
solve the problem with neat new methods. I want to especially thank Dr. Luigi
vii
Scaccabarozzi, who was my mentor at the beginning of my Ph.D.. He systemically
trained me on the optical devices simulations, fabrications, and measurements, which
I keep benefit from. I also want to thank other project members, Michelle, Jun, Sunil,
and Norbert for all their great work.
I would like to thank Professor Cui, Professor McIntyre, and Professor Wong
for the collaboration in nano-structured multi-junction solar cell and LED projects. I
really benefit from their insightful knowledge on nano-scale science. I would like to
especially thank Anjia who is the one of the most motivated people I have met. I
also would like to thank Dong, Evan, Yangsen, Antonio, Meiyueh, Xinyu, Erik, Shu,
Jia for their collaborations.
I would like to thank Professor Goldhaber-Gordon, Professor Zhang for their
advising on graphene and topological insulator. I would like to especially thank
Professor Hesjedal who is not only superior at the fundamental science but also good
at many engineering aspects. I also would like to thank Shuang and David for their
great collaborations.
I would like to especially thank our administrative assistant, Gail ChunCreech, who kindly and skillfully handles all the group daily affairs. She is one of
the nicest people I have met. Harris group is a large group with a few tens of group
members doing research in different areas. We are working as a huge family to help
each other and share the happiness. I would like to thank all the lab members that I
have not mentioned yet: Ofer, Xiaojun, Junxian, Seth, Homan, Mark, Xian, Hong,
Ke, Zhilong, Li, Rafael, Michael, Paul, Xiao Hann, Meredith, Thomas, Lele, Pascale,
viii
Larkhoon, Sonny, Altamash, Phani, and all the previous and future group members. I
also would like to thank all the SNF, SNL staffs to share their professional
knowledge with me. Making many great friends at Stanford is one of the best gifts
for me. I would like to thank you all for making my life at Stanford colorful.
Last, but not least, I thank my family for their tremendous support. My
presents, as a professor and a teacher in Tsinghua University, have given me the
heart to chase curiosity and science, the targets to pursue, the courage to achieve my
dreams, and their unconditional support from day to day. I can never adequately
thank them for their givings. And most of all, I must give my thanks to my wife,
Chen Wu, who is getting her Ph.D. degree from Stanford as well. She not only takes
care of me every day but also helps me with my studies. She always encourages me
to pursue higher targets when I am wandering.
Her positive attitude and fast
implementation ability always inspire me. She brings me many changes not only in
working style but more importantly in the living style. Thank you so much for your
love and support all the way. I will do my best to make us the happiest couple for
the rest of our lives.
ix
Dedication
To my Mom, Dad, and my wife Chen.
x
Contents
Abstract
iv
Acknowledgement
vi
Dedication
x
1
1
Introduction
1.1
On-Chip Interconnections . . . . . . . . . . . . . . . . . . . . . . . . .
1
1.1.2
Optical interconnects . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2
Key Devices for Si Photonics . . . . . . . . . . . . . . . . . . . . . . . . . . . .
1.3
Ge Band Structure Engineering . . . . . . . . . . . . . . . . . . . . . . . . . . . 6
2.2
3
1.3.1
Heavily n-Type Doping Germanium . . . . . . . . . . . . . . . . . 6
1.3.2
Biaxial Tensile Strained Germanium . . . . . . . . . . . . . . . . . 7
1.3.3
GeSn Alloy . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 9
Organization . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11
Growth Equipment and Characterization Techniques
2.1
1.1.1
1.2
1.4
2
On-Chip Interconnection Challenge and Solution . . . . . . . . . . . . . 1
16
Molecular Beam Epitaxy (MBE) . . . . . . . . . . . . . . . . . . . . . . . . . . 16
2.1.1
Introduction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 16
2.1.2
MBE System. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17
2.1.3
Material Growth Mechanism . . . . . . . . . . . . . . . . . . . . . . . 20
2.1.4
Advantages of MBE Systems. . . . . . . . . . . . . . . . . . . . . . . 20
Characterization Techniques. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 22
2.2.1
Scanning Electron Microscopy (SEM) . . . . . . . . . . . . . . . . 22
2.2.2
Scanning Probe Microscopy (SPM) . . . . . . . . . . . . . . . . . . 23
xi
2.3
3
2.2.3
X-Ray Diffraction (XRD) . . . . . . . . . . . . . . . . . . . . . . . . .
2.2.4
Transmission Electron Microscopy (TEM) . . . . . . . . . . . . 31
2.2.5
Raman Spectroscopy. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 36
2.2.6
Secondary Ion Mass Spectrometry (SIMS) . . . . . . . . . . . . 38
2.2.7
Photoluminescence (PL) . . . . . . . . . . . . . . . . . . . . . . . . . . . 39
Summary. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 44
Biaxial Tensile Strained Ge
3.1
25
46
Introduction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 46
3.1.1 Overview of Strained Ge. . . . . . . . . . . . . . . . . . . . . . . . . . . 46
3.2
3.1.2
Methodology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48
3.1.3
Section Arrangement. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50
Relaxed InGaAs Buffer Layers. . . . . . . . . . . . . . . . . . . . . . . . . . . . 50
3.2.1 Motivation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 50
3.2.2
Traditional InGaAs growth recipe. . . . . . . . . . . . . . . . . . . . 52
3.2.3
InGaAs buffer layer growth optimization. . . . . . . . . . . . . . 53
3.2.4
Discussion of Other Optimization and Growth Mechanisms for InGaAs Buffer Layers. . . . . . . . . . . . . . . . . . . . . . 62
3.2.5 Summary. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 68
3.3
3.4
4
Properties of Biaxial Tensile Strained Ge . . . . . . . . . . . . . . . . . . .
3.3.1
Critical Thickness for Strained Ge Layers . . . . . . . . . . . . . 69
3.3.2
Strained Ge Growth. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72
3.3.3
Strain Calibration in Thin Ge Layers. . . . . . . . . . . . . . . . . . 72
3.3.4
Material and Optical Characterization of Strained Ge. . . .
81
Strained Ge Summary. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 90
GeSn Alloys
94
4.1
Introduction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 94
4.2
GeSn Materials growth. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 95
4.2.1
69
Approach to GeSn Alloy Growth. . . . . . . . . . . . . . . . . . . . 95
xii
4.2.2
Sn Growth Calibration. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 97
4.2.3
GeSn alloy growth mechanism. . . . . . . . . . . . . . . . . . . . . . 100
4.2.4 GeSn Material Quality Characterization . . . . . . . . . . . . . .
4.2.5
5
103
Strained GeSn . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 105
4.3
GeSn Band Gap Energy Characterization . . . . . . . . . . . . . . . . . . .
4.4
Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 111
Summary and Future Work
108
113
5.1
Summary . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
113
5.2
Suggestions for Future Work . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 114
5.2.1
Ge Material Studies . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .
5.2.2
Ge Electro-Optics Device Studies . . . . . . . . . . . . . . . . . . . 116
5.2.3
On-chip Optical Interconnect System . . . . . . . . . . . . . . . . . 116
xiii
115
List of Tables
3.1
Biaxial tensile strain inside Ge with different InGaAs buffer layers. . . . . . 79
4.1
Summary of Sn composition in three different samples using SIMS
calibrated growth rates and XRD lattice constant with linear lattice
constant interpolation between Ge and Sn. . . . . . . . . . . . . . . . . . . . . . . . . . 99
4.2
GeSn growth samples and their 10 μm x 10 μm RMS surface roughness from AFM measurements. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 101
xiv
List of Figures
1.1
Relative delay vs. process technology node from 2007 ITRS. . . . . . . . . . . 2
1.2
Si photonics key devices, including integrated light source,
modulator, waveguide, and photo-detector. . . . . . . . . . . . . . . . . . . . . . . . .
1.3
4
Calculated valence and conduction band shifts at various symmetry
points in Ge as a function of in-plane biaxial strain. . . . . . . . . . . . . . . . . . . 7
1.4
Calculated electron and hole mobilities of (100) Ge as a function
of biaxial tensile strain. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 8
1.5
Calculation and experimental result of band crossing of GeSn
alloy as a function of Sn concentration. . . . . . . . . . . . . . . . . . . . . . . . . . . . 9
1.6
Ge-Sn phase diagram, showing less than 1% solid solubility of Sn in Ge. . 10
1.7
SEM image of Ge0.78Sn0.22 alloy grown at 200 °C with clear
Sn phase segregation. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 11
2.1
(a) A schematic drawing, and (b) a picture of an MBE chamber from side
view, showing material sources, beam flux gauge, and the RHEED system. 17
xv
2.2
Typical (a) tilted cross section SEM image and (b) top view SEM
image of GaAs nano-structured solar cells. . . . . . . . . . . . . . . . . . . . . . . . . . 23
2.3
A typical AFM image of a relaxed InGaAs buffer layer on a GaAs
substrate. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 24
2.4
(a) A photograph of Phillips X'Pert PRO diffractometer.
(b) Diffractometer plane geometry showing incident x-rays, sample
detector and all relevant angles. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25
2.5
Schematic diagram of a symmetric ω/2θ scan geometry of (00l) planes. . . 26
2.6
XRD 1D rocking scans for a sample with two InGaAs buffer layers on
a GaAs substrate. (a) No slit, (b) 1/8º, and (3) 1/32º receiving slits are
used in each scan. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 27
2.7
Schematic for RSM (004) symmetric and (224) asymmetric scans. The
figure shows substrate and films peaks under different scenarios: fully
strained, partly relaxed, fully relaxed and partly relaxed with lattice tilt. . . 29
2.8
(a) Schematic and (b) XRD (224) 2D RSM of SiGe quantum confined
Stark effect modulator sample. The SiGe buffer is totally relaxed, and
the Ge quantum wells are fully strained to the SiGe buffer layers. . . . . . . . 30
xvi
2.9
FEI Tecnai G2 F20 X-TWIN Microscopy at Stanford
Nanocharacterization Laboratory, Stanford University. . . . . . . . . . . . . . . . 31
2.10 Schematic of conventional high resolution TEM (HRTEM). . . . . . . . . . . . 32
2.11 Cross-section TEMs image of relaxed InGaAs buffer layers grown
on GaAs showing clear defects. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 34
2.12 Bright field TEM images of Ge multi-quantum wells, where the Ge
QW is 10 nm thick and the SiGe barrier is 19 nm thick. . . . . . . . . . . . . . . . 35
2.13 Schematic of photoluminescence of a direct band gap semiconductor. . . . . 40
2.14 Schematic of the photoluminescence setup. . . . . . . . . . . . . . . . . . . . . . . . . . 42
3.1
Schematic (a) and SEM image (b) of biaxial tensile strained Ge
sample structures. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 49
3.2
Defects in hetero-epitaxy material growth. . . . . . . . . . . . . . . . . . . . . . . . . . 52
3.3
Schematic (a) and SEM image of cross-section structure of (b) sample A
with high temperature growth method and (c) sample B with low
temperature growth and high temperature annealing method. . . . . . . . . . . . 55
3.4
AFM image of InGaAs buffer layer surface morphology of (a) sample A
with high temperature growth method and (b) sample B with low
temperature growth and high temperature annealing method. . . . . . . . . . . . 56
xvii
3.5
2D reciprocal space mapping of InGaAs buffer layers for two
samples with different growth recipes.. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57
3.6
Cross-section TEMs of InGaAs buffer layers grown on GaAs by two
different methods: (a) high temperature growth and (b) low temperature
growth and high temperature annealing. . . . . . . . . . . . . . . . . . . . . . . . . . . . 58
3.7
Cross-section TEM of sample B InGaAs buffer layers. . . . . . . . . . . . . . . . . 59
3.8
PL intensity spectrum for high temperature growth method (sample A,
blue dashed line) and low temperature growth high temperature
annealing method (sample B, black solid curve) . . . . . . . . . . . . . . . . . . . . . 60
3.9
Cross-section TEM image of two step InGaAs buffer layers on GaAs substrate
with threading dislocations propagation and bending. . . . . . . . . . . . . . . . . . 64
3.10 AFM of In0.3Ga0.7As buffer layers on GaAs without (a) and with
(b) in situ high temperature annealing after last InGaAs layer growth
in MBE system. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 65
3.11 PL spectra for In0.3Ga0.7As and In0.15Ga0.85As buffer layers with
530 ºC annealed sample (solid black line) and 580 ºC annealed
sample (dashed blue line) . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66
xviii
3.12 Schematic illustration of Matthews-Blakeslee model. Two conditions
are shown as (b) totally strained coherent growth and (c) totally
relaxed growth. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 70
3.13 Matthews-Blakeslee critical thickness for Ge at thermal equilibrium. . . . . 71
3.14 Simulated XRD 1D rocking curve for different thickness Ge on
InGaAs buffer layer and GaAs substrate. . . . . . . . . . . . . . . . . . . . . . . . . . . . 74
3.15 XRD rocking curve for InGaAs buffer layers on GaAs with (blue dashed
curve) and without (black solid curve) a strained Ge epitaxial layer. . . . . . 75
3.16 Reciprocal lattice mapping of strained Ge layer and InGaAs buffer
layers on GaAs substrate for both (004) and (224) scans. . . . . . . . . . . . . . . 76
3.17 Raman spectra for bulk Ge (red dotted curve), In0.3Ga0.7As buffer
layers (blue solid curve) and Ge on In0.3Ga0.7As buffer layers
(black dashed curve). . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78
3.18 Raman spectra for different tensile strained Ge layers. . . . . . . . . . . . . . . . . 78
3.19 The tensile strain inside Ge layers with different InGaAs buffer layers.
The black curve, blue triangles and red squares represent the strain
inside Ge calculated from theory, XRD measurement and Raman
spectroscopy, respectively. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 80
xix
3.20 Cross-section TEM image for strained Ge samples. (a) Schematic of
the sample structure, (b) cross-section TEM image of InGaAs buffer
layers and Ge QW, (c) cross-section TEM image of Ge QW, and
(d) high resolution TEM image of Ge QW region. . . . . . . . . . . . . . . . . . . . 82
3.21 Cross section TEM images with indication of selected area for
diffraction images. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 83
3.22 Selected area diffraction (a) at In0.3Ga0.7As and In0.15Ga0.85As buffer
layers as well as zoomed-in images (b) and (c). . . . . . . . . . . . . . . . . . . . . .
85
3.23 Selected area diffraction (a) at Ge quantum well and In0.3Ga0.7As buffer
layer as well as zoomed-in images (b) and (c). . . . . . . . . . . . . . . . . . . . . . . 85
3.24 Low temperature (5 K) PL of different tensile strained Ge layers. . . . . . . . 88
3.25 Temperature dependent PL intensity for different tensile strained Ge. . . . . 88
4.1
SIMS measurement for the calibration sample of 4 GeSn layers grown
directly on GaAs at 200 ºC. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 98
4.2
Sn growth rate vs. Sn evaporate source temperature. . . . . . . . . . . . . . . . . . 98
4.3
XRD (004) rocking curve for three different Sn concentration GeSn
layers on InGaAs buffer layers and GaAs substrate. . . . . . . . . . . . . . . . . . . 99
4.4
SEM images from Ge0.954Sn0.046 (a) and Ge0.93Sn0.07 (b) on In0.1Ga0.9As
buffer layer. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 100
xx
4.5
AFM images of 50 nm Ge0.954Sn0.046 grown on In0.1Ga0.9As buffer layer
at (a) 200 ºC, (b) 150 ºC, and (c) 100 ºC substrate temperature. . . . . . . . . . 102
4.6
AFM images of 50 nm Ge1-xSnx grown on InGaAs buffer layers at
100 ºC substrate temperature with Sn concentrations of (a) 4.6%,
(b) 7.4%, and (c) 9.2%.. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 102
4.7
Schematic (a) and cross section TEM image (b) of sample B-200,
which is Ge0.954Sn0.046 grown on an In0.1Ga0.9As buffer layer on a
GaAs (001) substrate. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 103
4.8
High resolution TEM image of sample B-200 (a) at GeSn and InGaAs
interface and (b) at the center of GeSn epitaxial layer. . . . . . . . . . . . . . . . . 104
4.9
Selected area diffraction pattern for Ge0.954Sn0.046 epitaxial layer. . . . . . . . 105
4.10 XRD 1D rocking curve for Ge0.95Sn0.05 epitaxial layer on three different
InGaAs buffer layers. Inset figure shows the schematic of the sample
layer structure. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 107
4.11 XRD 2D reciprocal lattice mapping for Ge0.95Sn0.05 - In0.18Ga0.82As
buffer layer sample. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 108
4.12 Transmission spectra for 150 nm thick Ge1-xSnx samples grown
on GaAs substrates. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 110
xxi
4.13 Approximate band gap energies of GeSn epitaxial layers with
different Sn compositions. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 110
5.1
Schematic of waveguide coupled GeSi electro-absorption modulator. . . . . 117
5.2
Schematic of on-chip an optical interconnect system with light
source, modulator, and photodetector. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 117
xxii
Chapter 1
Introduction
1.1
On-Chip Interconnection Challenge and Solution
1.1.1 On-Chip Interconnections
Integrated circuits (IC) have been a major driving force to revolutionize
electronic technology in the past few decades. With the success of Moore’s law,
which predicts that the number of transistors in an integrated circuit doubles every 2
years [1], society has benefited from this scaling. Unfortunately, anything with
exponential growth runs into limits and that is now beginning to happen with scaling
of ICs. The increase of clock speed for commercial computers has come to a halt [2].
The main reason is huge power dissipation, especially in the passive interconnect
component, leading to severe heating problems. Although new materials (copper
and low–K dielectric interconnects) and new architectures (multi-core) have been
applied to solve the speed limitation, the most fundamental RC limitation is still the
major bottleneck for the speed of ICs. As shown in figure 1.1, according to the 2007
International Technology Roadmap for Semiconductors (ITRS), with continued
technology node scaling, the relative delay for logic devices and local interconnects
decreases [3]. However, global interconnects, especially global interconnects without
repeaters, show a dramatic increase in delay time (ref). All of these arguments show
that the on-chip interconnect is one of the major challenges for the IC industry.
1
100
Gate Delay
Metal 1
Global W
Repeater
Global WO
Repeater
Relative Delay
10
Global interconnects
1
CMOS device
0.1
Local connects
250
190
130 90
66 45 32
Process Technology Node (nm)
Figure 1.1: Relative delay vs. process technology node from 2007 ITRS [3].
