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GaN nanostructures grown by selective area growth for solid-state lighting

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GaN nanostructures grown by selective area growth for solid-state lighting

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GAN NANOSTRUCTURES GROWN BY SELECTIVE AREA GROWTH FOR
SOLID-STATE LIGHTING

by

Ting-Wei Yeh

A Dissertation Presented to the
FACULTY OF THE USC GRADUATE SCHOOL
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY
(MATERIALS SCIENCE)

December 2012

Copyright 2012 Ting-Wei Yeh

ii
Dedication

To my dear wife, Yili Huang,
and to my lovely daughters, Danica A. Yeh, and Petra E. Yeh

iii
Acknowledgements
I would like to give my gratitude to my thesis advisor, Prof. P. Daniel Dapkus, for
his guidance and support during these years, as it has been my aspiration to study GaN
nanostructures since I joined Compound Semiconductor Laboratory (CSL). Dr. Dapkus’
passion towards the research and his profound experience in MOCVD growth further
inspired my enthusiasm and perseverance to work on exotic GaN nanostructure growth.
Through his mentorship, I have gained and learned hands-on experiences from the
projects that I engaged in, especially the modification of the GaN reactor and the
patterning technique. Based on these fundamentals, the growth of GaN nanostructures
can be explored and exciting results of NanoLED can be achieved.
In addition, I would like to thank Prof. Steven Nutt and Prof. Chongwu Zhou for
being the committee members for my defense and qualifying exam, and Prof. Edward
Goo and Prof. Atul Konkar for being the committee members for my qualifying exam.
The work in this dissertation could not be fully accomplished without the
collaboration with other groups. I would like to thank Raymond Sarkissian, who worked
with me on dry etching and cathodoluminescence measurement; Dr. Byungmin Ann, who
prepared some TEM samples; and Saptaparna Das, who did time-resolved
photoluminescence measurement with me.
For administrative support, no work could have been done smoothly without the
assistance from Jenny Lin, Eliza Aceves, and Loring Smith. In the cleanroom, Donghai

iv
Zhu helped me fix the equipments I needed. I would like to thank all of their supports
during these years.
I would like to thank my former CSL colleagues: Dr. Yuanming Deng, who
trained me on the operation of the GaAs reactor; Dr. Dawei Ren, who shared his
experience of the GaN growth and the assistance of pattern preparation by laser
interferomatric lithography; Dr. Hyung-Joon Chu, who trained me on using the Raith
eLine lithography system; Dr. Lawrence Stewart, who helped me in the cleanroom work
and collaborated on GaAs material growth; Dr. Ruijuan Li, who discussed and shared the
ideas in both research and raising kids; and Dr. Zhen Peng and Dr. Qi Yang for their help
in the cleanroom.
I would like to thank my current CSL colleagues, Chun-Yung Chi, Yen-Ting Lin,
and Maoqing Yao, for their assistance of maintaining the reactors and helps in the
experiments in the growth room. I also would like to thank Suzana Sburlan, Yunchu Li,
Yoshitake Nakajima, Krisna Bhargava for their support and discussion. I would like to
thank all my friends as well in letting me have an unforgettable Ph.D. life.
Finally, I would like to thank my father, who encourages and supports me on
pursuing my Ph.D. degree at USC. I also would like to thank my mother for the faith she
passed on to me to rely on when I faced frustrating challenges in both research and life.
Lastly, I would like to thank my wife, Yili, for her unconditional support and for taking
care of our two girls, Danica and Petra, when I focused on wrapping up my dissertation
and experiments in the last few months.

v
Table of Contents
Dedication........................................................................................................................... ii 
Acknowledgements............................................................................................................iii 
List of Tables ..................................................................................................................... ix 
List of Figures..................................................................................................................... x 
Abstract............................................................................................................................ xvi 
Chapter 1 Introduction ........................................................................................................ 1 
1.1 InGaN based LEDs ........................................................................................... 1 
1.2 InGaN-based LEDs as the engines for solid-state lighting............................... 2 
1.2.1 Efficient blue LEDs as pumps for phosphors to produce white light 3 
1.2.2 Blue, green, and red multicolor LEDs for lighting ............................ 4 
1.2.3 White light generated by photon recycling........................................ 5 
1.3 Current challenges for InGaN-based LED for lighting applications ................ 6 
1.3.1 High dislocation density caused by the growth on lattice-mismatched
substrates..................................................................................................... 6 
1.3.2 Strain induced piezoelectric fields in multiple quantum wells .......... 6 
1.3.3 Efficiency droop................................................................................. 7 
1.3.4 Nonpolar GaN substrate..................................................................... 9 
1.3.5 Light extraction efficiency................................................................. 9 
1.4 Nanostructure growth...................................................................................... 10 
1.4.1 Nanowires grown by vapor-liquid-solid (VLS)............................... 10 
1.4.2 Selective area growth (SAG) ........................................................... 15 
1.5 III-V materials grown by MOCVD system .................................................... 18 

vi
1.5.1 Epitaxial growth by MOCVD.......................................................... 18 
1.5.2 GaN reactor and system modification.............................................. 20 
1.6 Thesis outline.................................................................................................. 22 
1.7 Chapter references .......................................................................................... 23 
Chapter 2 Fabrication and Characterization of GaN Nanostructures Grown by Selective
Area Growth...................................................................................................................... 28 
2.1 SAG mask preparation.................................................................................... 28 
2.1.1 Photolithography.............................................................................. 28 
2.1.2 Electron-beam lithography............................................................... 30 
2.1.3 Laser interferometric lithography .................................................... 31 
2.1.4 Nanoimprint lithography.................................................................. 33 
2.1.5 Nanosphere lithography................................................................... 34 
2.1.6 Block-copolymer lithography .......................................................... 35 
2.2 SAG GaN nanostructures................................................................................ 37 
2.2.1 GaN nanostripes............................................................................... 37 
2.2.2 GaN nanopyramids .......................................................................... 41 
2.2.3 GaN nanorods .................................................................................. 43 
2.2.4 GaN nanosheets ............................................................................... 46 
2.3 GaN nanorod selective area growth mechanism............................................. 48 
2.3.1 Pulsed growth on (0001) Ga-polar plane......................................... 48 
2.3.2 Continuous growth on (0001) N-polar plane................................... 55 
2.4 Crystal structure analysis and sample preparation.......................................... 57 
2.4.1 TEM sample preparation by focused ion beam ............................... 57 
2.4.2 Crystal structure of GaN nanorods analyzed by transmission electron
microscope ................................................................................................ 59 
2.4.3 Heterostructures investigated by scanning transmission electron
microscope ................................................................................................ 62 

vii
2.5 Summary......................................................................................................... 63 
2.6 Chapter references .......................................................................................... 64 
Chapter 3 InGaN/GaN Multiple Quantum Wells Growth on GaN Nanorods.................. 68 
3.1 The dependence of mask design on GaN nanorod growth ............................. 69 
3.2 Multiple quantum well growth on GaN nanorods .......................................... 71 
3.3 Light emission from multiple quantum wells grown on GaN nanorods studied
by photoluminescence........................................................................................... 73 
3.4 Multiple quantum wells grown on GaN nanorods studied by transmission
electron microscopy.............................................................................................. 79 
3.5 Local emission of multiple quantum wells grown on GaN nanorods
investigated by cathodoluminescence................................................................... 83 
3.6 Summary......................................................................................................... 88 
3.7 Chapter references .......................................................................................... 88 
Chapter 4 InGaN/GaN Multiple Quantum Well Growth on GaN Nanosheets................. 91 
4.1 The dependence of mask design on GaN nanosheet growth .......................... 92 
4.2 Multiple quantum well growth on GaN nanosheets ....................................... 96 
4.3 Light emission from multiple quantum wells on GaN nanosheets studied by
photoluminescence................................................................................................ 97 
4.4 Multiple quantum wells grown on GaN nanosheets studied by transmission
electron microscopy............................................................................................ 101 
4.5 Local emission of multiple quantum wells grown on GaN nanosheets
investigated by cathodoluminescence................................................................. 104 
4.6 Summary....................................................................................................... 106 
4.7 Chapter references ........................................................................................ 106 
Chapter 5 InGaN/GaN Nano-scale Light Emitting Diodes ............................................ 108 

viii
5.1 Planar InGaN-based light emitting diode ..................................................... 108 
5.2 NanoLEDs grown on GaN nanorod arrays................................................... 120 
5.2.1 Dopant concentration estimated by capacitance-voltage measurement
................................................................................................................. 120 
5.2.2 Resistivity measurement of single GaN nanorod........................... 125 
5.2.3 NanoLED structures....................................................................... 127 
5.2.4 I-V characteristics of NanoLED grown on GaN nanorods............ 132 
5.2.5 Electroluminescence of NanoLED grown on GaN nanorods........ 133 
5.3 Summary....................................................................................................... 138 
5.4 Chapter references ........................................................................................ 138 
Chapter 6 Conclusions and Future work......................................................................... 142 
6.1 Conclusions................................................................................................... 142 
6.2 Future work................................................................................................... 144 
6.3 Chapter references ....................................................................................... 145 
Bibliography ................................................................................................................... 146 

ix
List of Tables
Table 2-1 Etch pit density comparison between two different GaN bulk materials......... 40 
Table 2-2 Plane estimation from the SEM images recorded at 0° and 45°. ..................... 42 
Table 5-1 The estimated indium composition of InGaN vs. growth temperatures......... 116 

x
List of Figures
Figure 1-1 Schematic diagram of spectrum and structure of a white LED. ....................... 4 
Figure 1-2 Schematic diagram of recombination mechanisms in an InGaN-based
LED.[19] ............................................................................................................................. 8 
Figure 1-3 A metal particle dispersed on a substrate surface. (b) A metal alloy formed in
a raised temperature to wet the sample surface. (c) Nanowire growth from a metal alloy
reaching its supersaturation condition. ............................................................................. 11 
Figure 1-4 1-nm-thick Au film annealed in H
2
/AsH
3
ambient for 10 mins to form metal
particles on a GaAs substrate surface. .............................................................................. 12 
Figure 1-5 (a) GaAs nanowires grown by VLS on a GaAs (111)B substrate. (b) InAs
nanowires grown on a GaAs (111)B substrates................................................................ 12 
Figure 1-6 (a) GaAs nanowires grown by 1 nm Au film. (b) GaAs nanowire growth by 40
nm Au nanoparticels using TBAs as arsenic precursor. ................................................... 13 
Figure 1-7 (a) Schematic illustration of patterned Au discs surrounded by a SiN
x
mask on
top of the GaAs substrate surface before nanowire growth. (b) A GaAs nanowire array is
formed by using position-controlled Au discs confined in openings of a SiN
x
film. ....... 14 
Figure 1-8 (a) A SiN
x
dielectric layer is deposited on top of a GaAs wafer. (b) The
dielectric layer is patterned by electron beam lithography. (c) GaAs nanowires are grown
from the openings of the mask.......................................................................................... 15 
Figure 1-9 (a) GaAs nanowires grown by selective area growth on a GaAs (111)B
substrate with small opening size. (b) Thicker GaAs nanowires grown on a mask with
larger hole diameters exhibit clear six {1-10} facets. The vertical sidewalls are {1-10}
planes. ............................................................................................................................... 16 
Figure 1-10 GaAs grown on a mask with stripe openings prepared along <11-2>-
orientation to form vertical sidewalls as a waveguide structure....................................... 17 
Figure 1-11 (a) 100 pairs of AlGaAs/GaAs superlattice are grown on a GaAs (100)
substrate. (b) Enlarged FE-SEM image of the superlattice demonstrates the uniform
control of layer thickness.................................................................................................. 19 
Figure 1-12 Schematic diagram of system layout............................................................. 20 

xi
Figure 2-1 (a) GaAs tetrahedral shaped structures grown in micron-size openings on a
GaAs (111)B surface. (b) A FE-SEM image of a GaAs tetrahedron recorded at higher
magnification. (c) GaN hexagons grown in hole opening arrays in micrometer scale on a
nitrogen-polar GaN surface. (d) A magnified FE-SEM image of a GaN hexagon with six
vertical {1-100} sidewalls. ............................................................................................... 29 
Figure 2-2 (a) Si-doped (0001) GaN bulk material is utilized to grow GaN nanostructures.
(b) SiN
x
is deposited on top of the GaN surface by plasma enhanced chemical vapor
deposition (PECVD). (c) A dot array is prepared by EBL and transferred into the SiN
x
30 
Figure 2-3 (a) AFM image of a stripe array pattern transferred into a SiN
x
mask. (b) A
line scan across the line indicated in (a) to estimate the depth of the mask, which is
around 30 nm. (c) FE-SEM image of the stripe pattern. The periodicity is 220 nm and the
opening is 63 nm. (d) FE-SEM image of a dot array pattern generated by exposure the
resist in two perpendicular directions. .............................................................................. 32 
Figure 2-4 Procedures of two NIL process: (a) Thermal NIL and (b) UV-NIL [3] ......... 34 
Figure 2-5 Schematic diagram of block copolymer lithography for GaAs quantum dots. (a)
Copolymer microdomain monolayer applied on a SiN
x
mask after annealing in vacuum.
(b) After ozone treatment. (c) During CF
4
plasma etching process. (d) hole array pattern
transferred to GaAs surface. (e) GaAs quantum dots selectively grown in the holes. (f)
GaAs dots after SiN
x
mask removal.[8]............................................................................ 36 
Figure 2-6 Stripe masks prepared with different spacings and different electron beam
dosage by EBL.................................................................................................................. 38 
Figure 2-7 (a), (b), and (c) shows the triangular stripes grown on different center-to-
center pitches, 500, 750, 1000 nm, respectively, before forming a coalescent GaN layer.
(e), (f), and (g) demonstrates coalescent GaN layer with different center-to-center
spacings............................................................................................................................. 39 
Figure 2-8 (a) A nanopyramid array is grown by SAG growth. (b) A top view FE-SEM
image of nanopyramids shows nanopyramids are formed by the six inclined planes...... 42 
Figure 2-9 (a) Schematic diagram of a GaN nanorod array grown vertically from its
substrate. The enlarged diagram shows two dominant facets formed by GaN nanorods. (b)
Uniform GaN nanorod array is grown by SAG by MOCVD. The arrows indicate the
polar and nonpolar facets of the nanorods. In the inset, a top view of hexagonal GaN
nanorods shows the six vertical sidewalls. Ordered GaN nanorod arrays grown on 250 nm,
500 nm, 750 nm, and 1 µm center-to-center spacings are demonstrated in (c−f),............ 45 
Figure 2-10 (a) A uniform nanosheet array with 500 nm center-to-center spacing grown
on a GaN/Al
2
O
3
buffer layer patterned with a stripe patterns. (b) Polar, semipolar, and
nonpolar planes are indicated in the schematic diagram of GaN nanosheets. (c) FESEM

xii
image taken with the sample rotated 90° with respect to (a). The nanosheets show vertical
and parallel sidewalls, which are {1-100} planes. (d) A nanosheet array grown for double
the growth cycles of (c). Each of the FESEM images was recorded at a 60° angle. The
scale bar is 500 nm in all figures.[14]............................................................................... 47 
Figure 2-11 Growth comparison between continuous and pulsed injection of gas flux. (a)
Vertical GaN nanorods structure can be formed by pulsed-mode GaN growth. (b) During
continuous growth, the slow growth rate of the {1-101} semipolar plane leads to form
GaN nanopyramids. (c) The growth scheme of the pulsed growth mode. T
1
and T
3
are the
injection time of TMG and ammonia, respectively. T
2
and T
4
are the growth interruption
between switching TMG and ammonia. ........................................................................... 50 
Figure 2-12 Schematic diagram of the pulsed growth mechanism for SAG GaN nanorod.
The yellow balls are the Ga adatoms. ............................................................................... 51 
Figure 2-13 GaN nanorods are grown on nitrogen-polar GaN surface. ........................... 56 
Figure 2-14 (a) Trenches prepared by FIB for lift-out process. (b) A sample is separated
from its substrate and mounted on a micro-manipulator to transfer to a TEM grid. (c) A
nanorod sample mounted on a TEM grid is sectioned along <11-20>-direction to study 58 
Figure 2-15 (a) A FESEM image of GaAs nanowires grown from low temperature GaAs
nuclei on a Si (111) substrate. (b) High density of SFs observed in the GaAs nanowire. 59 
Figure 2-16 TEM images taken at <11-20> zone axis. No SF is observed in both TEM
images recorded at (a) high magnification and (b) low magnification. (c) Selective area
diffraction pattern confirms the positioning of the nanorod as well as the single wurtzite
........................................................................................................................................... 60 
Figure 2-17 TEM images taken at <1-100> zone axis. No SF is observed in the TEM
images taken at (a) high magnification and (b) low magnification. Wurtzite crystal....... 61 
Figure 2-18 Comparison of STEM image between (a) bright field and (b) dark field..... 63 
Figure 3-1 (a) Nanorod diameter is proportional to hole diameter. (b) Nanorod height
changes with respect to its diameter. (c) m-plane surface area estimated from the product
of the height and diameter of GaN nanorods. (d) Nanorod height varies with the fill factor
of growth mask. The error bars are the standard deviations of sampling over 20 nanorods
for each point. ................................................................................................................... 70 
Figure 3-2 Comparison of the PL spectra before and after ICP etching. The inset shows
the c-plane and semipolar planes are removed on the tips of GaN nanorods................... 74 

xiii
Figure 3-3 The emission wavelengths measured from nanorod samples grown with
different opening sizes are shown in the spectra. The inset shows the emission
wavelength decreases as the opening size on the dielectric mask increases..................... 76 
Figure 3-4 (a) The decay curves comparison between GaN nanorods with and without
MQWs. (b) The fitted curve (black) includes the contribution from the instrument
response (red). Blue curve is the measured result and the other blue curve at the bottom is
the residuals of the fitting. ................................................................................................ 78 
Figure 3-5 (a) The nanorod is sectioned along <11-20>-direction to investigate the MQW
growth on three different planes. The inset shows the diffraction pattern of the <11-20>
zone axis and a beam blocker blocks the center beam. Three colored boxes indicate the
locations where magnified images are taken in Figure 3-5 (c), (d), and (e). (b)
InGaN/GaN MQWs are grown on three different planes as indicated in the arrows. (c)
TEM image shows thick a quantum well is grown on the c-plane, polar plane. (d) Thin
MQWs are grown on the semipolar plane. (e) MQWs are grown on the nonpolar plane. (f)
........................................................................................................................................... 81 
Figure 3-6 (a) Thickness variation of MQW grown along the axial direction of GaN
nanorod. The pitch spacing is 400 nm. (b) Decrease in thickness of MQWs from the inner
well to the outer well. The sample is sectioned from a GaN nanorod array with pitch of
500 nm. ............................................................................................................................. 83 
Figure 3-7 CL analysis from the GaN nanorod grown in the pattern with 400 nm center-
to-center spacing in trigonal arrangement. (a) CL spectra were taken at the spots
designated in the SEM image. (b) The comparison of CL spectra from the corresponding
spots in the nanorods. (c) The spectra comparison of peak positions with normalized
intensity from spot 1 to 9. ................................................................................................. 84 
Figure 3-8 CL analysis from the GaN nanorod grown in the pattern with 500 nm center-
to-center spacing in trigonal arrangement. (a) CL spectra were taken at the spots
designated in the SEM image. (b) The comparison of CL spectra from the corresponding
spots in the nanorods. (c) The spectra comparison of peak positions with normalized
intensity from spot 1 to 9. ................................................................................................. 85 
Figure 3-9 (a) A uniform array of GaN nanorods with sharp and terminated c-plane is
achieved after 10 minutes of continuous GaN growth. The inset shows the c-plane is
minimized or terminated on top of the nanorods. The FE-SEM image was recorded at a
45° angle. (b) CL spectrum is recorded at the boxed region in the inset. CL mapping
results at 430 and 480 nm are shown in (c) and (d), respectively..................................... 87 
Figure 4-1 Width and height of GaN nanosheets affected by the orientation of the growth
pattern. (a) The nanosheet width shows angle-dependent lateral growth within the small
........................................................................................................................................... 93 

xiv
Figure 4-2 Growth transition of nanosheet width varied by pattern orientations. Average
width of nanosheets varying with growth cycles is dependent on the orientation of the
stripe opening on the mask. The error bar is the standard deviation from the measured
result.................................................................................................................................. 95 
Figure 4-3 Room temperature photoluminescence measured from InGaN/GaN MQWs
grown on GaN nanosheet arrays. Photoluminescence (PL) of three pairs of MQWs grown
........................................................................................................................................... 98 
Figure 4-4 (a) Comparison of PL spectra of nanosheets before and after ICP etching. (b)
FE-SEM image of GaN nanosheets with three pairs of quantum wells before ICP etching.
(c) FE-SEM image of GaN nanosheets after ICP etching for 1 minute.......................... 100 
Figure 4-5 TEM bright field images taken from a sectioned nanosheet with MQWs. (a)
InGaN/GaN MQWs were grown on three different planes as indicated in the arrows. The
scale bar is 10 nm. (b) A vertical GaN nanosheet grown from its bulk material is shown
in the low magnification image. The magnified images taken in Figure 4-5 (c), (d), and (e)
are indicated in b. Nanosheet surface was covered with carbon to prevent ion beam
damage during sample preparation. The scale bar is 20 nm. (c) TEM image shows thick
MQWs are grown on the c-plane, polar plane. The scale bar is 10 nm. (d) Thin MQWs
are grown on the semipolar plane. (e) MQWs are grown on the nonpolar plane. The scale
bar is 2 nm in (d) and (e)................................................................................................. 103 
Figure 4-6 Dislocation bending towards the sidewalls of the GaN nanosheet. Dislocation
bending was observed in the nanosheet in the bright field TEM cross-sectional image.
Three pairs of quantum wells were grown after the GaN nanosheet growth. The sample
surface was coated with platinum in the FIB system to avoid ion beam damage. The scale
bar is 50 nm..................................................................................................................... 104 
Figure 4-7 CL spectrum of MQWs grown on GaN nanosheets. The inset shows an
overlap of a CL mapping with a FE-SEM image. .......................................................... 105 
Figure 5-1 Dependence of the FWHM of the HRXRD rocking curves for the (002)
reflections on the V/III ratio during the growth of GaN buffer layer. ............................ 111 
Figure 5-2 Dependence of the concentration and mobility on the V/III ratio during the
growth of the GaN buffer layer....................................................................................... 111 
Figure 5-3 The comparison of growth temperature with indium incorporation efficiency
investigated by HRXRD. ................................................................................................ 115 
Figure 5-4 Schematic structure of a InGaN-based LED................................................. 117 

