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Optical and electrical characterization of one-dimensional (1D) and two-dimensional (2D) nanostructures

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Optical and electrical characterization of one-dimensional (1D) and two-dimensional (2D) nanostructures

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Content Optical and Electrical
Characterization of One-dimensional
(1D) and Two-dimensional (2D)
Nanostructures

by

Shermin Arab

A Thesis Presented to
The Department of Ming Hsieh Electrical Engineering
Viterbi School of Engineering
UNIVERSITY OF SOUTHERN CALIFORNIA
In Partial Fulfillment of the
Requirements for the Degree
DOCTOR OF PHILOSOPHY

Thesis Approved by:
___________________________
Dr. Stephan. B. Cronin, Chair

___________________________
Dr. Aiichiro Nakano, Member

___________________________
Dr. Wei Wu, Member
August 2015

1

Dedication

This thesis is dedicated to my parents and my brother, who have supported me all
the way. It is also dedicated to my aunt, Sharareh, and grandmother, Taj.

2

Acknowledgment:
This material is based upon work supported as part of the Center for Energy
Nanoscience (CEN), an Energy Frontier Research Center (EFRC) funded by the
U.S. Department of Energy, Office of Science and Office of Basic Energy
Sciences, under Award DE-SC0001013.

3

Table of Contents
Dedication 1
Acknowledgment 2
List of Figures 5
List of Tables 8
Abstract 9

Chapter 1: Introduction 10
1.1 Introduction
1.1.1 One-dimensional Nanostructures (Nanowires)
1.1.2 Two-dimensional Nanostructures (Nanosheets)
10
1.2 Optical Characterization Methods 13
1.3 Electrical Characterization Methods 18
1.4 Growth and Synthesis of Nanostructures
1.4.1 Selected Area Growth Metal Organic Chemical Vapor Deposition
Method
1.4.2 Chemical Vapor Deposition Method
1.4.3 Exfoliation Method for 2D Materials
21

Chapter 2: Observation of Fabry-Perot Resonance from GaAs Nanowires 24
2.1 Observation of Fabry-Perot Resonance from bare GaAs Nanowires at 4K
and Room Temperature
24
2.2 Effects of Passivation on Optical Responses and Fabry-Perot Resonance 32

Chapter 3: Optical Characterization of Doped GaAs Nanowires 36
3.1 Study on the Concentration of Dopants in n-type GaAs Nanowires at 4K 36
3.2 Optical Responses from PIN (p-type, intrinsic, n-type) GaAs Nanowires 46

Chapter 4: Observation of Fabry-Perot Resonance from GaAs Nanosheets 48
4.1 Observation of Fabry-Perot Resonance from Bare GaAs Nanosheets at 4K
and Room Temperature
49
4.2 Effects of Passivation on Optical Responses from GaAs Nanosheets 60

4

Chapter 5: Optical Enhancement of GaAs Nanosheets after AlGaAs Passivation 62
5.1 Passivation Methodology and Measurement Setups 63
5.2 Steady-state and Time-resolved PL 67

Chapter 6: Electronic Characterization of GaAs Nanostructures 75
6.1 Challenges of Making Ohmic Contacts to GaAs Nanowires 75
6.2 I-V Characterization of GaAs Nanosheets 80
6.3 Photocurrent and Photo-spectroscopy Measurement of GaAs Nanosheets 87
6.4 Electron Beam Induced Current (EBIC) Measurement of GaAs Nanosheets 90

Chapter 7: Enhancement of III-V Semiconductor for Photo-Catalysis 94
7.1 Introduction 95
7.2 Photocatalysis Enhancement in III-V Semiconductor after TiO
2

Passivaiton
96
7.3 Comparison and Conclusion 100

Bibliography 106

5

List of Figures

1.1.1 GaAs nanowires grown using Selected Area Growth,
Metal Organic Chemical Vapor Deposition (SAG-MOCVD)
11
1.1.2.1 GaAs nanosheet grown using SAG-MOCVD 13
1.1.2.2 MoS
2
monolayer flake on Si/SiO
2
substrate 13
1.2.1 Photocurrent spectroscopy setup 15
1.2.2 Schematic representation of TRPL setup 17
1.3.1 Electron beam induced current mechanism 19
1.3.2 Photocurrent spectroscopy setup (room temperature) 19
1.3.3 X-ray Photoelectron Spectroscopy 21
1.4.2 CVD Graphene growth recipe 23
2.1.1 SEM images GaAs nanowires (a) as grown and (b) deposited on a Si/SiO
2

Substrate
28
2.1.2 Photoluminescence spectra of individual GaAs nanowires taken at
(a) 4K and (b) room temperature under 532nm excitation (3.5 µW), deposited
on Si/SiO
2
and Au substrates
29
2.1.3 Schematic diagrams illustrating the two simulated geometries:
(a) GaAs nanowire on top of Si/SiO
2
and
(b) GaAs nanowire on top of a 200 nm Au film
(c,d)Cross-sectional electric field intensity distributions for the respective
substrates.
30
2.1.4 Schematic diagrams of the two systems simulated for PL emission: (a) GaAs
nanowire on top of 300 nm thick SiO
2
and (b) GaAs nanowire on top of 200
nm of Au with 300 nm of SiO
2
. Dashed lines indicate flux monitors.
Incoherent summation results for (c) TDP and (d) EP of each competing
topology.
31
2.2.1 STM image of a GaAs surface grown by MBE. The density of surface defects
is ~ 5X10
19
cm
-3
.
33
2.2.2 1-Ethyl-3-methylimidazolium bis(trifluoromethylsulfonyl)imide 33
2.2.3 Room temperature PL spectra of (a) AlGaAs passivated and (b) EMIM TFSI
ionic liquid passivated GaAs nanowire at 8.5 µW power intensity.
34
2.2.4 PL images of GaAs nanowires: (a) Si substrate w/t passivation, (b) Si
substrate
w/ ionic liquid and c) Si substrate w/ core-shell passivation
35
3.1.1 SEM image of n-type doped GaAs nanowires (a) arrays of nanowires, (b)
individual nanowire on Si/SiO
2
substrate.
42
3.1.2 Photoluminescence response of n-type doped GaAs nanowires excited by a
CW 532 nm laser at 4K, (a) lightly doped, (b) moderately doped and (c)
heavily doped.
43
3.1.3 PL of n-type doped GaAs nanowires excited at 4K and 77K, (a) lightly doped
GaAs NW at 77K, (b) moderately doped GaAs NW at 77K and (c) heavily
doped GaAs NW at 77K. (d) lightly doped GaAs NW at 4K, (e) moderately
44
6