1.1.2 Optical interconnects
Light has been used to transmit signals for several thousands of years. In the
past three decades, photons have been widely used in optical fiber communication
systems, especially in long-distance communications. The fundamental reason for
the optical interconnect advantage is the zero rest mass of a photon, which can
greatly reduce the required energy. More specifically, there are many advantages of
optical fiber communications (OFC).
The first one is the high speed (high
information bandwidth) due to the lack of RC delay limitation, which is the main
constraint for electrical interconnections.
Secondly, photons with different
wavelengths do not interact with each other, to first order.
Thus, wavelength
multiplexing (WDM) technique can be applied in the optical communication channel,
which leads to much higher bandwidth capacity for each physical channel. Thirdly,
the optical fiber is the most transparent media for the signals. This is very critical for
2
long distance signal transmission. Finally, the cost of the silica optical fiber is much
lower than the cost of expensive metal wire. Due to all of these reasons, optical
transmission has replaced electrical transmission in all long haul telecommunication
systems. With this technology improvement, the optical interconnect approaches
have been gradually introduced into even shorter range communication systems. For
example, the board-to-board data communication in modern super computers is
largely dependent on optical fiber communication. Another example is the Light
Peak cable released by Intel Corp.
Optical transceivers and optical fiber are
integrated in this cable, which provides a 10 Gbits/second signal channel. Nowadays,
more interest is focused on how to use optical communication to solve the high
speed interconnection for even shorter distances between chips, and especially onchip.
However, many challenges need to be solved to achieve on-chip optical
interconnects. The first issue is to integrate all the optical devices with the silicon
microelectronic devices on silicon chips. All the discrete macroscopic components
used in optical fiber communication systems have to be redesigned to become onchip devices, and to be used in so-called electronic-photonic integrated circuits
(EPICs). Another challenge is materials used in each device. For traditional optical
components used in OFC, III-V material, such as GaAs and InP, are heavily used due
to their excellent optical properties. However, all the materials and fabrication
techniques in the current IC industry require complete compatibility with Si CMOS
IC technology for the on-chip interconnect solution. Thus, the development of
silicon-based photonics is absolutely necessary.
3
1.2
Key Devices for Si Photonics
Figure 1.2 shows several key Si photonics devices for on-chip
interconnections [4]. The on-chip interconnect system contains an integrated light
source (such as a continuous wave (CW) light), a modulator to transfer the electrical
signal to an optical signal, a waveguide or waveguide device to direct the optical
signal to the destination, and finally a photo-detector to convert the optical signal
back to an electrical signal. Most of these devices are already developed on a Si
platform with high bandwidth capability.
Integrated
Light Source
Waveguide
Modulator
Photo-detector
Figure 1.2: Si photonics key devices, including integrated light source,
modulator, waveguide, and photo-detector [4].
The only missing key device is the integrated CMOS compatible electrically
pumped laser source. Although many different methods have been investigated,
there is still no perfect solution for this device. In a first approach, III-V materials
are either epitaxially grown on or bonded onto the Si substrate [5, 6, 7]. These lasers
work at convenient wavelengths for optical communications and have high
efficiency. Unfortunately, the integration of III-V materials, which are dopants in Si,
creates major challenges and greatly limits their application. The second approach is
4
based on nonlinear effects in silicon. This type of optically pumped silicon laser is
based on stimulated Raman scattering [8]. However, the requirement of optical
injection and their low efficiency make a Si Raman lasers unsuitable for on-chip
integrated light sources. The third approach is to use a nano-structure material to
modify the band structure to achieve efficient light emission.
Porous Si, Si
nanostructures, and SiGe nanostructures have been studied [9]. However, the weak
optical gain of these approaches limits the lasing process in this type of material.
The fourth approach is to use erbium doped material as the gain medium. Erbium
doped Si, Si oxide, and Si nitride have been studied, and a high Q cavity laser has
been demonstrated [10, 11].
The limitation of erbium concentration, energy
conversion, and current injection prevents useful application of any of these
approaches as efficient light source devices.
From the material point of view, in order to achieve a CMOS compatible
light source, a group IV material (such as Si, Ge, or Sn) must be used. In order to
have an efficient light source, a direct band gap semiconductor is preferred. At least,
a local minimum at the Γ point of the conduction band is required to accumulate
electrons and achieve efficient radiative recombination. This requirement makes Si
nearly impossible to be an efficient light source. Fortunately, germanium, which is a
group IV material, has a local minimum at the Γ point of the conduction band. More
attractively is that the lowest energy point in the Ge conduction band is at the L point,
which is only 0.14 eV lower than the lowest energy at the Γ point at room
temperature. Therefore, Ge has the potential to be engineered to become a direct
band gap material and used as an on-chip integrated light source.
5
1.3
Ge Band Structure Engineering
As discussed in the previous section, Ge is the most interesting group IV
material for the light emitting process. However, achieving direct band gap in Ge
and improving the Ge light emitting efficiency are still huge challenges. Three
methods have been proposed to tune the Ge band structure to achieve a direct band
gap material, and they are discussed separately in the following sections.
1.3.1 Heavily n-Type Doping Germanium
N+ doped Germanium is capable of behaving like a direct band gap material
as discovered by the MIT groups [12]. In this approach, Ge was grown on Si in a
CVD system, and 0.2% tensile strain was achieved due to the different thermal
expansion coefficient of germanium compared to that of silicon. Different in-situ
doping and ion implantation techniques were applied to provide n+ doping in the Ge
layer. Room temperature light emitting diodes (LEDs) have been demonstrated with
an internal quantum efficiency of Ge LEDs at 10-3 [13-15].
Higher
electroluminescence (EL) is obtained at higher temperatures. This is because, with
heavy n-type doping, the excess electrons in the conduction band first fill the indirect
L band and then fill the direct Γ band of Ge. With higher temperature, there is more
possibility for electrons to thermally distribute in the Γ valley of the conduction band.
This approach itself has some difficulty to achieve an electrically pump laser due to
the high free carrier absorption as well as the low internal quantum efficiency.
Although the n+ doped Ge shows carrier radiative recombination at the direct band
gap, it’s not a truly direct band semiconductor.
6
1.3.2 Biaxial Tensile Strained Germanium
The second method to achieve direct band gap germanium is in-plane biaxial
tensile strain. In 1996, Fischetti and Laux from IBM first presented theoretical work
on strained Ge and predicted that, with about 1.75% biaxial tensile strain, Ge
becomes a direct band gap material, as shown in figure 1.3 [16]. The advantage of
direct band gap germanium is that most electrons stay in the Γ valley of the
conduction band.
Thus, the radiative recombination through the direct band
transitions is more efficient, and direct band gap germanium is preferable for
optoelectronic devices.
Figure 1.3: Calculated valence and conduction band shifts at various
symmetry points in Ge as a function of in-plane biaxial strain [16].
7
Tensile strain in Ge is not only beneficial for optical devices, but also good
for electrical devices.
Under in-plane biaxial tensile strain, both the electron
mobility and the hole mobility are greatly increased, as shown in figure 1.4.
Although tensile strained Ge has many advantages, how to achieve high-quality
tensile strained Ge is really a challenge and is discussed in detail in Chapter 3.
Figure 1.4: Calculated electron and hole mobilities of (100) Ge as a function
of biaxial tensile strain [17].
8
1.3.3 GeSn Alloy
The third method to achieve direct band gap Ge is to grow a GeSn alloy.
Many theoretical studies have predicted that a GeSn alloy can become a direct band
gap material [18]. However, the Sn concentration to achieve a direct bandgap is very
poorly defined and predictions vary from 6% to 20%, depending on the theoretical
model. A typical result from the Caltech group in 1997 is shown in figure 1.5 [19].
Figure 1.5: Calculation and experimental result of band crossing of GeSn
alloy as a function of Sn concentration [19].
In the study of these alloys, problems arise from the difficulty of material
growth. The most critical limitation is the low thermodynamic solid solubility of Sn
in Ge crystal, which is less than 1% as shown in figure 1.6. MBE growth has been
reported in previous work [20, 21]; however, the best quality samples with around
6.3% Sn concentration in Ge without Sn precipitation. Others have produced films
9
with much higher Sn concentrations. However, the material quality is questionable
due to the clear Sn phase segregation, as shown in figure 1.7 [22]. In Chapter 4, we
Temperature (ºC)
discuss our work on GeSn alloy growth and characterization.
Ge
Atomic percentage of Sn
Sn
Figure 1.6: Ge-Sn phase diagram, showing less than 1% solid solubility of Sn
in Ge [22].
Figure 1.7: SEM image of Ge0.78Sn0.22 alloy grown at 200 °C with clear Sn
phase segregation [23].
10
Ge1-xSnx alloys have many other good properties as well. Simulations have
shown that the electron mobility can exceed 105 cm2/V·s [24]. Also the increased
lattice constant (ranging from 5.66 Å to greater than 5.90 Å) of GeSn provides a
large lattice constant for strained Ge as well. All of these properties make GeSn
alloy an interesting research topic and a potential material for a future CMOS
compatible light source.
1.4
Organization
This thesis describes the growth, material characterization, and optical
properties of biaxial tensile strained Ge and GeSn alloys grown using molecular
beam epitaxy (MBE). Chapter 2 introduces the background for this thesis, including
the MBE material growth system and the special material and optical
characterization techniques.
Chapter 3 presents the results on in-plane biaxial tensile strained Ge epitaxial
layers.
It starts with the InGaAs relaxed thick buffer layer growth method
optimization. Then, the material properties of the coherently grown strained Ge
layers are examined. At the end of the chapter, the photoluminescence (PL) behavior
of strained Ge is studied, showing that a direct band gap semiconductor is achieved.
Chapter 4 describes the growth of GeSn alloys by low temperature MBE
growth. The growth method is optimized to prevent Sn phase segregation and
11
precipitation in the Ge crystal. The material quality is characterized, and the band
gap energy is examined as well.
Finally, Chapter 5 summarizes the contributions made in this dissertation and
makes recommendations for future material, device and system work needed in this
field.
References:
[1] G. Moore, "Cramming more components onto integrated circuits," Electronics 38,
144117 (1965).
[2] D. A. Muller, "A sound barrier for silicon?," Nature Mater. 4, 645-647 (2005).
[3] Semiconductor Industry Association, 2007 International technology Roadmap for
Semiconductors, (2007).
[4] www.intel.com
[5] B. Kunert, S. Reinhard, J. Koch, M. Lampalzer, K. Volz, and W. Stolz, “First
demonstration of electrical injection lasing in the novel dilute nitride Ga(NAsP)/GaPmaterial system,” Phys. Stat. Sol. (c) 3, No. 3, 614 (2006).
[6] Z. Mi, J. Yang, P. Bhattacharya, D. L. Huffaker, “Self-organised quantum dots as
dislocation filters: the case of GaAs-based lasers on silicon,” Electronics Letters,
vol.42, no.2, 121, (2006).
12
[7] A. W. Fang, H. Park, O. Cohen, R. Jones, M. J. Paniccia, and J. E. Bowers,
“Electrically pumped hybrid AlGaInAs-silicon evanescent laser,” Optics Express, Vol.
14, Issue 20, 9203 (2006).
[8] O. Boyraz and B. Jalali, "Demonstration of a silicon Raman laser," Opt. Express,
12 (21), 5269 (2004).
[9] N. Koshida and H. Koyama, "Visible electroluminescence from porous Si," Appl.
Phys. Lett. 60 (3), 347 (1992).
[10] M. Makarova, V. Sih, J. Warga, R. Li, L. D. Negro, and J. Vuckovic, "Enhanced
light emission in photonic crystal nanocavities with Erbium-doped silicon
nanocrystals," Appl. Phys. Lett. 92 (16), 161,107 (2008).
[11] A. Polman, B. Min, J. Kalkman, T. J. Kippenberg, and K. J. Vahala,
"Ultralowthreshold erbium-implanted toroidal microlaser on silicon," Appl. Phys. Lett.
84 (7), 1037 (2004).
[12] J. Liu, X. Sun, D. Pan, X. Wang, L. C. Kimerling, T. L. Koch, and J. Michel,
“Tensile-strained, n-type Ge as a gain medium for monolithic laser integration on Si,”
Optics Express, Vol. 15, No. 18, 11272, (2007).
[13] X. Sun, J. Liu, L. C. Kimerling, and J. Michel, “Room-temperature direct
bandgap electroluminesence from Ge-on-Si light-emitting diodes,” Optics Letters, Vol.
34, No. 8, (2009).
13
[14] S. Cheng, J. Lu, G. Shambat, H. Yu, K. Saraswat, J. Vuckovic, and Y. Nishi,
“Room temperature 1.6 μm electroluminescence from Ge light emitting diode on Si
substrate,” Optics Express, Vol. 17, No. 12, 10019 (2009).
[15] J. Liu, X. Sun, R. Camacho-Aguilera, L. C. Kimerling, and J. Michel, “Ge-on-Si
laser operating at room temperature,” Optics Letters, Vol. 35, Issue 5, pp. 679 (2010).
[16] M.V. Fischetti, and S.E. Laux, “Band structure, deformation potentials, and
carrier mobility in strained Si, Ge, and SiGe alloys,” J. Appl. Phys. 80, 2234 (1996)
[17] Y. Bai, K. E. Lee, C. Cheng, M. L. Lee, and E.A. Fitzgerald, “Growth of highly
tensile-strained Ge on relaxed InxGa1-xAs by metal-organic chemical vapor
deposition”, J. Applied Physics 104, 084518 (2008).
[18] J. Kouvetakis, J. Menendez, and A.V.G. Chizmeshya, “Tin-based group IV
semiconductors: new platforms for opto- and microelectronics on silicon,” Ann. Rev.
of Mater. Res. 36, 497 (2006).
[19] G. He and H. A. Atwater, “Interband transitions in SnxGe1-x alloys,” Physical
Review Letters, Vo. 79, 10, 1937 (1997).
[20] G. He and H.A. Atwater, “Synthesis of epitaxial SnxGe1-x alloy films by ionassisted molecular beam epitaxy,” Appl. Phys. Lett. 68 (5), (1996).
[21] Y. Shimura, N. Tsutsui, O. Nakatsuka, A. Sakai, S. Zaima, “Low temperature
growth of Ge1-xSnx buffer layers for tensile–strained Ge layers,” Thin Solid Films, 518,
(2010).
[22] http://www.crct.polymtl.ca/fact/phase_diagram.php?file=Ge-Sn.jpg&dir=SGTE
14
[23] Gang He, Ph.D dissertation, (1997).
[24] J. D. Sau, and M. L. Cohen, “Possibility of increased mobility in Ge-Sn alloy
system”, Physical Review B 75, 045208 (2007).
2
15
Chapter 2
Growth Equipment and
Characterization Techniques
Background information about material growth equipment is first introduced
in this chapter. Following that, many material quality characterization tools are
discussed, emphasizing their special applications for ultra-thin strained Ge and GeSn
epitaxial materials.
2.1 Molecular Beam Epitaxy (MBE)
2.1.1 Introduction
Epitaxial growth is a process in which layers of materials are deposited on a
substrate while conserving the substrate’s crystalline structure. There exist many
types of epitaxial growth techniques, ranging from liquid-phase epitaxy (LPE) to
vapor-phase epitaxy, such as metal-organic vapor-phase epitaxy (MOVPE) and
chemical vapor deposition (CVD).
Compared to all these epitaxial growth
techniques, the MBE growth technique shows superiority because it allows a wider
range of materials, precise growth control, the growth is farthest from thermal
equilibrium conditions, lack of memory effects as well as other advantages. The
16
development of MBE has been one of the most important achievements for the
advances of electronics and optoelectronics.
Nowadays, MBE systems have
achieved many improvements that provide much higher throughput and much greater
economic efficiency for large volume production.
2.1.2 MBE System
In this work, we use two coupled MBE systems. The first one is a group IIIV MBE system, which mainly grows GaAs, InGaAs, AlGaAs, and GaP materials
with silicon and beryllium as n and p-type dopants, respectively. The second system
is a group IV MBE system, which mainly grows silicon, germanium, and tin
materials with arsenic and boron as n and p-type dopants, respectively. The two
MBE systems are coupled together through a transfer tube which is under 1x10-9
Torr vacuum.
Substrate
RHEED gun Cells
Figure 2.1: (a) A schematic drawing, and (b) a picture of an MBE chamber
from side view, showing material sources, beam flux gauge, and the RHEED system.
17
Figure 2.1 shows a side view of an MBE growth chamber.
The main
component of an MBE system is an ultra-high vacuum (UHV) chamber, which has a
background pressure of 1x10-10 Torr. These chambers are pumped down to UHV by
using a combination of different vacuum pumps, including turbo pumps with
mechanical pumps as fore-line backing pumps, cryo-pumps, and ion pumps. A
liquid nitrogen filled cryo-shroud around the substrate holder is used to further pump
the region near the substrate and reduce memory effects, as discussed later in this
section.
Atomic or molecular beams are generated by heating the materials in various
individual effusion cells. The types of cells are determined by the specific material
properties. Typically Knudsen effusion cells are used for Ga, Al, In, Ge and Sn.
Effusion cells supply each constituent (e.g. gallium) atoms or molecules from an
ultra-pure (99.999% purity) ingot contained in an individually heated ceramic
crucible (typically pyrolytic boron nitride, PBN). The vapor pressure of each species
is controlled by setting the temperature of the effusion cell. The material growth rate
is calibrated and monitored by a beam-flux ion gauge at the sample region.
Mechanical shutters driven from outside of the vacuum chamber are used to switch
the beams on and off.
A valved-cracker is used to supply As2 flux; the cracker includes two parts: a
sublimator generates As4 vapor at equilibrium pressure, and a cracking heater which
18
cracks As4 to As2. The overall flux of As2 is controlled by a valve placed at the
cracking zone that produces precisely an arsenic flux into the sample region.
An electron-beam evaporator (e-gun) is used for silicon growth.
In the
conventional silicon MBE system, a Knudsen effusion cell is used to provide silicon
flux. However, since the silicon vapor pressure is extremely low even at relatively
high temperatures, the conventional effusion cell can only generate sufficient silicon
flux for doping purposes. On the other hand, the e-gun Si source only heats up a
small region of the tip of a Si bar and forms a mm-scale crucible which gives a
relatively high silicon flux and hence increases the growth rate.
In our MBE machines, all the cells (except for the Si e-beam cells) and
shutters are controlled by computers and can be programmed to produce quite
complex multi-layers, super-lattices and hetero-junction structures.