xv
Figure 5-5 Cross-sectional HRTEM images of MQWs with p-GaN epilayers grown at (a)
1050°C and (b) 900°C, and EDS indium composition profile across (c) MQW grown at
1050°C and (d) MQW gown at 900°C. [16]................................................................... 118 
Figure 5-6 (a) EL spectrum of a blue LED wafer. (b) A picture of a blue LED after turned
on by a forward bias........................................................................................................ 119 
Figure 5-7 (a) An example of a Schottky diode. (b) A schematic diagram of a Schottky
barrier formed by a metal with work function lager than that of a semiconductor. ....... 121 
Figure 5-8 (a) Schematic diagram of CV measurement for a GaN nanorod array. (b)
Exposed tips of GaN nanorods are shown at the Au contact boundary.......................... 124 
Figure 5-9 (a) A FE-SEM image of two contacts prepared on a GaN nanorod for........ 126 
Figure 5-10 (a) Schematic diagram of NanoLEDs. (b) A FE-SEM cross-sectional image
of the NanoLED. A coalescent p-GaN layer is grown on top of the nanorods. (c) Surface
morphology of a coalescent p-GaN over a trigonal nanorod array with 400 nm pitch
spacing. (d) Surface morphology of a coalescent p-GaN over a square nanorod array with
500 nm pitch. .................................................................................................................. 129 
Figure 5-11 (a) Top view FE-SEM image of non-planar NanoLED. (b) A connected p-
GaN layer covers the top of nanorods to form a non-planar surface.............................. 130 
Figure 5-12 (a) Top view FE-SEM image of NanoLED with a SiO
2
current blocking layer.
(b) FE-SEM image of the p-GaN grown laterally to link all the nanorods..................... 131 
Figure 5-13 (a) A NanoLED is sectioned by FIB to study the nanorods embedded in the
p-GaN layer. (b) A magnified FE-SEM image taken from the boxed region in (a)....... 132 
Figure 5-14 (a) A picture of a NanoLED after being turned on. (b) I-V curve of the
NanoLED. ....................................................................................................................... 133 
Figure 5-15 (a) EL spectra of a NanoLED under different injection current levels. (b) EL
spectrum measured at 90 mA is fitted by two Gaussian peaks. (c) EL emission peak shifts
with respect to the injection current................................................................................ 134 
Figure 5-16 (a) EL intensity versus injection current density. (b) Relative efficiency
varies with injection current density............................................................................... 136 
Figure 5-17 (a) Demonstration of different EL spectra from NanoLEDs (b) Pictures of
NanoLEDs with different emission wavelengths under forward bias. ........................... 137 

xvi
Abstract
There are several challenges for the InGaN-based light emitting diodes (LEDs) to
be used for lighting, such as high dislocation density resulting from the lattice mismatch
between GaN and sapphire substrates, strong piezoelectric fields induced by the strain in
InGaN/GaN multiple quantum wells (MQWs) grown on the polar direction, and fair
amount of light trapped inside GaN material due to the difference in refractive indices
between GaN and air. Also, the efficiency droop is an important issue for InGaN-based
LEDs to be driven at high current for solid-state lighting.
In this thesis, GaN nanostructures, especially GaN nanorods and GaN nanosheets,
were studied as potential solutions to overcome the current issues that the conventional
planar InGaN-based LEDs are facing. Nanostructures can help to release the strain from
the lattice mismatch between GaN and its conventional sapphire substrates. Dislocation
bending was observed in the nanostructures. Therefore, GaN nanostructures contribute to
forming dislocation-free templates for InGaN/GaN MQWs. Large surface area of
nonpolar planes was formed by these nanostructures to serve as growth templates for
InGaN/GaN MQWs to eliminate the piezoelectric fields. Both nanorod and nanosheet
structures can improve light extraction efficiency due to their nonplanar geometries for
reducing the total reflection loss at the GaN/air interface. The feasibility of growing
thicker MQWs on nonpolar substrate has been demonstrated and these thick MQWs
grown on the nonpolar facets of GaN nanostructures can contribute to increasing the
possibility of radiative recombination and reducing the carrier density in MQWs to

xvii
mitigate the efficiency droop. NanoLEDs are successfully fabricated on GaN nanorod
arrays and electrically driven up to 200 A/cm
2
with a relative maximum efficiency about
60 A/cm
2
. This result demonstrates the potential of using NanoLEDs as efficient emitters
for solid-state lighting.
1
Chapter 1 Introduction
1.1 InGaN based LEDs
A compound semiconductor light emitting diode (LED) was first reported by
Round in 1907. It was a SiC Schottky diode.[1] In the years since, several arsenide- and
phosphide-based III-V compound semiconductors have been studied for their
optoelectronic properties, such as GaAs, AlGaAs, GaP, GaAsP, and AlGaInP, etc.
However, these materials emitting at wavelengths in the range of visible wavelengths
restrict themselves to be used in lighting or full color displays due to their smaller energy
band gaps. Many wide band gap materials have been explored to emit short wavelengths
in ultraviolet (UV) or blue spectra. GaN is one of these materials and the emission
wavelength can be manipulated by the incorporation of indium. Single crystal GaN films
were first reported by Maruska and Tietjen in 1969.[2] The first electroluminescence
from GaN was announced in 1971 by Pankove et al.[3] However, the structure was a
metal-insulator-semiconductor (MIS) diode, not a conventional p-n junction, due to the
difficulty in achieving p-type GaN. The first p-type doping in GaN was reported by
Akasaki’s group in 1989 using low-energy electron beam irradiation.[4] They
demonstrated the first p-n homojunction LED emitting in the range of UV and blue in
1992.[5] Efficient blue LEDs and laser diodes with p-type GaN activated by high
temperature post-growth anneal were demonstrated by Nakamura et al.[6-10] This anneal
approach has been widely adopted in the fabrication of InGaN-based LEDs. This early
2
research opened the door for InGaN-based materials to be studied for application to many
optoelectronic devices.
InGaN-based LEDs are likely to greatly impact our daily lives due to their
compact size and high efficiency. These LEDs can be integrated into variety of
applications, such as cell phones, computers, TVs, traffic lights, and full color outdoor
displays in stadiums. The global need to conserve energy has created a strong potential
application for white light emitting InGaN-based LEDs as efficient, environmentally
friendly electronic devices. They are viewed as the likely candidates to substitute for the
conventional incandescent light bulbs and fluorescent lamps in general lighting
applications.

1.2 InGaN-based LEDs as the engines for solid-state lighting
Due to wide distribution of energy band gaps ranging from 0.7 to 3.4 eV, In
x
Ga
1-
x
N has drawn lots of attention for optoelectronic device applications. Ideally, this ternary
material covers the whole range of visible light. However, the lattice mismatch between
InGaN and GaN results in degradation of material quality for high In compositions, such
as the generation of dislocations and the reduction in radiative recombination caused by
the strain-induced piezoelectric fields. Therefore, the efficiency of InGaN-based LEDs
decreases with increasing emission wavelength. AlGaInP is a quaternary material suited
for high-brightness emission in red, orange and yellow. However, its efficiency drops
dramatically with decreasing emission wavelength because the energy band gap becomes
3
indirect with high Al and Ga content. This efficiency reduction phenomenon, usually
called “Green gap” becomes severe at green wavelengths for both InGaN and AlGaInP
materials. None of these materials are, at present, efficient enough to cover a wide range
of spectra for lighting. Therefore, efficient white light LEDs are usually achieved by a
combination of different color LEDs or wavelength converters.

1.2.1 Efficient blue LEDs as pumps for phosphors to produce white light
Luminous efficiency and color-rendering index of LEDs are the most important
factors for the general lighting applications. There are several approaches to fabricate
white LEDs for solid-state lighting. These approaches use either UV or blue LEDs to
excite one or several phosphors and they can be classified in dichromatic, trichromatic,
and tetrachromatic approaches.[11] The most widely adopted approach is using blue
LEDs to pump yellow phosphors that allows people to perceive white light from the
mixture of blue and yellow light. A schematic diagram of using a blue LED to pump
phosphors is shown in Figure 1-1.
4

Figure 1-1 Schematic diagram of spectrum and structure of a white LED.
In general, dichromatic white LEDs have a higher luminous efficiency than
several other types of white LEDs. The color-rendering capability is lowest for
dichromatic sources and it increased with the multi-chromaticity of the source. There is
another approach to fabricate white LEDs with improved color rendering by pumping
blue, green, and red phosphors with UV LEDs. However, the poorer efficiencies of UV
LEDs and the inherent wavelength-conversion loss hinder this approach from being
widely adopted for lighting. Therefore, white light generated by the combination of blue
LEDs and yellow phosphors is still most common approach for solid-state lighting.

1.2.2 Blue, green, and red multicolor LEDs for lighting
The color rendering has been a concern for the white LEDs using the combination
of blue LEDs and phosphors due to the relative weakness in the red portion of the
5
spectrum. To have more even luminous intensities in blue, green, and red colors, an
approach of combining the three primary colors has been carried out by adopting three
different kinds of LEDs to form white LEDs, emitting in blue, green, and red
wavelengths. Although the color rendering is improved for this form of white LED, the
spatial separation of these LED chips results in changes of color temperature when
viewed from different angles. Appropriate design and positioning the LED chips are
crucial to improve this issue. Using this methodology will also increase the cost and
reduce the yield of the white LED. Therefore, the combination of blue, green, and red
LEDs in a single package is mainly applied to outdoor displays.

1.2.3 White light generated by photon recycling
LEDs using semiconductor wavelength converters have been demonstrated by
Guo et al.[12] The structures consist of a InGaN-based blue LED pumping source and an
electrically passive AlGaInP epitaxial layer. The light emitted from the blue LED is
absorbed by a AlGaInP secondary active region which re-emits lower energy photons.
The calculated efficiency can exceed 300 lm/W for an ideal device with unity quantum
efficiency. Due to the narrow emission spectra of the two light sources, the photon-
recycling semiconductor LED demonstrates high luminous efficiency but low color-
rendering index.[12]
All the approaches mentioned above require efficient LEDs as pumps or light
sources for lighting applications.
6
1.3 Current challenges for InGaN-based LED for lighting applications
1.3.1 High dislocation density caused by the growth on lattice-mismatched
substrates
The large lattice mismatch between the GaN materials and the sapphire substrates
upon which they are grown results in strain and high dislocation densities. The high
dislocation density of GaN based materials, around 10
9
~10
10
cm
-2
, is attributed to the
16% lattice mismatch between GaN and sapphire. As compared to the other III-V
compounds, there is a small impact of such high dislocation densities on the radiative
efficiency of blue LEDs, possibly due to the formation of localized states formed by
indium fluctuations as recombination centers.[13] However, the efficiency of UV LEDs
is strongly dependent on the dislocation density, which still limits the application of the
GaN-based material for optoelectronic devices.

1.3.2 Strain induced piezoelectric fields in multiple quantum wells
InGaN/GaN multiple quantum well (MQW) active regions are usually employed
in LED devices to increase the capture of electrons and holes and to promote radiative
recombination in the active region. The presence of large spontaneous polarization fields
and piezoelectric fields caused by the strain inside the active region affects the efficiency
of quantum wells used for light emission. The piezoelectric field inside the multiple-
quantum-well active region causes spatial separation of the electron and hole
wavefunctions; consequently, thin quantum wells, approximately 3 nm thick, are grown
7
to increase the radiative recombination efficiency.[14, 15] The thin structure leads to
inefficient electron capture and high carrier concentrations at the operating current that
can lead to Auger recombination. Both of these effects have been implicated in the high
current efficiency “droop” observed in blue LEDs.[16-18]

1.3.3 Efficiency droop
For the lighting applications, InGaN-based LEDs have to be driven at relatively
high currents . However, a phenomenon, called “efficiency droop”, is observed when the
current density reaches tens ampere per square centimeter. A drop in the light emission
efficiency is only observed in the nitride-based LEDs that is not directly related to
heating of the device. The total current can be divided into two parts: (1) carriers that
generate photons in the QW (I
rad
) and (2) carriers that are lost to other processes (I
lost
).
The observation of efficiency droop is a measure related to the external quantum
efficiency, which is the product of internal quantum efficiency (IQE) and the light
extraction efficiency. Therefore, efficiency droop is dependent upon the IQE inside the
LEDs in general. IQE can be expressed as the following equation.
( 1-1)
Efficiency droop occurs when the I
lost
increases stronger than I
rad
. The
recombination mechanism can be illustrated in Figure 1-2. They can be categorized in
8
four different mechanisms: spontaneous emission reduction, defect-assisted mechanisms,
Auger recombination, and carrier leakage.

Figure 1-2 Schematic diagram of recombination mechanisms in an InGaN-based
LED.[19]
Carrier recombination mechanisms outside the QWs can be indicated as carrier
leakage (I
leak
). Total injection current can be simply expressed into four parts:
(1-2)
Proposed mechanisms resulting in the efficiency droop can relate to one of the
current components in the equation above. Collaborative experiments have also been
conducted at University of Southern California (USC) and University of California-
Santa Barbara (UCSB) to have better understanding of the droop behavior.[20, 21]
9
Although several droop mechanisms have been proposed, none of them are conclusive
enough to be generally accepted.

1.3.4 Nonpolar GaN substrate
Growth on nonpolar or semipolar GaN substrates have been explored as methods
for reducing the piezoelectric field and increasing the quantum well thickness to increase
the radiative recombination efficiency in the active region of InGaN/GaN LEDs.[22-24]
GaN substrates can be grown by different approaches, such as hydride vapor phase
epitaxy and ammonothermal method.[25-27] Nonpolar and semipolar GaN substrates can
be cut from specific orientations from the thick GaN bulk. However, the high cost of
these specialized substrates has prohibited widespread adoption.

1.3.5 Light extraction efficiency
The refractive index difference between GaN and air causes optical reflection at
the interface. Assume the refractive index of GaN material is 2.4 and that of air is 1. The
reflectance and transmittance at the interface can be calculated from the following
equations.
(1-3)
(1-4)
10
Only 83% of the light can transmit through the GaN/air interface from the normal
incident angle. Although these equations are probably too simplified for the actual case,
the reflection at GaN/air interface can be roughly estimated. Another loss is the critical
loss resulting from light coming out from a material with higher refractive index into
another material with lower refractive index. The critical angle can be calculated from
Snell’s law.
(1-5)
When θ
air
is equal to 90°, the critical angle is 24.6°. The result indicates large
portion of light is trapped inside GaN material. Therefore, several different surface
texturing techniques have been applied to alleviate this loss to enhance the performance
of LEDs.[28-30]

1.4 Nanostructure growth
Nanostructures have drawn lots of attention due to their appealing physical
properties. There are several approaches to grow III-V nanostructure materials. Two
major approaches are described in the following sections, vapor-liquid-solid (VLS) and
selective area growth (SAG), respectively.
1.4.1 Nanowires grown by vapor-liquid-solid (VLS)
Silicon nanowhiskers grown by the VLS mechanism were discovered and
demonstrated by Wagner and Ellis in 1964.[31] This approach utilizes metal particles to
11
form alloys with the semiconductor at an elevated temperature. Figure 1-3 shows a
schematic diagram of nanowire growth by VLS.

Figure 1-3 A metal particle dispersed on a substrate surface. (b) A metal alloy formed in
a raised temperature to wet the sample surface. (c) Nanowire growth from a metal alloy
reaching its supersaturation condition.
In the case of III-V nanowires, the growth temperature is slightly above the
eutectic point of the III-elements and the metal particles. Therefore, the metal alloys
usually form round balls in the liquid phase. During growth, the growth system keeps
supplying III-elements to let the metal balls stay in supersaturation condition. The III-
and V-elements will precipitate out from the metal balls and combine with the V-element
from the gas phase to form III-V materials. The growth orientation of the nanowires is
usually determined by the directions with slow growth rates on the sidewalls. Growth
along <111>-direction is the typical orientation to grow vertical nanowires in most
studies. In the preliminary study, thin Au film (1 nm) is used to form metal particles as
seeds to grow GaAs nanowires using VLS approach. The diameters of these metal
particles are widely distributed due to lack of controllability of this formation process. An
example is shown in Figure 1-4.
12

Figure 1-4 1-nm-thick Au film annealed in H
2
/AsH
3
ambient for 10 mins to form metal
particles on a GaAs substrate surface.
Figure 1-5 demonstrates vertical GaAs and InAs nanowires grown by VLS. The
diameter and height vary across the whole samples.

Figure 1-5 (a) GaAs nanowires grown by VLS on a GaAs (111)B substrate. (b) InAs
nanowires grown on a GaAs (111)B substrates.
To form a more uniform nanowires by using VLS approach, the dimension of the
metal particles has to be controlled. One of the approaches is to disperse size-controlled
Au particles immersed in a colloidal solution onto a substrate surface. The following
13
SEM images taken at the same magnification show the comparison between using a Au
film and Au particles to grow GaAs nanowires.

Figure 1-6 (a) GaAs nanowires grown by 1 nm Au film. (b) GaAs nanowire growth by 40
nm Au nanoparticels using TBAs as arsenic precursor.
Using size-controlled nanoparticles can control the uniformity of GaAs nanowires
in diameter and height. However, the random distribution of nanoparticles on the
substrate surface limits the accuracy in control of nanowire density and location.
Position-controlled GaAs nanowires can be achieved by growing nanowires in confined
regions to avoid the migration of the metal alloys on the substrate surface. Figure 1-7
demonstrates a GaAs nanowire array growing from openings of a SiN
x
mask patterned by
electron beam lithography. The schematic diagram shows a cross-sectional illustration of
patterned Au discs confined by SiN
x
on a GaAs substrate prior to GaAs nanowire growth.
The ratio of diameter and thickness of Au discs was found to be crucial to form single
nanowire in each opening.
14

Figure 1-7 (a) Schematic illustration of patterned Au discs surrounded by a SiN
x
mask on
top of the GaAs substrate surface before nanowire growth. (b) A GaAs nanowire array is
formed by using position-controlled Au discs confined in openings of a SiN
x
film.
Although this technique can achieve vertical nanowire growth, there are still some
issues about using VLS to grow semiconductor materials. The possible incorporation of
the metal particles inside the nanowires could create unwanted impurities inside the
nanowires. Also, the surface tension plays an important role in achieving heterojunctions
in axial directions.[32] The metals evaporated from the alloys can deposit on the
susceptors and reactor wall. The susceptor and the quartz liner need to be cleaned by
chemicals and dried by nitrogen to remove the metal deposition. More frequent cleaning
routines are expected when using the VLS growth technique.

15
1.4.2 Selective area growth (SAG)
Selective area growth is a growth technique to grow materials in preferential
regions created by growth masks, which are generally made of dielectric materials due to
their thermal and chemical stability. Also, the cracking efficiency of precursors on the
heterointerface results in the growth selectivity. Patterned masks can be prepared by
various kinds of lithography techniques, such as photolithography, electron beam
lithography, and block copolymer lithography, etc. The following Figure 1-8 shows a
schematic diagram of SAG GaAs nanowire growth.

Figure 1-8 (a) A SiN
x
dielectric layer is deposited on top of a GaAs wafer. (b) The
dielectric layer is patterned by electron beam lithography. (c) GaAs nanowires are grown
from the openings of the mask.
This approach has been applied in forming dislocation free GaN layers by lateral
epitaxial overgrowth of the masking layer for InGaN-based lasers. Vertical nanowire
arrays have also been grown using this growth technique. The following figure exhibits
vertical GaAs nanorod grown by SAG. This anisotropic growth is based on the growth
rate differences on different facets. It requires slow growth planes on the sidewalls to
16
promote the growth rate on the tips to form vertical structures. Therefore, substrates with
specific orientations are required for this kind of research. Similar to the VLS growth,
<111>-orientations are usually chosen for the SAG nanowire growth. Figure 1-9 shows
some SAG nanowires grown by this technique.

Figure 1-9 (a) GaAs nanowires grown by selective area growth on a GaAs (111)B
substrate with small opening size. (b) Thicker GaAs nanowires grown on a mask with
larger hole diameters exhibit clear six {1-10} facets. The vertical sidewalls are {1-10}
planes.
In Figure (b), the clear hexagons are formed due to the {1-10} facets with slowest
growth rate. The flat surface of the {1-10} can be utilized to form continuous vertical
sidewalls, as shown in Figure 1-10. A waveguide can be grown without surface damage
caused by dry etching process. Although the light can be confined by the parallel
sidewalls, it is challenging to grow a flat layer on (111)B plane with low refractive index
prior to the waveguide growth.
17

Figure 1-10 GaAs grown on a mask with stripe openings prepared along <11-2>-
orientation to form vertical sidewalls as a waveguide structure.
SAG technique is adopted in this study to avoid possible impurity incorporations
during growth. The growth conditions between GaAs and GaN are quite different. The
typical GaAs nanowire growth condition cannot be transferred to GaN nanorods directly.
Another growth scheme, the pulsed growth mode, is introduced to grow vertical GaN
nanorods or nanosheet. Its growth mechanism will be discussed in Chapter 2.

18
1.5 III-V materials grown by MOCVD system
The materials including planar LED structures and novel nanostructures studied in
this thesis are mainly grown in a close-coupled showerhead metalorganic chemical vapor
deposition (MOCVD) system.

1.5.1 Epitaxial growth by MOCVD
MOCVD is the epitaxial technique used to grow III-V crystals in this thesis. The
he unique aspect of this process is that relatively involatile metals can be transported in
the vapor phase to a reaction zone by the use of metalorganic compounds of interest (Ga
in the growth of GaN). The metalorganic compounds (usually liquids) have relatively
high vapor pressure that allows them to be carried in a carrier gas by passing the carrier
gas through the compounds in temperature- and pressure-controlled metalorganic vessels,
called bubblers. This process facilitates accurate control of the metalorganic sources
introduced into the growth chamber.[33] The gas control system, Epifold, allows quick
switching of the gaseous sources to form abrupt interfaces. Therefore, MOCVD can
provide a precise control of both compositions and film thickness with atomic layer
precision. An example of MOCVD-grown superlattice is shown in Figure 1-11. 70% Al
content of AlGaAs is estimated by X-ray diffraction. The dark contrast of AlGaAs was
enhanced by dipping the sample into buffered oxide etchant (BOE) for 5 minutes. The
AlGaAs and GaAs layers are 15 and 19 nm, respectively. This structure has been applied
to form self-assembled nanometer-scale wires of organic materials. [34]
19

Figure 1-11 (a) 100 pairs of AlGaAs/GaAs superlattice are grown on a GaAs (100)
substrate. (b) Enlarged FE-SEM image of the superlattice demonstrates the uniform
control of layer thickness.
Different metalorganic sources are available for the GaN-related material growth,
such as trimethylgallium (TMG), triethylgallium (TEG), trimethylaluminum (TMA), and
trimethylindium (TMIn), etc. Ammonia is the most commonly used hydride source for
the nitrogen. Disilane and biscyclopentadienylmagnesium (CP
2
Mg) are the dopants for n-
type and p-type doping for GaN. In the case of GaP- or GaAs-based materials, group V
metalorganic sources are also available, such as tertiarybutylarsine (TBA) and
tertiarybutylphosphine (TBP). Diethylzinc (DEZn) and tetrabromomethane (CBr
4
) are
precursors for p-type dopants. This epitaxial technique has been widely applied to III-V
material growth for optoelectronic devices. Also, the dimension of the reactor has been
scaled up to meet the requirement of reducing cost nowadays.
20
1.5.2 GaN reactor and system modification
GaN materials have been grown in the Compound Semiconductor Laboratory
(CSL) for several years. The previous GaN manifold system used in CSL was not
sophisticated enough for more complicated structure growth because there were many
limitations in the system design. For example, one pressure controller controlled all the
pressure of the bubblers. Thus, it had insufficient flexibility to test wider range of growth
conditions, in which metalorganic sources with different partial pressures are needed. The
reactor was modified by the author to grow GaN materials in Summer 2007. The design
is to install the previous GaN chamber into an existing GaAs reactor system for GaN
growth. The hydride gases used in the system are quite different; ammonia and arsine are
the precursors for GaN and GaAs, respectively. Also, the GaAs charcoal scrubber was
not suitable for ammonia. Thus, separate exhaust lines were needed for both chambers.
The layout of the system and modified items on different cabinets are listed below.