doped GaAs NW at 4K and (g) heavily doped GaAs NW at 4K.
3.1.4 Schematic for indirect band-acceptor transition 44
3.1.5 (a) Carrier concentrations estimated from photoluminescence FWHM, (b)
effective band gap estimated from both experimental photoluminescence
response and model.
45
3.2.1 SEM micrograph of the PIN GaAs nanowire 46
3.2.2 The PL spectra of the PIN GaAs nanowire as we traverse from one end of the
wire to the other.
47
3.2.3 The PL peak position corresponding to each position point on the PIN GaAs
nanowire
47
4.1.1 (a) SEM image of a GaAs nanosheet and (b) Photoluminescence spectrum
taken at 77K
56
4.1.2 Photoluminescence spectra from GaAs nanosheets on Si/SiO
2
and Si/SiO
2
/Au 57
4.1.3 Schematic diagrams illustrating the two simulated pump geometries: GaAs
nanosheet on top of (a) 300 nm of SiO
2
and a semi-infinite Si substrate and
(b) 200 nm of Au on top of 300 nm of SiO
2
and a semi-infinite Si substrate.
The nanosheet has a maximum side length of 6 m, obtuse angle of 119.5°,
minimum angle of 21.8°, and a thickness of 200 nm. The green spot in each
diagram represents a 532 nm optical pump with a beam waist of 1.5 m
focused onto the nanosheet
57
4.1.4 (a) Incoherently summed extracted power for the topologies shown in Figure
3. (b) Electric-field patterns for resonances in each material system; the
wavelength of each resonance is indicated by the dashed lines in (a). The
solid white lines indicate the triangular outline of the nanosheet, whereas the
green dots indicate the position of the dipole emitters.
58
4.1.5 (a) Incoherently summed extracted power (EP) results for three different
cross-sectional emitter positions. (d-f) Electric-field patterns for resonances in
each material system; the wavelength of each resonance is indicated by the
dashed lines in (a-c), respectively. The solid white lines indicate the triangular
outline of the nanosheet whereas the green dots indicate the position of the
dipole emitters. Complicated interference patterns are present in each figure
and depend sensitively on emitter position. The broadband extracted power
(EP) enhancement for each dipole position in (a-c) is 1.0X, 1.0X, and 1.1X,
respectively.
59
4.2.1 PL spectra of GaAs nanosheet after application of ionic liquid and bias
voltage
60
4.2.2 PL spectra of non-passivated and passivated n-type GaAs nanowires with
ionic liquid at 300K
61
5.1.1 SEM images of non-passivated GaAs nanosheet: (a) array of as-grown
nanosheets, (b) individual nanosheet deposited on Si/SiO
2
substrate
66
5.2.1 Photoluminescence spectra of an individual GaAs nanosheet with and without
AlGaAs passivation. Inset graph: Raman spectra of passivated GaAs
nanosheet.
70
5.2.2 Spatial maps of the PL intensity (a) PL spatial map for an as-grown and thus
unpassivated GaAs nanosheet, (b) Optical image of the nanosheet passivated
71
7

with a 20 nm AlGaAs layer, where the arrows show the direction of the line
scans used for spatially-resolved micro-Raman measurements shown later,(c)
PL spatial map for the passivated GaAs/AlGaAs nanosheet.
5.2.3 Selected Raman spectrum near the center of the GaAs/AlGaAs sample. GaAs
TO (267cm
-1
) and AlAs-like TO (360cm
-1
) and LO (388cm
-1
) phonon
structures are clearly identified, along with a weak GaAs LO response
(~290cm
-1
).
The weak structure at the lower energy side of the GaAs TO is presumably
GaAs-like TO response from AlGaAs.
72
5.2.4 Longitudinal and transverse Raman scattering scan analysis. The Al-
concentration is calculated based on AlAs-like LO phonon energy due to its
strong Al-concentration dependent nature. From both of the scans, we
observe a consistent Al-concentration 55% across the sample with a
deviation of ±1%.
73
5.2.5 Time resolved PL data from passivated and unpassivated GaAs nanosheets at
300 K. The non-passivated GaAs nanosheet PL lifetime is less than the
system response (30 ps). The passivated nanosheet PL exhibits a lifetime of
300 ps.
74
6.1.1 I-V characteristics of n-type GaAs wafer using, recipe A, the resistance is
1.75Ω
77
6.1.2 I-V characteristics of n-type GaAs wafer using, recipe B, the resistance is
1.42 Ω
78
6.1.3 I-V characteristics of n-type GaAs wafer using, recipe C, the resistance is
1.67 Ω
78
6.1.4 I-V characteristics of n-type GaAs nanowires (data retrieved for three
nanowires from the same batch)
79
6.2.1 SEM micrograph of nanosheet with possible electric contacts 81
6.2.2 SEM micrograph of nanosheet and corresponding I-V characteristics 82
6.2.3 SEM micrograph of GaAs nanosheet after RTA annealing at 375˚C , 380˚C,
385˚C, 390˚C and 395˚C
83
6.2.4 I-V characteristics of GaAs nanosheet after RTA annealing at 375˚C, 380˚C,
385˚C, 390˚C and 395˚C
83
6.2.5 RTA temperature dependence of contact resistivity for GaAs nanosheet 84
6.2.6 SEM micrograph of GaAs nanosheet after RTA annealing at 375˚C, 380˚C,
385˚C, 390˚C and 395˚C
85
6.2.7 SEM micrograph of GaAs nanosheet after RTA annealing at 375˚C, 380˚C,
385˚C, 390˚C and 395˚C
85
6.2.8 RTA temperature dependence of contact resistivity for GaAs nanosheet 86
6.3.1 Photocurrent measured for GaAs nanosheets 88
6.3.2 Photocurrent spectroscopy measured for GaAs nanosheets 89
6.4.1 Schematic EBIC setup located at CEMMA, USC 91
6.4.2 EBIC measurement on passivated vs. non-passsivated GaAs nanowire: (a)
EBIC micrograph, (b) SEM micrograph (c) extracted and exponentially fitted
signal, (d) EBIC concept
92
6.4.3 EBIC measurement of non-passsivated GaAs nanosheet: (a) SEM 93
8