A thin, crystalline substrate wafer is mounted on a substrate heater, such that
it can be brought to face the source crucibles used to evaporate the constituent atoms
or molecules. In-situ reflection high energy electron diffraction (RHEED) is built
into the MBE system to check the film quality during growth. The high energy
electron beam is incident on the substrate at a glancing angle. The diffraction from
the sample surface crystal structure creates a diffraction pattern on a phosphor-coated
window. By observing the electron beam diffraction pattern, we can check the oxide
desorption from the wafer surface, the 2D smooth single crystal epitaxial growth,
and the GaAs phase orientation etc.
19
2.1.3 Material Growth Mechanism
For typical III-V material growth, such as GaAs or InGaAs, beams of atoms
or molecules in an ultrahigh vacuum environment are incident upon a heated
substrate crystal and form a crystalline layer epitaxially without any other collisions.
At typical GaAs growth temperatures, the sticking coefficient of arsenic is much less
than sticking coefficient of gallium. Thus, we usually use about 15 times higher
arsenic beam-flux-equivalent-pressure (BEP) (overpressure) to Ga BEP during the
growth. The total growth rate is directly proportional to the sum of the group III
material beam fluxes, and the composition of the group III material is determined by
the ratios of the individual constituent beam fluxes.
For group IV materials, such as Si, Ge and Sn alloys, the growth is relatively
simple due to the unity sticking coefficient to each element. However, the substrate
temperature needs to be carefully chosen in order to prevent phase segregation and
achieve good crystal quality.
2.1.4 Advantages of MBE Systems
Several advantages of MBE systems are discussed in this section. The UHV
condition (~ 1x10-10 Torr) in the MBE system provides a low impurity/contamination
epitaxial growth environment. Furthermore, the low pressures also lead to long
20
mean free paths between collisions. As a result, there are no precursor interactions.
The atoms only react when they hit the substrate.
MBE systems decouple the substrate temperature and growth rate. Unlike
normal MOCVD and CVD systems, the material growth rate in MBE systems almost
totally depends on the constituent flux, which is controlled by each source
temperature. The substrate temperature can be separately controlled to optimize the
crystal quality, eliminate phase segregation, and control the phase of the epitaxial
crystal.
MBE systems can accurately control the layer structure with atomicly abrupt
layer changes. Under UHV conditions with liquid nitrogen cooled cryo-shroud and
chilled source cooling panel, the constituent beams make a single pass through the
chamber and hit the substrate or condense on the cold chamber walls.
This
eliminates any memory effects of what is previously grown. Also, the growth rate of
each material can be accurately controlled by the cell temperature.
Thus, the
material composition and doping level can be rapidly changed, producing crystalline
interfaces that are almost atomically abrupt. This accurate control allows production
of very complex structures by MBE, including quantum well devices, superlattices,
quantum cascade lasers, etc.
MBE systems use the highest purity source material, since all the materials
are elemental or molecular sources with more than 99.999% purity.
21
RHEED characterization in the MBE systems allows us to examine the
material properties in-situ during the growth and dynamically change our growth
recipes.
Because of all these advantages, MBE systems are the best tool for basic
materials research. MBE-grown device structures have the highest crystal quality
and are very close to the designed structures.
2.2 Characterization Techniques
Many types of characterization tools are required to measure both the
material as well as optical properties of our samples. In this section, we briefly
discuss the basic information obtained and the working principles for each of them.
More specific results are discussed later in each chapter.
2.2.1 Scanning Electron Microscopy (SEM)
An SEM is one of the most frequently used tools in semiconductor materials
and device research. In an SEM, a fine beam of electrons is scanned across the
surface of a specimen and a detector monitors the intensity of secondary electron
emission from the specimen. A spot is displayed on a screen; the spot is scanned in
synchronism with the scanning electron beam on the specimen, and the brightness of
the spot is controlled by the detected signal amplitude. When the intensity of the
22
emitted secondary signal changes across the specimen, the same contrast pattern is
shown in the SEM image. In this study, we use an FEI XL30 Sirion SEM with about
3nm resolution in the ultra-high resolution mode with normally 5 KeV electron beam
energy. Two typical SEM images are shown in figure 2.2 as the cross section image
and top view image of our GaAs nano-cone solar cell.
(a) (b)
Figure 2.2: Typical (a) tilted cross section SEM image and (b) top view SEM
image of GaAs nano-structured solar cells.
2.2.2 Scanning Probe Microscopy (SPM)
A scanning Probe Microscopy (SPM) (or, often referred to as atomic force
microscopy (AFM)) is a surface morphology characterization tool. It utilizes micromachined cantilever probes with sharp tips to scan the sample surface morphology.
In our study, we mainly use a Veeco multimode scanning probe microscope
equipped with a Quadrex Nanoscope IIIA controller. We use the tapping mode to do
the scan since our sample is relatively smooth and hard. In the tapping mode, the
cantilever is driven by a small piezoelectric transducer to oscillate up and down near
23
its resonance frequency. Due to the Van der Waals force and other forces acting on
the cantilever when the tip comes close to the surface, the amplitude of this
oscillation decreases. A laser is incident on the cantilever, and the light reflected to a
photodetector is a measure of the oscillation amplitude of the cantilever.
An
electronic piezoelectric actuator servo is used to adjust the distance between the
cantilever tip and the sample in order to maintain the cantilever oscillation amplitude.
Thus, a tapping AFM image is formed as the cantilever is scanned over the sample
surface. The root mean square (RMS) surface roughness is always calculated to
indicate whether a smooth crystal surface is obtained or not. A typical AFM scan
image of a relaxed InGaAs buffer layer on GaAs substrate is shown in figure 2.3.
With these AFM systems, we can scan a 1 μm x 1 μm or a 10 μm x 10 μm region
with almost atomic resolution in the vertical direction.
Figure 2.3: A typical AFM image of a relaxed InGaAs buffer layer on a GaAs
substrate.
24
2.2.3 X-Ray Diffraction (XRD)
High resolution XRD (HR-XRD) is one of the most useful techniques for
analyzing epitaxial films, such as the crystal lattice constant, film thickness, strain
relaxation and alloy composition. One of the major advantages of XRD is that we
can determine compositional information non-destructively without any complex
sample preparation.
For our study, we normally use a Phillips X'pert PRO
diffractometer as shown in figure 2.4 (a), and the diffractometer geometry is shown
in figure 2.4 (b).
(b) (a) Figure 2.4: (a) A photograph of Phillips X'Pert PRO diffractometer. (b)
Diffractometer plane geometry showing incident x-rays, sample detector and all
relevant angles [1].
The
most
common
HR-XRD
measurement
for
epitaxially
grown
semiconductors with cubic crystal lattices symmetry is a one-dimensional (1D) ω/2θ
rocking curve scan of the diffraction from (004) planes. These symmetric ω/2θ scans
are performed in order to determine the out-of-plane lattice constants of the films.
25
As shown in figure 2.5 (a), the symmetric (004) XRD reflections come from planes
that are parallel to the sample surface. The out-of-plane lattice spacing can be
calculated from Braggs Law:
nλ = 2dhklsin(θ)
(2-1)
where n is the diffraction order, λ is the wavelength of the incident x-rays, θ is the
scattering angle and dhkl is the spacing between the (hkl) planes. For cubic materials,
dhkl is related to the lattice constant of the crystal, a, by:
2
d hkl =
a2
h2 + k 2 + l 2
(2-2)
Figure 2.6 (b) shows a typical XRD 1D ω/2θ rocking curve scan around the
(004) diffraction point of a fully relaxed 250 nm thick Ge epitaxial layer on a Si (001)
substrate. The two diffraction peaks are related to the Si substrate and Ge epitaxial
layer. The out-of-plane lattice constant can be calculated from the diffraction angle.
For the GaAs substrate, the measured lattice constant is 5.653923 ± 0.00129 Å,
which is very close to the lattice constant of GaAs from the literature, 5.6533 Å with
~ 0.01% accuracy [2].
Si 4
10
Counts
Ge
2
10
0
10
(a) 32.5
33
33.5
34
34.5
Omega (degree)
(b) 35
Figure 2.5: Schematic diagram of a symmetric ω/2θ scan geometry of (00l) planes.
26
counts/s
10M
counts/s
10M
counts/s
1M
S
S
1M
1M
100K
100K
10K
10K
1K
1K
100
100
10
10
1
1
S
100K
10K
L
L
1K
0.1
31.0
100
31.5
32.0
32.5
33.0
33.5
34.0
Omega/2Theta (°)
0.1
31.0
10
1
31.5
32.0
(a) 32.5
(b) 33.0
33.5
34.0
Omega/2Theta (°)
0.1
31.0
31.5
32.0
32.5
33.0
33.5
34.0
Omega/2Theta (°)
(c) Figure 2.6: XRD 1D rocking scans for a sample with two InGaAs buffer
layers on a GaAs substrate. (a) No slit, (b) 1/8º, and (3) 1/32º receiving slits are used
in each scan.
A small receiving slit is used at the receiving arm to increase the angle
resolution. Three different receiving slits (none, 1/8º, and 1/32º) are used for a 1D
XRD rocking curve of the same sample with two InGaAs buffer layers on a GaAs
substrate, as shown in figure 2.6. Scan results clearly show that the angle resolution
is greatly improved by using a 1/8º receiving slit. However, with a 1/32º receiving
slit, the angle resolution does not change much, but the signal to noise ratio
decreases dramatically. Thus, most of the XRD 1D scans in this work are done with
1/8º, 1/16º, and 1/32º receiving slits.
The full width at half maximum of the
diffraction peaks is also determined by the epitaxial layer thickness. As discussed
later in section 3.3.3, in order to obtain a high enough signal-to-noise ratio, an
epitaxial layer about 10 nm thick is needed.
27
Although the XRD (004) 1D rocking curve is easy and quick to measure, it
provides quite useful information, however, the symmetry scan doesn’t provide
information about the in-plane lattice constant. Even more important, for alloy
crystalline materials, such as GeSn or InGaAs, the XRD 1D rocking curve scan is
determined by both the component composition and the lattice relaxation. Due to
these problems, two-dimensional (2D) asymmetric scans of diffraction patterns are
extremely important. This method is called XRD 2D reciprocal space mapping
(RSM) which is capable of unveiling the relative lattice tilt, the material composition,
and an estimate of the residual strain in epitaxial layers compared to the substrate [3].
Two types of XRD 2D RSM scan, symmetric (004) scans and asymmetric (224)
scans, are generally used in material characterization. The (004) scans are sensitive
for estimating the lattice tilt of layers to the substrate, and (224) scans are sensitive
for characterizing the relaxation and residual strain of the buffer layers. In the 2D
RSM, the out-of-plane (00l) direction is plotted as the y-axis and the in-plane (hk0)
direction is plotted as the x-axis. Figure 2.7 shows the RSM scheme of a substrate
(asub) and fully strained, partly relaxed and fully relaxed epitaxial layers with and
without lattice tilt. [4].
28
Figure 2.7: Schematic for RSM (004) symmetric and (224) asymmetric scans.
The figure shows substrate and films peaks under different scenarios: fully strained,
partly relaxed, fully relaxed and partly relaxed with lattice tilt [4].
An example of RSM around (224) for a Ge/SiGe quantum confined Stark
effect modulator is shown in figure 2.8. The thick SiGe buffer layer is first grown on
a Si substrate. As shown in the RSM plot, the diffraction points from the Si substrate
and SiGe buffer layers are aligned with the original point, which means that the thick
SiGe buffer layer is totally relaxed. Following that, 10 pairs of quantum wells (QW)
are grown on top of SiGe buffer layer. Although the Ge QWs and SiGe barriers have
29
different lattice constants, the RSM shows a vertical line of diffraction peaks, which
means the in-plane lattice constant is the same. This proves that the QWs/barriers
are fully strained to the totally relaxed SiGe buffer layer.
Si substrate
Quantum
well satellite
peaks
SiGe buffer
layer
Silicon Substrate
(b) (a) Figure 2.8: (a) Schematic and (b) XRD (224) 2D RSM of SiGe quantum
confined Stark effect modulator sample. The SiGe buffer is totally relaxed, and the
Ge quantum wells are fully strained to the SiGe buffer layers.
30
2.2.4 Transmission Electron Microscopy (TEM)
In this work, several different Transmission Electron Microscopy (TEM)
techniques are used as powerful tools in the structural characterization of Ge/InGaAs,
Si/Ge and GeSn structures at the atomic level. Cross-section TEM of the InGaAs
buffer layer interface is used to map the local strain and the distributions of misfit
dislocations and threading dislocations. Cross-section TEM also characterizes the
interface smoothness of Ge QWs as well as the crystalline quality of strained Ge
layers. Cross-section TEM also provides single crystal quality information, such as
Sn precipitation or phase segregation in a GeSn alloy. The diffraction pattern is also
used to check the crystalline quality and the coherent growth of strained Ge on
InGaAs buffer layers.
Figure 2.9: FEI Tecnai G2 F20 X-TWIN Microscopy at Stanford
Nanocharacterization Laboratory, Stanford University [5].
31
A few TEM systems are used in this work, and only one of these systems is
introduced in this section.
This TEM system is a FEI Tecnai G2 X-TWIN
Microscope with EDS for conventional bright field and dark field (BF/DF) TEM as
well as STEM, as shown in Figure 2.9. With a field emission gun (FEG), the FEI
Titan 300 kV FEG microscope is designed for both high resolution TEM/STEM and
analytical microscopy. The energy spread is 0.7 eV; the point resolution is 0.2 nm;
the HR-STEM resolution is 0.14 nm.
Figure 2.10: Schematic of conventional high resolution TEM (HRTEM) [6]
Figure 2.10 shows the schematic of a conventional HR-TEM.
The
transmission electron microscope uses a high-energy electron beam transmitted
through a very thin sample to image and analyze the microstructure of materials with
atomic scale resolution. The electrons are focused with electromagnetic lenses, and
the image is observed on a fluorescent screen, or recorded on film or a digital
cameras. Because a very thin sample still contains many layers of atoms, one does
not usually observe an image for each individual atom. Instead, the high resolution
32
imaging mode of the microscope images the crystal lattice of a material as an
interference pattern between the transmitted and diffracted electron beams. This
allows us to observe planar and linear defects, grain boundaries, interfaces etc. with
atomic scale resolution [7]. Because of the high resolution, HR-TEM is an important
tool for studying nanoscale properties of crystalline materials.
The BF/DF imaging modes of the microscope, which operate at intermediate
magnification, combined with electron diffraction, are also invaluable for giving
information about the morphology, crystal phases, and defects in a material.
The TEM is also capable of forming a focused electron probe, as small as 20
Å, which can be positioned on very fine features in the sample for micro-diffraction
information or analysis of x-rays for compositional information. The latter is the
same signal as used for energy-dispersive X-ray spectroscopy (EDAX) in SEM,
where the resolution is on the order of one micro-meter due to the electron beam
spreading in the bulk sample. The spatial resolution for this compositional analysis
in TEM is much higher and is similar to the probe size because the sample is very
thin. Conversely, the signal is much smaller and therefore less quantitative.
The most important information from a TEM image comes from different
types of intensity contrast between two adjacent areas. The basic reason for contrast
is the wave nature of electrons, causing amplitude and phase changes as electrons
pass through the specimen. Thus, there are mainly two mechanisms of contrast:
amplitude and phase contrast. Furthermore, amplitude contrast has two types of
contrast: diffraction contrast and mass thickness contrast.
33
Diffraction contrast is generally seen in crystalline samples with strong
diffraction patterns. It is a special form of amplitude contrast because scattering
occurs at special Bragg angles. To get strong diffraction contrast, we tilt specimen to
select a diffraction beam to form a dark field TEM image. Diffraction contrast is
generally used at low or intermediate magnification to image defects like
dislocations in crystals. Figure 2.11 shows an example of defects in InGaAs buffer
layers on a GaAs substrate. Because the diffraction contrast is a function of the
Bragg diffraction angle, the image strongly depends on the sample tilt. Thus, as the
sample is tilted, the image contrast at dislocations changes as well. In the extreme
case, if the sample is aligned correctly, some of the dislocations can become totally
invisible in the image.
200 nm
Figure 2.11: Cross-section TEMs image of relaxed InGaAs buffer layers
grown on GaAs showing clear defects.
34
Mass thickness contrast is observed in samples with varying thicknesses or
atom composition. Mass thickness contrast arises from incoherent elastic scattering
(Rutherford scattering) of electrons. Rutherford scattering is a strong function of
atomic number, i.e., the mass or the density as well as the thickness of the specimen.
The mass thickness contrast can be differentiated from the diffraction contrast by
tilting the sample about its axis, because the mass thickness contrast doesn’t change
much during the tilt. An example of mass thickness contrast is shown in figure 2.12
for a Ge/SiGe multi-QW sample. Another diffraction effect is related to thickness
variations in the sample causing periodic thickness fringes in the TEM image. The
thickness fringes arise due to multiple diffraction events occurring off a set of planes
through the sample thickness causing a periodic variation in the intensity of the beam
across the thickness gradient.
These thickness fringes are an artificial effect
occurring during the sample preparation.
10nm 19nm Figure 2.12: Bright field TEM images of Ge multi-quantum wells, where the
Ge QW is 10 nm thick and the SiGe barrier is 19 nm thick.
35
Phase contrast arises due to differences in the phases of the electron waves
scattered through a thin specimen. As a comparison, a phase contrast image requires
a selection of more than one beam, while amplitude contrast mechanisms require
selection of a single beam using an objective aperture. A lattice structure phase
contrast image can be obtained at the image plane from the interference between the
transmitted and diffracted beams.
This method is generally used to obtain
information regarding crystal plane spacing, structures of dislocations, etc.
2.2.5 Raman Spectroscopy
Raman spectroscopy is already widely used for stress measurement in silicon
and other materials [8]. In our study, Raman spectroscopy is a very important and
efficient method to measure the strain in thin Ge layers. We used several Raman
spectroscopy systems to measure our samples. Most of the Raman spectroscopy
discussed in this thesis is taken from a Renishaw RM Series System 1000 Raman
spectrometer system. A 50mW Ar+ laser at 514nm is used as the pump laser source.