Figure 1-12 Schematic diagram of system layout

21
Modified items:
• Ammonia cabinet
1. Anchor the cabinet
2. Run ammonia line to gas cabinet
3. Provide ventilation
• Gas control cabinet
1. Add ammonia line to process control system
2. Provide nitrogen line for bubbler sources
3. Modify three bubbler sources to provide nitrogen as optional carrier gas
4. Replace the gold seals and plungers for ammonia injection in the Epifold
• Purifier cabinet
1. Install power supply system to the heater for GaN growth
2. Separate vent lines for two growth chambers located in the reactor cabinet
• Reactor cabinet
1. Install nitride reactor (reactor 1)
2. Modify manual gas control panel
3. Provide pipe connection to reactors
4. Modify exhaust line
The phosphine line was partially converted into ammonia line for the Reactor 1 to
grow GaN material. After the modification of the MOCVD system, arsenide- and nitride-
based materials can be grown in the same system to make the system become versatile to
work on different projects. Although the system is capable of growing both nitride- and
22
arsenide-based materials in different chambers, sharing the same gas control system and
long purge time for switching chambers make the system hard to switching chambers
frequently. By virtue of the modification, GaN research continues. With the ammonia
injection through Epifold to the reactor facilitated the pulsed growth mode, also called
flow-rate modulation growth, used to grow GaN vertical nanostructures in this work. The
pulsed growth mode will be described in Chapter 2.

1.6 Thesis outline
Chapter1 describes the background of white LEDs and the current challenges for
the solid-state lighting using InGaN-based materials. The epitaxial tools and techniques
for nanostructure growth are also mentioned.
GaN nanorod and nanosheets are grown from opening hole and stripe arrays on
the growth masks by selective area growth. The dimension and uniformity of the
openings are crucial to grow ordered arrays of nanorods and nanosheets. Therefore, a
sophisticated technique, electron beam lithography, was utilized as the approach to
prepare the growth masks. The growth mechanism and growth mask preparation are
described in chapter 2. The crystal structure of GaN nanorod analyzed by transmission
electron microscopy (TEM) and TEM sample preparation by a focused ion beam system
are included in this chapter as well.
23
Chapter 3 presents the characterizations of InGaN/GaN MQWs grown on GaN
nanorods. TEM and cathodoluminescence (CL) verify the presence of MQWs
predominantly grown on non-ploar planes. Average light emission from the MQW is
studied by photoluminescence (PL) and local emission is analyzed by CL. The carrier
lifetimes in MQWs grown on GaN nanorod arrays are measured by time-resolved PL.
Chapter 4 exhibits the growth of MQWs on the unique GaN nanosheet structure.
Compared to the isolated GaN nanorod structure, the linearly connected structure of GaN
nanosheets provide a conducting path for device fabrication. TEM reveals the MQWs
grown on polar, semipolar, and non-polar planes. The variation of well width is based on
the growth rate on different planes. Light emission from the MQWs grown on GaN
nanosheet arrays is studied by PL and CL.
Chapter 5 demonstrates InGaN-based LEDs on GaN nanostructures. The planar
counterpart, LED structure growth on c-plane sapphire substrates, will be described first.
Electrical properties of GaN nanorods are studied by capacitance-voltage (CV) and
conductance measurement. Electroluminescence of GaN NanoLEDs is also demonstrated.
Conclusions and future work will be summarized in chapter 6.

1.7 Chapter references
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24

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[11] E. F. Schubert, Light-Emitting Diodes, 2003.

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25
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[15] Y. L. Li, R. Huang, and Y. H. Lai, "Efficiency droop behaviors of InGaN/GaN
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[16] M.-H. Kim, M. F. Schubert, Q. Dai, J. K. Kim, E. F. Schubert, J. Piprek, and Y.
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[17] Y. C. Shen, G. O. Mueller, S. Watanabe, N. F. Gardner, A. Munkholm, and M. R.
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[18] K. J. Vampola, M. Iza, S. Keller, S. P. DenBaars, and S. Nakamura,
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[19] J. Piprek, "Efficiency droop in nitride-based light-emitting diodes," Physica
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[20] R. Sarkissian, S. Roberts, T.-W. Yeh, S. Das, S. Bradforth, J. O'Brien, and P. D.
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(2012).

[21] R. Sarkissian, S. Roberts, T.-W. Yeh, S. Das, S. Bradforth, J. O'Brien, and P. D.
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[22] A. Chakraborty, B. A. Haskell, S. Keller, J. S. Speck, S. P. Denbaars, S.
Nakamura, and U. K. Mishra, "Demonstration of nonpolar m-plane InGaN/GaN
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[23] T. Koyama, T. Onuma, H. Masui, A. Chakraborty, B. A. Haskell, S. Keller, U. K.
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"Prospective emission efficiency and in-plane light polarization of nonpolar m-
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26
[24] H. Zhong, A. Tyagi, N. N. Fellows, F. Wu, R. B. Chung, M. Saito, K. Fujito, J. S.
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[25] K. Naniwae, S. Itoh, H. Amano, K. Itoh, K. Hiramatsu, and I. Akasaki,
"GROWTH OF SINGLE-CRYSTAL GAN SUBSTRATE USING HYDRIDE
VAPOR-PHASE EPITAXY," Journal of Crystal Growth, 99, 381, (1990).

[26] S. S. Park, I. W. Park, and S. H. Choh, "Free-standing GaN substrates by hydride
vapor phase epitaxy," Japanese Journal of Applied Physics Part 2-Letters, 39,
L1141, (2000).

[27] T. Hashimoto, F. Z. Wu, J. S. Speck, and S. Nakamura, "Growth of bulk GaN
with low dislocation density by the ammonothermal method using polycrystalline
GaN nutrient," Japanese Journal of Applied Physics Part 2-Letters & Express
Letters, 46, L525, (2007).

[28] C. M. Tsai, J. K. Sheu, W. C. Lai, Y. P. Hsu, P. T. Wang, C. T. Kuo, C. W. Kuo,
S. J. Chang, and Y. K. Su, "Enhanced output power in GaN-based LEDs with
naturally textured surface grown by MOCVD," Ieee Electron Device Letters, 26,
464, (2005).

[29] C. M. Tsai, J. K. Sheu, P. T. Wang, W. C. Lai, S. C. Shei, S. J. Chang, C. H. Kuo,
C. W. Kuo, and Y. K. Su, "High efficiency and improved ESD characteristics of
GaN-Based LEDs with naturally textured surface grown by MOCVD," Ieee
Photonics Technology Letters, 18, 1213, (2006).

[30] M.-Y. Ke, C.-Y. Wang, L.-Y. Chen, H.-H. Chen, H.-L. Chiang, Y.-W. Cheng,
M.-Y. Hsieh, C.-P. Chen, and J. Huang, "Application of Nanosphere Lithography
to LED Surface Texturing and to the Fabrication of Nanorod LED Arrays," Ieee
Journal of Selected Topics in Quantum Electronics, 15, 1242, (2009).

[31] R. S. Wagner and W. C. Ellis, "VAPOR-LIQUID-SOLID MECHANISM OF
SINGLE CRYSTAL GROWTH ( NEW METHOD GROWTH CATALYSIS
FROM IMPURITY WHISKER EPITAXIAL + LARGE CRYSTALS SI E ),"
Applied Physics Letters, 4, 89, (1964).

[32] K. A. Dick, S. Kodambaka, M. C. Reuter, K. Deppert, L. Samuelson, W. Seifert,
L. R. Wallenberg, and F. M. Ross, "The morphology of axial and branched
nanowire heterostructures," Nano Letters, 7, 1817, (2007).

[33] P. D. Dapkus, "METALORGANIC CHEMICAL VAPOR-DEPOSITION,"
Annual Review of Materials Science, 12, 243, (1982).
27

[34] H. Ohno, L. A. Nagahara, S. Gwo, W. Mizutani, and H. Tokumoto, "Nanometer-
scale wires of monolayer height alkanethiols on AlGaAs/GaAs heterostructures
by selective chemisorption," Japanese Journal of Applied Physics Part 2-Letters,
35, L512, (1996).

28
Chapter 2 Fabrication and Characterization of GaN
Nanostructures Grown by Selective Area Growth
This chapter demonstrates different GaN nanostructures grown by MOCVD. The
physical morphologies studied by field emission scanning microscopy (FE-SEM) and
crystal structures analyzed by transmission electron microscopy are investigated to
further understand these nanostructures.

2.1 SAG mask preparation
Prior to the nanostructure growth, different patterning techniques will be
discussed for the preparation of SAG growth masks. Each technique has its own
advantages for different purposes of nanostructure growth.

2.1.1 Photolithography
Photolithography is a mature and robust patterning technique, which has been
applied to variety of semiconductor processes in mass production. Especially in the
silicon industry, it has been used in high volume, large wafer scale production to increase
the volume of chips per wafer. By virtue of using photo masks, large area pattern designs
can be easily transferred to photoresists via UV exposure and development. Therefore,
photolithography is a low cost technique with high throughput. However, due to the
29
limitation of light diffraction, the smallest achievable dimension is typically limited to
somewhat less than 1 µm unless expensive deep UV projection lithography tools are used.
Therefore, it is not an appropriate technique to study the influence of patterns in
nanometer scale on the nanostructure growth. Figure 2-1 shows III-V materials grown in
the patterns with micrometer sized openings prepared by photolithography.

Figure 2-1 (a) GaAs tetrahedral shaped structures grown in micron-size openings on a
GaAs (111)B surface. (b) A FE-SEM image of a GaAs tetrahedron recorded at higher
magnification. (c) GaN hexagons grown in hole opening arrays in micrometer scale on a
nitrogen-polar GaN surface. (d) A magnified FE-SEM image of a GaN hexagon with six
vertical {1-100} sidewalls.
30
2.1.2 Electron-beam lithography
Electron-beam lithography (EBL) is a technique in which an electron beam is
used to exposed patterns on electron sensitive resists. Using the tools available at USC
(Raith eLine lithography system), the pattern is prepared in a 1 mm
2
area for most of the
masks used in nanostructure growth and it takes hours to fabricate a few samples.
Therefore, EBL is a time-consuming pattern transferring method to prepare growth masks
for nanostructures. Although it requires long electron beam writing time for mask
preparation, flexibility of mask design makes this technique attractive for research
purposes. Different opening sizes and patterns can be prepared to study the growth
behavior of nanostructures. The procedures to prepare a SAG mask are shown in Figure
2-2.

Figure 2-2 (a) Si-doped (0001) GaN bulk material is utilized to grow GaN nanostructures.
(b) SiN
x
is deposited on top of the GaN surface by plasma enhanced chemical vapor
deposition (PECVD). (c) A dot array is prepared by EBL and transferred into the SiN
x

31
mask by dry etching for nanorod growth. (d) A stripe arrays can be prepared by EBL for
nanosheet or nanostripe growth.
The influence of opening size on GaN nanorod growth will be discussed in
section 2.3. The effect of orientation of stripe openings on GaN nanosheet growth will be
shown in Chapter 4. Besides the SAG mask preparation, metal contacts can also prepared
and designed by EBL. The contact preparation and electrical properties of GaN nanorods
will be discussed in Chapter 5.

2.1.3 Laser interferometric lithography
Laser interferometric lithography utilizes the interference of laser beams to form
periodic patterns on the transferring resist. The coherency of lasers can generate a stable
fringe pattern for this technique. Once a photosensitive resist is exposed to this periodic
pattern, its pattern of interference fringes could be recorded in the resist material. After
photoresist development, this pattern can be transferred to the substrate using dry etch
processes. Figure 2-3 demonstrates SAG growth masks prepared by laser interferometric
lithography. The pattern can be prepared in large scale.
32

Figure 2-3 (a) AFM image of a stripe array pattern transferred into a SiN
x
mask. (b) A
line scan across the line indicated in (a) to estimate the depth of the mask, which is
around 30 nm. (c) FE-SEM image of the stripe pattern. The periodicity is 220 nm and the
opening is 63 nm. (d) FE-SEM image of a dot array pattern generated by exposure the
resist in two perpendicular directions.

33
2.1.4 Nanoimprint lithography
Nanoimprint lithography (NIL) was proposed and demonstrated by Stephen Chou
et al. in 1995.[1] It is an approach to manufacture patterns at the nanometer scale with
low cost, high throughput, and high resolution. NIL utilizes direct contact between the
mold and resist causing the deformation of the resist to transfer the design pattern from
the mold to the resist. Unlike the photolithography or EBL, which create patterns though
the use of photons or electrons to change the chemical and physical properties of the
resist. NIL is based on direct mechanical deformation of the resist using a mold or stamp
containing nanopatterns to achieve resolutions beyond the limitations of light diffraction
or electron beam scattering encountered in photolithography or EBL. There are two
fundamental processes, which are thermal nanoimprint lithography (T-NIL) and UV-
based nanoimprint lithography (UV-NIL). Both them have demonstrated a sub-10nm
resolution capability.
In a T-NIL process, a thin layer of imprint resist, thermoplastic polymer, is coated
on a substrate by spin-coating, followed by the contact of the resist with a mold under
certain pressure. After heating up the resist above its glass transition temperature, the
feature pattern on the mold is pressed and transferred into the melt resist. After the
temperature cools down, the mold is separated from the substrate and the resist with the
pattern is left on the substrate. The pattern can be transferred to the underneath substrate
or mask by a etch process. In UV-NIL process, a UV-curable photopolymer is applied to
the substrate and a mold is brought in contact with the UV resist. After the cavities in the
mold are filled with the resist, the resist is cured by UV light and becomes solid. For this
34
UV curing process, either the mold or substrate must be transparent for UV light to cure
the resist. After the mold is separated from the resist, the pattern on the resist can be
transferred to the substrate underneath or a mask by a dry or wet etch process.[2] Figure
2-4 shows the schematic diagram of two NIL processes.

Figure 2-4 Procedures of two NIL process: (a) Thermal NIL and (b) UV-NIL [3]
Once an optimum pattern is decided for a specific nanostructure, the imprint mold
can be made to prepare high throughput growth mask in large area for the growth of the
nanostructure.

2.1.5 Nanosphere lithography
Nanosphere lithography (NSL) is also a high throughput, low cost technique to
prepare large pattern areas. The hole size and spacing are controlled by the diameter of
the nanospheres. They can be used as etch masks to transfer ordered patterns to the
35
substrates underneath or as evaporation masks to prepare inverse patterns after the gaps
between nanospheres are fill with different materials. The transfer patterns can be
designed by a single layer or double layers of nanospheres.[4] Nanospheres can be
applied to a substrate using different approaches, such as spin-coating, dip-coating or
Langmuir-Blodgett (LB) patterning. The LB technique is a easily integrated method to
deposit well-arranged arrays of nanoparticles or nanowires onto substrates.[5] Nanowires
prepared by either SAG or etch process have been demonstrated using NSL.[6, 7]

2.1.6 Block-copolymer lithography
Block copolymer lithography is another patterning technique, which can
overcome some limitations of lithographic techniques for patterning on the nanometer
scale. For the purpose of this patterning technique, a block copolymer is synthesized,
which self-assembles into an ordered lattice of spheres whose dimensions are determined
by the length of the polymer chain. An example of employing polystyrene(PI)-
polyisoprene(PS) copolymer to form a dense hole array for the GaAs quantum dot growth
is shown in the Figure 2-5.
36

Figure 2-5 Schematic diagram of block copolymer lithography for GaAs quantum dots. (a)
Copolymer microdomain monolayer applied on a SiN
x
mask after annealing in vacuum.
(b) After ozone treatment. (c) During CF
4
plasma etching process. (d) hole array pattern
transferred to GaAs surface. (e) GaAs quantum dots selectively grown in the holes. (f)
GaAs dots after SiN
x
mask removal.[8]
Block copolymer lithography can prepare a large area and close-packed pattern to
grow nanostructures. Dense quantum dots arrays have been demonstrated by Li et al. to
show the feasibility of applying this technique to the optoelectronic devices.[9]
All of these techniques have their own pros and cons. The choice of these
nanopatterning techniques depends on the growth of nanostructures. For example, core-
shell structures need a patterning technique to control the pitch spacing of growth sites to
37
grow heterostructure laterally. In the case of dense arrays, a technique, which can
generate close-packed pattern, is preferable.

2.2 SAG GaN nanostructures
GaN nanostructures have been explored using different kinds approaches. This
section will demonstrate the four different kinds of nanostructures, nanostripe,
nanopyramid, nanorod, and nanosheet, respectively. All of them are grown by selective
area growth.

2.2.1 GaN nanostripes
Due to the lattice mismatch between sapphire and GaN, a high density of
dislocations is generated to release the strain. However, the high dislocation density
affects the device performance, especially laser diodes. Dislocation-free GaN bulk can be
achieved by lateral epitaxial overgrowth (LEO), also called epitaxial lateral overgrowth
(ELOG), to grow high quality materials as growth substrate for lasers.[10, 11] The
dislocation density can be reduced by two to three orders in the mask-covered regions,
where dislocations bend laterally toward to the adjacent growth fronts. However, the
growth mask prepared for LEO is typically on the micrometer scale due to the limitation
of photolithography. The formation of a coalescent layer over the growth masks requires
growth that is several micrometers thick. If the growth mask can be reduced to the
38
nanometer-scale, the growth time and layer thickness to achieve coalescence can be
shortened dramatically. The decrease in the layer thickness not only reduces the cost, but
also has the potential to grow crack-free GaN layer on foreign substrates, such as silicon
substrates. Figure 2-6 shows stripe masks prepared along <1-100> direction with
different spacings, 500, 750, and 1000 nm, respectively. The dark stripes in the figure are
SiN
x
masks and the bright stripes are the exposed GaN surface.

Figure 2-6 Stripe masks prepared with different spacings and different electron beam
dosage by EBL.
Figure 2-7 demonstrates GaN growth on a masked substrate with different
opening spacings. To form a coalescent film and bend the dislocation laterally, two
growth steps were adopted. In the first step, as shown in Figure 2-7 (a), (b), and (c),
triangular stripes are grown for 3 minutes at growth temperature setting around 1030°C,
which is about 120°C lower than the planar bulk growth temperature. The TMG flow rate
39
and ammonia flow rate are 8.85 µmole/min and 107 mmole/min. This step results in
bending of the dislocation laterally towards the adjacent stripes. Figure 2-6 (d), (e), and
(f) shows the planar GaN coalescent layer formed by the nanostripes after the 1030°C
growth for 4 minutes, followed by a high temperature growth, 1150°C, for 20 minutes.
Coalescent GaN layers less than 1 µm are demonstrated by this approach.

Figure 2-7 (a), (b), and (c) shows the triangular stripes grown on different center-to-
center pitches, 500, 750, 1000 nm, respectively, before forming a coalescent GaN layer.
(e), (f), and (g) demonstrates coalescent GaN layer with different center-to-center
spacings.
To estimate the dislocation reduction via this nanostripe technique, a chemical
etch was carried out to count etch pit density (EPD) on the sample surface. The samples
!"#$"%&$'&("#$"%)
*+,(-#./011#2)
!"#$"%&$'&("#$"%)
*+,(-#./301#2)
!"#$"%&$'&("#$"%)
*+,(-#./4111#2)
5,6) 576) 5(6)
!"#$"%&$'&("#$"%)
*+,(-#./011#2)
!"#$"%&$'&("#$"%)
*+,(-#./301#2)
!"#$"%&$'&("#$"%)
*+,(-#./4111#2)
5"6) 586) 5.6)
40
are etched in a mixed acid (H
3
PO
4
: H
2
SO
4
=3:1) at 180°C for 10 minutes. The etch pit
density is estimated from the FE-SEM images in the masked region with 500 nm spacing.
Two GaN samples with different crystal quality evaluated by the full width of half
maximum (FWHM) of GaN (002) rocking curve are prepared for comparison. Besides
FE-SEM images, a more accurate estimation can be carried out using an atomic force
microscopy, which is more sensitive to depth. FE-SEM images are beneficial to
investigate wider arranges of surface area and the relative difference in etch pit density
can still be estimated as well. Table 2-1 shows the results for two different GaN samples.
Before immersing the samples in the hot acid, SiN
x
masks were removed to compare the
EPD between nanostripe regions and planar bulk regions.

Table 2-1 Etch pit density comparison between two different GaN bulk materials.
The number of EPD is reduced in the coalescent regions formed by the
nanostripes for both samples. The dislocation reduction is only about 4 to 5 times, rather
than 2 or 3 orders. A different estimation approach might lead to different results. The
dislocation-free region estimated by typical LEO technique is evaluated at the mask-
covered region where the dislocations bend laterally over the growth masks. The EPD
estimated in the GaN layer formed by nanostripes is an average result covering 8 to 9
nanostripes, including the opening windows of the mask. Due to the short spacing
41
between each nanostripe, the dislocation-free region is barely observed on the surface.
Although nanostripes are not suitable to obtain a large dislocation-free region, dislocation
density can still be reduced to form a layer with thickness about 1 µm via this approach. It
is a potential technique to grow a GaN layer on a silicon substrate with moderate
reduction in dislocation density.

2.2.2 GaN nanopyramids
GaN is typically grown at high V/III ratio due to the poor homogeneous
decomposition reaction (cracking) efficiency of ammonia gas. In the typical growth
ambient, {1-101} surfaces tend to form in non-restricted openings of SAG masks. If
openings are prepared in nanometer-scale and symmetric, GaN can grow from the
openings and form a self-confined structure, called nanopyramid, under a wide range of
growth conditions. Figure 2-8 demonstrates an example of GaN nanopyramids grown by
SAG from openings in nanometer-size. The c-plane is pinched off to form a sharp tip on
the top of each nanopyramid.
42

Figure 2-8 (a) A nanopyramid array is grown by SAG growth. (b) A top view FE-SEM
image of nanopyramids shows nanopyramids are formed by the six inclined planes.
If these inclined planes are unknown, the orientation of the inclined planes can be
estimated from FE-SEM image, as shown Table 2-2.

Table 2-2 Plane estimation from the SEM images recorded at 0° and 45°.
43
(2-1)
Where
The above approach is valid for a 45° SEM stage tilt angle. From the substrate
orientation, these unknown planes belong to {1-10n} group. The n in the {1-10n} plane is
1 from the comparison of y/y’ ratio measured from the FE-SEM images and calculated
from trigonometry. The semipolar {1-101} plane is the slowest growth plane under most
of growth conditions, possibly due to hydrogen passivation of the surface bonds.[12] The
angle between the semipolar plane and c-plane is around 62°. The maximum surface area
increased by the GaN nanopyramids is only about a factor of 2 to reduce the piezoelectric
fields if the InGa/GaN MQWs grow on these semipolar planes. Therefore, this structure
has its limitation to serve as growth templates for the subsequent heterostructures.