micrograph of GaAs nanosheet, (b) EBIC micrograph of GaAs nanosheet, (c)
EBIC micrograph of GaAs nanosheet with opposite polarity
7.2.1 Photo I-V measurement setup, 3-terminal potentiostat is used with Pt counter
electrodes and Ag/AgCl reference electrode. The illumination source is a 532
nm CW laser.
97
7.2.2 Photocurrent diagrams of III-V with and without TiO
2
passivation: (a) GaAs,
(b) GaP and (c) InP
102
7.2.3 Photoluminescence spectra of III-V with and without TiO
2
passivation: (a)
GaAs, (b) InP
103
7.2.4 XPS measurement results of III-V with and without TiO
2
passivation: (a)
GaAs, (b)GaP, (c) InP
105

List of Tables

4.1.1 Comparison between the properties of the optical cavity in the nanowires vs.
nanosheets
52
6.1.1 Ohmic cont Ohmic contact recipe for n-type GaAs 76
6.2.1 Resistance vs. RTA temperature (nanosheet sample1) 84
6.2.2 Resistance vs. RTA temperature (nanosheet sample2) 86
7.2.1 Shift of overpotential required for TiO
2
-passivated samples with respect to
their bare counterparts
99

9

Abstract

In this thesis, we study the optical and electrical properties of one-dimensional and two-
dimensional nanostructures used for different energy applications. These applications would
include solar cells, LEDs, nanolasers and water splitting cells. Different characterization
methodologies are used here; inducing but not limited to static and time-resolved
photoluminescence at room temperature and low temperatures, Raman spectroscopy,
photocurrent spectroscopy, scanning electron microscopy, electron beam induced current, etc.
Our goal is to not only identify the optoelectronic properties of low-dimensional nanostructures
but also modify and enhance such properties. Therefore, in this thesis, we present new methods
to enhance the optical properties of such materials and characterization of such enhancements.

10

Chapter 1

Introduction

In this chapter some of the main optical and electrical measurement methods which are used
in this research will be described in detail. The nature and importance of optoelectronic
characterizations of one and two-dimensional nanostructures and their applications will be
presented. Finally, different synthesis and fabrication methods used throughout this work are
provided.

1.1.1 One-dimensional Nanostructures (Nanowires)
One-dimensional (1D) materials including nanowires, quantumwires and nanotubes have
received a lot of attention in the recent years. Both experimental and theoretical studies show that
1D material demonstrates outstanding qualities in comparison to their bulk form. These
improvements may be but not limited to higher luminescent efficiency, faster electrical transport
response, compact structure and high surface to volume ratio. 1D materials are the best
candidates for several electronic and optical applications such as nanowire lasers, nanowire
transistors and solar cells[1-4] [1, 3-5]. The application of nanowires especially in the energy-
harvesting field has stimulated several research groups [6-10]. It should be mentioned that
semiconductor nanowires with high surface-to-volume ratios poses several advantageous and
outstanding optoelectronic properties, such as enhanced light absorption [11, 12], electrostatic
gate tenability [2], and enhanced carrier collection [13], that are beneficial for applications
11

including: detectors [14, 15], sensors [16, 17], energy storage [18], and lasers [19]. Among
different forms of semiconductor nanowires, III-V nanowire and especially GaAs nanowires
have been studied for photovoltaic applications. Direct band gap of GaAs nanowires make them
quite attractive for optical applications. However, minority carrier lifetime (τ) and mobility (µ),
which are two of the most important characteristics of solar cells, are poor in GaAs nanowires in
comparison to planer GaAs. Nevertheless, due to higher surface to volume ratio in nanowires,
better optical efficiency is observed in these structures. In this work, we focus on GaAs
nanowires (with or without passivation). There are different methodologies to grow GaAs
nanowires among which Vapor-liquid-solid (VLS) method[20, 21], Metal Organic Vapor Phase
Epitaxy (MOVPE)[22, 23] and Selected Area Growth, Metal Organic Chemical Vapor
Deposition (SAG-MOCVD)[12, 24]have received more attentions. Here, Selected Area Growth,
Metal Organic Chemical Vapor Deposition (SAG-MOCVD) is used; the details of the growth
will be discussed in the future chapters. An image of the GaAs nanowire grown by SAG-
MOCVD is shown in Figure.1.1.1. Higher crystalline quality, more abrupt junction interfaces
and lack of metal catalyst are the main advantages of this method.

Figure 1.1.1 GaAs Nanowire grown using Selected Area Growth, Metal Organic Chemical Vapor
Deposition, GaAs nanowire on GaAs (111)B substrate.
12

1.1.2 Two-dimensional Nanostructures (Nanosheets)
After the discovery of Graphene [25] a new path to 2D materials or single atomic layer of
materials was introduced. There have been so many different systems introduced under this
family including: Graphene, MoS
2
, WSe
2
, etc. Under a more general concept 2D materials are
referred to any structure with one dimension confinement. Here, we both focus on single (or few)
atomic-layer materials and on GaAs nanosheets. In Figure 1.1.2.1 an image of GaAs nanosheet is
depicted. The proposed GaAs nanosheets are grown using SAG-MOCVD. As for GaAs
nanowires, the main application considered for nanosheet are solar cells and nanolaser (optical
cavities). In the later chapters we will talk about Van de Waals materials including MoS
2
, WSe
2
,
BN, Black phosphorous, etc. These 2D materials are retrieved through exfoliation of their bulk
form. In Figure 1.1.2.2 optical images of MoS
2
and WSe
2
samples are shown.

Atomic Layer Deposition (ALD) can be used to obtain thin layers of different materials
including TiO
2
, HfO, etc. During ALD different gas phase precursors (TMA, H
2
O, TiCl
4
, etc.)
are exposed; hence react with the surface of the substrate in a sequential form, which leads to
formation of a thin film.