The vibrations of a crystal are described not in terms of the vibrations of
individual atoms but in terms of collective motions in the form of waves, called
lattice vibrations or phonons. Raman spectroscopy measures this phonon energy,
which is related to the crystal lattice constant and strain in the crystal, even when the
crystal layer is only 10 nm thick. When monochromatic light of frequency ωi is
incident on a crystal, it induces an electric moment, P, described by [9]:
36
(2-3)
where, ε0 is the vacuum permittivity, χ0 is the first order approximation of
susceptibility, E0 is the electric field, ki is the wave vector in the direction of incident
light, and ω is the light frequency. The vibration amplitude at position r is described
by Qj which is the normal coordinate of the vibration, and Aj is a constant. In this
equation, the first term, with frequency component ωi, is called Rayleigh scattering,
and the second and third terms with frequencies ωi + ωj and ωi - ωj are called antiStokes and Stokes Raman scattering, respectively.
When considering the
polarization of incident and reflected light, the crystal orientation, the setup geometry,
and the Raman tensor, for a back scattering setup from a (001) cubic crystal surface,
only longitudinal optical (LO) phonons provide an active Raman mode.
When the crystal is under strain, the lattice constant changes and
consequently phonon energy changes. The change of the Raman shift under the
biaxial stresses, σxx and σyy, is given by [9, 10]
Δω = [(pS12 + q(S11 + S12)](σxx + σyy) / (2ωj)
(2-4)
where ωj is the frequency of the Raman shift without biaxial stress, Sij are the
elastic compliance tensor elements of Ge, and p and q are phonon deformation
potentials for Ge. In the case of biaxial tensile strain, where σxx = σyy, equation (2-4)
can be simplified as Δω = bε‖ , where ε‖ is the biaxial tensile strain. As calibrated
by Dr. Y. Fang et al. [11], for germanium, the Raman peak is at 300 cm-1
wavenumbers, b is equal to -415±40 cm-1. Thus, by measuring the change of the
37
germanium Raman peak shift, we can calibrate the strain in a Ge layer with an
accuracy of 0.05%.
2.2.6 Secondary Ion Mass Spectrometry (SIMS)
SIMS is a characterization technique that detects material composition with
very high accuracy. SIMS uses an ion beam to strike the sample. The secondary
ions formed during the sputtering are extracted and analyzed using a mass
spectrometer. The concentration of these secondary ions can range from matrix
levels down to 1 ppm. Thus, SIMS is widely used in semiconductor research to
determine the material composition, low level dopant concentrations and impurity
levels. SIMS is also used with ion beam sputtering to provide elemental depth
profiles with about 1 nm depth resolution.
In our study, the SIMS data were
acquired from a Cameca SC-Ultra SIMS system at IBM and a Physical Electronics
Model 6600 SIMS system at Charles Evans and Associates.
38
2.2.7 Photoluminescence (PL)
Luminescence is the term used to describe the emission of radiation from a
solid when it is supplied with energy [ref: Vincent_gambin.pdf86]. The different
types of luminescence are distinguished based on the method of excitation used to
create excess electron hole pairs. For example, PL arises from the absorption of
photons, cathodoluminescence (CL) is from bombardment with a beam of electrons,
while electroluminescence (EL) results from the injection of carriers across a PN
junction.
PL has been the most important characterization technique, and is the most
sensitive method to determine the optical properties of direct bandgap materials (for
example, GaAs). Good crystal quality indicated by SEM, AFM, XRD, and/or TEM
does not guarantee good optical properties, which can be characterized by PL. Even
more importantly, PL results and laser performance have a very strong positive
correlation [13, 14].
Thus, in many laser materials growth optimizations, PL
optimization is the most useful method to develop the best laser materials. PL is also
important to determination whether a semiconductor material is a direct band gap
material or not. PL from a direct band gap material is very bright. However, PL
from an indirect bandgap material is normally too weak to be detected, as described
later in this section.
39
Figure 2.13: Schematic of photoluminescence of a direct band gap
semiconductor.
In a PL measurement of a direct band gap material, the pump source photon
energy is larger than the band gap energy.
The absorption of photons creates
electrons in the conduction band and holes in the valence band. Both electrons and
holes have higher energy than the band minimum energy. Then, the carriers quickly
(~0.1ps) non-radiatively relax to the lowest energy state of their respective bands
through interactions with phonons.
After that, electrons at the bottom of the
conduction band recombine with holes at the top of the valence band and emit a
photon. This whole process is shown in figure 2.13. The energy of the emitted
photons is the same as the band gap energy of the direct band gap material to obey
the energy conservation law. In a PL process the total momentum of the system
must also be conserved. The photon wave vector is given by 2π/λ, where λ is the
wavelength. The electron wave vectors can be anywhere between the first Brillouin
40
zone boundaries, between -π/a and +π/a, where a is the crystal lattice spacing. Using
light with wavelength λ of 500 nm and a crystal lattice spacing around 5 Å as an
example, the photon wave vector (k) is several orders of magnitude smaller than the
possible electron wave vectors (k). Thus, in an E-k diagram, as shown in figure 2.13,
the radiative recombination is only allowed in the vertical directions without phonon
assistance. For a direct band gap semiconductor, e. g. GaAs, the global minimum of
the conduction band and the global maximum of the valence band are aligned
vertically in momentum space (k-space). Thus, the radiative recombination in GaAs
happens with only two particles (electron and hole), and the PL signal is strong.
It is also possible to have non-vertical radiative recombination processes in an
indirect bandgap system, such as Ge. However, in order to conserve the system
momentum, at least one phonon must be either created or absorbed simultaneously.
However, because this non-vertical transition process involves at least three particles,
instead of the two particles needed for the vertical transition, radiative recombination
in an indirect band gap material has much lower probability than in a direct band gap
material. Therefore, the PL signals in Si and Ge are very weak. On the other hand,
PL serves as a good indicator to tell whether the material is a direct band material or
not.
The excess free carriers may also recombine non-radiatively, most notably,
through Shockley-Read-Hall (SRH), or defect recombination. In a high quality
direct band gap material, where the impurity and defect densities are low, the nonradiative recombination probability is limited, and the PL signal is strong. On the
other hand, if the material quality is low, most of the excess carriers recombine non 41
radiatively through defects.
Thus, the PL signal from low quality material is
significantly reduced. Therefore, the intensity of the PL correlates positively with
the material quality, and serves as an indicator of material optical quality.
Filter
Figure 2.14: Schematic of the photoluminescence setup. (Courtesy of Dr. Maria
Makarova)
Several PL setups were used in this study. Figure 2.14 shows a micro-PL
setup. A 473 nm continuous wave laser is used as a pump laser with 5 mW or less
power. The laser is first transmitted through a 470 nm band pass filter, which only
allows the 473 nm blue laser to go through. Then the laser is reflected by a cold
mirror, which reflects all the visible light but transmits all the infrared light. Next,
the laser is fed into a 100X objective lens to obtain a focus point at the diffraction
limit. When the pump laser is focused on the surface of the sample, a strong PL
signal is collected and collimated by the same objective lens. After that, the light
passes through a couple different filters to eliminate the 473 nm pump laser and keep
42
the entire useful PL signal. Then, an iris is used at the focal plane to remove all the
ambient light. At the end, the light is either transmitted through a polarizing beam
splitter and imaged by a CCD camera, or reflected to a spectrometer to measure the
PL spectrum. An array of Ge detectors is used in the spectrometer to collect the PL
signal. Another PL setup employed in this work uses a chopper to modulate the
pump laser.
The emission signal passes through a monochrometer to select a
specific wavelength light. The selected light is then detected by a photo detector
from material which absorbs in the PL signal spectral range. The detected signal is
fed into a lock-in amplifier to enhance the signal-to-noise ratio. For both setups, the
system response over the whole spectral range is not flat because of the unequal
spectral response of the optical parts, the monochromator grating and the photo
detectors. Thus, each PL measurement is normalized by dividing by the system
response, which is obtained from the system response to a white light (known,
broadband) source. Also, a standard sample may be used to eliminate day-to-day
variations. The sample is mounted in air to take PL at room temperature. However,
if low temperature PL is preferred, the sample is placed in a small vacuum chamber
with an optical window. The sample stage is cooled by liquid He2 and heated by an
electrical heater to maintain a desired temperature.
In our cryo-stage, the
temperature can be steadily controlled from 5 K to room temperature.
43
2.3 Summary
In this chapter, we have described the two MBE systems used in this study as
well as their advantages for materials growth of novel materials. Several materials
property characterization tools were described, especially their applications to ultrathin Ge epitaxial layers.
References:
[1] http://www.stanford.edu/group/glam/xlab/XPert1/XPert1.htm
[2] S. Adachi, Properties of group-IV, III-V and II-VI semiconductors, Wiley (2004).
[3] J. A. Olsen, E. L. Hu, S. R. Lee, I. J. Fritz, A. J. Howard, B. E. Hammons, and J. Y.
Tsao, “X-ray reciprocal-space mapping of strain relaxation and tilting in linearly
graded InAlAs buffers”, J. Appl. Phys., 79 (7) 3578-3584, (1996).
[4] J. M. Chauveau, Y. Androussi, A. Lefebvre, J. Di Persio, Y. Cordier, “Indium
content measurement in metamorphic high electron mobility transistor structures by
combination of X-ray reciprocal space mapping and transmission electron
microscopy”, J. Appl. Phys., 93 (7) 4219-4225, (2003).
[5] http://www.stanford.edu/group/snl/
[6] Donghun Choi, Ph.D. thesis, (2008).
[7] H. H. Rose, “Optics of high-performance electron microscopes," Science and
Technology of Advanced Materials, vol. 9, no. 1, p. 014107, (2008).
44
[8] E. Anastassakis, A. Pinczuk, E. Burstein, F.H. Pollak and M. Cardona, “Effect of
static uniaxial stress on the Raman spectrum of silicon”, Solid State Communications,
Vol. 8, p 133, (1970).
[9] Ingrid De Wolf, “Micro-Raman spectroscopy to study local mechanical stress in
silicon integrated circuits,” Semicond. Sci. Technol. 11, 139 (1996).
[10] Y. Bai, K. E. Lee, C. Cheng, M. L. Lee, and E. A. Fitzgerald, “Growth of highly
tensile-strained Ge on relaxed InxGa1-xAs by metalorganic chemical vapor deposition,”
Journal of Applied Physics 104, 084518 (2008).
[11] Y.-Y. Fang, J. Tolle, R. Roucka, A. V. G. Chizmeshya, John Kouvetakisa, V. R.
D’Costa and José Menéndez, “Perfectly tetragonal, tensile-strained Ge on Ge1-ySny
buffered Si(100),” Applied Physics Letters 90, 061915 (2007).
[12] S. Wilson, C. R. Brundel, and C. Evans, “Encyclopedia of Materials
Characterization: Surfaces, Interfaces, Thin Films,” Butterworth-Heinemann, 1992.
[13] G. Jaschke, R. Averbeck, L. Geelhaar, and H. Riechert, “Low threshold
InGaAsN/GaAs lasers beyond 1500 nm,” Journal of Crystal Growth, vol. 278, pp.
224–228, May 2005.
[14] S. R. Bank, B. Hopil, L. L. Goddard, H. B. Yuen, M. A. Wistey, R. Kudrawiec,
and J. S. Harris, “Recent progress on 1.55- μm dilute-Nitride Lasers,” Quantum
Electronics, IEEE Journal of, vol. 43, no. 9, pp. 773, 2007.
3
45
Chapter 3
Biaxial Tensile Strained Ge
3.1 Introduction
3.1.1 Overview of Strained Ge
As discussed in Chapter 1, Germanium (Ge), already in use with modern Si
MOSFETs, has recently gained much interest for photonics due to its near direct
band gap property and its compatibility with Si CMOS. The demonstration of Ge
detectors and modulators [1, 2] based upon direct band gap absorption suggests the
possibility of using Ge as a gain medium for Group-IV-based lasers, LEDs, and
optical amplifiers.
Especially, research in tensile-strained Ge has made much
progress in recent years through the investigation of integrated optoelectronics on
silicon.
In 1996, Fischetti and Laux from IBM first presented theoretical work on
strained Ge and predicted that with about 1.75% biaxial tensile strain, Ge becomes a
direct bandgap material due to the Γ band-edge decreasing more rapidly than the L
band-edge with increasing tensile-strain [3].
In direct bandgap Ge, radiative
recombination becomes far more likely than non-radiative recombination that
dominates indirect gap materials.
Thus, electrons have a higher probability of
photon emission through recombination via direct transitions, making direct-bandgap
46
Ge an exciting material for optoelectronic devices. It has already been demonstrated
that the performance of many Ge-based optoelectronic devices, such as detectors and
modulators, is enhanced by inducing even a small strain inside the Ge layers (2-4).
For example, the Ge detector response range can cover the whole C-band in optical
communications with only 0.2% tensile strain in the Ge layer.
Although the advantages of tensile-strained Ge can be easily demonstrated,
the amount of strain needed to reach the direct-bandgap transition is difficult to
achieve due to 1) a lack of larger than Ge lattice constant substrate material, 2) the
need to maintain coherent Ge growth before dislocations are generated and the lattice
strain is relaxed.. Still, much work has been done to achieve high strain and modify
the direct-bandgap properties of Ge. Bai et al. from MIT used relaxed In0.11Ga0.89As
buffer layers grown by MOCVD to achieve 0.5% biaxial tensile strain in Ge layers
[APL 5]. Larger strain resulted in the formation of Ge quantum dots with tensilestrain of 1.37%. Hoshina et al. also used a similar method with MBE to achieve thin
layers of 1.55% tensile-strained Ge [apl6]. However, no optical property
characterization was been reported for these samples.
However, achieving more than 1.75% biaxial tensile strain in Ge layers is
critical to achieve a direct bandgap material, which is crucial to demonstrate an
efficient, electrically pumped CMOS compatible laser source. Thus, how to achieve
more than 1.75% tensile strain inside Ge with good material quality is a critical
problem which has not yet been solved.
47
3.1.2 Methodology
The most important factor for achieving highly tensile strained Ge layers is
the buffer layer choice. Other than good material and optical qualities of the buffer
layer, it must have a larger lattice constant than Ge’s lattice constant in order to
provide in-plane biaxial tensile strain in the Ge layers. With this requirement, the
appropriate buffer layer materials are limited. In the ideal case, a relaxed GeSn
buffer layer would first be grown on a Si substrate, and then a the tensile strained Ge
layer can be grown on top of the GeSn buffer layer. However, the growth of GeSn is
not trivial due to the low solid-solubility of Sn in Ge. GeSn growth is discussed
separately later in Chapter 4. In order to study the physical properties of tensile
strained Ge, several binary or ternary III-V semiconductors can be used as buffer
layers. InxGa1-xAs buffer layers are discussed in detail here, while other possible
buffers are described later. In this work, we grow and characterize thin Ge layers
which are highly-tensile strained (up to 2.33%) using lattice relaxed InxGa1-xAs
buffer layers with indium concentrations of up to x = 0.4. The layer structure is
shown in figure 3.1. By controlling the indium composition in each buffer layer, we
can control the crystal in-plane lattice constant and thus control the biaxial tensile
strain inside the Ge layer. To our knowledge, this work demonstrates the highest
amount of in-plane biaxial tensile strain in epitaxial Ge layers. Also for the first time,
these samples exhibit excellent optical luminescent properties at low temperatures
48
with high tensile strain. These results suggest that a direct-bandgap Ge material has
been achieved with a good possibility for use as a CMOS compatible light source.
In0.3Ga0.7As
175 nm
11.2 nm
tensile strained Ge
In0.3Ga0.7As
In0.15Ga0.85As
275 nm
GaAs substrate
438 nm
(a) (b)
Figure 3.1: Schematic (a) and SEM image (b) of biaxial tensile strained Ge sample
structures
A particularly critical issue for highly strained Ge is the maximum layer
thickness before dislocation generation and lattice relaxation occurs. In order to
achieve high tensile strain in a Ge layer, the epitaxially grown Ge layer cannot be
very thick.
As described later in section 3.3.1, through the critical thickness
calculation, the critical thickness of Ge is only a few nanometers for 2% tensile strain
under thermal-equilibrium conditions. Since low temperature MBE growth produces
a highly non-thermal-equilibrium condition, we may achieve about a 10 nm thick Ge
layer with 2% tensile strain without lattice relaxation. Although the 10 nm thick
strained Ge layer is thick enough for the active gain region of a quantum well edgeemitting laser, it still generates many difficulties for adequate the material and
optical characterization.
Thus, some special characterization techniques and
methods are required to investigate the properties of these thin layers of highlystrained Ge.
49
3.1.3 Section Arrangement
In section 3.2, the growth of lattice relaxed thick InxGa1-xAs layers are first
discussed. The material properties from SEM, XRD, TEM, and PL characterization
are also discussed to prove that a good buffer layer growth method was achieved.
Next, in section 3.3, the material and optical properties of different tensile strained
Ge are described and analyzed. Section 3.4 contains the final summary about the
strained Ge material.
3.2 Relaxed InGaAs Buffer Layers
3.2.1 Motivation
InxGa1-xAs growth on GaAs has many applications in both electrical and
optical devices. For the traditional III-V electro-optic device, the band gap energy of
GaAs, which is 1.424 eV (or 870 nm in terms of photon wavelength), is too large for
many IR applications.
Although many 1.3 μm and 1.55 μm (the most useful
wavelength range in optical fiber communication systems) optical devices have been
fabricated on InP substrates, devices on GaAs substrates still offer several
advantages. Such advantages include better material quality, less-expensive
substrates, high refractive index contrast in AlAs/GaAs lattice matched material
systems, the avoidance of dangerous phosphorous growth in the chamber, etc..
Currently, the extended wavelength InGaAs photodetectors can operate at
wavelengths longer than 2 μm and are commercially available [6].
50
Equally
importantly, the high electron mobility in InGaAs makes it attractive for high-speed
electronic devices and most of the low-noise, high electron mobility transistors
(HEMT) in cell phones utilize strained InGaAs channel layers.
Furthermore, relaxed InxGa1-xAs buffer layers on GaAs substrates are also
very important in our biaxial tensile strained Ge study. InxGa1-xAs can provide a
good virtual substrate with different desired lattice constant for the following Ge
layers by simply varying the indium content. Therefore the strain status of the
following Ge layers can be accurately controlled. The key criteria for the InxGa1-xAs
buffer layers are low threading dislocation density (TDD), efficient strain relief,
smooth surface, and strong PL. All these factors are further discussed in the
following sections.
The major challenge to grow InxGa1-xAs on a GaAs substrate is the lattice
mismatch. The lattice constant of GaAs is 5.6533 Å. And, the lattice constant of
InxGa1-xAs is given by
a = 5.6533 + 0.4050x
(3-1)
where x is the indium composition. For example, In0.3Ga0.7As has a lattice constant
2.15% larger than the GaAs lattice constant. As-grown InxGa1-xAs buffer layers with
significant lattice-mismatch usually have rough surfaces and a very high defect density.