2.2.3 GaN nanorods
The idea of growing vertical nanorod arrays is to explore the interesting
physical properties of the nonpolar facets exposed on the GaN nanorods. Also, the
diameter of nanometer scale rods contributes to accommodating more strain caused by
the lattice mismatch from either core-shell or axial heterostructures. A schematic diagram
showing the exposure of nonpolar m-planes on individual rods of a nanorod array is
shown in Figure 2-8 (a). A FE-SEM image of GaN nanorods grown from a hole array
prepared on a SiN
x
mask by MOCVD is demonstrated in Figure 2-8 (b). The facet
44
orientations are deduced from the GaN bulk material and the sapphire substrate. The
vertical sidewalls are nonpolar planes, which can be used as growth templates for
InGaN/GaN MQWs to eliminate the strain-induced piezoelectric field caused by the
lattice mismatch between InGaN and GaN. The advantage of this nanostructure is the
controllability of nonpolar surface area, proportional to the height of the nanorod, and it
contributes to the optimal control the device structure as well. The opening diameter and
pitch can be designed using EBL. Nanorod arrays with different center-to-center spacings
are shown in Figure 2-8 (c)-(f).
45

Figure 2-9 (a) Schematic diagram of a GaN nanorod array grown vertically from its
substrate. The enlarged diagram shows two dominant facets formed by GaN nanorods. (b)
Uniform GaN nanorod array is grown by SAG by MOCVD. The arrows indicate the
polar and nonpolar facets of the nanorods. In the inset, a top view of hexagonal GaN
nanorods shows the six vertical sidewalls. Ordered GaN nanorod arrays grown on 250 nm,
500 nm, 750 nm, and 1 µm center-to-center spacings are demonstrated in (c−f),
46
respectively. The scale bars are 500 nm in all figures. All the FE-SEM images were
recorded at a 45° angle.[13]
The exposed nonpolar surface area can be controlled by the diameter and the
height of GaN nanorods. The growth of core-shell InGaN/GaN MQW will be discussed
in Chapter 3.

2.2.4 GaN nanosheets
GaN nanosheets are unique structures with long connected nonpolar sidewalls that
can be considered as a series of linked nanorods. To form this nanosheet structure, a SAG
mask with stripe patterns aligned along <11-20> directions is prepared to form {1-100}
nonpolar planes. Figure 2-10 (a) shows the GaN growth along <11-20> with limited
growth along the <1-100> direction due to the growth behavior and mask design. The
facets exposed in the nanosheets are indicated in Figure 2-10 (b). The comparison of
height with different growth cycles are shown in Figure 2-10 (c) and (d). The nonpolar
surface exposed from GaN nanosheets can also serve as growth template for InGaN/GaN
MQWs to eliminate the strain-induced piezoelectric fields. The detailed structure and
characteristics of the heterostructures will be discussed in Chapter 4.

47

Figure 2-10 (a) A uniform nanosheet array with 500 nm center-to-center spacing grown
on a GaN/Al
2
O
3
buffer layer patterned with a stripe patterns. (b) Polar, semipolar, and
nonpolar planes are indicated in the schematic diagram of GaN nanosheets. (c) FESEM
image taken with the sample rotated 90° with respect to (a). The nanosheets show vertical
and parallel sidewalls, which are {1-100} planes. (d) A nanosheet array grown for double
the growth cycles of (c). Each of the FESEM images was recorded at a 60° angle. The
scale bar is 500 nm in all figures.[14]

48
2.3 GaN nanorod selective area growth mechanism
GaN nanowires/nanorods have been grown by various growth approaches,
including VLS and SAG by MOCVD.[15, 16] Different growth techniques have also
demonstrated GaN nanowire growth, such as laser-assisted catalytic growth and
molecular beam epitaxy (MBE).[17, 18] However, general growth conditions under high
V/III ratios in MOCVD are not suitable to grow vertical hexagonal wires along <0001>
direction with six {1-100} sidewalls. The most stable {1-101} planes, which have the
slowest growth plane under most of the growth ambient, prevent the definition during
growth of {1-100} planes. Therefore, GaN pyramids with six surrounding {1-101}
semipolar planes are the most common feature observed under high V/III ratio conditions.
Therefore, either different growth approaches or a different growth surface must be
adopted to grow vertical GaN nanorods.

2.3.1 Pulsed growth on (0001) Ga-polar plane
Hersee et al. first demonstrated a uniform GaN nanorod array on a
GaN/sapphire substrate by introducing flow rate modulation growth method, or called
pulsed growth in 2006.[19] The hole array mask is prepared by laser interferometic
lithography on a masked GaN bulk substrate. TMG and ammonia are introduced into
growth chamber alternatively. Vertical GaN nanorods can be grown on the GaN template
under these conditions. There are some special features of vertical GaN nanorods grown
along c-axis. For example, the small bases of GaN nanorods can be utilized to reduce
49
dislocations originated from the GaN bulk material. The defining sidewalls of the vertical
nanorods are nonpolar m-planes, on which the piezoelectric fields can be eliminated
when the InGaN/GaN MQW structures are grown.. All the MOCVD-grown vertical
nanorod arrays reported earlier were grown in high rotation speed, disk reactors and the
nanowire diameters were over 200 nm.[19-21] GaN nanorod arrays grown in a close-
coupled showerhead reactor was first demonstrated in this work.[13] Also, the diameter
of the GaN, smaller than 150 nm, is demonstrated by using masks prepared by EBL. In
our research, there are several parameters affecting the growth of GaN nanowires, such as
temperature, gas ambient, pressure, exposure durations of reactants during pulsed growth,
gas flow rate, and the design of growth masks. Uniform GaN nanorod arrays have been
achieved by optimizing these growth parameters.
Under conventional MOCVD growth conditions, precursor compounds are
simultaneously injected into the growth chamber and GaN pyramids are formed as a
result of high V/III ratios. A comparison of GaN grown on dot array patterns using pulsed
and continuous injection of TMG and ammonia are shown in Figure 2-11 (a) and (b).
Vertical GaN nanorods grown by the pulsed mode are defined by six nonpolar {1-100}
facets; however, the slow growth rate on semipolar {1-101} planes forms hexagonal
pyramids under the continuous growth mode. Although the nonpolar {1-100} planes
disappeared under the continuous gas flux injection due to faster growth rate on those
planes than on the {1-101} planes, core-shell InGaN/GaN MQW structures can be
achieved by utilizing this feature.
50

Figure 2-11 Growth comparison between continuous and pulsed injection of gas flux. (a)
Vertical GaN nanorods structure can be formed by pulsed-mode GaN growth. (b) During
continuous growth, the slow growth rate of the {1-101} semipolar plane leads to form
GaN nanopyramids. (c) The growth scheme of the pulsed growth mode. T
1
and T
3
are the
injection time of TMG and ammonia, respectively. T
2
and T
4
are the growth interruption
between switching TMG and ammonia.
Due to the different facets exposed in these two structures, the growth
mechanism of the GaN vertical nanorods must be different from the typical planar growth,
which is typically performed under continuous high V/III ratio conditions. To grow
vertical nanorods along <0001>-direction, the growth rate on the tip must be faster than
those of sidewalls, especially the inclined semipolar planes. The growth scheme of the
51
alternative injection of precursors into the reactor chamber and the growth interruptions
are shown in Figure 2-11 (c). The period covers from T
1
to T
4
is called a growth cycle.
The pulsed growth mode contributes to the enhancement of growth rate on the tips of
nanorods, resulting in the exposure of {1-100} planes. A schematic growth mechanism is
demonstrated in Figure 2-12.

Figure 2-12 Schematic diagram of the pulsed growth mechanism for SAG GaN nanorod.
The yellow balls are the Ga adatoms.
The group V element is typically much more volatile than the group III element.
Thus, the growth mechanism mentioned here is based on the group III element, Ga.
Vertical nanorod growth can be briefly described in the following steps.
a. Ga directly impinges on the top of the nanorod.
b. Ga diffuses from the gas phase to the mask surface.
c. Ga migrates to the nanorod growth site.
!"
#"
$"
%"
&'(
)
"*!+,"
-!("#./,"
52
d. Ga migrates from the bottom of the nanorod to the top of the rod.
The surface migration is affected by several growth parameters, such as
temperature, gas ambient, pressure, durations of pulsed growth, gas flow rate, and mask
design. The influence of these parameters will be discussed in the following sequence.

• Temperature
Temperature controls the kinetic energy of the gas phase diffusion and surface
migration of adatoms. Also, it is involved in controlling the thermal etching in
hydrogen ambient. High growth temperature results in a decrease in the growth rate in
the <0001> direction. Low growth temperature reduces the energy of adatoms
migrating to the top of the GaN nanords and decreases the desorption rate of Ga
adatoms especially on (0001) and {1-100} planes. The nanorods grown at low
temperature are usually wider with sharper tips.

• Gas ambient
In a hydrogen growth ambient, the desorption rate of GaN is fast at high
temperature. If the pulsed growth mode is applied to the GaN nanorod growth, the
strong desorption rate results in not only no crystal growth but also strong dissociation
of GaN under the parts of the cycle during which there is no overpressure of the group
V source. Thus, a nitrogen-rich ambient is used during the pulsed growth and the
desorption rate reduces dramatically. However, the low desorption rate of Ga under
53
these conditions leads to the deposition of GaN on the mask surface and non-uniform
growth of GaN nanorods. Therefore, the TMG flow rate and the design of the mask
pattern have to be taken into account to grow GaN SAG nanorods in the nitrogen-rich
ambient.

• Pressure
Pressure affects the diffusion length of metalorganic sources as well as the
desorption rate. It is found the growth pressure around 200~250 Torr is an appropriate
growth condition for the current nanorod growth. Also, low growth pressure is
beneficial for the maintenance of the reactor because it results in less deposition on the
growth chamber, as compared to atmospheric growth condition.

• Durations of pulsed growth
Flow–rate modulation (FM) growth, or pulsed growth, was used to control the
layer growth precisely in the atomic layer range.[22] In our study, it was found the
duration of gas flow and the timing of switching between TMG and ammonia
influenced the nanorod growth rate and morphology. In our early work on nanorod
growth, the typical duration of TMG flow was 20 sec and that of ammonia was 30 sec.
A 15-second growth interruption was introduced between switching from TMG to
ammonia injection. The growth interruption duration time affects the depletion of
either one of the sources, Ga or N, during the pulsed growth to avoid the overlap of
54
both sources. If the high concentrations of both elements coexist, the nanorods tend to
grow laterally and form pyramidal shapes. The goal of pulsed growth is somewhat
different from the original purpose of FM growth because the duration of each source
is longer than that of traditional FM growth and the layer thickness grown for each
growth cycle is more than 10 monolayers, rather than single monolayer. Long duration
of source interruption results in the long surface migration of Ga to the tips of nanorods.
If the migration length is shortened by the existence of ammonia, Ga may not be able
to migrate to the top of nanorods and will deposit on the sidewalls, resulting in the
lateral growth of nanorods and low growth rate in c-direction. It is found that the
growth interruption is crucial to the reduction of hydrogen passivation on the {1-101}
planes to enhance the growth rate on the tips of GaN nanorods. The detailed growth
mechanisms are reported by Lin et al.[23]

• Gas flow rate
Both TMG and ammonia flow rates affect the nanorod growth. As compared to
ammonia, the flow rate of TMG shows more obvious effects on the nanorod growth,
probably due to the excessive supply of ammonia gas under most of the growth
conditions. At low TMG flow rate, the growth rate of GaN nanorods is typically very
low because the Ga source dominates the growth. At high TMG flow rate, excessive
Ga will grow non-uniformly on the nanorods and forms unwanted surface deposits on
the mask. Those deposits work as sinks for source materials, consuming metalorganic
sources, and affects growth uniformity. Therefore, gas flow rates of both TMG and
55
ammonia need to be adjusted to grow uniform nanorod arrays with acceptable growth
rates.

• Mask design
The patterns on the growth mask impact the nanostructure growth, including the
opening sizes, pitch spacings, and orientations. Detailed demonstration of the influence
from the growth masks for both GaN nanorods and nanosheets will be shown in
Chapter 3 and 4, respectively.

2.3.2 Continuous growth on (0001) N-polar plane
The growth of vertical GaN nanorods under continuous growth mode on nitrogen-
polar GaN surface by MOCVD has been reported.[24, 25] A nitridation step is introduced
to prepare a nitrogen-polar surface on sapphire substrates to grow GaN nanorod arrays.
Nanorod growth on this surface results in a nitrogen-polar GaN surface grown on a
sapphire substrate is shown in Figure 2-13. Although the growth can be achieved without
using the sophisticated pulsed growth condition, the uniformity of the nanorods are
typically not good enough to serve as growth templates for MQWs. Also, it is reported
the MQWs tend to grow on the inclined semipolar planes around the tips of nanorods
grown on nitrogen-polar surface.[25, 26] Therefore, the large nonpolar surface area
exposed from the nanorods is not utilized well as the growth templates for InGaN/GaN
56
MQWs. Therefore, GaN nanostructures grown on Ga-polar GaN bulk material are mainly
investigated in this research.

Figure 2-13 GaN nanorods are grown on nitrogen-polar GaN surface.

57
2.4 Crystal structure analysis and sample preparation
Different GaN nanostructures, including nanostripe, nanopyramid, nanorod, and
nanosheet, are successfully grown by SAG in our laboratory. To understand the crystal
structure of these GaN nanostructures, using general analysis tools for planar structures,
such as X-ray diffraction, are not appropriate methods because of the small and isolated
features of these nanostructures. Therefore, crystal analysis tools are limited for the
nanostructures. The following sections will discuss the crystal analysis by transmission
electron microscope (TEM) and the sample preparation by focused ion beam (FIB).

2.4.1 TEM sample preparation by focused ion beam
Transmission electron microscopy requires thin sample for effective use of the
tool. When the thickness of a sample is, in general, thicker than 100 nm, the sample must
be thinned by either mechanical lapping and polishing or ion beam milling to let electrons
penetrate the sample to form images on a projection screen in a TEM system. Focused
ion beam (FIB) is a tool to section a sample of interest and transfer it to a TEM grid for
structure analysis. The FIB milling system used in this study is made by JEOL
(MultiBeam JIB-4500) and the resulting slices are transferred to a TEM grid by a lift-out
process using a micromanipulator (Omniprobe Autoprobe 200). The TEM sample
preparation procedures by FIB are shown in Figure 2-14. A carbon sacrificial layer is
deposited on the sectioned region to protect the sample from the damage caused by the
ion beams, followed by trench milling to create the thin web of material to be examined
58
by TEM and to create enough space for the lift-out process, as shown in Figure 2-14 (a).
In Figure 2-14 (b), a sample is separated from its substrate and transferred to a TEM grid.
The positioning of the mounted sample on a TEM grid depends on the orientations in
which the cross-sectional nanostructures are investigated. Examples are shown in the
Figure 2-14 (c) and (d). The insets in both (c) and (d) are the schematic diagrams of the
cross-sectional structures of nanorods observed from <11-20> and <0001> directions,
respectively.

Figure 2-14 (a) Trenches prepared by FIB for lift-out process. (b) A sample is separated
from its substrate and mounted on a micro-manipulator to transfer to a TEM grid. (c) A
nanorod sample mounted on a TEM grid is sectioned along <11-20>-direction to study
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(a) (b)
(c) (d)
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59
the MQWs grown on the nonpolar planes. (d) A TEM sample is sectioned along <0001>-
direction to investigate the core-shell MQW structure.

2.4.2 Crystal structure of GaN nanorods analyzed by transmission electron
microscope
Transmission electron microscope (TEM) is a useful tool to understand the crystal
structure and orientation of nanostructures. Due to the small dimension of these structures,
they can be removed from their substrates and dispersed on TEM grids for direct TEM
observation. High density of stacking faults (SFs) has been observed in GaAs, InAs, or
InP nanowires grown by SAG.[27-29] They tend to form dark band-like contrasts along
the <111> axial direction of nanowires. Figure 2-15 demonstrates an example of dense
SFs observed in a GaAs nanowire grown directly on a Si (111) substrate without any
growth mask.

Figure 2-15 (a) A FESEM image of GaAs nanowires grown from low temperature GaAs
nuclei on a Si (111) substrate. (b) High density of SFs observed in the GaAs nanowire.
60
For the crystal investigation of GaN nanorods, individual nanorods were
immersed in isopropanol in a small container and then removed from the GaN/sapphire
substrate in the solution by scratching or sonication. GaN nanorods suspended in
isopropanol were dispersed on a TEM grid using a pipette. The crystal structure of the
GaN nanorods was investigated by a TEM (JEOL 2100LB). Figure 2-16 shows a wurtzite
crystal structure is observed from the TEM images of the GaN nanorod taken at <11-20>
zone axis.

Figure 2-16 TEM images taken at <11-20> zone axis. No SF is observed in both TEM
images recorded at (a) high magnification and (b) low magnification. (c) Selective area
diffraction pattern confirms the positioning of the nanorod as well as the single wurtzite
61
crystal structure of the nanorod. (d) Wurtzite lattice fringes recoded by the high
resolution TEM.
To confirm the crystal structure again, TEM images of another GaN nanorod
taken along a different zone axis was investigated. Figure 2-17 demonstrates the TEM
images of wurtzite crystal structure recorded along the <1-100> zone axis.

Figure 2-17 TEM images taken at <1-100> zone axis. No SF is observed in the TEM
images taken at (a) high magnification and (b) low magnification. Wurtzite crystal
62
structure is confirmed by both (c) diffraction pattern, and (d) lattice fringes taken by the
high resolution TEM.
No SFs are observed in the TEM images taken along both zone axes. This result
suggests that the GaN nanorods are potential candidates for device fabrication without the
any influences on their electrical property caused by the crystal discontinuity defect
structure.

2.4.3 Heterostructures investigated by scanning transmission electron microscope
Scanning transmission electron microscopy (STEM) is similar to SEM in that its
images are formed by scanning a fine and highly focused beam of electrons over a thin
specimen. Usually, STEM is performed in a conventional TEM equipped with additional
scanning coils and the needed circuitry to drive the focused electron beam and with
detectors for collecting the different kinds signals, such as energy dispersive X-ray
spectroscopy (EDS), electron energy loss spectroscopy (EELS), and annular dark-field
imaging (ADF). The rastering of the beam across the sample enables the microscope to
perform different analysis techniques simultaneously. The image contrast, z-contrast
image, recorded by a high angle annular dark field (HAADF) detector is directly related
to the atomic number of the materials. The z-contrast images can facilitate the structure
analysis of layers with different compositions.[30-32] Figure 2-18 demonstrates clear
interfaces between GaN barriers and InGaN wells in the MQWs in the STEM dark field
image. This MQW structure grown on the GaN nanorods is investigated by a TEM
63
equipped with STEM (JEOL-JEM- 2100F). The result shows that MQWs are grown
predominantly on the sidewalls. The c-plane surface is pinched-off after a continuous
GaN growth for 10 minutes prior to the growth of MQWs. Detailed structure of the core-
shell MQWs will be discussed in Chapter 3.

Figure 2-18 Comparison of STEM image between (a) bright field and (b) dark field.

2.5 Summary
Different mask preparation techniques for nanostructures have been discussed
in this chapter. The choice of these nanopatterning techniques depends on the purpose of
nanostructure growth. EBL is the main approach used in this study to have the flexibility
of controlling mask design for different kinds of nanostructures. Four different kinds of
GaN nanostructure have been demonstrated using the SAG technique. Single wurtzite
64
crystal structure of GaN nanorods is confirmed by TEM without the observation of SFs
or dislocations. The result suggests the GaN nanostructures can be used as pristine
growth templates for the subsequent heterostructure growth.

2.6 Chapter references
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[4] C. L. Haynes and R. P. Van Duyne, "Nanosphere lithography: A versatile
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[13] T.-W. Yeh, Y.-T. Lin, L. S. Stewart, P. D. Dapkus, R. Sarkissian, J. D. O'Brien, B.
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[15] G. T. Wang, A. A. Talin, D. J. Werder, J. R. Creighton, E. Lai, R. J. Anderson,
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[17] X. F. Duan and C. M. Lieber, "Laser-assisted catalytic growth of single crystal
GaN nanowires," Journal of the American Chemical Society, 122, 188, (2000).
[18] M. Yoshizawa, A. Kikuchi, M. Mori, N. Fujita, and K. Kishino, "Growth of self-
organized GaN nanostructures on Al2O3(0001) by RF-radical source molecular
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66
[19] S. D. Hersee, X. Sun, and X. Wang, "The controlled growth of GaN nanowires,"
Nano Letters, 6, 1808, (2006).
[20] Y.-S. Chen, W.-Y. Shiao, T.-Y. Tang, W.-M. Chang, C.-H. Liao, C.-H. Lin, K.-C.
Shen, C. C. Yang, M.-C. Hsu, J.-H. Yeh, and T.-C. Hsu, "Threading dislocation
evolution in patterned GaN nanocolumn growth and coalescence overgrowth,"
Journal of Applied Physics, 106, 023521, (2009).
[21] T.-Y. Tang, C.-H. Lin, Y.-S. Chen, W.-Y. Shiao, W.-M. Chang, C.-H. Liao, K.-C.
Shen, C.-C. Yang, M.-C. Hsu, J.-H. Yeh, and T.-C. Hsu, "Nitride Nanocolumns
for the Development of Light-Emitting Diode," Ieee Transactions on Electron
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[22] N. Kobayashi, T. Makimoto, and Y. Horikoshi, "FLOW-RATE MODULATION
EPITAXY OF GAAS," Japanese Journal of Applied Physics Part 2-Letters, 24,
L962, (1985).
[23] Y.-T. Lin, T.-W. Yeh, and P. D. Dapkus, "Mechanism of selective area growth of
GaN nanorods by pulsed mode metalorganic chemical vapor deposition,"
Nanotechnology, accepted, (2012).
[24] W. Bergbauer, M. Strassburg, C. Koelper, N. Linder, C. Roder, J. Laehnemann, A.
Trampert, S. Fuendling, S. F. Li, H. H. Wehmann, and A. Waag, "N-face GaN
nanorods: Continuous-flux MOVPE growth and morphological properties,"
Journal of Crystal Growth, 315, 164, (2011).
[25] W. Bergbauer, M. Strassburg, C. Koelper, N. Linder, C. Roder, J. Laehnemann, A.
Trampert, S. Fuendling, S. F. Li, H. H. Wehmann, and A. Waag, "Continuous-
flux MOVPE growth of position-controlled N-face GaN nanorods and embedded
InGaN quantum wells," Nanotechnology, 21, (2010).
[26] A. Waag, X. Wang, S. Fuendling, J. Ledig, M. Erenburg, R. Neumann, M. Al
Suleiman, S. Merzsch, J. Wei, S. Li, H. H. Wehmann, W. Bergbauer, M.
Strassburg, A. Trampert, U. Jahn, and H. Riechert, "The nanorod approach: GaN
NanoLEDs for solid state lighting," in Physica Status Solidi C: Current Topics in
Solid State Physics, Vol 8, No 7-8. vol. 8, 2011.
[27] K. Tomioka, Y. Kobayashi, J. Motohisa, S. Hara, and T. Fukui, "Selective-area
growth of vertically aligned GaAs and GaAs/AlGaAs core-shell nanowires on
Si(111) substrate," Nanotechnology, 20, (2009).
[28] K. Tomioka, J. Motohisa, S. Hara, and T. Fukui, "Crystallographic structure of
InAs nanowires studied by transmission electron microscopy," Japanese Journal
of Applied Physics Part 2-Letters & Express Letters, 46, L1102, (2007).
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[29] H. J. Chu, T. W. Yeh, L. Stewart, and P. D. Dapkus, "Wurtzite InP nanowire
arrays grown by selective area MOCVD," in Physica Status Solidi C: Current
Topics in Solid State Physics, Vol 7, No 10. vol. 7, P. Bhattacharya, et al., Eds.,
2010.
[30] Y. Li, J. Xiang, F. Qian, S. Gradecak, Y. Wu, H. Yan, D. A. Blom, and C. M.
Lieber, "Dopant-free GaN/AlN/AlGaN radial nanowire heterostructures as high
electron mobility transistors," Nano Letters, 6, 1468, (2006).
[31] F. Qian, S. Gradecak, Y. Li, C. Y. Wen, and C. M. Lieber, "Core/multishell
nanowire heterostructures as multicolor, high-efficiency light-emitting diodes,"
Nano Letters, 5, 2287, (2005).
[32] F. Qian, Y. Li, S. Gradecak, H. G. Park, Y. J. Dong, Y. Ding, Z. L. Wang, and C.
M. Lieber, "Multi-quantum-well nanowire heterostructures for wavelength-
controlled lasers," Nature Materials, 7, 701, (2008).