13

Figure 1.1.2.1 GaAs nanosheet grown using SAG-MOCVD

Figure 1.1.2.2 MoS
2
monolayer flake on Si/SiO
2
substrate

1.2 Optical Measurement Methods
A large portion of the characterization performed here is focused on optical measurements.
Studying the optical response of different nanostructures allows us to validate the feasibility and
efficiency of using such systems for optical applications including energy harvesting and lasers.
14

Among the optical characterizations performed we focus on Photoluminescence
(temperature dependent, steady state or time-resolved), Raman Spectroscopy and photocurrent
measurement. Below we describe each of these methods.

Photoluminescence (Steady State)
Photoluminescent (PL) is referred to a photo emission phenomenon from a matter. This
photo emission process occurs due to excitation of the sample by photons. During PL
measurement the semiconductor is excited by a light source (usually a continuous wavelength
laser), the emitted photons from the light source excite and separate the electrons and holes;
hence, electrons which gain energy from the photon move to the conduction band and holes stay
in valance band. The electron and hole pair has the tendency to go back to the relaxation state
and finally recombine and emit photons. The emitted photons can be measured and used as an
indication of different energy states and the band gap (E
g
) of the semiconductor.
PL measurement is an important characterization method to investigate the band gap of
different materials. This method is particularly useful for measuring the purity, defect points and
crystalline quality of the material (especially GaAs and InP).
PL measurements can be performed at different temperatures with different excitation powers.
Apart from PL measurement, Time-resolved Photoluminescent (TRPL) is another branch of PL
characterization which is used to measure the carrier lifetime (τ) in semiconductors. In this
method, the sample is excited with a light pulse (pulse laser) and the decay in the PL response is
measured as the function of decay time (in units of femtoseconds or picoseconds), which reveals
the lifetime of the carriers in the sample. Knowing the lifetime of carriers one can estimate the
15

mobility (using the formula bellow) of the carries which is another important factor both for
solar cell applications and electronic circuit applications.

In the above equation µ is the carrier mobility, q electron unit charge, m
*
is the effective mass of
the electrons and τ is the lifetime.
In Figure 1.2.1 a simple PL setup is depicted. The main components of such setup are
light source (pulse or continuous laser), detector (Si, Ge or InGaAs detector based on the band
gap of the material) and a controllable stage to hold and move the sample. It should be indicated
that temperature-dependent PL requires a more complex setup (cryostat to hold the sample at
specific temperature, pump to vacuum the sample environment and liquid N
2
or He to cool down
the sample).

Figure 1.2.1 Photoluminescent setup (for low temperature measurements).

Photoluminescence (Time-Resolved)
Another important and critical characterization is the identification of carrier’s lifetime, τ,
(how long it takes for separated electron-holes to recombine). This can be done, using a method
called Time-Resolved Photoluminescent (TRPL). In the previous chapters, we talked about PL;
16

in the PL setup the sample is excited using continues-wave laser (CW), and the emitted light
from the sample is detected and graphed versus the energy or the wavelength. In TRPL, the
dynamics of the photoexcited carriers are studied, the method which is used here is excite-and-
probe. For this measurement a pulse laser is required. Pulse laser is cable of generating lasers
pulses with small bandwidths in the range of picoseconds or even femtoseconds. The setup also
needs a fast detector to determine the emitted light from the sample as the function of time. The
pulse laser can be used for both exciting the sample and probing. The pulse laser excites the
sample and the signal from the sample is focused on a crystal, on the other hand, a part of the
laser’s incident beam is also focused (after passing a variable delay generator) on the same
crystal. The converted beam is sent to spectrometer. Based on the required sensitivity and
resolution one can use Time Correlated Single Photon Counting (TCSPC) (picoseconds or
slower), Streak Camera (picoseconds or slower) or Intensified CCD camera (picoseconds or
slower). As for our GaAs nanowires and nanosheets we need a resolution of picoseconds. The
longer it takes the emission signal to die the higher is the quality of the sample. This shows that
defect points, crystalline lack of quality or any other collision point is less in the sample. Similar
to PL this measurement can be done both at room temperatures and low temperatures. A part of
the proposal is to perform TRPL measurements for our passivated and non-passivated samples
(AlGaAs or ionic liquid passivation), not only to find the carrier’s lifetime of our samples but
also to validate the affectivity of our passivation method and quality.
17

Figure 1.2.2 Schematic representation of TRPL setup (Courtesy of University of Cincinnati).

Raman Spectroscopy:
Raman spectroscopy is used to study the vibrational, rotational or other low-frequency
modes of the materials. In this method the material is excited by a laser light, the laser will
interact with the molecular vibration or phonons; this interaction which is known as Raman
scattering leads to energy shift (up or down). The shift in energy gives information about the
vibrational modes in the system.

Photocurrent Measurement
Photocurrent measurement is used to study the current generated as the result of
photoexcitation of p-n junction or a Schottky barrier. The sample should have metal contacts
18

(Ohmic or Schottky) so we can apply bias voltages. In this setup a light source (laser) is used to
excite the sample at the metal-semiconductor interface. Meanwhile voltage (varying) is applied
to the sample and the current generated through the sample is measured. The final result is in a
form of a Current-Voltage diagram.

1.3 Electrical Measurements
Electrical measurements are another set of important characterizations which give more
insight into the properties of nanostructures. Electrical measurements cover a larger variety of
methods including: Current-Voltage relations (I-V), Capacitance-Voltage relation (C-V),
Electron beam Induced Current (EBIC), Photocurrent (mentioned in the previous section) and
Photocurrent Spectroscopy. Both I-V and C-V measurements are applicable methods for
calculating the dopant concentration. Measuring the dopant concentration is quite an important
factor in order to investigate the growth mechanism and properties of a semiconductor. EBIC
measurement is a way to investigate the diffusion length of minority carriers in a semiconductor.
Here, we shortly describe each method.