Many types of defects incorporate into the relaxed buffer layers, including misfit
dislocations at the material interface, threading dislocations propagating to the outer
surface of the layers, and other defects resulting from the strain and growth conditions,
as shown in figure 3.2 [7]. Among all of these defects, misfit dislocations (MD) and
51
threading dislocations (TD) are of great interest to both theoretical and experimental
research because they may propagate into subsequent layers, which impact the device
performance. The defect density can be reduced by optimizing the growth parameters.
Figure 3.2: Defects in hetero-epitaxy material growth [7].
3.2.2 Traditional InGaAs growth recipe
In order to achieve a good InGaAs quantum well (QW) laser, the growth
recipe of thin (10 nm) strained InGaAs is optimized based on optical properties,
especially PL results. However, this recipe is not suitable for thick, strain relaxed
InxGa1-xAs buffer layer growth. MBE growth of thick InxGa1-xAs buffer layers has
also been studied by former students in our lab. Lord investigated the temperature
dependence of InxGa1-xAs growth and found that if the substrate temperature is
higher than 550 ºC, indium atoms desorb from the surface during growth [8].
52
Linearly graded InGaAs buffer layers were also studied, and TEM shows they have
fewer defects than step graded InGaAs buffer layer samples, although they require
thicker buffer layers.
Fu later found that low temperature growth with high
temperature annealing decreases the defect density in the relaxed InGaAs buffer
layers [7]. However, the final defect density is still high, and some material and
optical properties, which are very critical for the strained Ge layer investigation,
were not characterized.
3.2.3 InGaAs buffer layer growth optimization
In our study of MBE growth of InGaAs buffer layers, step-graded InxGa1-xAs
buffer layers on GaAs (100) substrates have one, two or three InxGa1-xAs layers with
different indium concentrations. In order to discuss the recipe optimization clearly,
two very different In0.3Ga0.7As buffer growth methods are first compared. Other
growth recipes are described afterwards.
For all of our samples, the surface oxide on the GaAs wafer is first blown off
around 680°C (all the substrate temperature readings in this thesis refer to the
reading of the thermocouple in the MBE substrate holder) in the III-V MBE chamber.
The substrate temperature is gradually increased until a clear (2×4) RHEED pattern
is observed. Then, the sample is kept at the same temperature for 15 minutes more
to totally remove the oxide from the surface. After that, a 50 nm to 100 nm GaAs
seed layer is grown on a GaAs substrate at the oxide blow off temperature with 15
53
times arsenic over pressure. This GaAs seed layer growth gives a good and
repeatable crystal surface for InGaAs buffer layer growth.
In the first method (Sample A, high temperature growth method), InxGa1-xAs
is grown with substrate temperature from 520 ºC to 540 ºC. For a higher indium
concentration, the substrate temperature is lower in order to prevent desorption of
indium atoms. The arsenic over pressure is 13 times the gallium flux pressure. In
the second method (Sample B, low temperature growth and high temperature anneal
method), InGaAs is first grown at 360 ºC to 380 ºC with 10 times arsenic over
pressure. Then the sample is quickly ramped to 520-540 ºC and annealed for 20
minutes. After that, the sample is cooled back to 360-380 ºC and the second InGaAs
buffer layer growth starts. Normally, each buffer layer is around 200 nm thick in
order to achieve maximum relaxation. In the end, an InGaAs cap layer (from 10 nm
to 200 nm thick depending on the purpose) with the same indium content may be
grown without subsequent annealing.
In both methods, the indium concentration increases by about 10%~18% in
each successive layer by varying the In:(In+Ga) flux ratio for this step graded
InGaAs buffer layer growth. In order to achieve different indium concentrations in
the InGaAs buffer layers, the indium evaporator temperature is changed to obtain
different indium flux rates based on the calculations from the calibration done before
each sample growth.
The gallium source temperature is fixed to maintain a
relatively small change in the crystal growth rate. The needle valve position of the
arsenic source is changed in each step to provide sufficient group V element over
pressure all the time.
54
Figure 3.3 shows SEM cross section images for these two methods. As
shown in figure 3.3 (b), in the high temperature growth method (sample A), the
defect density is so high that even the interface of the InGaAs buffer layers is not
smooth after cleaving. On the other hand, with the low temperature growth high
temperature annealing method (sample B), the defect density is dramatically
decreased so that the interface is continuous and smooth. Only a few cracks show up
in the cross section of sample B, as shown in figure 3.3 (c).
In0.3Ga0.7As
In0.15Ga0.85As
GaAs
(a) 329 nm
192 nm
191 nm
178 nm
(c) (b) Figure 3.3: Schematic (a) and SEM image of cross-section structure of (b) sample A
with high temperature growth method and (c) sample B with low temperature growth
and high temperature annealing method.
55
Surface roughness is particularly critical in our research. Since the strained
Ge layer is only 10 nm thick and forms a quantum well, large surface roughness
changes the quantum well thickness and broadens the energy distribution of carrier
states associated with quantum confinement effects. The surface morphology is
measured by AFM. As shown in figure 3.4, a typical InGaAs thick buffer layer
growth on GaAs substrate always has a cross-hatch pattern. The RMS surface
roughness in a 1x1 μm2 region is decreased from 1.5 nm for sample A with the high
temperature growth method to 0.62 nm for sample B with the low temperature
growth high temperature annealing method. Sample B with a smoother surface is a
better buffer layer for the following Ge layer epitaxial growth.
X: 0.2 μm/div
X: 0.2 μm/div (a) (b) Figure 3.4 AFM image of InGaAs buffer layer surface morphology of (a)
sample A with high temperature growth method and (b) sample B with low
temperature growth and high temperature annealing method.
56
XRD 2D reciprocal space mapping is measured for both symmetric (004) and
asymmetric (224) diffraction peaks. From 2D reciprocal space mapping results, the
indium concentration is calculated and is close to our designed indium concentration.
The average relaxation for each InGaAs layer is about 85%. The 2D scans at (224)
for two samples with different InGaAs growth methods are shown in figure 3.5. As
we can see, the diffraction pattern of the In0.3Ga0.7As buffer layer is very different for
the two samples. The FWHM of the In0.3Ga0.7As along the qx is different by 4 times
(0.0043 rlu for sample A and 0.0012 rlu for sample B). In general, the threading
dislocations can reduce the in-plane coherence length inside a single crystal. Thus,
sample A with the high temperature growth method has many more threading
dislocations than sample B [9].
(a) (b) Figure 3.5: 2D reciprocal space mapping of InGaAs buffer layers for two samples
with different growth recipes. (Courtesy of Hai Lin)
57
TEM is one of the most powerful instruments to characterize the material
qualities of our samples. (110) oriented cross-section TEM images for the two
methods are shown in figure 3.6. For the high temperature growth method (sample
A), an increasing density of threading dislocations are shown in the InGaAs buffer
layers as the indium concentration increases. A high concentration of defects behave
as recombination centers to capture the minority carriers and dramatically reduce the
radiative recombination process.
However, in figure 3.6 (b), with the low
temperature growth high temperature anneal method (sample B), the threading
dislocations are much fewer. Most of the defects terminate at the first interface
between the step-graded InGaAs buffer layers and leave the top buffer layer almost
defect free in cross-section TEM view. The horizontal dark wide line is a thickness
fringe from sample preparation.
30% InGaAs
15% InGaAs
GaAs
200 nm
200 nm
(a) (b) Figure 3.6: Cross-section TEMs of InGaAs buffer layers grown on GaAs by
two different methods: (a) high temperature growth and (b) low temperature growth
and high temperature annealing.
58
The defects are more clearly observed by selecting a diffraction beam to
image the sample cross section. As show in figure 3.7, most of the defects are misfit
dislocations at the interface which induce local strain. A few threading dislocations
are shown in the In0.15Ga0.85As first buffer layer. However, the second In0.3Ga0.7As
buffer layer is defect free. By the limitation of the amount of material sampled by
cross-section TEM, the defect density in the top layer is estimated to be less than or
within the 1*108 cm-2 range.
Figure 3.7: Cross-section TEM of sample B InGaAs buffer layers.
As discussed at the beginning of this section, in addition to the material
quality, the optical properties of the buffer layer are most important, especially the
PL intensity. The material direct band gap energy can be calculated from the PL
spectrum. For a bulk layer of InxGa1-xAs at room temperature, the band gap energy
is given by:
Eg bluk = 1.424 -1.614x + 0.540x2
(3-2)
where x is the indium concentration in the material. Figure 3.8 shows the PL
spectrum of these two samples with two different growth methods at room
59
temperature. The two curves in this figure clearly show the difference in optical
properties between the two different growth methods.
The two luminescence intensity peaks occur at 1.0 μm and 1.2 μm and arise
from the In0.15Ga0.85As and In0.3Ga0.7As buffer layers, respectively.
For
In0.15Ga0.85As, the band gap energy is 1.19 eV, corresponding to 1.04 μm. So the PL
intensity starts to increase at 1040 nm and achieve the maximum at 1.0 μm, which is
due to free carrier band filling from the pumping laser. Similarly, for In0.3Ga0.7As,
the band gap energy is 0.99 eV; the PL intensity starts to increase around 1.25 μm
and has the highest luminescence at 1.2 μm.
In0.3Ga0.7As
PL intensity (a.u.)
1
In0.15Ga0.85As
LT+Anneal
0.1 -1
10
HT
900 1000 1100 1200 1300 1400 1500
Wavelength (nm)
Figure 3.8: PL intensity spectrum for high temperature growth method
(sample A, blue dashed line) and low temperature growth high temperature
annealing method (sample B, black solid curve)
60
The PL spectrum difference between the two samples grown with different
methods is very obvious. For Sample A (high temperature growth method), the blue
dashed curve in figure 3.8 shows two small PL intensity peaks around 1.0 μm and
1.1 μm. The PL signal from the In0.3Ga0.7As buffer layer almost totally disappears,
which means the material quality of the second In0.3Ga0.7As buffer layer is very poor.
However, with sample B, when the buffer layers are grown at low temperature and
annealed at high temperature, as the black curve shows in figure 3.8, the PL signal
from the In0.3Ga0.7As buffer layer is more than 10 times stronger than sample A’s
In0.3Ga0.7As buffer layer PL signal. One point that needs to be clarified is that, in the
black curve, the PL signal from the first (bottom) In0.15Ga0.85As buffer layer is much
smaller than the signal from the top In0.3Ga0.7As buffer layer. This is because both
the pumping laser and luminescence light need to be transmitted through the
In0.3Ga0.7As buffer layer and they have been absorbed. The FWHM of the PL from
the In0.3Ga0.7As buffer layer is about 50 nm, which is about 50meV or 2kT. The
energy spreading is coming from both the carrier thermal distribution and the nonuniform material band gap energy. This is because the non-uniform relaxation of
InGaAs buffer layers gives non-uniform strain inside the buffer layers, and the
bandgap energy changes accordingly.
Comparing the PL spectra from these two samples, the buffer layer in sample
B (low temperature growth high temperature annealing) has a much higher radiative
recombination probability than sample A. This is also reasonable from all material
characterization predictions. The buffer layers following the low temperature growth
61
high temperature anneal method, can provide a good substrate to strain Ge layers by
having achieved both excellent material and optical properties.
3.2.4 Discussion of Other Optimization and Growth Mechanisms for
InGaAs Buffer Layers
Two typical growth examples (sample A and B) were described in the
previous section. In this section, we provide further discussion about the growth
mechanism as well as additional studies about other samples with different InGaAs
step graded buffer layer growth methods.
Growth temperature is one of the most important parameters in material
growth. As discussed before, if the substrate growth temperature is higher than 550
ºC, indium atoms may desorb from the crystal. On the other hand, if the substrate
temperature is higher, the column III atoms have greater surface mobility and have
more chance to find a perfect crystal lattice point upon which to grow. Thus, if the
substrate is a perfect crystal with the same lattice constant as the epitaxially grown
layer lattice constant, a better InGaAs crystal quality may be achieved with higher
substrate temperature as long as the substrate temperature is lower than 550 ºC.
However, if the lattice constant is mismatched, many kinds of defects can be
generated at the interface, especially at high temperature. Also, at lower temperature,
the energy for defect migration is low as well. Thus, most of misfit defects can be
confined near the hetero-epitaxial growth interface. Therefore, a lower substrate
temperature is preferred during the hetero-epitaxial growth with lattice mismatch.
62
However, if the growth temperature is too low, adatoms don’t have enough surface
mobility to find a good lattice point. Thus many local defects may occur and the
material quality decreases. In the extreme low temperature case, amorphous material
is grown instead of single crystal material. As a result, low temperature growth is
good for defect confinement with material quality decrease as a tradeoff.
In order to confine the misfit locations at the hetero-epitaxial growth interface
and decrease the threading dislocation density, low temperature growth is used as the
first step. When the InGaAs buffer layers are grown at a relatively low temperature,
the material quality is less than perfect. The local atom disorder is distributed
throughout the epitaxial layer, because atoms do not have enough energy to move
around to the right lattice point. In order to fix this problem, in-situ high temperature
annealing is used after each InGaAs buffer layer growth. During the annealing step,
the high substrate temperature provides enough energy to move the atoms locally to
the lowest energy point, which forms a closer to perfect lattice structure as well.
Thus, in this low temperature growth and high temperature anneal method, most of
misfit dislocations are confined at the interface region and the material qualities as
well as the optical properties are significantly improved.
Annealing not only benefits the bulk material but also affects defect
migration, defect termination at the top surface, top surface morphology, and other
properties in different ways.
63
During the annealing stage, the high temperature provides more energy to the
defects as well, which allows the defects to travel more.
For the threading
dislocations, with more energy, they have a greater chance to slide to the edge of
sample if they are not stuck by other defects. Thus, annealing can help to reduce the
threading dislocation density.
At high annealing temperatures, the atoms at the surface have a greater
chance to move around to cure defects at the surface or change the direction of the
defects to travel in-plane to the edge. As shown in the figure 3.9, some threading
dislocations clearly change directions or are terminated at the interfaces.
GaAs Top surface Figure 3.9: Cross-section TEM image of two step InGaAs buffer layers on
GaAs substrate with threading dislocations propagation and bending.
64
(a) (b) Figure 3.10: AFM of In0.3Ga0.7As buffer layers on GaAs without (a) and with
(b) in situ high temperature annealing after last InGaAs layer growth in MBE system.
During the annealing step, the surface morphology may change noticeably as
well. For example, if the Ge surface is very rough, the RMS roughness can be
reduced from 20 nm to 2 nm in a 1x1 μm2 region [10]. Sometimes, in an H2
environment, hydrogen atoms can significantly increase the Si atom surface mobility.
Thus, high temperature H2 annealing can cause the Si surface to reflow and result in
a very smooth surface. In our case, our sample surfaces are normally quite smooth.
Even in the worst case, the surface roughness is normally smaller than 1nm. In this
case, during the annealing step, the InGaAs surface becomes slightly rough as shown
in figures 3.10. The RMS surface roughness increases from 0.56 nm to 0.70 nm in a
1x1 μm2 region. Thus, the annealing step may not be preferred after the final top
InGaAs buffer layer is grown.
The annealing temperature is studied as well. Ideally, a higher temperature
improves the bulk material quality.
However, the surface may become a little
rougher. At an even higher annealing temperature, a more serious problem occurs:
65
the indium atoms may evaporate from the crystal, which causes the indium
concentration to decrease. As shown in figure 3.11, although the In0.3Ga0.7As PL
peak intensity (dashed blue line) is increased after annealing at 580 ºC and shows a
better material optical properties, the peak is blue shifted in wavelength after 580 ºC
annealing for 20 minutes. The PL peak at 1064 nm is reflection from the pumping
laser and can be ignored. This proves that the indium concentration is greatly
reduced during very high temperature annealing. Thus annealing at 520 ºC to 540 ºC
is used for most samples.
PL intensity (a.u.)
-12
10
-13
10
900
1000
1100
1200
Wavelength (nm)
1300
Figure 3.11: PL spectra for In0.3Ga0.7As and In0.15Ga0.85As buffer layers with
530 ºC annealed sample (solid black line) and 580 ºC annealed sample (dashed blue
line)
66
Many other growth recipes were explored as well. We have tried to vary the
indium concentration change between two adjacent buffer layer steps. If the indium
composition change is more than 18% in one step, the lattice mismatch is so large
that far more defects are generated at the interface and too many threading
dislocations propagate through the buffer layers. The final sample has much rougher
surface and the threading dislocation defects are dramatically increased as well. If
the indium composition change is less than 10% per step, the lattice mismatch is less
than 0.7%. Under this small lattice mismatch condition, the critical thickness at
thermal equilibrium increases almost exponentially and only partial lattice relaxation
occurs, thus making the layers not useful as substrates for strained Ge growth. Due
to our MBE epitaxy growth under highly non-thermal-equilibrium conditions, the
critical thickness is even larger to start the relaxation process. However, with such
thick epitaxial layers, it is much harder to achieve 100% full relaxation.
As
mentioned before, characterization by 2D XRD reciprocal lattice mapping show that
we normally have about 80% relaxation on 15% indium concentration step graded
buffer layers. However, we achieve less than 50% relaxation when the indium
concentration step grading is less than 10% in each step.
67
3.2.5 Summary
InGaAs relaxed thick buffer layers are systemically studied. Low temperature
growth with high temperature in-situ annealing proved to be the best method for
thick buffer layer growth. About 80% relaxed, up to 40% indium composition
InGaAs buffer layers were grown on GaAs (001) substrates. The top surface is
smooth (less than 1 nm RMS roughness in 1x1 μm2). The threading dislocation
density is in the 1x108 cm-2 range as observed by cross-section TEM. The PL signal
is increased by more than 10 times compared to the traditional growth recipe. Good
InGaAs buffer layers were prepared by this process for our investigations of strained
Ge epitaxial layer growth.
Both the growth mechanism and the annealing mechanism were studied
under different conditions. Other buffer layer materials, such as InAlAs, InGaP and
GaAsSb were studied as well. They are attractive for quantum confinement due to
their larger band gap energies. However, those studies are limited by source material
availability or material growth quality currently and are not discussed in detail in this
thesis.
68
3.3 Properties of Biaxial Tensile Strained Ge
We have a unique MBE facility in having a coupled two-chamber system
with one chamber devoted to Group III-V materials and the second to Si-Ge-Sn
Group IV materials.