68
Chapter 3 InGaN/GaN Multiple Quantum Wells Growth
on GaN Nanorods
Core-shell InGaN/GaN MQWs have been demonstrated on GaN
nanowires/nanorods for application to LEDs and lasers.[1-3] Electroluminescence has
been observed in single nanowire LEDs with the core-shell structure.[4, 5] Also,
InGaN/GaN MQWs embedded in nanowires/nanorods along or around axial directions
have been applied to LEDs.[6-9] However, less emphasis has been placed on the
elimination of piezoelectric fields in the MQWs resulting from the growth orientations or
different epitaxial growth techniques. In this study, semipolar and nonpolar facets are
exposed on GaN nanorods grown vertically from the basal plane of the polar substrates.
These exposed semipolar/nonpolar facets can be utilized as growth templates for InGaN-
based LED active regions to reduce or eliminate the piezoelectric fields in the active
regions. The chapter starts from the impact of dot array patterns to the nanorod growth,
followed by InGaN/GaN MQWs growth on these nanorods. Light emission will be
characterized by photoluminescence (PL) and cathodoluminescence (CL). The growth of
MQW predominantly on the nonpolar planes will be verified by CL mappings and cross-
sectional TEM images.

69
3.1 The dependence of mask design on GaN nanorod growth
To grow MQWs on the nonpolar planes created by GaN nanorods, the structure
initiated from the growth of GaN nanorod arrays. It is found that the growth masks play
an important role in the nanorod growth. In the case of dot arrays for nanorod growth, the
hole dimension can affect the both the diameter and height of GaN nanorod. The
influence of the hole size of the growth mask on the nanorod dimension is shown in
Figure 3-1. The number of growth cycles is 120 and the pitch spacing is 500 nm. All the
data points are measured from FE-SEM images. The diameter of the nanorod is
proportional to the hole diameter as shown in Figure 3-1 (a). The relation between the
height and diameter of the nanorods is demonstrated in Figure 3-1 (b). From the second
order polynomial fitting curve, the maximum height, 447 nm, occurs at the nanorod
diameter of 187 nm. Assume the nanorod diameter is linearly proportional to the hole
diameter in Figure 3-1 (a), the maximum height of the nanorod can be achieved by
preparing the corresponding hole diameter, 143 nm, on the mask. Surface area of
nonpolar plane, m-plane, is closely related to the height and diameter of the nanorods. If
the total m-plane surface area can be expressed as 2πhd, where h and d are the height and
diameter of GaN nanoord, the maximum surface area can be derived from the relation of
height and diameter in Figure 3-1 (b). The maximum m-plane surface area will occur at
the nanorods with the diameter of 206 nm and the corresponding diameter of the hole
opening on the mask is 165 nm. Also, the m-plane surface area increased from the
nanorod structure grown for 120 cycles will be 2.26 times larger than the planar surface
70
area of c-plane. Therefore, the growth mask can be designed to meet the experimental
purpose based on these relations.

Figure 3-1 (a) Nanorod diameter is proportional to hole diameter. (b) Nanorod height
changes with respect to its diameter. (c) m-plane surface area estimated from the product
of the height and diameter of GaN nanorods. (d) Nanorod height varies with the fill factor
of growth mask. The error bars are the standard deviations of sampling over 20 nanorods
for each point.
The relation between nanorod height and fill factor is plotted in Figure 3-1 (d). In
general, a negative linear slope is expected if the height of nanorod is inverse to the hole
!"#$ !%#$
!&#$ !'#$
71
diameter square due to the concentration gradient in the gas phase. However, the height
of nanorod is the shortest in the regions with the smallest fill factor and increases
dramatically up to a fill factor of 3%. This result suggests that the reaction surface area is
too small to grow nanorods and might indicate that the formation of GaN is dependent on
a heterogeneous decomposition of NH
3
during the NH
3
exposure in the pulsed growth
mode. This phenomenon possibly occurs on the sidewalls of the nanorods. The results
shown in Figure 3-1 apply only to the specific growth conditions. These curves might
vary when the growth parameters, not limited to those mentioned above, are changed. It
is also found that the growth parameters are correlated. Changing one parameter
sometimes requires an adjustment of another one to grow uniform nanorod arrays.

3.2 Multiple quantum well growth on GaN nanorods
GaN substrates consisting of an epitaxial layer of c-plane GaN grown on sapphire
were prepared for nanorod growth by patterning the substrate with a dielectric film
containing a dot array of nanoscale openings prepared by electron beam lithography and
reactive ion etching. A 20 nm-thick SiN
x
layer was deposited onto Si-doped GaN layers
by plasma enhanced chemical vapor deposition (PECVD). PMMA was used as an
electron beam resist on the dielectric mask prior to electron beam writing. The hole
diameter was controlled between 70 nm and 250 nm by varying the writing parameters,
and dot arrays with 250 nm ~ 1 µm center-to-center spacing were prepared for GaN
nanorod growth. The pattern was transferred from the PMMA resist into the SiN
x

72
dielectric mask using a CF
4
based reactive ion etching process, and the surface was then
cleaned with solvents and further treated with oxygen plasma to remove any resist
residue. The samples were then loaded into a close-coupled showerhead MOCVD system
for GaN nanorod growth.
Semipolar or nonpolar planes have also been demonstrated in micrometer-scale
stripe structures grown by modulating NH
3
injection to enhance lateral growth.[10, 11]
To achieve isolated nanorod structure with nonpolar surfaces, a pulsed growth mode was
applied to enhance the vertical growth along <0001> directions.[12] Trimethylgallium
(TMG) and ammonia (NH
3
) were used as the precursors for the nanorod growth. The
growth pressure was 200 Torr for the entire pulsed growth mode process. Various growth
parameters, including growth temperature, gas flow rate, and precursor flow period, were
adjusted to achieve GaN nanorods with high aspect ratios. The flow rates of TMG and
NH
3
were 17.7 µmole/min and 67 mmole/min, respectively. InGaN/GaN multiple
quantum wells were grown subsequently on the GaN nanorods using trimethylindium
(TMI) and triethylgallium (TEG) as the indium and gallium sources, respectively. The
flow rate of TEG was 20 µmole/min and that of TMI ranged between 3.4 and 17.0
µmole/min for different MQW growth tests. The growth pressure was increased to 300
Torr in nitrogen ambient to grow InGaN quantum wells and GaN barriers under
continuous growth conditions. The growth temperature varied between 730°C and 815°C
for different experimental purposes. NH
3
was injected into the chamber until the
temperature decreased below 400°C. After nanorod growth, the surface morphology was
evaluated by scanning electron microscopy (Hitachi S-4800 field emission FE-SEM) and
73
the facet orientations of the GaN nanorod were deduced from the orientation of the
underlying GaN material.

3.3 Light emission from multiple quantum wells grown on GaN
nanorods studied by photoluminescence
Photoluminescence measurements were performed by illuminating the samples
with a focused 325 nm HeCd laser source to measure the light emission from the MQWs.
Typically, 500 nm height GaN nanorods are grown over a period of 120 growth cycles.
The nonpolar surface formed by their vertical sidewalls is exposed for subsequent
InGaN/GaN MQW growth. Three pairs of MQWs were grown on a GaN nanorod array
to study the light emission from the nanostructure. PL emission from the MQWs grown
on the nonpolar planes has been verified by the removeal of c-plane and semipolar planes
by a dry etching process. This selective etching involved in filling the gaps with
hydrogen silsesquioxane (HSQ) to protect the MQWs grown on nonpolar planes,
followed by a selective etching using inductively coupled plasma reactive ion etching. PL
spectra before and after the removal of the c-plane are shown in Figure 3-2. The hollow
nanorods after the dry ethcing are demonstrated in FE-SEM image in Figure 3-2. The
growth of MQWs on the nonpolar sidewalls of nanorods is verified by the PL emission
peaks located at the same wavelength with/withot the presence of MQWs grown on c-
74
plane. The increase in PL intensity might result from the removal of light aborbing
MQWs grown on the c-plane.

Figure 3-2 Comparison of the PL spectra before and after ICP etching. The inset shows
the c-plane and semipolar planes are removed on the tips of GaN nanorods.
Changes in emission wavelengths with respect to opening diameters in the mask
are observed in the MQWs grown on GaN nanorods. The growth temperate was
performed at 770°C with 17.0 µmole/min TMI injection during MQW growth. By
varying the opening size in the patterned dielectric growth mask, emission wavelengths
from 432 nm to 498 nm are achieved, as shown in Figure 3-3. The variation of the peak
wavelength with opening size is shown in the inset. The emission wavelengths are closely
related to the indium compositon and the quantum well width. Varying the opening size
affects the enhancement of the growth rate that occurs in selective area growth which
75
may affect both the indium composition and the QW thickness. However, the MQWs
thickness cannot be determined directly from the changes of the nanorod diameters, and
extensive TEM studies of sectioned samples are required to determine the cause of the
wavelength change. The portions of the sample from which these spectra are taken are
grown simultaneously on different regions of the same wafer. The apparent reduction of
the efficiency with increasing wavelength may be a consequence of unoptimized growth
conditions for the longer wavelength portions of the sample rather than a fundamental
effect. In spite of the need for further studies, these results suggest that the emission
wavelength may be controllable by careful mask design and control of the nonpolar
plane surface area. They also suggest that precise design will be necessary to control the
wavelength of emission in LEDs.
76

Figure 3-3 The emission wavelengths measured from nanorod samples grown with
different opening sizes are shown in the spectra. The inset shows the emission
wavelength decreases as the opening size on the dielectric mask increases.
Since the MQWs are predominantly grown on nonpolar planes from the results
mentioned above, the radiative recombination efficiency should be increased due to the
absence of the piezoelectric fields inside MQWs. Time-resolve PL was carried out to
study the carrier lifetime in MQWs grown on nonpolar planes formed by GaN nanorods.
A nanorod array with MQWs emitting at 450 nm was prepared for the measurement. The
emission peak was confirmed by the steady state PL measurement. A pulsed Ti:Sapphire
laser with pulse width of 100 fs and repetition rate of 250 kHz was used as the excitation
77
source and the beam spot was focused down to 300 um in diameter on the sample surface.
The excitation wavelength was set to 400 nm to excite the carriers inside the quantum
wells only and the excitation power, 21 mW, was measured before the experiment. A
nanorod sample without any quantum wells grown on it was also prepared for
comparison. The results are shown in Figure 3-4 (a). Neutral density filters were
introduced before the detector for the measurement to reduce the light intensity from the
MQWs grown on GaN nanorods to ensure the measurement was single photon counting.
Therefore, the result of the MQWs grown on nanorods in Figure 3-4 (a) shows poor
signal-to-noise ratio at low intensity. It is obvious to see that the MQWs grown on GaN
nanorods exhibit faster decay than the sample without MQWs. Two decay time constants,
205 ps and 1.024 ns, are derived from fitting the curve of GaN nanorods without MQWs.
These decays might relate to weak and broad light emission from the defect-related
yellow band originating from the GaN material. After deconvoluting of the instrument
response and applying the decay constants obtained from the GaN nanorods to the MQW,
the estimated carrier lifetime in the quantum well is 96 ps. The fitted curve and the
instrument response are black and red curves, respectively, as shown in Figure 3-4 (b).
This result might suggest the MQWs grown on the nonpolar planes formed by the
nanorods have the potential to improve the radiative recombination efficiency for InGaN-
based LEDs.

78

Figure 3-4 (a) The decay curves comparison between GaN nanorods with and without
MQWs. (b) The fitted curve (black) includes the contribution from the instrument
response (red). Blue curve is the measured result and the other blue curve at the bottom is
the residuals of the fitting.

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79
3.4 Multiple quantum wells grown on GaN nanorods studied by
transmission electron microscopy
GaN nanorods with InGaN/GaN multiple quantum wells were also investigated
by transmission electron microscope (TEM, JEOL 2100LB). Due to the thickness of the
nanorods, samples were sectioned by a focused ion beam milling system (JEOL
MultiBeam JIB-4500), and the resulting slices were transferred to a TEM grid by a lift-
out process (Omniprobe Autoprobe 200). GaN nanorods with three pairs of MQWs were
sectioned by focused ion beam and transferred onto a TEM grid for structure analysis.
GaN nanorods are sliced in two different orientations to demonstrate the core-shell
structure of InGaN/GaN MQWs. Bright field TEM images of sectioned GaN nanorods
with MQWs are shown in Figure 3-5. The surfaces of the nanorods were covered with a
carbon layer to protect them from damage from the focused ion beam. The GaN nanorods
sectioned along <11-20> are shown in Figure 3-5 (a), and three boxed regions indicate
the locations where magnified images were recorded and shown in Figures 3-5 (c), (d),
and (e). Three pairs of MQWs are clearly distinguishable on nonpolar planes. The
exposure of three main facets in the nanorod resulted in different growth rates on the
various planes.
Figure 3-5 (b) designates the orientations of MQWs grown on polar, semipolar,
and nonpolar planes. The MQWs grown at the edge between semipolar and nonpolar are
slightly thicker possibly due to the migration of growth species from the semipolar plane
to the nonpolar plane. Because of the hexagonal pyramid shapes of the GaN nanorod tips
80
after MQW growth and the nanometer-scale dimension of the tip, interfaces between
barriers and wells are difficult to distinguish on this plane. In addition, the interface
between well and barrier is not as abrupt as for other orientations presumably due to the
fast growth rate on the c-plane and possibly due to migration of species from the slow
growth semipolar facet. In Figure 3-5 (c), only the first well is clearly distinguishable at
the nanorod tip. The QW grown on the c-plane is considerably thicker than typically
used for efficient light emitters. This results in poor emission efficiency due to the large
spatial separation between electron and hole wavefunctions.
MQWs grown on the semipolar and nonpolar planes are shown in Figure 3-5 (d)
and (e), respectively. In Figure 3-5 (d), the MQWs grown on semipolar planes are
perhaps too thin to capture the electrons and holes efficiently for radiative recombination.
Similar to the results for GaN nanosheet structures, the growth rates for each plane were
estimated from TEM images, resulting in the following sequence: polar (0001) >
nonpolar {1-100} > semipolar {1-101}. The nanorod sectioned along <0001>-direction is
shown in Figure 3-5 (f). Three pairs of MQW are clearly seen and are equally distributed
on the six nonpolar planes. These results verify the InGaN/GaN MQWs are grown mainly
on the nonpolar planes, which are benificial for InGaN-based LEDs to eliminate the
peizoelectric fields in the active region.
81

Figure 3-5 (a) The nanorod is sectioned along <11-20>-direction to investigate the MQW
growth on three different planes. The inset shows the diffraction pattern of the <11-20>
zone axis and a beam blocker blocks the center beam. Three colored boxes indicate the
locations where magnified images are taken in Figure 3-5 (c), (d), and (e). (b)
InGaN/GaN MQWs are grown on three different planes as indicated in the arrows. (c)
TEM image shows thick a quantum well is grown on the c-plane, polar plane. (d) Thin
MQWs are grown on the semipolar plane. (e) MQWs are grown on the nonpolar plane. (f)
82
The nanorod is sectioned along <0001>-direction to investigate the MQW growth on the
periphery of the GaN nanorod. Three pairs of concentric hexagonal rings of InGaN are
clearly seen in the image. The inset shows the diffraction pattern of the <0001> zone axis.
Since the growth of MQWs on those nonpolar planes are confirmed by the cross-
sectional TEM images, the widths of quantum wells, within the critical thickness, can be
grown much thicker without the concern of the piezoelectric fields. Therefore, thicker
quantum wells were grown on GaN nanorods and their cross-sectional STEM dark field
images are shown in Figure 3-6. The MQWs grown on a GaN nanorod array with 400 nm
pitch spacing show thickness variation from the bottom to the top, as shown in Figure 3-6
(a). The thickness variation is more obvious as the aspect ratio of the gaps between
nanorods becomes higher, probably resulting from the capability of the source species to
reach to the bottom of GaN nanorods. Similar result is also reported by other group.[13]
It is found this issue is less severe as the pitch spacings between nanorods are larger,
resulting in smaller aspect ratios of gaps. The variation in thickness along the radial
direction of nanorods is also observed, as shown in Figure 3-6 (b). This phenomenon
might relate to the increase in the nonpolar surface area, leading to the reduction of the
growth rate on those m-planes, as the diameter of nanorods becomes wider. This issue
could be alleviated by increasing the growth rate for each layer to compensate the
decrease in thickness. These results suggest the mask pattern and the structure need to be
carefully designed to achieve uniform growth of InGaN/GaN MQWs on GaN nanorods.
The uniformity of MQW thickness could be improved by adjusting the growth conditions
to enhance the surface migration as well.
83

Figure 3-6 (a) Thickness variation of MQW grown along the axial direction of GaN
nanorod. The pitch spacing is 400 nm. (b) Decrease in thickness of MQWs from the inner
well to the outer well. The sample is sectioned from a GaN nanorod array with pitch of
500 nm.

3.5 Local emission of multiple quantum wells grown on GaN nanorods
investigated by cathodoluminescence
Local emission from MQWs grown on GaN nanorods was studied by
cathodoluminescence (CL) measurements (Horiba Scientific CLUE series equipped in
the FE-SEM). To understand the local emission spectra from the MQWs on GaN
nanorods, a nanorod with three pairs of quantum wells, lying on its substrate, is
investigated, as shown in Figure 3-7. It is clear to see the CL intensity decreases
dramatically from top to the bottom in Figure 3-7 (b). This might relate to the variation in
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84
MQW thickness along the axial direction of the nanorod, as demonstrated in the FE-SEM
image in Figure 3-7 (a). From both Figure 3-7 (b) and (c), there is an additional peak
close to the tip of the nanorod, possibly resulting from the light emission from the MQWs
grown on the c-plane.

Figure 3-7 CL analysis from the GaN nanorod grown in the pattern with 400 nm center-
to-center spacing in trigonal arrangement. (a) CL spectra were taken at the spots
designated in the SEM image. (b) The comparison of CL spectra from the corresponding
spots in the nanorods. (c) The spectra comparison of peak positions with normalized
intensity from spot 1 to 9.
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85
Another CL measurement was done on a sample with 500 nm ptich spacing. The
result shows relatively uniform CL intensity from the MQWs grown on the nonpolar
plane. The results also suggest the uniform light emission on the nonpolar can be
achieved by optimizing the design of growth mask.

Figure 3-8 CL analysis from the GaN nanorod grown in the pattern with 500 nm center-
to-center spacing in trigonal arrangement. (a) CL spectra were taken at the spots
designated in the SEM image. (b) The comparison of CL spectra from the corresponding
spots in the nanorods. (c) The spectra comparison of peak positions with normalized
intensity from spot 1 to 9.
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86
MQWs grown on c-plane may affect the light emission when the c-plane is
present on the nanorod template. Both samples in Figure 3-7 and Figure 3-8 demonstrate
light emssion from the c-plane. To avoid or minimize MQWs grown on the polar plane,
c-plane surface area must be eliminated. Difference in growth rates on the polar,
semipolar, and nonpolar planes is utilized, and growth on the c-plane is terminated by
applying a continous GaN growth for 10 minutes after GaN nanorod growth. The slow
growth rate of inclined semipolar planes also contributes to enhanced c-plane growth.
Therefore, c-plane is teminated due to its faster growth rate over 10 mins of continuous
growth, as shown in Figure 3-9 (a). The sharp tips of the nanorods are shown in the inset.
After the sample was cleaved to observe the emission from the sidewalls, a CL spectrum
was recorded in the region outlined in the image of Figure 3-9 (b). CL mappings recorded
at 430 nm and 480 nm are shown in Figure 3-9 (c) and (d). The peak emission at 430 nm
is attributed to the MQWs grown on the nanorod sidewalls. The longer emission
wavelength is especially observed at the edges between semipolar and nonpolar planes.
This result is consistent with observation from the cross-sectional TEM image in Figure
3-5 (b). No light emission is observed from the nanorod tips due to exlusion of MQWs
grown on the c-plane. In this way, piezoelectric fields in the MQWs can be greatly
reduced or eliminated due to only the presence of semipolar and nonpolar planes from the
GaN nanorods serving as growth templates for MQWs.
87

Figure 3-9 (a) A uniform array of GaN nanorods with sharp and terminated c-plane is
achieved after 10 minutes of continuous GaN growth. The inset shows the c-plane is
minimized or terminated on top of the nanorods. The FE-SEM image was recorded at a
45° angle. (b) CL spectrum is recorded at the boxed region in the inset. CL mapping
results at 430 and 480 nm are shown in (c) and (d), respectively.

500 nm
430 nm (peak)
480 nm
500 nm
500 nm
(a) (b)
(c)
(d)
1 µm
300 nm
88
3.6 Summary
InGaN/GaN MQWs on uniformly grown GaN nanorod arrays are produced by
selective area growth by MOCVD. Tunable emission over a 60 nm wavelength span,
from near UV to blue green, is demonstrated by varying the mask design. From cross-
sectional TEM images, the InGaN/GaN MQWs are predominantly grown on the {1-100}
nonpolar planes. Also, variation in thickness of MQWs can be minimized by careful
design of growth mask and the MQW structure. The growth of MQWs on c-plane can be
minimized or eliminated by terminating the c-plane before MQW growth. No light
emission from c-plane under these conditions is observed in the CL mapping, confirming
that MQWs grow only on semipolar and nonpolar planes. The elimination of
piezoelectric fields is expected to improve radiative recombination. Therefore, the
exposed nonpolar sidewalls of GaN nanorods are potential templates for the InGaN/GaN
MQWs growth to reduce the strain-induced piezoelectric field in the active regions of
InGaN/GaN light emitting diodes. Reduction of the piezoelectric fields may mitigate the
causes of efficiency droop in LEDs.

3.7 Chapter references
[1] F. Qian, Y. Li, S. Gradecak, D. L. Wang, C. J. Barrelet, and C. M. Lieber,
"Gallium nitride-based nanowire radial heterostructures for nanophotonics," Nano
Letters, 4, 1975, (2004).

[2] W. Bergbauer, M. Strassburg, C. Koelper, N. Linder, C. Roder, J. Laehnemann, A.
Trampert, S. Fuendling, S. F. Li, H. H. Wehmann, and A. Waag, "Continuous-
89
flux MOVPE growth of position-controlled N-face GaN nanorods and embedded
InGaN quantum wells," Nanotechnology, 21, 305201, (2010).