Electron Beam Induced Current (EBIC)
This technique is based on the presence of the local electric field in the sample. Such a
field may be due to a p-n junction, Schottky barrier or even an electrical field induced by the
crystal defects[26, 27]. Having this field, the electron beam (usually in the Scanning Electron
Microscopy setup) separate electron and hole pairs. This separation avoids the recombination of
electrons and holes and forms a current. The magnitude of such a current is quite small and one
19

needs current amplifiers to amplify such currents. This method allows us to measure the minority
carrier lifetime and surface recombination velocities (Figure 1.3.1).

Figure 1.3.1 Electron beam Induced Current mechanism.

Photocurrent Spectroscopy
In this method, the sample with a p-n junction or a Schottky barrier is exited by a pulse
laser, the wavelength of the light source changes (here we use a Fianium, ultra fast fiber laser,
SC450) and hence a current is induced across the junction. The variation of such current versus
the excitation wavelength is an indication of band gap and sub band gap state. In this experiment
an additional bias voltage may be applied. In Figure 1.3.2 the photocurrent spectroscopy setup is
demonstrated.

Figure 1.3.2 Photocurrent Spectroscopy setup (room temperature).
p n
l
min_diff
e
+
e
-
ħ w ε
metal
e-beam
e
-
e
+
n
l
min_diff
20

X-ray Photoelectron Spectroscopy
X-ray photoelectron spectroscopy (XPS) is used to study the elemental composition on the
very top surface of materials. This measurement provides information about the chemical state,
oxidation state and elements’ count. Here, an X-ray irradiates the sample and both the number of
the electrons leaving the surface and the kinetic energies are measured. This measurement is
accurate for the 1-10nm thickness of the surface.

21

Figure 1.3.3 X-ray Photoelectron Spectroscopy setup (courtesy of California Technology Institute)

1.4 Synthesis Methods
1.4.1 Selected Area Growth Metal Organic Chemical Vapor Deposition
Method
In the work presented here, GaAs nanostructures (nanowires or nanosheets) are synthesized
by metal organic chemical vapor deposition (MOCVD) with selective area growth (SAG).
Trimethylgallium (TMGa), trimethylaluminum (TMAl), and arsine are used as precursors for Ga,
Al, and As deposition, respectively, and disiline is used as a precursor for Si doping. A thermally
grown silicon oxide layer is used as a mask for the SAG growth. An array of holes (1 mm  1
mm) with 600 nm pitch are patterned with silicon nitride using electron beam lithography (EBL)
and wet chemical etching. Nanowires are grown in a Thomas Swan MOCVD system at a
pressure of 0.1 atm. The total flow rate of the carrier gas is 7 SLM, and the partial pressure of
TMG, arsine, and disilane are 7.56x10
-7
atm, 2.14x10
-4
atm, and 1.43x10
-8
atm during the growth
of the nanowire core. During the growth of AlGaAs passivation layer, arsine partial pressure was
increased to 1.07x10
-3
atm, and the partial pressures of TMG and TMAl were kept at 3.78x10
-7

atm and 1.51x10
-6
atm. The growth temperature was fixed at 760˚C degrees for both the
nanowire core and passivation layer. These nanowires are grown vertically along the (111)B
22

direction, sonicated in isopropyl alcohol, and then deposited onto various substrates including Si
substrates with 300 nm of SiO
2
and Si/SiO
2
substrates with 200 nm of Au.

1.4.2 Chemical Vapor Deposition Method for 2D Materials
In this method different precursors are introduced to the substrate in vacuum or specific
environment. During the growth, temperature and partial pressure of the chamber changes and
governs the direction and specification of the growth. This method has been widely used for
MoS
2
and graphene. This method provides easily scalable and controllable growth of 2D
materials[28, 29].
The CVD growth mechanism for graphene is quite well-known. The growth occurs at low
pressures and under methane flow. The literature shows that graphene growth on Cu-substrate
provides better control on the morphology and quality of the graphene. Therefore, in this work
we use CVD-growth mechanism for graphane on Cu substrate. The recipe used for CVD growth
is depicted in Figure 1.4.2. Here, the Cu film (99.99%, Sigma) is cleaned with IPA and DI water;
then situated in the 2.5” furnace tube. During the first 40 min, H
2
gas is flowed at 7 sccm and the
pressure is kept at 40 mTorre while the temperature is increased from 25˚C to 1000˚ C. The
temperature is kept constant for 20 min and then the H
2
and CH
4
are flowed at 20 and 7 sccm.
The tube pressure is increased to 150 mTorre. The growth occurs during this window for 30 min.
Then the temperature is decreased to room temperature; when temperature reaches 400˚C the
CH
4
is turned off and at 100˚C the H
2
is turned off.

23

Figure 1.4.2 CVD graphene growth recipe.

1.4.3 Exfoliation Method for 2D Materials
Mechanical exfoliation (scotch-tape) of 2D materials (MoS
2
, WSe
2
, BN, etc.) was the first
method introduced in this field. This method provides high quality and pristine single flakes,
which are perfectly suitable for fundamental studies and fabrication of single devices; however,
this method is not scalable and the size and thickness control of flakes is hard [30-32].

24

Chapter 2

Observation of Fabry-Perot Resonance from GaAs
Nanowires

In this chapter we will describe the procedures and outcomes of optical measurements
performed on GaAs nanowires. Our work can be categorized into photoluminescent
measurement of GaAs nanowires, formation of Fabry-Perot microcavity in GaAs nanowires and
study of the passivation schemes to improve the optical activity of GaAs nanostructures and the
efficiency of each method.