These chambers are coupled by a high-vacuum transfer
tube/trolley such that the strained Ge layers can be grown on an InGaAs buffer layer
which has never been exposed to air. After a relaxed InGaAs buffer layer is grown,
a germanium layer is epitaxially grown on top of the InGaAs buffer layer. The
strained Ge layer growth and material properties are discussed in this section.
3.3.1 Critical Thickness for Strained Ge Layers
Strained layer epitaxy is a rapidly growing area for both research and real
applications. In electronics, strained Si is widely applied to increase the carrier
mobility in order to improve the device performance [11].
In electro-optics
applications, strained III-V epitaxy layers are widely used in the quantum well region
to achieve good laser behavior [12]. Strain is also the fundamental driving force in
producing quantum dots, which are now widely used in low threshold lasers [13].
Lattice mismatched materials can be stable and may not have very high
defect density if the epitaxial layers are thinner than a given thickness, which is
called the critical thickness. The most widely used model for lattice mismatched
epitaxy is known as the Matthews-Blakeslee model [14] and is shown in figure 3.12.
When the epitaxial layer thickness is smaller than the critical thickness, the mismatch
69
between substrate and epitaxial layer is accommodated by strain in the film. As
shown in figure 3.12 (b), the epitaxial layer in this example is compressively strained
in-plane.
Also, following the Poisson ratio, the out-of-plane lattice constant is
enlarged. In this case, the quality of the epitaxial layer is good since no defects are
introduced. If the epitaxial layer thickness is larger than the critical thickness, the
film can no longer accommodate such high strain. Thus misfit dislocations are
introduced, as shown in figure 3.12(c).
At regions far from the interface, the
epitaxial layer is almost totally relaxed. Some of the misfit dislocations are confined
near the interface. However, in some cases, the defects travel through the whole
epitaxial layer and degrade the material quality.
Epitaxial film (b) h < hc Substrate (c) h > hc (a) Figure 3.12: Schematic illustration of Matthews-Blakeslee model.
Two
conditions are shown as (b) totally strained coherent growth and (c) totally relaxed
growth.
70
It is difficult to provide an exact critical thickness for given materials
combination since this is a thermal equilibrium calculation and nearly all
mismatched heteroepitaxy is done under non-equilibrium growth conditions. A good
reference is the Matthews-Blakeslee critical thickness, which is given by
ε=
2hc
a (1 −ν / 4)
+ 1)
(ln
a
2hc 2π (1 + ν )
(3-3)
where ε is the strain, a is the lattice constant, ν is the Poisson’s ratio, and hc is the
critical thickness. Equation (3-3) can be plotted as shown in figure 3.13. From this
figure, we find that with 2% strain in the Ge layer, the critical thickness of the Ge
layer is about 7 nm. This means, in order to have enough strain in a Ge layer to
achieve direct band gap material, the Ge layer thickness cannot appreciably exceed
10nm.
This 10nm or less thick Ge layer creates many difficulties for material
characterization as well as optical properties measurements.
Ge critical thickness (nm)
50
40
30
20
10
0.5
1
1.5
2
2.5
Ge tensile strain (%)
3
Figure 3.13: Matthews-Blakeslee critical thickness for Ge at thermal equilibrium.
71
3.3.2 Strained Ge Growth
The thin layer of Ge is grown in a group IV Varian Gen II modified MBE
system. The samples with InGaAs buffer layers are first transferred through a
transfer tube, which is under ultra-high vacuum (~10-9 torr) to this group IV MBE
chamber. As discussed in the previous section, the Ge layer thickness cannot be
thicker than 10 nm, otherwise the Ge layer is relaxed by many misfit defects and the
Ge crystal quality degrades. Thus, normally, we only coherently grow 10 nm thick
Ge layers on the InGaAs buffer layers at a substrate temperature of 400 ºC to avoid
generation of more defects in the Ge layer. The growth rate is around 1 nm/minute.
If the substrate temperature is too high 3D Ge quantum dot growth occurs. If the
substrate temperature is too low, the Ge crystal quality is degraded due to
insufficient atom surface mobility. As a result, this recipe gives a good strained
germanium epitaxial layer and these epitaxial Ge layers are first characterized and
can then be applied to many applications.
3.3.3 Strain Calibration in Thin Ge Layers
The calibration of 10 nm thick Ge epitaxial layers is very challenging. This is
because the Ge material volume is much smaller than the material volume of the
substrate and buffer layers so that the Ge layer signal intensity in most techniques is
overwhelmed by the substrate background noise intensity in many characterization
techniques. Some very useful techniques for our case are discussed here, starting
with strain characterization and followed by material and optical characterizations in
the next section.
72
In order to measure the biaxial tensile strain inside the germanium epitaxial
layer, XRD is first used. Our samples were first examined in the XRD system at the
Stanford Nanocharacterization Laboratory at Stanford University.
However, we
could not observe a clear Ge signal in either XRD 1D rocking curve scan or 2D
reciprocal lattice mapping. This is due to both the relatively low detection dynamic
range of the system and the Ge layer thickness. As we know, the diffraction of xrays is based on the coherent interaction of the x-rays reflected from different atom
layers. The more atom layers we have, the narrower the angular distribution of the
diffracted beam. This relationship is described by the Debye-Scherrer equation as
Dhkl = kλ/βcosθ
(3-4)
where Dhkl is the film thickness in the (hkl) direction, k is the Scherrer coefficient
and is normally set to 0.89, λ is the wavelength of the x-rays which is 0.15406 nm
for the Cu Kα line for our instrument, θ is the Bragg diffraction angle, and β is the
full width at half maximum (FWHM) in radians. For a 10 nm thick Ge epitaxial
layer even with a perfect crystal lattice, the FWHM of the 004 diffraction peak is as
large as 1 degree. Figure 3.14 shows a simulation of the diffraction pattern of
different Ge layers ranging in thickness from 0 to 50 nm. As we see, the FWHM of
10 nm epitaxial layer diffraction signal is very large. Thus, in order to see the 10 nm
thick Ge diffraction signal, the background noise needs to be very low. However,
the diffraction peak of Ge is very close to the GaAs substrate diffraction peak and the
signal to noise ratio for our XRD instrument is not high enough, and in most of cases,
the Ge epitaxial layer signal is very hard to observe.
73
GaAs InGaAs Strained Ge
Buffer layer
15% InGaAs
Ge GaAs
Figure 3.14: Simulated XRD 1D rocking curve for different thickness Ge on
InGaAs buffer layer and GaAs substrate.
In order to have better signal to noise ratio, our samples were further
measured at the Stanford Synchrotron Radiation (SSRL) at the SLAC national
accelerator laboratory. The x-ray source is much brighter. Also, three different
attenuation films are automatically controlled to achieve a much wider detection
dynamic range. Figure 3.15 is a 1D XRD rocking curve scan for samples with and
without a Ge epitaxial layer on In0.3Ga0.7As and In0.15Ga0.85As buffer layers and
GaAs substrate. The different decay curves on the right side shoulder of the GaAs
peak show the Ge diffraction signal. Thus the Ge diffraction signal can be calculated
from the difference between these two measurements and are plotted in the inset.
The Ge diffraction peak intensity is at 34 degrees which proves the Ge is under
biaxial tensile strain.
74
0
Diffraction intensity (a.u.)
10
Intensity
0.5
0.4
0.3
0.2
0.1
0.0
0.1
-0.2
-0.3
-2
10
-4
34.0
10
With Ge
-6
10
Without Ge
32
32.5
33
33.5
34
Diffraction angle (degree)
34.5
Figure 3.15: 1D XRD rocking curve for InGaAs buffer layers on GaAs with
(blue dashed curve) and without (black solid curve) a strained Ge epitaxial layer.
However, the information from a 1D rocking curve is limited, and the beam
time at SSRL is limited. Thus this method with direct reading from the Ge diffraction
signal cannot be used all the time. Instead, if we assume the thin Ge layer is always
coherently grown on the InGaAs buffer layers during non-equilibrium growth, we
can calculate the Ge strain from the in-plane lattice constant of the InGaAs buffers.
The 2D XRD reciprocal lattice mapping at the (004) and (224) diffraction points for
a sample with Ge on an In0.4Ga0.6As buffer layer on a GaAs substrate are shown in
figure 3.16. From the 2D reciprocal lattice mapping result, we calculate that the
In0.4Ga0.6As buffer layer is 85% relaxed, the indium concentration is 39.9%, and the
in-plane lattice constant of the In0.4Ga0.6As buffer layer is 5.788 Å. So the biaxial
tensile strain in the Ge epitaxial layer is 2.31%.
75
GaAs
In0.4Ga0.6As
(224)
(004)
Figure 3.16: Reciprocal lattice mapping of strained Ge layer and InGaAs
buffer layers on GaAs substrate for both (004) and (224) scans.
Although XRD is a very powerful characterization instrument in terms of Ge
strain calibration, XRD measurement is still not the most convenient way to measure
the strain in Ge epitaxial layers. The 2D XRD reciprocal lattice mapping takes hours
for one sample measurement and XRD in SSRL is hard to access. Instead, Raman
spectroscopy is a very quick and accurate method for strain characterization.
The fundamental physics background of Raman spectroscopy has been
discussed in Chapter 2. Since GaAs/InGaAs have a similar lattice constant and atom
mass as Ge, it is important to first check if the Ge Raman signal is clearly separate
from the background signal of the GaAs substrate and InGaAs buffer layers. By
choice of short wavelength pumping laser, the penetration depth in InGaAs/GaAs is
only a few hundred nanometers. This is important so that the signal from GaAs
76
substrate is almost totally eliminated, which is totally different from the XRD
measurement where all layers, especially the GaAs substrate is the largest
contributor to the final x-ray diffraction pattern.
Figure 3.17 shows the Raman spectra of three samples. The first sample is a
bulk Ge sample. The second sample is In0.3Ga0.7As and In0.15Ga0.85As thick buffer
layer on top of a GaAs substrate. And, the third sample is a thin layer (10nm) of Ge
coherently grown on an In0.3Ga0.7As thick buffer layer which is the same as the
second sample. As we observe in the figure, the Raman shift of bulk Ge is around
302 cm-1. Some bias of Raman shift may happen in this measurement due to the
instrument calibration. Thus a bulk Ge substrate is always used to calibrate the
system. For the second sample, In0.3Ga0.7As Raman signal shift is around 282 cm-1.
There is some unclear background noise at the left side of the main Raman shift
signal. However, the signal around 300 cm-1 is very clean. In the third sample,
where the Ge is coherently grown on top of an InGaAs buffer layer, the InGaAs
Raman signal is not shifted but greatly reduced. The signal from Ge is shifted from
302 cm-1 to 294 cm-1, which means the in-plane lattice constant of Ge is changed.
77
Normalized intensity (a.u.) 1
Strained
Ge
0.8
InGaAs
0.6
0.4
Bulk Ge
0.2
0
260
270
280 290 300
Raman shift (cm‐1)
310
320
Figure 3.17: Raman spectra for bulk Ge (red dotted curve), In0.3Ga0.7As
buffer layers (blue solid curve) and Ge on In0.3Ga0.7As buffer layers (black dashed
curve).
The strain inside the Ge layers can be calculated from the change of Raman
shift using the equation Δω = bε‖, where ε‖ is the biaxial strain in the Ge layers and
b is equal to -415 ± 40 cm-1 [4].
1
Normalized intensity (a.u.)
In0.1Ga0.9As 0.8
0.6
In0.2Ga0.8As In Ga As
0.4
0.2
0
260
280
300
raman shift (cm-1)
320
Figure 3.18: Raman spectra for different tensile strained Ge layers.
78
A set of strained Ge samples is prepared and measured by Raman
spectroscopy. Each sample has the same thickness of Ge epitaxial layer, however
different biaxial tensile strain is applied to these Ge layers by using different indium
concentrations in the buffer layers. The Raman spectra are plotted in figure 3.18. A
clear shift of the Ge Raman signal is observed. The calculated biaxial tensile strain in
Ge is listed in table 3.1. The biaxial tensile strain in the Ge layer vs. the indium
concentration in the InGaAs buffer layer is plotted in figure 3.19 as well.
A
theoretical prediction is calculated based on the InGaAs lattice constant and the
assumption of totally coherent growth of Ge.
The Ge biaxial tensile strains
measured by Raman spectroscopy and by XRD 2D reciprocal lattice mapping are
marked by red squares and blue triangles respectively in this figure. The results from
these two methods are very close to each other. Because of the partial relaxation of
InGaAs buffer layers (normally 80% relaxed), the measured results are always
smaller than the theoretical prediction.
Indium
concentration
10%
20%
30%
40%
Theoretical Tensile strain of Ge
Ge strain
Raman
XRD
0.64%
0.26%
0.42%
1.35%
0.91%
0.93%
2.07%
1.78%
1.84%
2.79%
2.35%
2.31%
Table 3.1: Biaxial tensile strain inside Ge with different InGaAs buffer layers.
79
3%
Raman
XRD
Theory
Strain
2%
1%
0%
0
0.1
0.2
0.3
In concentration
0.4
Figure 3.19: The tensile strain inside Ge layers with different InGaAs buffer
layers. The black curve, blue triangles and red squares represent the strain inside Ge
calculated from theory, XRD measurement and Raman spectroscopy, respectively.
All the strain characterizations show that, we have successfully achieved up
to 2.3% in-plane biaxial tensile strain in Ge layers. To our knowledge, this is the
highest strained epitaxial Ge layer achieved.
80
3.3.4 Material and Optical Characterization of Strained Ge
Many of our samples have been inspected by TEM as well. Samples are
prepared by manual polishing and ion beam milling, or directly with focused ion
beam. Different TEM techniques, including cross-section bright field, dark field,
high resolution TEM, and selected area diffraction, are applied to our sample
inspections.
Figure 3.20 shows a cross section TEM image of a sample containing a 10
nm Ge layer grown between an In0.3Ga0.7As buffer layer and an In0.3Ga0.7As cap
layer.
Figure 3.20 (b) shows that most misfit dislocations are confined at the
interface between the In0.3Ga0.7As and the In0.15Ga0.85As buffer layers, and some
strain fringes appear around the defects due to the local strain induced by defects. A
faded line is shown in the middle of the In0.3Ga0.7As buffer layer. This is due to an
annealing step after the first 200 nm of In0.3Ga0.7As growth. Another 100 nm of
In0.3Ga0.7As is grown after annealing. The 10 nm Ge quantum well is grown on top
of this InGaAs buffer layer, then another 10 nm In0.3Ga0.7As is grown on top of the
Ge quantum well as a cap layer. In the zoom-in TEM image shown in figure 3.20 (c),
the two interfaces between In0.3Ga0.7As and Ge epitaxial layers are very clear and
smooth. The Ge epitaxial layer is about 9 nm thick and very uniform. The 10 nm
In0.3Ga0.7As cap layer on top of the Ge layer has many defects. These defects are
caused by anti-phase domains (APD) in the InGaAs cap layer which is due to the
polar material growth on non-polar material. The surface roughness is also induced
by these APD. The high resolution cross-section TEM image is shown in figure
81
3.20(d). The image shows a clear atom column fringe around the Ge QW region.
Although the intensity contrast between Ge and InGaAs is not very strong in high
resolution TEM mode, defects in the Ge layer should be clearly observed if they are
present. By examining the cross section of this sample, we confirm that the defects
density is in the 1x108 cm-2 range.
In0.3Ga0.7As 10 nm
Tensile strained Ge 9 nm
In0.3Ga0.7As 300 nm
In0.15Ga0.85As 200 nm
100 nm
GaAs substrate
(a) (b) Ge
Ge
20 nm
5 nm
(c) (d) Figure 3.20 Cross-section TEM image for strained Ge samples. (a) Schematic
of the sample structure, (b) cross-section TEM image of InGaAs buffer layers and
Ge QW, (c) cross-section TEM image of Ge QW, and (d) high resolution TEM
image of Ge QW region.
82
The biaxial tensile strain inside the Ge layer can be measured from high
resolution TEM images as well. The Ge in-plane lattice constant is calculated by
counting the number of atom columns in a certain length. The local tensile strain for
this sample is around 2%, however, with a relatively large variance. The large strain
variance is because this method only provides a measure of strain in a local region.
The selected area diffraction method gives some unique information about
the strain in a larger region. The advantage of diffraction in TEM compared to the
XRD method is that, in TEM, the electron beam is incident parallel to the top surface,
thus sampling an effectively thick region of the sample. Especially in the selected
area diffraction mode, the diffraction pattern is only generated from a small region.
Thus, there is no huge background diffraction peak from the GaAs substrate. Figure
3.21 shows the cross section of our sample. The dashed circles indicate the selected
regions from which the diffraction patterns are observed in the two cases.
Second case: In0.3Ga0.7As 100 nm
First case: In0.3Ga0.7As Figure 3.21: Cross section TEM images with indication of selected area for
diffraction images.
83
In the first case, the selected area is chosen at the interface of the In0.3Ga0.7As
and In0.15Ga0.85As buffer layers. A standard diffraction pattern is shown in figure
3.22 (a). However, in this diffraction pattern, each diffraction spot actually contains
two small diffraction spots: one from the In0.3Ga0.7As buffer layer, and the other one
from the In0.15Ga0.85As buffer layer. As shown in figure 3.22 (b) and (c), each pair of
two small spots are aligned together and pointed to the original point of the
diffraction pattern. This means all the buffer layers are totally relaxed so that the
out-of-plane lattice constant is equal to the in-plane lattice constant.
In the second case, when the selected area is on the top surface of the sample,
only the In0.3Ga0.7As buffer layer and strained Ge layer are selected. A similar
diffraction pattern is obtained as shown in figure 3.23 (a). However, in this second
case, the two spots are aligned together and all of them point to the [001] direction
which is the direction normal to the sample surface as shown in figure 3.23 (b) and
(c). This means the Ge epitaxial layer and In0.3Ga0.7As buffer layer have the same
in-plane lattice constant and different out-of-plane lattice constants. Thus, the
selected area diffraction images also prove that a totally relaxed InGaAs layer with
biaxial tensile strained Ge layer is achieved.
84
(c) 15%InGaAs
30%InGaAs
(a) (b) Figure 3.22: Selected area diffraction (a) at In0.3Ga0.7As and In0.15Ga0.85As
buffer layers as well as zoomed-in images (b) and (c).