[3] F. Qian, Y. Li, S. Gradecak, H.-G. Park, Y. Dong, Y. Ding, Z. L. Wang, and C. M.
Lieber, "Multi-quantum-well nanowire heterostructures for wavelength-controlled
lasers," Nature Materials, 7, 701, (2008).

[4] F. Qian, S. Gradecak, Y. Li, C. Y. Wen, and C. M. Lieber, "Core/multishell
nanowire heterostructures as multicolor, high-efficiency light-emitting diodes,"
Nano Letters, 5, 2287, (2005).

[5] R. Koester, J.-S. Hwang, D. Salomon, X. Chen, C. Bougerol, J.-P. Barnes, D. L. S.
Dang, L. Rigutti, A. d. L. Bugallo, G. Jacopin, M. Tchernycheva, C. Durand, and
J. Eymery, "M-Plane Core-Shell InGaN/GaN Multiple-Quantum-Wells on GaN
Wires for Electroluminescent Devices," Nano Letters, 11, 4839, (2011).

[6] H. M. Kim, Y. H. Cho, H. Lee, S. I. Kim, S. R. Ryu, D. Y. Kim, T. W. Kang, and
K. S. Chung, "High-brightness light emitting diodes using dislocation-free indium
gallium nitride/gallium nitride multiquantum-well nanorod arrays," Nano Letters,
4, 1059, (2004).

[7] Y.-J. Lee, S.-Y. Lin, C.-H. Chiu, T.-C. Lu, H.-C. Kuo, S.-C. Wang, S. Chhajed, J.
K. Kim, and E. F. Schubert, "High output power density from GaN-based two-
dimensional nanorod light-emitting diode arrays," Applied Physics Letters, 94,
141111, (2009).

[8] Y. L. Chang, J. L. Wang, F. Li, and Z. Mi, "High efficiency green, yellow, and
amber emission from InGaN/GaN dot-in-a-wire heterostructures on Si(111),"
Applied Physics Letters, 96, 013106, (2010).

[9] H. P. T. Nguyen, S. Zhang, K. Cui, X. Han, S. Fathololoumi, M. Couillard, G. A.
Botton, and Z. Mi, "p-Type Modulation Doped InGaN/GaN Dot-in-a-Wire White-
Light-Emitting Diodes Monolithically Grown on Si(111)," Nano Letters, 11,
1919, (2011).

[10] X. Zhang, P. D. Dapkus, and D. H. Rich, "Lateral epitaxy overgrowth of GaN
with NH3 flow rate modulation," Applied Physics Letters, 77, 1496, (2000).

[11] R. S. Q. Fareed, J. W. Yang, J. P. Zhang, V. Adivarahan, V. Chaturvedi, and M.
A. Khan, "Vertically faceted lateral overgrowth of GaN on SiC with conducting
buffer layers using pulsed metalorganic chemical vapor deposition," Applied
Physics Letters, 77, 2343, (2000).

90
[12] S. D. Hersee, X. Sun, and X. Wang, "The controlled growth of GaN nanowires,"
Nano Letters, 6, 1808, (2006).

[13] J. J. Wierer, Jr., Q. Li, D. D. Koleske, S. R. Lee, and G. T. Wang, "III- nitride
core-shell nanowire arrayed solar cells," Nanotechnology, 23, (2012).

91
Chapter 4 InGaN/GaN Multiple Quantum Well Growth
on GaN Nanosheets
Recently, GaN nanostructures have been studied due to their promising properties,
including large surface-to-volume ratio and the exposure of facets other than the typical
polar (0001), basal plane. These nanostructures are grown on easily accessible substrates
mainly by two approaches: vapor-liquid-solid growth (VLS), and selective area growth
(SAG) by metal organic chemical vapor deposition (MOCVD).[1, 2] In this work, we
demonstrate that semipolar and nonpolar facets can be exposed in GaN nanosheets grown
vertically from the basal plane of the polar substrates. These exposed semipolar/nonpolar
facets can be utilized as growth templates for InGaN-based LEDs to reduce or eliminate
the piezoelectric fields in the active regions. GaN nanowire arrays have been
demonstrated for application to LED devices, but the active area of the resulting material
is limited by the aspect ratio and fill factor of structures and the vertical and isolated
features require sophisticated procedures to prepare contacts for devices.[3-5] MQWs
grown on semipolar planes of GaN nanowires/nanorods have been demonstrated by
MOCVD, but the nonpolar surface has not been utilized as the dominant growth planes in
these studies.[6, 7]
In this work, GaN nanosheets are formed and confined by two parallel {1-100}
planes by applying a pulsed GaN growth mode on a dielectric mask with <11-20>-
oriented stripe patterns prepared on GaN bulk material in a close-coupled showerhead
92
MOCVD system. InGaN/GaN MQWs are then grown on these nanostructures by
changing the growth scheme from pulsed to continuous injection of the gas flux, resulting
in nanostructures with MQWs predominantly grown on the exposed GaN nonpolar
surface. By virtue of their linearly connected structure, GaN nanosheets provide a
simplified electrically conducting path for LED fabrication.

4.1 The dependence of mask design on GaN nanosheet growth
Electron beam lithography was used in this work to define stripe patterns and to
explore the orientation-dependent growth of the nanosheets. In particular, the dependence
of the lateral and vertical growth rate as a function of small angular misalignment to the
<11-20> direction helped to define the sensitivity of the structure shape to the alignment.
The design of the stripe patterns nominally aligned to the <11-20>
GaN
was varied in the
range of ±3° to test this sensitivity. The width was observed to increase as the orientation
of the stripe patterns deviated from the <11-20>
GaN
orientation, as shown in Figure 4-1
(a). The exaggerated illustration showing the approach to estimate angles and widths of
the FE-SEM top view images of nanosheets is illustrated in the upper left corner of
Figure 4-1 (a). Some representative top-view images are also included in Figure 4-1 (a).
The red dotted line is a linear fit to these data, yielding a slope of ~60 nm/degree. The y-
intercept of the fitted curve showed a minimum width of approximately 131 nm would be
achieved if the mask orientation was aligned exactly to <11-20>
GaN
. This corresponds to
the designed width of the stripe pattern. The dependence of the nanosheet width on
93
growth mask orientation indicates that the width of the nanosheet can be controlled either
by the width of the design pattern or by the misorientation angles. Conversely, control of
reproducibility will require precise alignment control.

Figure 4-1 Width and height of GaN nanosheets affected by the orientation of the growth
pattern. (a) The nanosheet width shows angle-dependent lateral growth within the small
H1
H2
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Width
94
range of misoriented patterns. The scale bar is 1 µm in all SEM top view images. (b) The
nanosheet height also shows orientation-dependent growth within the small angles. The
inset is a schematic diagram of the nanosheet cross-section. H1 indicates the height of
nanopolar planes and H2 shows the height measured from the c-plane top surface.
The height and shape of the nanosheet is related to the stripe orientation as well.
The illustration of the nanosheet cross-section shown in the upper left corner of Figure 4-
1 (b) indicates the average height of the non-polar plane (H1) and the average height of
the nanosheet (H2), respectively. Within the 3° misorientation range, H1 shows a slightly
decreasing rate of 0.7 nm/degree by a linear fit. The height of non-polar planes, H1, is
less sensitive to the mask orientation, and also implies that the non-polar surface utilized
for the subsequent MQW growth does not vary greatly within the small misorientation
range in the mask. However, H2 shows a more pronounced increasing rate in height than
H1. This may result from the larger exposed surface area of {1-101} surface, a slow
growth plane, enhancing the growth rate on the c-plane. From the results shown in Figure
4-1 (a), the greater the deviation of the stripe from the <11-20>
GaN
direction, the wider is
the nanosheet. Thus, the lateral growth rate of the nanosheets with larger misorientation
increases more rapidly than those with smaller misorientation before the nanosheets form
continuous vertical {1-100} sidewalls.
To observe the growth transition of the nanosheets, an experiment was conducted
by varying growth cycles on samples with different stripe orientations on the mask, and
the results are shown in Figure 4-2. These results are interpreted to indicate that a larger
deviation from the <11-20>
GaN
orientation results in more growth steps created on the
95
sidewalls of the nanosheets in the early growth stage. Therefore, the lateral growth
increased dramatically in the nanosheets grown on the stripe pattern with larger
misoriented angles within the first 60 growth cycles. The growth steps on the sidewalls
diminished as the sidewalls became parallel to the {1-100} nonpolar planes, resulting in a
saturation of the lateral growth, especially for those nanosheets with larger misorientation.

Figure 4-2 Growth transition of nanosheet width varied by pattern orientations. Average
width of nanosheets varying with growth cycles is dependent on the orientation of the
stripe opening on the mask. The error bar is the standard deviation from the measured
result.

96
4.2 Multiple quantum well growth on GaN nanosheets
Similar to the sample preparation for GaN nanorod growth, Si-doped GaN grown
on a c-plane sapphire substrate was used, onto which a 20 nm-thick SiN
x
layer was
deposited by plasma-enhanced chemical vapor deposition. Polymethylmethacrylate
(PMMA), an electron beam resist, was coated on top of the dielectric mask prior to
electron beam writing. A specific orientation of stripe pattern along <11-20>
GaN
was
prepared for nanosheet growth. The pattern was transferred from the PMMA resist into
the SiN
x
dielectric mask using a CF
4
-based reactive ion etching process, and the surface
was then cleaned with solvents and further treated with oxygen plasma to remove any
resist residue. The sample was then loaded into the MOCVD chamber for GaN nanosheet
growth. GaN nanosheet arrays were grown in a close-coupled showerhead MOCVD
system. Trimethylgallium (TMG) and ammonia (NH
3
) were used as the precursors for the
nanosheet growth. The growth pressure was 200 Torr for the entire pulsed growth mode
process. Various growth parameters, including growth temperature, gas flow rates, and
precursor flow periods, were adjusted to achieve GaN nanosheets with high aspect ratios.
The flow rates of TMG and NH
3
are 17.7 µmole/min and 67 mmole/min, respectively.
The surface morphology was evaluated by a Hitachi S-4800 field emission
scanning electron microscope (FE-SEM) and the facet orientations of the GaN nanosheet
were deduced from the orientation of the underlying GaN bulk material. No evidence of
Ga droplets was ever observed, indicating that the vertical nanosheets were the result of
kinetically controlled selective growth of the structures, as opposed to VLS growth
97
processes. Near-vertical semi-polar/non-polar sidewalls grown on <1-100>-orientated
stripes have been reported previously using alternate exposure growth.[8] However, the
low vertical-to-lateral ratio achieved in that work resulted in a smaller area of exposed
nonpolar surface owing to faster growth rates on the {11-20} plane than on the {1-100}
plane that occurred under the growth conditions employed. InGaN/GaN MQWs were
grown sequentially on the GaN nanosheets using trimethylindium (TMI) and
triethylgallium (TEG) as precursors as the indium and gallium sources, respectively. The
growth pressure was increased to 300 Torr in nitrogen ambient to grow InGaN quantum
wells and GaN barriers. NH
3
was injected into the chamber until the temperature
decreased below 400°C.

4.3 Light emission from multiple quantum wells on GaN nanosheets
studied by photoluminescence
Photoluminescence measurements were performed by illuminating the samples
with a focused 325 nm laser beam from a HeCd laser to measure the light emission from
the MQWs. Light emission from the MQWs is confirmed by photoluminescence (PL)
measurements, as shown in Figure 4-3. A dominant QW emission peak, around 436 nm,
and a peak from the GaN band edge are clearly observed in the spectrum. A typical
defect emission, around 550 nm, is also observed in Figure 4-3, which may come from
the underlying GaN bulk, or GaN nanosheets, or both. The inset of Figure 4-3 shows a
slight decrease in peak wavelengths of the QW emission with increasing mask
98
misorientation. The rate is about -1.71 nm/degree, a linear fit. The negative slope on this
data may result from the larger surface area of the inclined semipolar planes in the stripe
pattern with larger misoriented angles. The energy difference between the shortest and
the longest peaks is only 30 meV, possibly the result of the independence of the nonpolar
surface area on the orientation of the stripe described earlier. This result suggests the
geometry of GaN nanosheets provides a viable growth template for MQWs to emit
consistent wavelengths, even with small misorientation in the mask design. From
consideration of the cross-sectional TEM images and the trends we observe with stripe
misorientation, we believe the light emission is predominantly from the MQWs grown on
the vertical sidewalls of the nanosheets.

Figure 4-3 Room temperature photoluminescence measured from InGaN/GaN MQWs
grown on GaN nanosheet arrays. Photoluminescence (PL) of three pairs of MQWs grown
99
on a GaN nanosheet array. The inset shows PL results of emission peaks versus
misoriented angles of a stripe mask.
PL emission from the MQWs grown on the nonpolar planes of nanosheets has also been
verified by the removeal of c-plane and semipolar planes by a selective dry etching
process, which involves filling the gaps between nanosheets with HSQ to protect the
MQWs grown on nonpolar planes, followed by a selective etching using inductively
coupled plasma reactive ion etching. PL spectra before and after the removal of the c-
plane are shown in Figure 4-4 (a). This experiemnt was conducted on the same sample.
Therefore, it is clear to compare the surface morphology of nanosheets before and after
the dry ethcing in FE-SEM images in Figure 4-4 (b) and (c). No obvious shift in PL
emission peaks with/withot the presence of MQWs grown on c-plane verifys the growth
of MQWs on the nonpolar planes of nanosheets. The slight increase in PL intensity might
result from the removal of light aborbing MQWs grown on the c-plane.

100

Figure 4-4 (a) Comparison of PL spectra of nanosheets before and after ICP etching. (b)
FE-SEM image of GaN nanosheets with three pairs of quantum wells before ICP etching.
(c) FE-SEM image of GaN nanosheets after ICP etching for 1 minute.

101
4.4 Multiple quantum wells grown on GaN nanosheets studied by
transmission electron microscopy
GaN nanosheets with InGaN/GaN multiple quantum wells were examined using a
JEOL 2100LB transmission electron microscope (TEM) system. Samples were sectioned
using a focused ion beam milling system, JEOL MultiBeam JIB-4500, and the resulting
slices were transferred to a TEM grid by a lift-off process using an Omniprobe
micromanipulator. In Figure 4-5, a cross sectional TEM bright field image is shown of a
nanosheet aligned to the <11-20>
GaN
direction on which three pairs of InGaN/GaN
MQWs were grown. The TEM image confirms that MQWs were grown on the GaN
nanosheet. The exposure of three main facets in the nanosheet resulted in growth
competition leading to thickness variation of MQWs on different planes. Figure 4b also
designates the regions where MQWs grew on polar, semipolar, and nonpolar planes, as
show in Figure 4-5 (c), (d), and (e), respectively. The growth rates of MQWs on each
plane resulted in the following sequence: polar (0001) > nonpolar {1-100} > semipolar
{1-101}. This resulted in quantum well (QW) thicknesses of 5.7 nm, 2.7 nm, and 0.7 nm,
respectively. The QWs grown on the c-plane shown in Figure 4-5(c) are considerably
thicker than typically used for efficient light emitters and result in poor emission
efficiency due to the large spatial separation between electron and hole wavefunctions.
Furthermore, we observe that the interface between well and barrier is not as abrupt as for
other orientations, presumably due to the fast growth rate on the c-plane and possibly due
to migration of species from the slow growth semipolar facet. In Figure 4-5(d), we
102
observe that the MQWs grown on semipolar planes are extremely thin – perhaps too thin
to capture the electrons and holes for radiative recombination.

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GaN nanosheet
growth direction
!"#$%
&'(%
&)(%
&*(%
&+(%
364%
!"#$%
&'(%
)*++,-+.!/01#$%
2-33.40/#$%
)*++,-+%
2-33%
374%
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*)++$
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*)++,0.1"#$
234$
384%
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1&2$
314%
103
Figure 4-5 TEM bright field images taken from a sectioned nanosheet with MQWs. (a)
InGaN/GaN MQWs were grown on three different planes as indicated in the arrows. The
scale bar is 10 nm. (b) A vertical GaN nanosheet grown from its bulk material is shown
in the low magnification image. The magnified images taken in Figure 4-5 (c), (d), and (e)
are indicated in b. Nanosheet surface was covered with carbon to prevent ion beam
damage during sample preparation. The scale bar is 20 nm. (c) TEM image shows thick
MQWs are grown on the c-plane, polar plane. The scale bar is 10 nm. (d) Thin MQWs
are grown on the semipolar plane. (e) MQWs are grown on the nonpolar plane. The scale
bar is 2 nm in (d) and (e).
GaN nanostructures exhibit unique physical properties that relax the strain resulting from
growth on the non-lattice matched sapphire substrate which, in turn, reduces the
dislocation density in the nanostructures.[9, 10] Although the linked structure of GaN
nanosheet limits the dislocation bending to the surface in the <11-20> direction, the large
exposed surface area of GaN nanosheets still enables the threading dislocations, which
originate from the GaN bulk material, to bend to the surface, as shown in Figure 4-6.
Therefore, the geometry of GaN nanosheets contributes to the dislocation reduction.
104

Figure 4-6 Dislocation bending towards the sidewalls of the GaN nanosheet. Dislocation
bending was observed in the nanosheet in the bright field TEM cross-sectional image.
Three pairs of quantum wells were grown after the GaN nanosheet growth. The sample
surface was coated with platinum in the FIB system to avoid ion beam damage. The scale
bar is 50 nm.

4.5 Local emission of multiple quantum wells grown on GaN nanosheets
investigated by cathodoluminescence
Local emission from MQWs grown on GaN nanorods was studied by
cathodoluminescence (CL) measurements (Horiba Scientific CLUE series equipped in
the FE-SEM). The CL spectrum of GaN nanosheets with MQWs grown them is shown in
Figure 4-7. Three distinguishable emission peaks from GaN band edge, MQWs, and
105
yellow band are observed in the spectrum. CL mapping was also conducted to understand
the origin of the emission peak from MQWs. From the inset of Figure 4-7, light emission
from the nonpolar sidewalls contributes to the dominant emission. The result is consistent
with the results from PL measurements and the observation from TEM.

Figure 4-7 CL spectrum of MQWs grown on GaN nanosheets. The inset shows an
overlap of a CL mapping with a FE-SEM image.

106
4.6 Summary
Large nonpolar GaN surface areas are present in uniform GaN nanosheet arrays
grown vertically from the GaN bulk material, which serves as a growth template for
InGaN/GaN MQWs. The width and height of GaN nanosheets are closely related to the
misorientation of the stripe opening of the dielectric masks on which the nanosheets are
grown. InGaN/GaN MQWs grown on polar, semipolar, and nonpolar planes are revealed
by cross-sectional TEM analysis. The strong PL peak indicates that the nonpolar planes
of GaN nanosheets are potential candidates for InGaN/GaN MQW growth to reduce the
piezoelectric fields inside the active regions of light emitting diodes. The aspect ratio of
well-aligned nanosheets may also allow the fabrication of LEDs with large three-
dimensional active areas per chip area that may help to further mitigate efficiency droop
in LEDs.

4.7 Chapter references
[1] X. F. Duan and C. M. Lieber, "Laser-assisted catalytic growth of single crystal
GaN nanowires," Journal of the American Chemical Society, 122, 188, (2000).

[2] S. D. Hersee, X. Sun, and X. Wang, "The controlled growth of GaN nanowires,"
Nano Letters, 6, 1808, (2006).

[3] S. D. Hersee, M. Fairchild, A. K. Rishinaramangalam, M. S. Ferdous, L. Zhang, P.
M. Varangis, B. S. Swartzentruber, and A. A. Talin, "GaN nanowire light
emitting diodes based on templated and scalable nanowire growth process,"
Electronics Letters, 45, 75, (2009).

107
[4] Y.-Y. Huang, L.-Y. Chen, C.-H. Chang, Y.-H. Sun, Y.-W. Cheng, M.-Y. Ke, Y.-
H. Lu, H.-C. Kuo, and J. Huang, "Investigation of low-temperature
electroluminescence of InGaN/GaN based nanorod light emitting arrays,"
Nanotechnology, 22, (2011).

[5] H. M. Kim, Y. H. Cho, H. Lee, S. I. Kim, S. R. Ryu, D. Y. Kim, T. W. Kang, and
K. S. Chung, "High-brightness light emitting diodes using dislocation-free indium
gallium nitride/gallium nitride multiquantum-well nanorod arrays," Nano Letters,
4, 1059, (2004).

[6] W. Bergbauer, M. Strassburg, C. Koelper, N. Linder, C. Roder, J. Laehnemann, A.
Trampert, S. Fuendling, S. F. Li, H. H. Wehmann, and A. Waag, "N-face GaN
nanorods: Continuous-flux MOVPE growth and morphological properties,"
Journal of Crystal Growth, 315, 164, (2011).

[7] F. Qian, S. Gradecak, Y. Li, C. Y. Wen, and C. M. Lieber, "Core/multishell
nanowire heterostructures as multicolor, high-efficiency light-emitting diodes,"
Nano Letters, 5, 2287, (2005).

[8] X. Zhang, P. D. Dapkus, and D. H. Rich, "Lateral epitaxy overgrowth of GaN
with NH3 flow rate modulation," Applied Physics Letters, 77, 1496, (2000).

[9] Y.-S. Chen, W.-Y. Shiao, T.-Y. Tang, W.-M. Chang, C.-H. Liao, C.-H. Lin, K.-C.
Shen, C. C. Yang, M.-C. Hsu, J.-H. Yeh, and T.-C. Hsu, "Threading dislocation
evolution in patterned GaN nanocolumn growth and coalescence overgrowth,"
Journal of Applied Physics, 106, (2009).

[10] R. Colby, Z. Liang, I. H. Wildeson, D. A. Ewoldt, T. D. Sands, R. E. Garcia, and
E. A. Stach, "Dislocation Filtering in GaN Nanostructures," Nano Letters, 10,
1568, (2010).

108
Chapter 5 InGaN/GaN Nano-scale Light Emitting Diodes
InGaN/GaN LED is a structure containing multiple epitaxial layers. Therefore,
planar layers are studied first, including nucleation, doping, and ternary InGaN material
growth, to understand the growth of etch layer in the structure of InGaN/GaN LED. The
growth conditions for the nano-scale light emitting diodes (NanoLEDs) are based on
those from the planar LEDs. Chapter 5 will start with the growth of the planar LED
structure, followed by the fabrication and characteristics of the NanoLED devices.

5.1 Planar InGaN-based light emitting diode
The growth of planar structure is to calibrate the GaN growth conditions after the
system modification, as described in Chapter 1. The growth will be discussed in the
following sequence based on an InGaN-based LED structure from bottom GaN buffer to
the topmost p-type GaN.