2.1 Observation of Fabry-Perot Resonance from Bare GaAs Nanowires at 4K
and Room Temperature
As it was mentioned in the previous chapters, semiconductor nanowires with high
surface-to-volume ratios possess several advantageous optoelectronic properties; such as
enhanced light absorption, electrostatic gate tunability and enhanced carrier collection that are
beneficial for applications including: detectors, sensors, energy storage and lasers. Nanowire
lasers based on binary semiconductor materials are of particular interest because of their wide
range of band gaps. GaN, GaAs, ZnO, and CdS are a few of the materials widely studied as
nanolasers. The first electrically pumped nanowire laser was presented by Duan et al., where the
electroluminescence spectra and lasing of an n-type CdS nanowire were observed under an
applied current of 200 µA [5]. ZnO electroluminescence and photoluminescence were reported at
room temperature by Chu et al. [33]. Fukui and coworkers reported the observation of Fabry-
25

Perot resonances and lasing of InGaAs/GaAs core-shell nanopillars [34, 35]. One of the main
challenges associated with GaAs nanowire-based optoelectronic and photonic devices is the non-
radiative recombination associated with surface states, which is particularly pronounced in GaAs
nanowires with high surface-to-volume ratios[36]. While many studies discuss the effects of
surface passivation of bulk GaAs[37-39] and GaAs nanowires [36], none have investigated these
effects on the Fabry-Perot (FP) resonance.
Here, as shown in Figure 2.1.1.a , the standing GaAs nanowires are sonicated in
isopropyl alcohol, and then deposited onto various substrates including Si substrates with 300 nm
of SiO
2
and Si/SiO
2
substrates with 200 nm of Au (Figure 2.1.1.b). Micro-Raman and
photoluminescence spectroscopy were collected using a 40X objective lens, an 1800 l/mm
grating, and a silicon CCD detector to detect the PL in the range of 750 nm to 950 nm. Moderate
to low power excitation (10
2
-10
5
mW/cm
2
) using a 532 nm laser was used to avoid optical
heating.
Figure 2.1.2.a shows the PL spectra of individual GaAs nanowires deposited on Si/SiO
2

and Au substrates. In Figure 2.1.2.b, spectra from both substrates were taken at 4K and exhibit
strong Fabry-Perot peaks. The spectrum taken on the Au film is significantly enhanced with
respect to that of the Si/SiO
2
. The integrated areal intensity of the dominant PL peak is 5X larger
for the nanowire on the Au film, whereas the relative strength of the FP peaks are also enhanced
by a factor of 373%. Figure 2.3.2.b shows the PL spectra of nanowires at room temperature.
Here, the Au substrate results in strong FP resonance peaks, which are completely absent from
the nanowire deposited on the Si/SiO
2
substrate. As observed in the 4K dataset, the integrated
intensity of the main PL peak increases by a factor of 5X due to the presence of the Au substrate.
The increased PL intensity is attributed to the local field enhancement from the underlying Au
26

substrate. In the absence of the Au film, the high dielectric constant of the silicon dioxide
substrate causes light to be preferentially emitted downwards, thus reducing the observed PL
intensity. The Q-factor of the FP peaks of the nanowires deposited on the Au film is 140, and
181 for nanowires on Si/SiO
2
. This decrease in Q-factor may be due to the lossy nature of Au
substrate. The mode spacing of the FP peaks can be calculated by the relation:

For the nanowire deposited in Si/SiO
2
, we observe a mode spacing of 8-10 nm (at 4K) for
this 270nm diameter and 7µm long nanowire, which agrees with the calculated mode spacing of
11.5nm, assuming n=3.655, and dn/dx= -0.822/µm. A similar agreement was obtained for the
nanowire on Au substrate. In order to understand the underlying mechanism of enhancement,
finite-difference time-domain (FDTD) simulations are carried out on a GaAs nanowire on top of
both Si/SiO
2
and Au substrates, as illustrated in Figures 2.1.3.a and 2.1.3.b. Figure 2.1.3.a shows
a GaAs nanowire on top of 300 nm of SiO
2
and a semi-infinite layer of Si, whereas the sample in
Figure 2.1.3.b has an additional 200 nm layer of Au below the nanowire. The optical pump is
simulated using a Gaussian beam (1.5μm in diameter) focused onto the center of the GaAs
nanowire. The electric-field intensity distributions are plotted in Figures 2.1.3.c and 2.1.3.d,
corresponding to the structures shown in 2.1.3.a and 2.1.3.b, respectively. As evident from the
Figures, the gold substrate produces larger electric-field intensity within the nanowire.
We can calculate the relative enhancement in the optical absorption of the nanowire due
to the presence of the metallic substrate by integrating the electric-field intensity inside the
hexagonal nanowire cross-section of each system and taking the ratio:
27

For an incident wavelength of 532 nm, the total electric field enhancement is 2.0, indicating the
gold substrate concentrates 2.0 times more total electric field intensity within the nanowire than
the alternative structure. However, the absorption enhancement varies as a function of nanowire
diameter.
PL emission is modeled in the GaAs nanowire using electric dipoles, as shown in
Figures 2.1.4.a and 2.1.4.b. The bandwidth of each dipole is taken as 200 nm, to resemble the PL
spectra collected experimentally. The extracted power (EP) was calculated at the flux planes
indicated by the upper dashed lines in Figures 2.1.4.a and 2.1.4.b. The dashed box surrounding
each dipole is used to calculate the total dipole power (TDP). The power emitted by each dipole
is affected by the photonic density of states and therefore will vary depending on its position
within the nanowire. Figures 2.1.4.c and 2.1.4.d show the TDP and EP results for both structures.
Fabry-Perot fringes are clearly visible in the simulated spectra. The average peak-to-peak
spacing and quality factors of the resonances closely resemble those found experimentally. By
taking the integral of both EP spectrums for the 260nm diameter nanowire, we find that the gold
substrate produces 0.74 times as much EP as the alternative geometry. This slight decrease in
emitted power is attributed to losses in the Au film. We, therefore, attribute the enhancements
observed experimentally to increased absorption as described above in Figure 2.1.3.
In previous works, local field enhancement of Raman spectroscopy has been observed on
nanowires after Au nanoparticle decoration, as reported by Cao et al. In our work, we also
explored Au nanoparticles deposited on top of GaAs nanowires, which resulted in an increased
28

PL intensity; however the FP resonance was lost even at low temperatures. Here, the Au
nanoparticles cause significant out-coupling of light from the nanowire, thereby, reducing the Q-
factor of the FP cavity.