(c) [001] Ge
30%InGaAs
(a) (b) Figure 3.23: Selected area diffraction (a) at Ge quantum well and In0.3Ga0.7As
buffer layer as well as zoomed-in images (b) and (c).
85
The optical properties of thin strained Ge layer are the most important as well
as the most difficult to characterize. Due to the small volume of the Ge epitaxial
layer, most of the transmission and reflection methods are hard to use. Fortunately,
PL can separately detect the signal generated from Ge epitaxial layers and InGaAs
buffer layers instead of combining these signals with the background signal from the
GaAs substrate. Also, PL is one of the most important characterization techniques
for electron luminescence and lasing. PL is also important for determining whether
the sample is a direct band gap material as discussed in chapter 2.
The low temperature (5 K) PL spectra of a set of Ge samples with different
biaxial tensile strain are plotted in figure 3.24. In these measurements, due to the cut
off of our InGaAs detector, only PL signals with wavelengths shorter than 1600 nm
are detected. In this set of Ge samples, a 10 nm thick InGaAs cap layer is grown on
top of Ge epitaxial layer to prevent the Ge exposure to air and eliminate the surface
states which decrease the PL intensity significantly. In this measurement, we can
decouple the PL signals from the InGaAs buffer layer materials from the Ge PL
signal. The bandgap of GaAs to In0.4Ga0.6As at room temperature changes from
1.420 eV to 0.893 eV (corresponding to emission anywhere between 870 nm to 1400
nm). Furthermore, as the sample is cooled to low temperatures (5 K), the bandgap
energy of InGaAs increases. Thus, the PL signals from the InGaAs buffer layers
correspond to even shorter emission wavelengths (the longest wavelength is around
1300nm at 5K). This simple analysis shows that the PL from InGaAs buffer layers
at 5K is shorter than 1300 nm, which is about 300 nm away from the peaks of the
86
epitaxially grown Ge layer PL spectrum. Consequently, the PL signals plotted in
figure 3.24 are generated in the Ge epitaxial layers only.
At 5 K, the indirect band gap energy of Ge increases from 0.66 eV to 0.8 eV.
Therefore, for the unstrained Ge sample, the photo luminescence signal at 1.6 μm is
from phonon assisted indirect band gap radiative recombination processes. The
sample with 0.3% and 0.9% tensile strained Ge is also measured with the same setup.
These two samples are Ge grown on In0.1Ga0.9As and In0.2Ga0.8As buffer layers
respectively. However, no PL signal other than the background noise is detected for
these two samples, which means the Ge epitaxial material quality is degraded,
because of the imperfect crystal quality of relaxed InGaAs buffer layers. The Ge
epitaxial layer grown on In0.3Ga0.7As buffer layer achieves 1.8% biaxial tensile strain.
A strong PL signal is detected for this highly tensile strained sample. The FWHM is
about 98 meV. For the Ge epitaxial layer on In0.4Ga0.6As buffer layer, the tensile
strain in the Ge reaches 2.3%. An even stronger PL signal is detected with peak
intensity at 1560 nm. The FWHM is about 53 meV. The broadening of the PL
peaks is due to non-uniform strain in the sample.
87
6000
2.33% tensile
strained Ge
5000
Intensity (a.u.)
4000
3000
1.81% tensile
strained Ge
2000
1000
Unstrained Ge
0
1300
1400
1500
λ (nm)
1600
0.34% & 0.92% tensile
strained Ge
Figure 3.24: Low temperature (5 K) PL of different tensile strained Ge layers.
The PL intensity changes for this set of samples at different temperatures are
also measured and plotted in figure 3.25. For the unstrained Ge sample, the PL
intensity does not change much. However, for Ge samples with large biaxial tensile
strain, especially the 2.33% tensile strained one, the PL signal increases by more
than 20 times when the sample is cooled from room temperature down to 5 K.
3
Intensity (a.u.) 2.5
2.33% tensile strained Ge 2
1.5
1.81% tensile strained Ge 1
0.5
Unstrained Ge 0
0
50
100
150
200
250
300
Temperature (K) Figure 3.25: Temperature dependent PL intensity for different tensile strained Ge.
88
Much information can be derived from these two PL measurement results.
The most important issue is the band gap energy in Ge epitaxial layers. In our case,
the problem is quite complicated since it combines the consideration of energy shifts
due to strain, temperature shift, quantum confinement, optical pump intensity and
many other effects. For Ge, when the temperature changes from 300 K to 5 K, the
bandgap energy is increased by 88 meV. In addition, we can also consider quantum
confinement within the Ge and InGaAs layers. Quantum confinement plays a strong
role in changing the ground state energy since the Ge epitaxial layer is so thin. For
the valence band and assuming an infinite well, the energy of the first energy state is
around 0.13 eV below the top of valence band energy. For 2% biaxial tensile
strained bulk Ge at room temperature, by using the data from the MIT group [15] as
well as from IBM [3], the direct band gap energy is about 0.5eV. Adding these
factors together, the direct band gap transition energy for 2% tensile strained Ge QW
at 5 K is about 0.72 eV. This energy corresponds to a photon wavelength at 1.7 μm,
which fits well with our experimental results. One more effect to be considered is
the PL peak wavelength change as a function of pumping laser energy. With optical
excitation, the quasi-Fermi level inside the Ge layer increases, and the luminescence
energy increases correspondingly. The detailed calculations for band gap energy are
not included here because it is very difficult to accurately determine the band gap
energy shift with many unknown parameters.
However, this is the first
characterization of the bandgap energy for such highly strained Ge epitaxial layers
and it provides much useful information for future study.
89
Although it is difficult to give a precise calculation of the band gap energy for
highly strained Ge, the intensity change during the cooling process is obvious and
can be clearly interpreted. For highly tensile strained Ge, when the sample is cooled
down to 5 K, the peak intensity increases by 20 times. At this lower temperature, all
the free carriers follow the Fermi-Dirac distribution and condense to the lowest
possible states. Thus, for a direct band gap material, all electrons fall into the gamma
point of the conduction band which is also the lowest energy point in the conduction
band. As a result, the PL signal increases quickly for direct band gap material during
the cooling process. On the other hand, a material with a quasi-direct band gap, for
example, heavily n-type doped Ge material [16], needs higher temperatures to
thermally excite the electron from the L-valley to the Gamma point. So the PL in
this indirect material increases during the heating process. Therefore, the strong
increase of the PL at low temperatures proves that direct band gap Ge is achieved
under high biaxial tensile strain.
3.4 Strained Ge Summary
In summary, we have used MBE to grow 2.33% biaxial tensile strained Ge on
InxGa1-xAs buffer layers, producing the highest reported strain in Ge to the best of
our knowledge. Low growth temperatures and low growth rates enable the 2D
growth of coherent tensile-strained Ge.
Material qualities and strain are
characterized using TEM, AFM, XRD, and Raman spectroscopy.
90
The strains
measured in the Ge from different methods are all in very good agreement with each
other and match theoretical predictions. Furthermore, optical characterization using
low-temperature PL shows a strong increase in PL intensity with higher strain and
lower temperature. Strong increase in PL intensity for a highly tensile strained Ge
epitaxial layer at low temperature is observed. Summing all the properties of tensilestrained Ge, a direct-bandgap tensile-strained Ge has been achieved and may be a
strong candidate for a group-IV laser. With the ability to synthesize highly-strained
layers of Ge, we can further investigate optical and transport properties of a material
that may have a place in future devices and optoelectronics on Si.
Reference:
[1] J. Oh, J. C. Campbell, S. G. Thomas, S. Bharatan, R. Thoma, C. Jasper, R. E.
Jones, and T. E.Zirkle, “Interdigitated Ge p-i-n photodetectors fabricated on a Si
substrate using graded SiGe buffer layers,” IEEE J. Quantum Electron. 38, 1238
(2002).
[2] Y.-H. Kuo, Y.K. Lee, Y. Ge, S. Ren, J.E. Roth, T.I. Kamins, D.A.B. Miller, and
J.S. Harris, “Strong quantum-confined Stark effect in germanium quantum-well
structures on silicon,” Nature 437, 1334-1336 (2005).
[3] M.V. Fischetti, and S.E. Laux, “Band structure, deformation potentials, and carrier
mobility in strained Si, Ge, and SiGe alloys,” J. Appl. Phys. 80, 2234 (1996).
91
[4] Y. Bai, K. E. Lee, C. Cheng, M. L. Lee, and E.A. Fitzgerald, “Growth of highly
tensile-strained Ge on relaxed InxGa1-xAs by metal-organic chemical vapor
deposition”, J. Applied Physics 104, 084518 (2008).
[5] Y. Hoshina, A. Yamanda, M. Konagai, “Growth and characterization of highly
tensile-strained Ge on InxGa1-xAs virtual substrate by solid source molecular beam
epitaxy,” Jpn. J. Appl.Phys. 48, 111102 (2009).
[6] http://www.thorlabs.us/newgrouppage9.cfm?objectgroup_id=2907
[7] Junxian Fu, Ph.D. thesis, (2005).
[8] Susan Marie Lord, Ph.D. thesis, (1993).
[9] H. Lin, Y. Huo, Y. Rong, R. Chen, T. I. Kamins, and J. S. Harris, “X-ray
diffraction analysis of step-graded InxGa1-xAs buffer layers grown by MBE,” 16th
International conference on Molecular Beam Epitaxy, Berlin, Germany(2010).
[10] D. Choi, Y. Ge, J. S. Harris, J. Cagnon, and S. Stemmer, “Low surface roughness
and threading dislocation density Ge growth on Si (001),” Journal of Crystal Growth
310, 4273–4279 (2008).
[11] T. Mizuno, S. Takagi, N. Sugiyama, H. Satake, A. Kurobe, and A. Toriumi,
“Electron and hole mobility enhancement in strained-Si MOSFET’s on SiGe-oninsulator substrates fabricated by SIMOX Technology,” IEEE Electron Device Letters,
Vol. 21, No. 5, (2000).
[12] D. Ahn, and S. Chuang, “Optical gain in a strained-layer quantum-well laser,”
IEEE Journal of Quantum Electronics, Vol 24. No 12, (1988).
92
[13] D. Leonard, M. Krishnamurthy, C. M. Reaves, S. P. Denbaars, P. M. Petroff, ,
“Direct formation of quantum-sized dots from uniform coherent islands of InGaAs on
GaAs surfaces,” Applied Physics Letters, Vol 63, 23, 3203 (1993)
[14] J. W. Matthews and A. E. Blakeslee, “Defects in epitaxial multilayers: I. Misfit
dislocations,” J. Cryst. Growth 27, 118 (1974)
[15] J. Liu, D. D. Cannon, K. Wada, Y. Ishikawa, S. Jongthammanurak, D. T.
Danielson, J. Michel, and L. C. Kimerling, “Silicidation-induced band gap shrinkage
in Ge epitaxial films on Si,” Applied Physics Letters, Vol 84, 5, (2004)
[16] S. Cheng, J. Lu, G. Shambat, H. Yu, K. Saraswat, J. Vuckovic, and Y. Nishi,
“Room temperature 1.6 μm electroluminescence from Ge light emitting diode on Si
substrate,” Optics Express, Vol. 17, No. 12, 10019 (2009).
4
93
Chapter 4
GeSn Alloys
4.1
Introduction
As mentioned in Chapter 1, different methods can be used to create a direct
bandgap Ge material. The strained Ge approach was discussed in Chapter 3. In this
chapter, we focus on forming a GeSn alloy and the materials growth and
characterization.
In recent years, Ge1-xSnx alloys have become very exciting research materials
due to their predicted properties of high carrier mobility (in excess of 105 cm2/V·s
electron mobility) [1], direct bandgap crossover (around x = 0.15) [2, 3], and
increased lattice constant (ranging from 5.66 Å to greater than 5.90 Å, enabling
strain as an energy band engineering tool for Ge). Moreover, GeSn alloys, which are
comprised of group IV atoms, can be grown on silicon, making them very useful for
monolithic integration of photonic devices on silicon wafers.
Although many
theoretical investigations have been done to study these alloys, strong experimental
evidence confirming such predictions has been limited because of the difficulty of
materials growth.
The most critical limitation is the low thermodynamic solid
solubility of Sn in Ge, which is less than 1%. MBE growth has been reported in
previous work [2, 4, 5]; however, the best quality samples with around 6.7% Sn
concentration in Ge exhibited Sn precipitation [5]. Others have produced films with
94
up to 34% Sn concentration, although the films have been either polycrystalline or
amorphous, making them unusable for electronic or photonic devices [4]. In this
chapter, we discuss our work on Sn growth rate calibration, GeSn alloy growth, and
characterization of high-quality Ge1-xSnx alloys with up to 9.2% Sn concentration in
GeSn alloys.
4.2
GeSn Materials growth
4.2.1 Approach to GeSn Alloy Growth
GeSn alloys have a much larger lattice constant than the Si lattice constant.
Therefore, two different growth methods have been used in this study. In the first
method, a Ge epitaxial layer is grown on a Si substrate by CVD or MBE, and then a
GeSn layer is grown on Ge in the MBE chamber. Ge growth on the Si (001)
substrates in the CVD system follows the technique of Choi et al. [6]. After three
cycles of the three-step process including low temperature growth, high temperature
growth, and high temperature annealing, the 1.4 μm thick Ge layer has 0.6 nm RMS
roughness over a 10 μm x 10 μm scan field. The threading dislocation density
measured by plan-view TEM is in the 1x107 cm-2 range. The Ge growth on a Si
wafer in the MBE system can produce good Ge cryatl quality in a much thinner
epitaxial structure. A 30 to 50 nm Ge layer is first grown on a Si (001) substrate at a
low temperature (around 300 ºC) to avoid 3D growth due to the 4% lattice
mismatching between Si and Ge. Then a second layer of 50 nm Ge is grown at 600
ºC to obtain good crystal quality. So, after a combined 100 nm growth of Ge on Si,
95
the smoothest Ge surface achieves 0.6 nm RMS roughness over a 1 μm x 1 μm scan
field. Although good Ge epitaxial layers can be achieved on a Si substrate, the
background signal from Ge layers affects many material and optical characterization
signals from a GeSn epitaxial layer on top of the Ge buffer layer. Thus, the second
method is discussed and applied in this chapter.
In the second method, layers are grown on (001) GaAs wafers. The GaAs
substrate wafers are first thermally treated in a III-V MBE chamber to remove all the
surface oxide and either a GaAs or InGaAs buffer layer is grown as discussed before.
Then, the wafer is transported under UHV (< 1x10-9 Torr) to a group IV MBE
chamber where Ge1-xSnx alloys are grown. Both Ge and Sn sources are thermal
effusion cells. The Ge growth rate is controlled from 0.6 nm/min to 1 nm/min.
During Ge1-xSnx growth, we used low substrate temperatures (100 ºC to 200 ºC) to
reduce Sn segregation and precipitation in Ge when the Sn concentration exceeds the
equilibrium solid solubility. To investigate the growth of high-quality Ge1-xSnx
alloys under unstrained or even tensile strained conditions, thick InGaAs buffer
layers were also used to provide a lattice constant larger than the Ge. For example,
Ge0.95Sn0.05 is grown on In0.10Ga0.90As buffer layers to achieve a nominally
unstrained GeSn epitaxial layer.
96
4.2.2 Sn Growth Calibration
Sn concentration in the sample needs to be calibrated first. However, the Sn
growth rate calibration is not trivial due to the small Sn flux pressure at low growth
rate. An ion gauge is mounted on the back side of the substrate holder and can be
flipped to face the Ge and Sn sources. Both Ge and Sn source flux pressures can be
monitored by this ion gauge and then used to calibrate to the growth rate and Sn
concentration from these data. Unfortunately, the ion gauge is not very consistent
during day to day running, therefore the Sn growth rate cannot be measured in all
cases. Thus, instead, the source growth rate is controlled by the source temperature
directly. In order to do so, the calibration samples are first grown in our chamber.
A 100 nm or thicker Ge epitaxial film is first grown on a GaAs (001)
substrate and measured by XRD and SEM to determine Ge growth rate. Ge growth
rates ranging from 0.75 nm/min to 0.96 nm/min are used in our sample growth. The
Sn growth rate calibration is done next using SIMS. The sample contains 4 layers of
Ge1-xSnx films with different Sn concentrations.
In each layer, the Ge source
temperature is fixed and the Sn source temperature increases from 880 ºC to 940 ºC.
Figure 4.1 presents the SIMS results for this sample, which shows discrete steps in
Sn composition corresponding to each layer. Using the previously determined Ge
growth rates, the Sn growth rates are calculated vs. source temperature and plotted in
figure 4.2.
97
940oC
920oC
900oC
880oC Figure 4.1: SIMS measurement for the calibration sample of 4 GeSn layers
grown directly on GaAs at 200 ºC.
Sn Growth Rate (nm/min)
0.08
0.06
0.04
0.02
0
1.05
1.1
1000/T (K
(K--11))
Figure 4.2: Sn growth rate vs. Sn evaporate source temperature.
The Sn growth rate and Sn concentration in GeSn can also be calculated from
XRD 1D rocking curve measurements of the lattice constant of a GeSn alloys.
Figure 4.3 presents three rocking curve measurements, showing the (004) GaAs
substrate, InGaAs buffer layers, and Ge1-xSnx epitaxial layer peaks.
The Sn
concentration calculated using the lattice constant from XRD and linear interpolation
of the Ge1-xSnx lattice constant is shown in Table 4.1. The growth rates from SIMS
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calibrations, which have similar results, are also listed in table 4.1. Rutherford
backscattering spectrometry (RBS) is one of the best tools to calibrate the Sn
concentration from first principles and will be used in the future. As a summary, all
the Sn concentrations mentioned in this section are derived from the SIMS calibrated
growth rate instead of XRD or RBS calibrated growth rates.
Figure 4.3 XRD (004) rocking curve for three different Sn concentration
GeSn layers on InGaAs buffer layers and GaAs substrate. (Courtesy of Hai Lin)
Sn concentration
from SIMS
Sn concentration from XRD,
Linear Interpolation
2.8%
2.7%
4.6%
3.6%
1.9%
1.4%
Table 4.1: Summary of Sn composition in three different samples using SIMS
calibrated growth rates and XRD lattice constant with linear lattice constant
interpolation between Ge and Sn.