• Low temperature GaN buffer
GaN can be grown on sapphire, silicon carbine, silicon, and GaN substrates. Due to
the concern of cost and crystal quality, sapphire substrates are still widely used in GaN
material growth. However, there are some difficulties for the GaN to grow on sapphire
substrates. Lattice mismatch between GaN and sapphire, around 14%, result in the
generation of strain between these two materials. Also, the difference in thermal
109
expansion coefficients between GaN and sapphire is large. Therefore, the generation of
high dislocation density is the consequence of the strain release when GaN is grown on a
sapphire substrate. Due to low ammonia cracking efficiency, high temperature (>1000°C)
was usually performed during GaN growth. However, high growth temperature leads to
the reevaporation of group III elements, Ga, on the substrate surface, resulting in sparse
island or pyramidal growth on the substrate surface due to lack of nucleation sites.
Generally, it is difficult to grow GaN directly onto the sapphire substrate at high
temperature. To avoid the shortcoming of heteroepitaxy between a sapphire substrate and
GaN material, a low temperature GaN buffer is introduced between the substrate and high
temperature GaN. The low temperature GaN could be either amorphous or
polycrystalline. After the low temperature buffer growth, the susceptor was heated up to
high temperature, around 1050°C, prior to high temperature GaN growth. This process is
called recrystallization. Then, high temperature GaN can grow directly on the dense
nucleation sites formed by the low temperature buffer. Some growth conditions before
and after low temperature GaN growth are reported to influence the crystal quality of
subsequent undoped GaN growth, for example, the heat treatment before low temperature
GaN growth, growth rate[1], thickness[2], and annealing rate during recrystallization.[3]
Hydrogen preannealing is used to remove surface damage on the substrate and also
contributes to create an Al-rich surface, which contributes to AlN buffer growth. From
the result of the full width at half maximum (FWHM) of X-ray diffraction measurement,
it is suggested 900°C is the minimum temperature which allows for high GaN crystal
quality.[4] In our case, 1070°C was used for the hydrogen preannealing before the low
110
temperature GaN buffer growth. Higher temperature could also remove the residual
coating of GaN from the previous growth. To see the influence of the V/III ratio effect on
the LT GaN buffer, the V/III ratio was changed from 1920 to 2824 and the corresponding
TMG flow rate was varied from 55.8 µmole/min to 37.9 µmole/min while the ammonia
flow rate was fixed. Then, a 700 second annealing for recrystallization was carried out
and 2µm undoped GaN was grown on the buffer layer. The growth pressure was kept at
200 Torr during the whole growth procedure. The growth temperature of the LT GaN
was kept at 550°C and that of undoped GaN was 1050°C. After growth, the rocking curve
of (002) high resolution X-ray diffraction (HRXRD) and Hall Effect measurement were
performed to check the crystal quality of GaN bulk material. The results are shown in
Figure 5-1 and Figure 5-2. The results show that higher mobility and lower background
concentration was achieved when a higher V/III ratio was applied to the LT buffer
growth.

111

Figure 5-1 Dependence of the FWHM of the HRXRD rocking curves for the (002)
reflections on the V/III ratio during the growth of GaN buffer layer.

Figure 5-2 Dependence of the concentration and mobility on the V/III ratio during the
growth of the GaN buffer layer.
112
The dependence of V/III ratio and crystal characteristics could result from the
influence of the growth rate and/or buffer layer thickness. The influence of the ammonia
cracking efficiency can be excluded because the ammonia flow and the growth
temperature were kept the same during the growth. Based on the HRXRD and Hall Effect
results, the highest V/III ratio, 2824, was chosen as the general condition for the buffer
layer. The ramping rate of the annealing for the low temperature GaN buffer before the
subsequent high temperature GaN growth has shown some influence on the GaN quality.
An optimum ramping rate is suggested around 20°C/min.[3] In our case, the ramping
time from 550°C to 1050°C is increased from 700 sec to 1200 sec. The corresponding
ramping rate was decreased from 43°C/min to 25°C/min. After the temperature ramped
up to 1050°C, the TMG flow was increased from 37.6 µmole/min to 70.8 µmole/min. The
ammonia flow, the same as LT GaN buffer, was fixed. The corresponding V/III ratio was
decreased to 1500. Based on the low ramping rate of the annealing condition, the
narrower FWHM of the rocking curve of the (002) reflection, 308 arcsec, was obtained.
After the Hall Effect measurement, the background concentration, mobility, and
resistivity were 2.1×10
17
cm
-3
, 309 cm
2
/s V, and 0.0912 ohm-cm, respectively. The
typical background concentration of the non-intentionally doped GaN should be around
mid 10
17
cm
-3
. The slightly high doping could result from the growth chamber condition
because several growth experiments have been performed without cleaning it. The
background doping might affect the Mg-doped GaN due to doping compensation from
the background carrier concentration. Thus, p-type GaN could be hard to achieve when
the background doping concentration is high.
113
• Si-doped GaN
For the n-type GaN, disilane is the precursor to provide Si dopants during growth.
2 µm thick Si-doped GaN was grown on 1µm non-intentionally doped GaN. The same
growth condition of the non-intentionally doped GaN is applied to n-type GaN, except for
the introduction of disilane. The disilane flow rate was 1.83 nmole/sec. Hall Effect
measurement was carried out to evaluate the electrical properties of n-type GaN. The
resulting concentration, mobility, and resistivity are 5.12×10
18
cm
-3
, 209 cm
2
/s V, 0.00584
ohm-cm, respectively. Thus, the growth condition is applied to grow blue LEDs.

• InGaN growth
InGaN based material has been widely used in optoelectronics devices, such as
LEDs, LDs, etc. due to its direct and wide range of energy band gaps. The band gap of
In
x
Ga
1-x
N ternary material ranges from 3.4eV to 0.7eV as x varies from 0 to 1. Thus,
InGaN based material ideally can be applied to light emitting or light absorbing devices
from UV to IR. However, it is challenging to grow InGaN materials. The large difference
in the interatomic spacing between InN and GaN results in a solid phase miscibility gap.
[5] The relative high vapor pressure of indium as compared to the vapor pressure of
gallium leads to low indium incorporation efficiency in the InGaN materials. These
problems can be minimized by optimizing the growth parameters, such as growth
temperature, pressure, V/III ratio, and growth rate, etc.[6] Triethylgallium (TEG) was
used as the group III source for InGaN growth due to its high purity and low vapor
pressure. Lower vapor pressure might lead to lower growth rate, which contributes to
114
improving InGaN crystal and optical properties. Ammonia was used as group V source
but its cracking efficiency is low at low temperature. Thus, the ammonia flow rate was
increased from 2.4 slm to 5 slm and the corresponding V/III ratio was increased from
1500 to 13373 at the same time for InGaN growth. Indium incorporation in InGaN bulk
was found to be affected by the growth rate.[7] The reevaporation of indium species from
the surface will be suppressed by higher growth rate due to the indium species trapped by
the growing layer. However, the improvement in the indium incorporation rate is not the
only concern during the InGaN growth. High quality InGaN could be obtained by
reducing the growth rate, since low growth rate allows adatoms on the surface to have a
longer time to arrive at the two-dimensional step edges of the growth front, which
enhances the optical and crystal quality.[8] Therefore, low growth rate, less than 1 Å/sec,
was carried out during InGaN growth. The pressure during InGaN growth influences the
growth rate of InGaN because mass transportation of precursor gases through the
boundary layer on the substrate is enhanced at lower growth pressure.[9] The higher the
growth pressure, the lower the growth rate. From the PL and HRXRD measurement, the
influence of growth pressure is similar to that of growth rate. Lower growth rate leads to
better crystal quality. However, high pressure also results in parasitic gas phase reactions.
Thus, the pressure, 300 Torr, was maintained during the InGaN growth. Nitrogen gas
ambient was also an important factor which affects the indium incorporation
efficiency.[10] Therefore, the gas ambient was changed from hydrogen to nitrogen during
InGaN growth. The growth temperature plays an important role in the indium
incorporation efficiency.[11] The desorption of indium from the surface is substantially
115
decreased at lower growth temperatures. However, lower growth temperature could lead
to high indium composition in InGaN which may cause indium droplets, phase separation,
and composition inhomogeneity.[6] Also, low ammonia cracking efficiency is another
factor which should be taken into account. Therefore, there is a trade-off during InGaN
growth. Generally, the growth temperature of InGaN was around 750°C. Based on those
conditions, the InGaN bulk materials were grown on 1 µm-thick non-intentionally doped
GaN and the growth temperatures of InGaN were between 710°C to 760°C. Figure 5-3
shows the results of these InGaN bulk materials from ω-2θ scan in HRXRD.

Figure 5-3 The comparison of growth temperature with indium incorporation efficiency
investigated by HRXRD.

116

Table 5-1 The estimated indium composition of InGaN vs. growth temperatures.

From the results shown in Table 5-1, the indium composition roughly estimated
by Vagard’s law can be achieved from 16.8% to 41.6% within this temperature range.

• Mg-doped GaN
Magnesium was used as the dopant for p-type GaN. The growth temperature and
pressure for p-GaN are 1030°C and 100 Torr. TMG flow rate and CP
2
Mg flow were 17.7
µmole/sec and 0.77 µmole/sec, respectively. The V/III ratio was around 3000. Nakamura
et al. demonstrated thermal annealing to activate p-type dopants.[12] After growth, the
sample was annealed at 750°C in nitrogen ambient for 15mins in a furnace. From the
results of the Hall Effect measurement, the concentration, mobility, and resistivity are
3.4×10
17
cm
-3
, 16.2 cm
2
/s V, and 2.39 ohm-cm. Therefore, p-type GaN can be grown,
even though the background concentration was around 2×10
17
cm
-3
. Then, a LED
structure can be grown in the stack of the undoped GaN, Si-doped GaN, InGaN-QW, and
Mg-doped GaN.

117
• Blue LED
To achieve bright blue LEDs, the growth of the active region, MQWs, is very
important. Therefore, some growth conditions should be taken into account, for example,
high barrier growth temperature [13], well protection layers [14], and the number of well-
barrier pairs.[15] Better GaN quality should be grown at high temperature due to the low
cracking efficiency of ammonia. However, the indium incorporation rate is strongly
influenced by the growth temperature. The lower the growth temperature, the higher the
indium incorporation rate. Thus, there is a trade-off between growth temperature and
GaN quality. The growth temperature difference between the well and barrier, around
100°C, was carried out to improve the material quality during MQW growth. Also, a well
protection layer is added to protect the quantum well from thermal damage and then the
temperature can be raised to higher temperature for barrier growth. The number of well-
barrier pairs is also a factor, which will affect the internal quantum efficiency of LEDs.
Thus, six pairs of well and barrier are chosen as the structure for MQWs. The LED is a
structure composed of buffer, undoped GaN, Si-doped GaN, MQWs, and p-GaN. The
structure is shown in Figure 5-4.

Figure 5-4 Schematic structure of a InGaN-based LED.
118
As mentioned before, InGaN growth is influenced by the growth temperature. If
the subsequent Mg-doped GaN is grown at the temperature above 1000°C, the growth
temperature is too high to deteriorate the InGaN inside the MQW. The influence of p-
GaN growth temperature is shown in Figure 5-5 from Oh et al.

Figure 5-5 Cross-sectional HRTEM images of MQWs with p-GaN epilayers grown at (a)
1050°C and (b) 900°C, and EDS indium composition profile across (c) MQW grown at
1050°C and (d) MQW gown at 900°C. [16]
Therefore, the p-GaN growth temperature was fined-tuned from 1030°C to 950°C
to reduce the damage to the InGaN layer. Figure 5-6 demonstrates EL spectrum and a
picture of a blue LED wafer after turned on by a forward bias.
119

Figure 5-6 (a) EL spectrum of a blue LED wafer. (b) A picture of a blue LED after turned
on by a forward bias.

Several GaN growth conditions were tested in the planar structures, such as buffer,
undoped GaN, Si-doped GaN and Mg-doped GaN. Although the background
concentration was around 2×10
17
cm
-3
, p-GaN still can be grown based on the background
concentration. High indium concentration of InGaN can be achieved by low growth
temperature, low growth rate, high V/III ratio, and nitrogen ambient. The fact the blue
LEDs can be grown in the GaN reactor validates the success of reactor modification.
From the results of bulk materials and LED devices, the GaN reactor is in good shape and
can be used to explore GaN-based advanced materials.

!"#$
!%#$
120
5.2 NanoLEDs grown on GaN nanorod arrays
NanoLEDs grown on GaN nanorod arrays have been demonstrated by using
different epitaxial techniques, such as MBE,[17] MOCVD,[18] and HVPE.[19] Single
nanowire LED with core-shell MQWs also has been demonstrated by catalyst and
catalyst-free growth.[20, 21] Anisotropic dry etching techniques can shape planar LED
wafers into nanorod arrays to fabricate NanoLEDs.[22, 23] The surface damages from the
dry etching process require a regrowth step or additional surface treatment to recover the
surface for subsequent heterostructure growth. A combination of dry etching and
regrowth has been explored to fabricate NanoLEDs.[24, 25] None of these results
demonstrates a monolithic growth of the whole structure with the full utilization of the
nonpolar facets formed by the GaN nanostructures as a solution to eliminate the
piezoelectric fields created by the InGaN/GaN MQWs. This section will focus on the
NanoLED structure and characteristics of their performances.

5.2.1 Dopant concentration estimated by capacitance-voltage measurement
Capacitance-voltage measurement (C-V) is an approach to estimate the
concentration of dopants in semiconductor materials by measuring the capacitance
changes with respect to reverse biases applied on a Schottky contact. The difference
between work function of chosen metals and the electron affinities of semiconductors
will form Schottky barriers. When an ohmic contact is prepared on the other side of the
121
semiconductor material, the structure is a Schottky diode. A depletion region on the
Schottky contact will be formed on the semiconductor side. The depletion width changes
with respect to the applied reverse bias will lead to changes in capacitance. Therefore, the
doping level of a semiconductor can be derived from the variation of capacitances. The
following Figure 5-7 (a) shows a Schottky diode structure.

Figure 5-7 (a) An example of a Schottky diode. (b) A schematic diagram of a Schottky
barrier formed by a metal with work function lager than that of a semiconductor.
An illustration of a Schottky barrier formed on a n-type semiconductor is shown
in Figure 5-7 (b). The work functions of metal and semiconductor are Φ
m
and Φ
s
,
respectively, before they contact to each other. W is the depletion width and χ is the
electron affinity of the semiconductor. The electron affinity of GaN is 4.1 eV and the
work function of Au is 5.1 eV. Therefore, the barrier height, Φ
B
, is 1 eV. The capacitance
changes with respect to the depletion width. The following common expression, which
charge is a liner function of voltage, is no longer feasible.
( 5-1)
!
C =
Q
V
Semiconductor
Ohmic contact
Schottky contact
!"#$ !%#$
!"#$%&
'
()&
'
(*&
'
+&
'
,&
-&
./
0
1.2/
)
345&
.2/
)
3/
6
51.,
7&
122
A more general expression definition is shown below.

(5-2)

The charge Q varies nonlinearly with the applied voltage. N
a
and N
d
are the
acceptor and donor concentrations and the depletion width, W, can be expressed in the
following equation.

(5-3)

Where q is the electron charge and ε is dielectric constant of the material. V
0
is the built-
in potential.

(5-4)

(5-5)

x
n0
and x
p0
are the depletion widths on the n-type and p-type side, respectively.

(5-6)
(5-7)

(5-8)

where A is the Schottky contact area.

!
C =
dQ
dV
!!
!
W =[
""V
#
q
(
N
a
+N
d
N
a
N
d
)]
$/"
!!
!
W =[
""(V
#
#V)
q
(
N
a
+N
d
N
a
N
d
)]
$/"
!!
!
Q =qAx
n"
N
d
=qAx
p"
N
a
!!
!
x
n"
=
N
a
N
a
+N
d
W
!!
!
x
p"
=
N
d
N
a
+N
d
W
!!
!
Q =qA
N
a
N
d
N
a
+N
d
W =A "q"(V
#
#V)
N
a
N
d
N
a
+N
d
$
%
&
'
(
)
$/"
123
(5-9)
(5-10)

In a Schottky contact of n-type semiconductor, assumed to be p
+
-n, N
a
>>N
d
and x
n0
≈ W
The capacitance can be derived from the following equation.
(5-11)
Based on the relation between the changes of capacitance with respect to the applied
voltage, the dopant concentration can be estimated.
To study the dopant concentration of the n-type GaN nanorod, GaN nanorod
arrays are grown for the C-V measurements, instead of single nanorod. The length of the
nanorod is about 1 µm long. A schematic diagram of the sample structure for C-V
measurement is shown in Figure 5-8 (a) and GaN nanorods with exposed tips are
demonstrated in Figure 5-8 (b).

!!
!
C
j
=
dQ
d(V
"
"V)
=
A
#
#q#
(V
"
"V)
N
a
N
d
N
a
+N
d
$
%
&
'
(
)
$/#
!!
!
C
j
="A[
q
""(V
#
#V)
N
a
N
d
N
a
+N
d
]
$/"
=
"A
W
!!
!
C
j
=
A
"
[
"q"
(V
#
#V)
N
d
]
$/"
124

Figure 5-8 (a) Schematic diagram of CV measurement for a GaN nanorod array. (b)
Exposed tips of GaN nanorods are shown at the Au contact boundary.
The sample preparation procedures are shown in the following sequence.
1. GaN nanorods were grown by MOCVD.
2. Gaps between GaN nanorods were infilled with HSQ. Spinning speed (1.2k rpm,
Soft baking 180°C for 1.5 minutes and hard baking 320°C for 20mins)
3. Au Schottky contact regions were defined by electron-beam lithography.
4. A 60-nm-thick Au metal contact was deposited by a metal evaporator, followed
by a lift-off process.
5. Ohmic contact is prepared by pressing indium metal directly on n-type GaN
surface.
Three different types of samples, non-intentionally doping, pulsed doping,
continuous doping, has been prepared for the purpose of this experiment. The disilane
flow rate was fixed at 0.1 sccm for the doping experiment. The doping level of the pulsed
N-type GaN
GaN nanorods
Au contact
Indium
contact
HSQ
Sapphire substrate
!"#$%&'($'#
)(*# )+*#
125
doping sample is around 5×10
18
cm
-3
and that of the continuous doping is about 7×10
18

cm
-3
. These are typically doping levels for the InGaN-based LED structure. However, the
result of the non-intentionally doped sample is not consistent with the others and it is
found the infilling of HSQ affects the result dramatically. The estimated doping level can
be two orders of magnitude different if the sidewalls of nanorods are not fully covered by
HSQ. The calculated dopant concentration is inverse proportional to the square of contact
area, resulting in the strong dependence of the contact surface area to the dopant
concentration. Another possible junction formed at the base of nanorods could influence
the doping estimation as well.

5.2.2 Resistivity measurement of single GaN nanorod
For device fabrication, studying the resistivity of individual nanorod can help to
understand the incorporation of dopants during the pulsed growth. First of all, the
nanorods have to be long enough for the contacts preparation. Generally, nanorods longer
than 2 µm were grown for the purpose of this measurement. After growth, GaN nanorods
were removed from the GaN/sapphire substrate in an isopropanol (IPA) solution, the
same preparation procedure for TEM samples. The nanorods suspended in IPA were
dispersed on a silicon substrate with 300-nm-thick thermally oxidized SiO
2
on top of it.
The density of these GaN nanorods on the silicon substrate depends on the quantity of
nanorods grown on the GaN/sapphire substrate and the number of IPA drops on the
silicon substrate. After nanorods dispersed on the Si substrate, an oxygen plasma
treatment was performed to clean the sample surface. The location of GaN nanorods can
126
be recorded from the coordinates in the electron beam writing system. To prepare
contacts on the nanorods, PMMA 495 C5 was spun on the silicon substrate under 3000
rpm for 60 sec and baked on a hot plate for 1.5 minutes. The sample was loaded into the
electron beam writing system to write the contact patterns on the silicon substrate. After
development, the whole sample was treated with oxygen plasma again for 6 seconds to
remove any possible residues in the contact pattern regions. Before metal evaporation, the
sample surface was cleaned with a dilute hydrochloric acid (HCl:H
2
O=1:10) for 3
seconds. Ti/Au (25 nm/250 nm) was deposited on GaN nanorods as the electrodes. After
metal lift-off process, the contacts prepared on a nanorod are shown in Figure 5-9 (a).
The separation between two contacts is about 1 µm. The I-V characteristic curves of four
different GaN nanorods are shown in Figure 5-9 (b). The legend in Figure 5-9 (b)
indicates the designated number for each nanorod. These nanorods were grown in a
pulsed growth mode with continuous disilane injection. The disilane flow rate was 0.1
sccm.
Figure 5-9 (a) A FE-SEM image of two contacts prepared on a GaN nanorod for
!"#$ !%#$
127
resistivity measurement. (b) I-V curves of four different nanorods. The curves show
Schottky behavior.
An average resistivity of these four nanorods is estimated form the linear regions
of both positive and negative biases, which is 0.067 ohm-cm. This number has been
calibrated by the dimension of each nanorod from its FE-SEM image. The corresponding
concentration in planar n-type GaN is estimated to be around mid 10
17
cm
-3
. The result
indicates the doping level is not high enough for the device fabrication or the carrier
concentration is underestimated due to the Schottky contacts. Surface treatment prior to
the contact deposition and appropriate choice of metals for ohmic contacts are crucial for
the measurement.

5.2.3 NanoLED structures
In chapter 3, the core-shell MQW structure grown on GaN nanorod arrays has
been revealed by TEM images and the light emission predominantly from the MQWs
grown on the nonpolar planes has also confirmed by CL. Therefore, the last step to
fabricate a NanoLED is to grow a p-type GaN to form a p-n junction for the diode. Three
different kinds of p-GaN will be discussed in the following sequence, planer p-GaN, non-
planar p-GaN, and current blocking p-GaN (CBP).

128
• Planar p-GaN
A schematic diagram of NanoLED with planar p-GaN is shown in Figure 5-10 (a).
The growth template for NanoLEDs can be formed by n-type GaN nanorods or GaN
nanosheets. To understand the coalescent behavior, nanorods have been studied first to
grow and fabricate NanoLEDs. Figure 5-10 (b) demonstrates a planar p-GaN formed by a
coalescent process. GaN nanorods embedded in the p-GaN can be observed in the cross-
sectional FE-SEM image as well. Two different growth steps are involved in the planar
p-GaN growth. The first step is to grow a p-GaN in nitrogen-rich ambient to enhance the
lateral growth of nanorods. TMG and ammonia flow rates were 17.7 µmole/sec and 53.6
mmole/sec, respectively. The growth pressure was 200 Torr and the growth temperature
setting was 1050°C. After the nanorods connected to each other via their {11-20} planes,
a second growth step in hydrogen ambient was introduced to planarize the p-GaN layer.
TMG and ammonia flow rates were 13.3 µmole/sec and 53.6 mmole/sec, respectively.
The growth pressure was 100 Torr and the growth temperature setting was 1070°C. Low
growth rate and high temperature lead to the planarization of p-GaN. The growth mask of
the NanoLED shown in Figure 5-10 (b) is a dot array with 400 nm pitch spacing in a
trigonal arrangement. The surface morphology studied by FE-SEM is shown in Figure 5-
10 (c). The trace of the nanorod positions and coalescent fronts can be seen on the p-GaN
surface. A nanorod array with square arrangement after the coalescence of p-GaN is
demonstrated in Figure 5-10 (d).
129

Figure 5-10 (a) Schematic diagram of NanoLEDs. (b) A FE-SEM cross-sectional image
of the NanoLED. A coalescent p-GaN layer is grown on top of the nanorods. (c) Surface
morphology of a coalescent p-GaN over a trigonal nanorod array with 400 nm pitch
spacing. (d) Surface morphology of a coalescent p-GaN over a square nanorod array with
500 nm pitch.
The surface morphology shown in Figure 5-10 (c) and (d) is flat enough for the
deposition of a p-type contact to fabricate NanoLED devices.