Figure 2.1.1 SEM images GaAs nanowires (a) as grown and (b) deposited on a Si/SiO
2

substrate[40]
(a)
(b)
2 µm
Laser
spot
1 µm
29

750 800 850 900 950
0.0
2.0x10
4
4.0x10
4
6.0x10
4
8.0x10
4
1.0x10
5
1.2x10
5
1.4x10
5
PL Intensity (Counts)
Wavelength(nm)
4K

750 800 850 900 950
0.0
8.0x10
3
1.6x10
4
2.4x10
4
3.2x10
4
4.0x10
4
300K
PL Intensity (Counts)
Wavelength(nm)

Figure 2.1.2 Photoluminescence spectra of individual GaAs nanowires taken at (a) 4K and (b)
room temperature under 532nm excitation (3.5 µW), deposited on Si/SiO
2
and Au substrates.

(a)
(b)
Si/SiO
2

Au
Au
Si/SiO
2

30

Figure 2.1.3 Schematic diagrams illustrating the two simulated geometries: (a) GaAs nanowire
on top of Si/SiO
2
and (b) GaAs nanowire on top of a 200 nm Au film. (c,d) Cross-sectional
electric field intensity distributions for the respective substrates. The white outline in each plot
indicates the hexagonal boundary of the nanowire.

(a)
(b)
31

Figure 2.1.4 Schematic diagrams of the two systems simulated for PL emission: (a) GaAs
nanowire on top of 300 nm thick SiO
2
and (b) GaAs nanowire on top of 200 nm of Au with 300
nm of SiO
2
. Dashed lines indicate flux monitors. Incoherent summation results for (c) TDP and
(d) EP of each competing topology.

32

2.2 Effects of Passivation on Optical Responses and Fabry-Perot Resonance
GaAs nanowires are well-known for the large number of surface states that they possess.
These surface states behave as an insulating layer; the separated electron and holes during any
optical excitation would recombine with these states and lead to degraded optical activity of
these nanostructures (Figure 2.2.1) [41]. The formation of localized surface states may lead to
high surface recombination velocity, short minority carrier diffusion length and short minority
carrier lifetime. A passivation layer (AlGaAs, ionic liquid, etc.) will behave as a barrier and
avoid the recombination of the photoexcited carries to recombine with the surface state and
hence improves the optical activity. In this chapter we describe two forms of passivations:
AlGaAs passivation and the ionic liquid passivation. AlGaAs passivation can be formed as an
additional layer over the GaAs nanowire during the growth process. This will form a core-shell
structure where the GaAs nanowire is the core and a shell of AlGaAs (about 10nm) is
surrounding it. The growth mechanism was described in the first chapter.
Ionic liquid is referred to salts in liquid forms at room temperature, a solution consisting
of ions and short-lived ion pairs. There are different classes of ionic liquids and here we mainly
focus on ethyl-methylimidazolium bis(trifluoromethyl-sulfonyl) imide (EMIM TFSI) (Figure
2.2.2). After transferring the GaAs nanowire on substrate, the nanowire is covered with ionic
liquid, which behaves as a passivation layer. We also investigate the effects of surface
passivation on the PL intensity and observation of FP resonance at room temperature. Here, we
consider two forms of passivation: AlGaAs passivation, which leads to a core-shell structure, and
ionic liquid passivation. In the former case, after the GaAs nanowire growth, a layer of AlGaAs
is deposited as a passivation layer. The PL spectra of AlGaAs-passivated nanowires at room
33

temperature are shown in Figure 2.2.3.a exhibiting the FP resonance peaks. In the latter case, we
deposit a few drops of ionic liquid ethyl-methylimidazolium bis(trifluoromethyl-sulfonyl) imide
(EMIM TFSI)[42] to passivate the surface states in the nanowire. This form of passivation leads
to a 12X enhancement in the PL intensity without applying any potential to the nanowire and
enables us to observe the FP resonance at room temperature (Figure 2.2.3.b). However, as it can
be observed the quality factor is smaller in comparison to that of AlGaAs passivated case.

Figure 2.2.1 STM image of a GaAs surface grown by MBE. The density of surface defects is ~
5X10
19
cm
-3
.

Figure 2.2.2 1-Ethyl-3-methylimidazolium bis(trifluoromethylsulfonyl)imide

[EMIM] [TFSI] ionic
liquid

-
-
-
-
-
-
-
-
-

-
-
-
-
Localized surface
states

Conducting
channel
Depletion region
34

750 780 810 840 870 900 930
0.0
5.0x10
4
1.0x10
5
1.5x10
5
2.0x10
5
2.5x10
5
PL Intensity (Counts)
Wavelength(nm)
3.5 uW
8.5 uW
32.5 uW
90 uW

300K
AlGaAs Passivation
(a)

750 800 850 900 950
0.0
2.0x10
3
4.0x10
3
6.0x10
3
8.0x10
3
1.0x10
4
1.2x10
4
1.4x10
4
1.6x10
4
PL Intensity (Counts)
Wavelength(nm)
w/o ionic liquid
w/ ionic liquid
300K

(b) Ionic liquid passivation

Figure 2.2.3 Room temperature PL spectra of (a) AlGaAs passivated and (b) EMIM TFSI ionic
liquid passivated GaAs nanowire at 8.5 µW power intensity.

35

Figure 2.2.4 PL images of GaAs nanowires: a) Si substrate w/t passivation, b) Si substrate w/
ionic liquid and c) Si substrate w/ core-shell passivation

0.05% power
100X
10% power
40X
10% power
100X
(C)
(B)
(A)
36

Chapter 3

Low Temperature Photoluminescence and Doping
Concentration Dependence of n-type GaAs Nanowires

In this chapter, we aim to estimate the carrier concentration of doped GaAs nanowires
using PL measurements at low temperatures (4K and 77K). We also study the PL peak energy
shift as a function of carriers’ concentration and type (n -type or p-type).