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4.2.3 GeSn alloy growth mechanism
GeSn samples are grown with different recipes and are characterized by SEM
and AFM first. The most critical parameter is the Sn concentration. For example,
two typical samples are grown with almost the same recipe but different Sn
concentrations. In the first sample, 50 nm Ge0.954Sn0.046 is grown on an In0.1Ga0.9As
buffer layer at 200 ºC substrate temperature as shown in figure 4.4 (a). The surface
RMS roughness is about 0.5 nm in a 10 x 10 μm region. However, in the second
sample with 50 nm Ge0.93Sn0.07 on an In0.1Ga0.9As buffer layer also at a 200 ºC
substrate temperature, Sn phase segregation is clearly observed as shown in figure
4.4 (b). The surface RMS roughness is about 10 nm in a 10 μm x10 μm scan field.
Absence of Sn phase segregation in both SEM and AFM suggests that 200 ºC is
good for Ge0.954Sn0.046, however, it is not good for Ge0.93Sn0.07 epitaxial growth in the
MBE system.
(a)
(b)
Figure 4.4: SEM images from Ge0.954Sn0.046 (a) and Ge0.93Sn0.07 (b) on
In0.1Ga0.9As buffer layer.
100
Sample
Sn % (based on GR)
Substrate T (°C)
10 x 10 μm2 RMS
A
2.8%
200
0.451nm
B-200
4.6%
200
0.403nm
B-150
4.6%
150
0.428nm
B-100
4.6%
100
0.524nm
C
7.4%
100
0.385nm
D
9.2%
100
0.626nm
Table 4.2: GeSn growth samples and their 10 μm x 10 μm RMS surface
roughness from AFM measurements.
Many of our samples are characterized by non-contact mode AFM to ensure
smooth Ge1-xSnx epitaxial layer growth and give more quantitative characterization.
Our results in Table 4.2 show all the growth parameters and surface roughnesses. A
set of Ge0.954Sn0.046 epitaxial layers is grown at different substrate temperatures. As
shown in table 4.2 and figure 4.5, all the samples have quite smooth surfaces. The
surface roughness is 0.403 nm, 0.428 nm, and 0.524 nm for GeSn grown at 200 ºC,
150 ºC, and 100 ºC, respectively. Thus, a slightly smoother surface is obtained for a
200 ºC grown sample compared to a 100 ºC grown sample. Also, a set of GeSn
alloys with different Sn concentrations and the same 100 ºC sample growth
temperature are shown in table 4.2 and figure 4.6. With Sn concentration varied
from 4.6% to 9.2%, the sample surfaces are always smooth. Sample D is expected to
be rougher because we use two step-graded InGaAs buffer layers for lattice matching.
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(a)
(b)
(c)
Figure 4.5: AFM images of 50 nm Ge0.954Sn0.046 grown on In0.1Ga0.9As buffer
layer at (a) 200 ºC, (b) 150 ºC, and (c) 100 ºC substrate temperature.
(a)
(b)
(c)
Figure 4.6: AFM images of 50 nm Ge1-xSnx grown on InGaAs buffer layers at
100 ºC substrate temperature with Sn concentrations of (a) 4.6%, (b) 7.4%, and (c)
9.2%.
As a conclusion, GeSn epitaxial layers are grown at 100 ºC ~ 200 ºC substrate
temperature with 0.6 to 1 nm/minute growth rate.
The highest achieved Sn
concentration in GeSn alloys with a smooth surface is 9.2%.
For high Sn
concentration samples, such as 9.2% Sn concentration, the substrate temperature
must be decreased to avoid phase segregation. However, the low temperature may
reduce the crystalline quality and may limit the GeSn epitaxial layer thickness.
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4.2.4 GeSn Material Quality Characterization
To inspect the material quality of the Ge1-xSnx epitaxial layer in detail, we
utilize cross-section TEM on sample B-200 as listed in table 4.2, which is
Ge0.954Sn0.046 grown on an In0.1Ga0.9As buffer layer at 200 ºC substrate temperature.
In the cross section TEM image shown is figure 4.7, no Sn precipitation or phase
segregation was observed in any of the observed areas, which should both be visible
in TEM images if present [5]. Also, the absence of visible threading dislocations in
the cross section TEM suggests a threading dislocation density lower than or in the
1x108 cm-2 range.
GeSn 10% InGaAs GaAs Figure 4.7: Schematic (a) and cross section TEM image (b) of sample B-200,
which is Ge0.954Sn0.046 grown on an In0.1Ga0.9As buffer layer on a GaAs (001)
substrate.
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The high resolution TEM shows clear fringes from atom columns for this B200 sample. At the interface of the Ge0.954Sn0.046 layer and In0.1Ga0.9As buffer layer,
where the dark region is the InGaAs layer and the bright region is the GeSn layer, no
misfit dislocation are found in most of the areas examined as shown in figure 4.8 (a).
Figure 4.8 (b) shows clear GeSn diamond lattice structures without any defects, Sn
precipitation, or phase segregation. Figure 4.9 shows the selected area diffraction
pattern for the Ge0.954Sn0.046 region. The diffraction pattern shows a perfect diamond
lattice with small and sharp diffraction spots, which also proves that there is no Sn
phase segregation in this sample.
(b)
(a)
Figure 4.8: High resolution TEM image of sample B-200 (a) at GeSn and
InGaAs interface and (b) at the center of GeSn epitaxial layer.
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Figure 4.9: Selected area diffraction pattern for Ge0.954Sn0.046 epitaxial layer.
4.2.5 Strained GeSn
Strain is also important in GeSn band structure engineering. Since in-plane
biaxial tensile strained GeSn requires both less strain than pure strained Ge or less Sn
than lattice matched GeSn, it is easier to produce a direct bandgap GeSn alloy by
growing under tensile strain. Lattice relaxed InGaAs buffer layers which have larger
lattice constants than our chosen GeSn alloy lattice constant are used as the substrate.
A set of GeSn samples was grown on InGaAs buffer layers to achieve different
tensile strain in GeSn layers with the same alloy composisiton. In these three
samples, the same 35nm of Ge0.95Sn0.05 are coherently grown on three different
In0.12Ga0.88As, In0.15Ga0.85As, and In0.18Ga0.82As buffer layers, and named sample A,
B, and C, respectively. Theoretically, in sample A, the Ge0.95Sn0.05 is lattice matched
to the In0.12Ga0.88As buffer layer. This is also confirmed by the XRD 1D rocking
curve scan. The magenta curve in figure 4.10 shows only one diffraction peak other
than the GaAs substrate peak. This proves that the In0.12Ga0.88As has a similar out 105
of-plane lattice constant to that of the Ge0.95Sn0.05. For sample B, which is the GeSn
on In0.15Ga0.85As buffer layer sample, the yellow curve shows a larger out-of-plane
lattice constant for the InGaAs buffer layer as expected. Also, in the same yellow
curve, a shoulder shows up on the right side of the InGaAs diffraction peak, which
comes from the Ge0.95Sn0.05 epitaxial layer. The out-of-plane lattice constant for
sample B Ge0.95Sn0.05 is not only smaller than the In0.15Ga0.85As’s out-of-plane lattice
constant, but also smaller than the out-of-plane lattice constant of sample A with a
Ge0.95Sn0.05 layer on an In0.12Ga0.88As buffer layer. This phenomenon can be more
clearly observed in sample C, which is Ge0.95Sn0.05 on top of an In0.18Ga0.82As buffer
layer; the corresponding XRD rocking curve is plotted as the blue curve in figure
4.10. In this case, the out-of-plane lattice constant of Ge0.95Sn0.05 gets even smaller.
Through the Poisson ratio relationship, we can conclude that the Ge0.95Sn0.05
epitaxial layer is under biaxial tensile strain with In0.15Ga0.85As or In0.18Ga0.82As
buffer layers.
106
1000000
5% GeSn
InxGa1-xAs buffer layer
100000
GaAs
GaAs substrate
10000
InGaAs
1000
100
10
GeSn
1
64
64.5
65
65.5
66
66.5
67
67.5
Figure 4.10: XRD 1D rocking curve for Ge0.95Sn0.05 epitaxial layer on three
different InGaAs buffer layers. Inset figure shows the schematic of the sample layer
structure.
In order to correctly calibrate the strain in GeSn epitaxial layers, more XRD
characterization was done for the tensile strained GeSn samples on InGaAs buffer
layers. As shown in figure 4.11, sample C, which is Ge0.95Sn0.05 on top of an
In0.18Ga0.82As buffer layer, is scanned for XRD 2D reciprocal lattice mapping. As
shown in this figure, the InGaAs buffer layer is almost totally relaxed, and the GeSn
is fully coherently grown on top of the InGaAs buffer layer. Most importantly, the
Ge0.95Sn0.05 epitaxial layer is about 0.2% tensile strained in this sample.
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Strain-line
GaAs
GeSn
InGaAs
Relaxationline
Figure 4.11: XRD 2D reciprocal lattice mapping for Ge0.95Sn0.05 In0.18Ga0.82As buffer layer sample.
4.3
GeSn Band Gap Energy Characterization
As discussed in the previous sections, good quality and biaxial tensile
strained GeSn epitaxial layers can be grown in our MBE system and are ready for
band gap energy characterization. Since InGaAs and GaAs buffer layers are used
instead of Ge or GeSn buffer layers, the signal is cleaner and easier to analyze.
However, the relatively low band gap energy of GeSn layers limits our choice of low
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noise photo detectors and our measurement techniques. Here we only discuss the
transmission spectrum results from spectro-photometer measurements.
Although the amount of Sn incorporated in Ge is below the predicted amount
for direct-bandgap GeSn, we conducted optical transmission measurements to
characterize the shifts in the absorption edge. In order to have enough absorption in
the GeSn layers, 150 nm of lattice-relaxed GeSn epitaxial layers with different Sn
concentrations were grown on GaAs substrates for this measurement. Figure 4.12
shows the optical transmission spectra for three GeSn samples with 1.5%, 3.5%, and
5.7% Sn concentration, respectively. The decrease in transmission around 870 nm is
due to the GaAs substrate absorption, and wavelength dependent reflections are not
accounted for in this transmission measurement. We assume that the energy of the
transmission curve bending point represents the band gap energy of the GeSn layers.
Although this is not entirely accurate, it does show the general trend in the band gap
energy as the Sn composition changes. This figure shows the band gap energy
decrease with increasing Sn composition, which is also plotted in figure 4.13. This
trend is consistent with the predicted band gap shift to lower energies with increasing
Sn concentration.
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Transmission
0.5
0.4
0.3
Sn concentration
0.2
1.5%
3.5%
5.7%
0.1
0
800
1000
1200 1400 1600
Wavelength (nm)
1800
2000
Figure 4.12: Transmission spectra for 150 nm thick Ge1-xSnx samples grown
on GaAs substrates.
Direct band gap energy (eV)
0.8
0.78
0.76
0.74
0.72
0
1
2
3
4
Sn concentration (%)
5
Figure 4.13: Approximate band gap energies of GeSn epitaxial layers with
different Sn compositions.
110
4.4
Summary
In conclusion, we have grown high Sn composition (up to 9.2%) Ge1-xSnx
alloys using low-temperature MBE growth. Growth rates are calibrated using SIMS,
XRD, and SEM. The Sn compositions calculated from growth rates match fairly
well with those back calculated from lattice constant measurements in XRD. The
parameter space for GeSn epitaxial layer growth was systemically explored. The
final GeSn samples have excellent material quality as determined by SEM, AFM and
TEM. Strained GeSn layers were also characterized by XRD 1D and 2D scans. This
leads us to believe that GeSn layers with about 10% or more Sn and small tensile
strained are very possible to grow by MBE and these can lead to direct band gap
Ge1-xSnx and become a very useful material system for group IV based photonics
devices.
Reference:
[1] J. D. Sau, and M. L. Cohen, “Possibility of increased mobility in Ge-Sn alloy
system”, Physical Review B 75, 045208 (2007).
[2] J. Kouvetakis, J. Menendez, and A.V.G. Chizmeshya, “Tin-based group IV
semiconductors: new platforms for opto- and microelectronics on silicon,” Ann. Rev.
of Mater. Res. 36, 497 (2006).
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[3] A. Moontragoon, Z. Ikonic and P. Harrison, “Band structure calculations of Si–
Ge–Sn alloys: achieving direct band gap materials,” Journal Semiconductor Science
and Technology, 22, 7 (2007).
[4] Gang He and Harry A. Atwater, “Synthesis of epitaxial SnxGe1-x alloy films by
ion-assisted molecular beam epitaxy,” Nuclear Instruments and Methods in Physics
Research Section B: Beam Interactions with Materials and Atoms, Volume 106, Issues
1-4, 126 (1995).
[5] S. Takeuchi, Y. Shimura, O. Nakatsuka, S. Zaima, M. Ogawa, and A. Sakai,
“Growth of highly strain-relaxed Ge1-xSnx/virtual Ge by a Sn precipitation controlled
compositionally step-graded method,” Applied Physics Letters 92, 231916 (2008)
[6] D. Choi, Y. Ge, J. S. Harris, J. Cagnon, and S. Stemmer, “Low surface roughness
and threading dislocation density Ge growth on Si (001),” Journal of Crystal Growth
310, 4273–4279 (2008).
5
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Chapter 5
Summary and Future Work
5.1 Summary
This dissertation describes an investigation of the growth and characterization
of strained Ge and GeSn material aimed at achieving a group IV direct band gap
semiconductor material for realizing CMOS compatible electro-optical devices.
The on-chip interconnect bandwidth limitation is becoming an ever greater
critical challenge to device scaling. Silicon photonics is capable of solving this
emerging problem because of its high speed, high bandwidth, and low power
capabilities.
Most of the key devices in Si photonic ICs have already been
extensively studied and applied in real applications. However, a practical siliconcompatible coherent light source is still a major challenge.
This thesis mainly
focuses on the band structure engineering of Ge to produce a direct band gap
semiconductor.
The background of silicon photonics, material growth and
characterization techniques are briefly reviewed in Chapters 1 and 2.
In Chapter 3, we first discuss our work on highly biaxial tensile strained Ge
on top of fully-relaxed InGaAs buffer layers grown in our coupled Group IV and IIIV MBE systems. A low temperature growth with high temperature annealing of step
graded InGaAs buffer layers was investigated.
113
InGaAs buffer layers were
characterized and proven to have smooth surfaces (0.6 nm RMS surface roughness in
1x1 μm2 scanning field), with up to 40% indium concentration, and about 85% strain
relaxation within several hundreds of nanometers of buffer layer growth. These
InGaAs buffer layers have a threading dislocation density less than 1x108 cm-2, and
strong PL signals. Up to 2.3% in-plane biaxial tensile strained thin Ge epitaxial
layers were coherently grown on top of InGaAs buffer layers. The strain-dependent
and temperature-dependent PL intensity from Ge epitaxial layers suggests that a
direct band gap semiconductor is achieved.
In Chapter 4, the growth and characterization of GeSn were discussed. Using
a low temperature MBE growth method, more than 9% Sn incorporation in Ge was
achieved, which is 10 times greater than the solid solubility limit of Sn in Ge. Good
crystalline quality as well as no Sn precipitation or phase segregation were
demonstrated by SEM, AFM, XRD, and TEM. The direct band gap narrowing with
increasing Sn percentage was observed by optical transmission measurements, which
is consistent with theoretical prediction.
This dissertation work provides a
foundation for realizing an efficient germanium based CMOS compatible integrated
coherent light sources.
5.2 Suggestions for Future Work
While the investigations in this thesis have provided a foundation toward
realizing a Group IV based photonic source, a number of interesting areas remain for
future research are suggested by the work presented in this dissertation, including
materials, device design and system integration.
114
5.2.1 Ge Material Studies
The initial band engineering effort focused on tensile strained Ge grown on
InGaAs buffer layers and was discussed in Chapter 3 of this thesis. There are also
other good buffer layer materials. AlGaAs buffer layers are interesting because of
their wider band gap of AlGaAs, which can provide stronger quantum confinement.
The totally CMOS compatible solution is to grow relaxed GeSn buffer layers on top
of Si substrates. The study of epitaxial Ge buffer layers on Si substrates and relaxed
GeSn buffer layers on GaAs has already been done. The growth optimization and
strained Ge/GeSn on Si substrates growth need further investigations.
The second band engineering effort was investigation of GeSn alloys and up
to 9% Sn incorporation and up to 0.2% tensile strained GeSn alloys were achieved.
However, no strong PL was observed. Further material growth with higher Sn
composition and higher strain and crystalline quality improvements by annealing are
important directions to investigate further. More materials property characterization,
such as Raman shift dependence on Sn composition and strain, direct and indirect
band energy crossing, and carrier mobility should also be pursued.
Ge quantum dots (QD) are also interesting for integrated light sources since
the 3-dimensional-confinement gives better light emitting properties. More studies
on the material growth, optical properties and electrical carrier injection for tensile
strained Ge QDs are crucial.
115
5.2.2 Ge Electro-Optical Device Studies
A photodiode is not only one of the most fundamental optical devices, but a
particularly useful characterization tool on the path to developing a useful, Si-based
optical source. Both metal-semiconductor-metal photo detectors and vertical PIN
photo detectors are useful for strained SiGe, Ge, or GeSn layers bandgap energy
characterization. GeSn photo detectors can also be used for long wavelength (~2 μm)
light detection.
Achieving highly efficient electroluminescence from both strained Ge and
GeSn is one of the critical milestones for realizing an on-chip integrated coherent
light source. More studies need to be done on the MBE growth of strained GeSn to
first realize a useful LED. Achieving a useful LED, one can then incorporate a
micro-disk, photonic crystal or a distributed Bragg reflector cavity to provide optical
feedback; finally achieving a CMOS compatible, integrated, electrically pumped
laser.
5.2.3 On-chip Optical Interconnect System
The next step is to integrate all the electro-optical devices on a Si chip to
demonstrate an on-chip high-speed optical interconnection. Our group has already
successfully demonstrated a vertical GeSi quantum confined Stark effect electroabsorption modulators. A novel waveguide modulator is under development now,
116
and the structure is shown in figure 5.1. Furthermore, based on the GeSn/SiGeSn
materials system, the waveguide coupled light source and photo detector can be
integrated together, as shown in figure 5.2. This architecture could yield a complete
on-chip optical communication system.
Ge Quantum
N-SiGe
W
P-SiGe
P-Si
id
SiO2
Figure 5.1: Schematic of waveguide coupled GeSi electro-absorption
modulator.
Light source
Modulator
Photodetector
SiGeSn cap layer
GeSn QWs
SiGeSn buffer layer
Si waveguide
SiO2
Si substrate
Figure 5.2: Schematic of on-chip an optical interconnect system with light
source, modulator, and photodetector.
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