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130
• Non-planar p-GaN
If only the first growth step mentioned in the planar p-GaN is applied to the
NanoLED, the p-GaN will form a non-planar surface. The non-planar p-GaN connects
all the nanorods through a-planes, as shown in Figure 5-11 (a). The coverage of p-GaN
layer over all the nanorods is clearly seen in Figure 5-11 (b). This non-planar surface
morphology has the potential to enhance the light extraction to improve the efficiency of
NanoLEDs. A thicker contact is required to form a continuous layer to improve current
spreading due to the thin p-GaN layer and non-planar feature.

Figure 5-11 (a) Top view FE-SEM image of non-planar NanoLED. (b) A connected p-
GaN layer covers the top of nanorods to form a non-planar surface.

• Current blocking p-GaN (CBP)
Direct growth of p-GaN on NanoLED structure could result in carrier
recombination at different facets. Even though the c-plane is pinched-off, the inclined
semipolar planes still coexist with the nonpolar planes. Therefore, a current blocking
!"#$ !%#$
131
layer, 10 nm thick SiO
2
, deposited on top of the nanorods by a dielectric evaporator can
serve as the growth mask to prevent the p-GaN from growing directly on the tips of GaN
nanorods after MQW growth. The result of CBP is shown in Figure 5-12. From both top
view and 45°-tilt FE-SEM images, the directional deposition of SiO
2
results in the p-GaN
growth only occurring on the sidewalls. However, this CBP structure involves the growth
interruption and possible SiO
2
deposition on the sidewalls. Therefore, this structure has
not been extensively studied.

Figure 5-12 (a) Top view FE-SEM image of NanoLED with a SiO
2
current blocking layer.
(b) FE-SEM image of the p-GaN grown laterally to link all the nanorods.
An advantage of this CBP structure is capable of investigating the dimension and
location the nanorods embedded in p-GaN. Figure 5-13 shows a p-GaN growth condition,
similar to the second growth step mentioned in the planar p-GaN, was applied to the CBP
to form a coalescent p-GaN. The sample was milled by a FIB to study the cross-sectional
structure. A carbon sacrificial layer was deposited on top of the region for investigation,
as shown in Figure 5-13 (a). Due to the charging effect caused by the dielectric materials,
!"#$ !%#$
132
SiN
x
and SiO
2
, the location and the shape of the nanorod can be easily identified in the
embedded structure, as demonstrated in the white dashed lines in Figure 5-13 (b).

Figure 5-13 (a) A NanoLED is sectioned by FIB to study the nanorods embedded in the
p-GaN layer. (b) A magnified FE-SEM image taken from the boxed region in (a).

5.2.4 I-V characteristics of NanoLED grown on GaN nanorods
In general, a planar structure facilitates the device fabrication. The study of
NanoLED initiated from the structure with planar p-GaN. A transparent metal contact,
Ni/Au (5 nm/5 nm), was deposited on top of the planar p-GaN to form an ohmic contact
after annealed at 500°C for 3 minutes in a nitrogen ambient. Figure 5-14 (a) demonstrates
light uniformly distributes across a 300 µm ×300 µm chip area under a forward bias. The
I-V curve is shown in Figure 5-14 (b). High resistivity and low breakdown voltage could
imply the structure and/or the contact of NanoLED need to be optimized to improve the
device performance.
!"#$%"%&'&()$*+,-$./0)$
1&"23)43%,$56!"#$
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133

Figure 5-14 (a) A picture of a NanoLED after being turned on. (b) I-V curve of the
NanoLED.

5.2.5 Electroluminescence of NanoLED grown on GaN nanorods
The EL spectra under different injection current levels have been studied using
the setup, similar to that of PL, with a current source. All the recorded spectra are plotted
in Figure 5-15 (a). The injection current varies from 5 mA to 90 mA. The device failed at
the current of 100 mA. It is clear to observe two emission peaks contributing to these
spectra. The shorter emission peak could result from the non-uniformity in current
spreading or inhomogeneous growth at local regions. These spectra were fitted with two
Gaussian peaks to study the changes of emission peaks with respect to the injection
current. An example of fitted curves at an injection current of 90 mA is shown in Figure
5-15 (b). Peak 1 and Peak 2 are designated the emission peaks at 486 nm and 438 nm,
!"#$ !%#$
&'()'*$('*)*+",-'.$/010$234$
134
respectively. After all the spectra are fitted with two Gaussian peaks, the resulting data
are plotted in Figure 5-15 (c).

Figure 5-15 (a) EL spectra of a NanoLED under different injection current levels. (b) EL
spectrum measured at 90 mA is fitted by two Gaussian peaks. (c) EL emission peak shifts
with respect to the injection current.
The emission peak, Peak 1, changes from 494 nm to 486 nm from 5 to 50 mA.
The emission peak fixed at 486 nm after the injection current of 50 mA. The blue shift of
Peak 2, from 444 to 438 nm, is also observed from 5 to 50 mA. After passing the
!"#$ !%#$
!&#$
135
injection current of 50 mA, the emission peak stays at 438 nm. Both peaks show the same
trend within the current range. There are some possibilities leading to the blue shift. First,
the poor p-GaN current distribution results in the radiative recombination occurring at
different regions as the injection current varies. This phenomenon has been reported by
Hong et al.[18] Second, the blue shift was also observed in the planar LED grown on m-
plane GaN substrate.[26] The blue shift could be attributed to the band filling at localized
states due to the composition fluctuation of InGaN. Third, the internal electric field in the
p-n junction could cause the band bending of MQWs, which results in the reduction of
the energy under no or low forward bias.
The integral of intensity divided by the total points taken in each spectrum at
different injection current can be plotted as a L-I characteristic measurement, as shown in
Figure 5-16 (a). The x-axis is also plotted as current density defined by the current over
the p-GaN surface area. The value of each point plotted in Figure 5-15 (a) divided by the
current density can be drawn as a relative efficiency plot, as shown in Figure 5-16 (b).
This plot can be seen as number of photons generated by carriers per unit area.

136

Figure 5-16 (a) EL intensity versus injection current density. (b) Relative efficiency
varies with injection current density.
The maximum efficiency occurs at 60 A/cm
2
in the preliminary NanoLED result.
There was no heat-controlled stage for the device measurement and each spectrum was
taken at continuous current injection for at least 5 minutes. Therefore, heat generated
during the measurement could also lead to the drop in efficiency. No clear red shift in
emission wavelength, as shown in Figure 5-15 (c), suggests light emission from the
MQW was not obviously affected by the heat. The NanoLED has the potential to be
operated at higher current density. Device structure and process need to be optimized for
the fabrication of efficient NanoLEDs. Also, different emission wavelengths have been
demonstrated from NanoLEDs, as shown in Figure 5-17. The emission wavelength can
be controlled by the TMI flow rate, resulting in the difference in In/Ga ratio in gas phase.
Also, the nonpolar surface area formed by GaN nanorods could lead to the variation of
both composition and thickness due to the difference in growth rates. The control of
137
growing of long emission wavelengths needs further study to fabricate NanoLEDs
capable of emitting a wide range of colors.

Figure 5-17 (a) Demonstration of different EL spectra from NanoLEDs (b) Pictures of
NanoLEDs with different emission wavelengths under forward bias.

!"#$
!%#$
138
5.3 Summary
Planar LED structure has been demonstrated to show the success of the reactor
modification for GaN growth after the calibration of growth conditions for each layer in a
LED structure. Based on the planar LED growth condition, NanoLEDs are also
successfully grown on GaN nanorods. The doping levels of the n-type nanorods are
estimated by C-V and resistivity measurement. Sample preparation greatly affects the
accuracy of both measurements. Also, the surface recombination in nanorods could also
add an additional effect on the doping estimation. Therefore, there is no conclusive result
about the doping level inside the n-type nanorods at this moment. The NanoLED
structure grown on the nonpolar planes formed by GaN nanorods can be fabricated by
growing a coalescent p-type GaN. Three different kinds of p-GaN layer have been
explored in this study. Finally, NanoLEDs are characterized by I-V, L-I, and EL
measurement and the promising results have been shown in the preliminary NanoLEDs.
Further studies are required to achieve efficient NanoLEDs as a solution for solid-state
lighting.

5.4 Chapter references
[1] K. S. Kim, C. S. Oh, K. J. Lee, G. M. Yang, C. H. Hong, K. Y. Lim, H. J. Lee,
and A. Yoshikawa, "Effects of growth rate of a GaN buffer layer on the properties
of GaN on a sapphire substrate," Journal of Applied Physics, 85, 8441, (1999).

[2] J. N. Kuznia, M. A. Khan, D. T. Olson, R. Kaplan, and J. Freitas, "INFLUENCE
OF BUFFER LAYERS ON THE DEPOSITION OF HIGH-QUALITY SINGLE-
139
CRYSTAL GAN OVER SAPPHIRE SUBSTRATES," Journal of Applied
Physics, 73, 4700, (1993).

[3] C. F. Lin, G. C. Chi, M. S. Feng, J. D. Guo, J. S. Tsang, and J. M. H. Hong, "The
dependence of the electrical characteristics of the GaN epitaxial layer on the
thermal treatment of the GaN buffer layer," Applied Physics Letters, 68, 3758,
(1996).

[4] M. Tsuda, K. Watanabe, S. Kamiyama, H. Amano, I. Akasaki, R. Liu, A. Bell,
and F. A. Ponce, "Mechanism of H-2 pre-annealing on the growth of GaN on
sapphire by MOVPE," Applied Surface Science, 216, 585, (2003).

[5] I. H. Ho and G. B. Stringfellow, "Solubility of nitrogen in binary III-V systems,"
Journal of Crystal Growth, 178, 1, (1997).

[6] F. K. Yam and Z. Hassan, "InGaN: An overview of the growth kinetics, physical
properties and emission mechanisms," Superlattices and Microstructures, 43,
(2008).

[7] S. Keller, B. P. Keller, D. Kapolnek, A. C. Abare, H. Masui, L. A. Coldren, U. K.
Mishra, and S. P. DenBaars, "Growth and characterization of bulk InGaN films
and quantum wells," Applied Physics Letters, 68, 3147, (1996).

[8] S. Keller, B. P. Keller, D. Kapolnek, U. K. Mishra, S. P. DenBaars, I. K. Shmagin,
R. M. Kolbas, and S. Krishnankutty, "Growth of bulk InGaN films and quantum
wells by atmospheric pressure metalorganic chemical vapour deposition," Journal
of Crystal Growth, 170, 349, (1997).

[9] D. J. Kim, Y. T. Moon, K. M. Song, I. W. Lee, and S. J. Park, "Effect of growth
pressure on indium incorporation during the growth of InGaN by MOCVD,"
Journal of Electronic Materials, 30, 99, (2001).

[10] M. Bosi and R. Fornari, "A study of Indium incorporation efficiency in InGaN
grown by MOVPE," Journal of Crystal Growth, 265, 434, (2004).

[11] S. Nakamura and T. Mukai, "HIGH-QUALITY INGAN FILMS GROWN ON
GAN FILMS," Japanese Journal of Applied Physics Part 2-Letters, 31, L1457,
(1992).

[12] S. Nakamura, N. Iwasa, M. Senoh, and T. Mukai, "HOLE COMPENSATION
MECHANISM OF P-TYPE GAN FILMS," Japanese Journal of Applied Physics
Part 1-Regular Papers Short Notes & Review Papers, 31, 1258, (1992).

140
[13] T. C. Wen, S. J. Chang, Y. K. Su, L. W. Wu, C. H. Kuo, W. C. Lai, J. K. Sheu,
and T. Y. Tsai, "InGaN/GaN multiple quantum well green light-emitting diodes
prepared by temperature ramping," Journal of Electronic Materials, 32, 419,
(2003).
[14] J.-W. Ju, H.-S. Kim, L.-W. Jang, J. H. Baek, D.-C. Shin, and I.-H. Lee, "A well
protection layer as a novel pathway to increase indium composition: a route
towards green emission from a blue InGaN/GaN multiple quantum well,"
Nanotechnology, 18, (2007).
[15] M. G. Cheong, E. K. Suh, and H. J. Lee, "High-quality In0.3Ga0.7N/GaN
quantum well growth and their optical and structural properties," Semiconductor
Science and Technology, 16, 783, (2001).
[16] M. S. Oh, M. K. Kwon, I. K. Park, S. H. Baek, S. J. Park, S. H. Lee, and J. J. Jung,
"Improvement of green LED by growing p-GaN on In0.25GaN/GaN MQWs at
low temperature," Journal of Crystal Growth, 289, 107, (2006).

[17] A. Kikuchi, M. Kawai, M. Tada, and K. Kishino, "InGaN/GaN multiple quantum
disk nanocolumn light-emitting diodes grown on (111)Si substrate," Japanese
Journal of Applied Physics Part 2-Letters & Express Letters, 43, L1524, (2004).
[18] Y. J. Hong, C.-H. Lee, A. Yoon, M. Kim, H.-K. Seong, H. J. Chung, C. Sone, Y. J.
Park, and G.-C. Yi, "Visible-Color-Tunable Light-Emitting Diodes," Advanced
Materials, 23, 3284, (2011).

[19] H. M. Kim, Y. H. Cho, H. Lee, S. I. Kim, S. R. Ryu, D. Y. Kim, T. W. Kang, and
K. S. Chung, "High-brightness light emitting diodes using dislocation-free indium
gallium nitride/gallium nitride multiquantum-well nanorod arrays," Nano Letters,
4, 1059, (2004).

[20] F. Qian, S. Gradecak, Y. Li, C. Y. Wen, and C. M. Lieber, "Core/multishell
nanowire heterostructures as multicolor, high-efficiency light-emitting diodes,"
Nano Letters, 5, 2287, (2005).

[21] R. Koester, J.-S. Hwang, D. Salomon, X. Chen, C. Bougerol, J.-P. Barnes, D. L. S.
Dang, L. Rigutti, A. d. L. Bugallo, G. Jacopin, M. Tchernycheva, C. Durand, and
J. Eymery, "M-Plane Core-Shell InGaN/GaN Multiple-Quantum-Wells on GaN
Wires for Electroluminescent Devices," Nano Letters, 11, 4839, (2011).

[22] H. W. Huang, C. C. Kao, T. H. Hsueh, C. C. Yu, C. F. Lin, J. T. Chu, H. C. Kuo,
and S. C. Wang, "Fabrication of GaN-based nanorod light emitting diodes using
self-assemble nickel nano-mask and inductively coupled plasma reactive ion
etching," Materials Science and Engineering B-Solid State Materials for
Advanced Technology, 113, 125, (2004).

[23] C.-Y. Wang, L.-Y. Chen, C.-P. Chen, Y.-W. Cheng, M.-Y. Ke, M.-Y. Hsieh, H.-
M. Wu, L.-H. Peng, and J. Huang, "GaN nanorod light emitting diode arrays with
141
a nearly constant electroluminescent peak wavelength," Optics Express, 16,
10549, (2008).

[24] J. R. Chang, S. P. Chang, Y. J. Li, Y. J. Cheng, K. P. Sou, J. K. Huang, H. C. Kuo,
and C. Y. Chang, "Fabrication and luminescent properties of core-shell
InGaN/GaN multiple quantum wells on GaN nanopillars," Applied Physics
Letters, 100, (2012).

[25] Q. Li, K. R. Westlake, M. H. Crawford, S. R. Lee, D. D. Koleske, J. J. Figiel, K.
C. Cross, S. Fathololoumi, Z. Mi, and G. T. Wang, "Optical performance of top-
down fabricated InGaN/GaN nanorod light emitting diode arrays," Optics Express,
19, 25528, (2011).

[26] A. Chakraborty, B. A. Haskell, S. Keller, J. S. Speck, S. P. Denbaars, S.
Nakamura, and U. K. Mishra, "Demonstration of nonpolar m-plane InGaN/GaN
light-emitting diodes on free-standing m-plane GaN substrates," Japanese Journal
of Applied Physics Part 2-Letters & Express Letters, 44, L173, (2005).

142
Chapter 6 Conclusions and Future work

6.1 Conclusions
• GaN nanostructures
GaN nanostructures have been successfully grown by selective area growth,
including nanostripes, nanopyramids, nanorods, and nanosheets. The growth masks for
these nanostructures are prepared by EBL, which has the flexibility of mask designs for
different kinds of nanostructures. The patterns on the mask demonstrate great impact on
the nanostructure growth, especially the opening diameter and the orientation of the stripe
for nanorods and nanosheets, respectively. Both GaN nanorods and nanosheets are grown
vertically from GaN/sapphire substrates and exhibit nonpolar sidewalls, which can serve
as the growth template for InGaN/GaN MQWs to eliminate the piezoelectric fields in the
quantum wells. Single wurtzite crystal structure of GaN nanorods is confirmed by TEM
without the observation of stacking faults or dislocations. It was found that dislocations,
originating from the GaN/sapphire substrates, bend toward the sidewalls of
nanostructures. The results suggest that the GaN nanostructures can provide dislocation-
free growth templates for the subsequent heterostructure growth.

143
• MQW grown on GaN nanostructures
InGaN/GaN MQWs on uniformly grown GaN nanorod and nanosheet arrays are
produced by selective area growth by MOCVD. CL mappings and cross-sectional TEM
images have verified the growth of MQW predominantly on the nonpolar {1-100} planes.
The surface area formed by nanorods and nanosheets affects PL emission peaks. In the
case of GaN nanorod, PL emission peak varies with the diameter of hole opening on the
growth masks. Emission wavelengths are stabilized in the nanosheet arrays due to slight
changes of sidewall surface area formed by the nanosheets grown within ±3°
misorientation angles. The growth of MQWs on c-plane can be minimized or eliminated
by pitching off the c-plane before MQW growth. No light emission from c-plane under
these conditions is observed in the CL mapping, confirming that MQWs grow only on
semipolar and nonpolar planes. The elimination of piezoelectric fields is expected to
improve radiative recombination and mitigate the causes of efficiency droop in LEDs.

• NanoLED
The structure of NanoLED grown on the nonpolar planes formed by GaN
nanorods has been demonstrated by growing a coalescent p-type GaN over GaN nanorod
arrays. Three different kinds of p-GaN layers have been explored in this research. InGaN-
based NanoLEDs have been characterized by I-V, L-I, and EL measurements and
electrically driven to demonstrate a wide range of emission wavelengths. These
NanoLEDs are expected to be efficient emitters for solid-state lighting.
144
6.2 Future work
• Nanostructure growth
EBL has its limitations in fabricating growth masks in large scale. Therefore,
different patterning techniques have to be explored to increase the throughput of the
growth masks, such as nanosphere lithography, nanoimprint lithography, laser
interferometric lithography, block-copolymer lithography, and etc. Also, growth
conditions might need to be adjusted to grow nanostructures on larger patterns. Growing
nanostructures on Si substrates is another approach to reduce the cost, and has the
capability to work on a larger wafer scale.

• InGaN/GaN growth on nanostructures
From the FE-SEM and TEM results, the variation of MQW thickness might lead
to broad emission spectra and non-uniformity in spatial distribution of light emission.
This issue can be minimized by the design of the mask together with the growth
conditions.

• Device fabrication
Thick p-type GaN layer in the NanoLED structure absorbs light emitting from the
MQWs and increases the resistance of the whole device. Therefore, the nonplanar p-type
GaN combined with a transparent contact layer could be a solution to improve NanoLED
performance. The doping levels of the NanoLED are still not clear. For p-type GaN, the
magnesium incorporation could be different on different growth planes.[1] Therefore,
145
exploring and searching for appropriate analysis technique might help to know the doping
profile of the p-type dopant. For example, atom probe tomography measurement has
demonstrated the capability of analyzing the impurity distribution.[2] In the case of n-
type GaN, the improvement in surface treatment and metal contact preparation can help
the conductivity measurement for single nanorod, and allow one to understand the dopant
incorporation inside the nanorods grown by pulsed growth. If the doping levels are
known and can be controlled, NanoLEDs can be operated at lower bias to reduce heat
generation, which affects the LED performance at high driving current. Once the
NanoLED fabrication is optimized, NanoLEDs can be well characterized to have further
understanding of efficiency droop in these nanostructures.

6.3 Chapter references
[1] D. W. Ren and P. D. Dapkus, "Anisotropic Mg incorporation in GaN growth on
nonplanar templates," Applied Physics Letters, 86, (2005).

[2] D. E. Perea, E. Wijaya, J. L. Lensch-Falk, E. R. Hemesath, and L. J. Lauhon,
"Tomographic analysis of dilute impurities in semiconductor nanostructures,"
Journal of Solid State Chemistry, 181, 1642, (2008).

146
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Abstract (if available)

Abstract There are several challenges for the InGaN-based light emitting diodes (LEDs) to be used for lighting, such as high dislocation density resulting from the lattice mismatch between GaN and sapphire substrates, strong piezoelectric fields induced by the strain in InGaN/GaN multiple quantum wells (MQWs) grown on the polar direction, and fair amount of light trapped inside GaN material due to the difference in refractive indices between GaN and air. Also, the efficiency droop is an important issue for InGaN-based LEDs to be driven at high current for solid-state lighting. ❧ In this thesis, GaN nanostructures, especially GaN nanorods and GaN nanosheets, were studied as potential solutions to overcome the current issues that the conventional planar InGaN-based LEDs are facing. Nanostructures can help to release the strain from the lattice mismatch between GaN and its conventional sapphire substrates. Dislocation bending was observed in the nanostructures. Therefore, GaN nanostructures contribute to forming dislocation-free templates for InGaN/GaN MQWs. Large surface area of nonpolar planes was formed by these nanostructures to serve as growth templates for InGaN/GaN MQWs to eliminate the piezoelectric fields. Both nanorod and nanosheet structures can improve light extraction efficiency due to their nonplanar geometries for reducing the total reflection loss at the GaN/air interface. The feasibility of growing thicker MQWs on nonpolar substrate has been demonstrated and these thick MQWs grown on the nonpolar facets of GaN nanostructures can contribute to increasing the possibility of radiative recombination and reducing the carrier density in MQWs to mitigate the efficiency droop. NanoLEDs are successfully fabricated on GaN nanorod arrays and electrically driven up to 200 A/cm2 with a relative maximum efficiency about 60 A/cm2. This result demonstrates the potential of using NanoLEDs as efficient emitters for solid-state lighting.

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University of Southern California Dissertations and Theses

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Asset Metadata

Creator Yeh, Ting-Wei (author)

Core Title GaN nanostructures grown by selective area growth for solid-state lighting

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Materials Science

Publication Date 05/21/2013

Defense Date 10/16/2012

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag GaN nanorod,GaN nanosheet,GaN nanostructures,LED,nonpolar,OAI-PMH Harvest

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Dapkus, Paul Daniel (committee chair), Nutt, Steven R. (committee member), Zhou, Chongwu (committee member)

Creator Email twpeteryeh@gmail.com,twpeteryeh@yahoo.com

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c3-118807

Unique identifier UC11292740

Identifier usctheses-c3-118807 (legacy record id)

Legacy Identifier etd-YehTingWei-1332.pdf

Dmrecord 118807

Document Type Dissertation

Rights Yeh, Ting-Wei

Type texts

Source University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the a...

Repository Name University of Southern California Digital Library

Repository Location USC Digital Library, University of Southern California, University Park Campus MC 2810, 3434 South Grand Avenue, 2nd Floor, Los Angeles, California 90089-2810, USA