3.1 Study on the Concentration of Dopants in n-type GaAs Nanowires at 4K
In this section, the photoluminescence spectra of n-type doped GaAs nanowires, grown
by the metal organic chemical vapor deposition (MOCVD) method are measured at 4K. Our
measurements indicate that increase in carrier concentration leads to increase in the complexity
of the doping mechanism; which, we attribute to the formation of different recombination
centers. At high carrier concentrations, we observe a blueshift of the effective band gap energies
by up to 25meV due to the Burstein-Moss shift. Based on the FWHM of the photoluminescence
peaks, we estimate the carrier concentrations for these nanowires which varies from 6x1017 cm-
3 (lightly doped), to 1.5x1018 cm-3 (moderately doped), to 3.5x1018 cm-3 (heavily doped) as
the partial pressure of the disilane is varied from 0.01 sccm to 1 sccm during the growth process.
We find that the growth temperature variation does not affect the radiative recombination
mechanism; however, it does lead to a slight enhancement in the optical emission intensities.
GaAs nanowires’ direct band gap a nd large surface-to-volume ratio have attracted considerable
attention for their potential use in high efficiency solar cells, energy storage, and lasers[1-4].
37

Measuring the concentrations of the dopants in III-V nanowires and specifically GaAs nanowires
is an important problem since it governs the mobility, minority carrier diffusion length/lifetime,
and conductivity. This fact becomes even more important when one designs PN, PIN, or core-
shell nanowires where concentrations of dopants and doping profiles directly affect the
performance of the junction. While accurate measurement of carrier concentration in the GaAs
nanowires is of great importance, precise measurement techniques of these carrier concentrations
have not been developed yet.

In previous literature, both contact and non-contact methods have been employed for the
estimation of carrier concentration in III-V nanowires. Field effect measurements are one way to
obtain the field effect mobility and carrier concentration[43]. However, this method presents
several limitations and technical challenges, for example formation of Ohmic contacts. Atom
probe microscopy and secondary ion mass spectroscopy are two alternative methods used for
establishing the doping profiles in nanowires; however, they are time consuming, expensive and
destructive to the sample being measured[44]. Capacitance-voltage measurements are another
approach that is commonly used[45, 46]; although, this method also suffers from contact
fabrication difficulties. Strom et al. reported Hall effect measurements on core-shell InP
nanowires[47]. Schaper et al. also used Hall measurements with four contacts to obtain the
carrier concentrations in InAs nanowires, as well as the electron concentrations in the surface
accumulation layer of the InAs nanowires[48]. It should be noted that the Hall method has not
been implemented in GaAs nanowires due to difficulties in forming Ohmic contacts. Ketterer et
al. reported a non-contact method that correlates the carrier concentration of p-type GaAs
nanowires with transmission Raman spectroscopy[49]. This method is based on forward Raman
38

scattering, and it is unlikely that nanowires with large diameters (larger than 250nm) will not
provide enough signal intensity. Chen et al. studied the temperature effects on the
photoluminescence (PL) spectra of the bulk metal organic vapour phase epitaxy grown
(MOVPE) GaAs films[50]. The photoluminscence spectra of MBE grown bulk GaAs was
studied at 4.2K, as reported in reference[51]. Both n-type doped GaAs, Te-doped[52] and p-type
doped[53, 54] were studied. N. Y. Lee et al. performed a room temperature PL analysis and
proposed a model where PL peak energy can be estimated from the carried concentrations[52].
S.I Kim et al. performed similar analysis for C-doped samples at 12K and proposed imperial
relations for hole conceteration estimation based on the FWHM of PL peak[53]. P.W Yu et al.
performed a temperature dependence study on p-type GaAs[54]. Here, we present a systematic
study of the carrier concentration of n-type doped GaAs nanowires grown by MOCVD using
low-temperature photoluminescence spectroscopy. The advantage of this method, especially for
n-type GaAs nanowires with large diameter is that it is contact-less; therefore, there is no need
for complex fabrication steps.

In this work, nanowires are grown using MOCVD with selective area growth (SAG)[55]. These
nanowires are intentionally doped with different dopant concentrations and deposited on Si/SiO
2

substrates. The measurement is performed at 4K using a micro-PL[40]. In the MOCVD process,
trimethylgallium (TMGa), trimethylaluminum (TMAl), and arsine are used as precursors for Ga,
Al, and As deposition, respectively, and disiline is used as a precursor for Si doping. Nanowires
are grown in a Thomas Swan MOCVD system at a pressure of 0.1 atm. The total flow rate of the
carrier gas is 7 SLM, and the partial pressure of TMG, arsine, and disilane are 7.56x10-7atm,
2.14x10-4atm, and 1.43x10-8atm during the growth of the nanowire. The partial pressure of
39

disilane is varied from 0.01 sccm to 1 sccm in order to increase the amount of Si dopant during
different growth runs. Two sets of samples were grown under these conditions at 760˚C and
780˚C. These nanowires are grown vertically along the (111 )B direction, as shown in Figure
3.1.1a, sonicated in isopropyl alcohol, and then deposited onto Si substrates with 300nm of SiO
2

(Figure 3.1.1b). Micro-Raman and photoluminescence spectroscopy were collected using a 40X
objective lens, a 1800 l/mm grating, and a silicon CCD detector. A 532nm continuous-wave laser
is used to excite the samples and spectra are collected over 750 nm to 950 nm wavelength range.
Low to moderate power excitation (102-105 mWcm-2) was used to avoid optical heating. The
measurement was performed at 4K in a continuous flow liquid helium optical cryostat.

Figure 3.1.1 shows SEM micrographs of GaAs nanowires arrays and individual GaAs nanowires
deposited on a Si/SiO2 substrate. These nanowires are doped with Si; where, Si behaves as a
donor. Here, the diameters of the nanowires lie in the range of 250-270 nm. Non-doped GaAs is
known to have p-type behavior due to residual C impurities from MOCVD process, which serve
as the main acceptor species[56]. Radiative recombination in GaAs nanowires can arise from
mechanisms various including: direct band-band transition with momentum conservation,
indirect band-band transition without momentum conservation, and indirect band-acceptor
transition without momentum conservation. The concentration of dopants governs which
mechanism will be dominant in the radiative recombination process. In figure 3.1.2, PL spectra
of three cases of lightly, moderately, and heavily doped GaAs nanowires are shown. Since these
PL spectra are complex, we attempt to assign the peaks to corresponding transitions. In figure
3.1.2a, two major peaks can be observed. The peak at 1.497 eV is due to the exciton,[57-59] and
the 1.49 eV peak is characteristic of a band-acceptor transition due to carbon contamination[60].
40

Wood et al. attributed the 1.454 eV peak to the phonon sideband of the band-to-acceptor
transition at 1.49 eV[60]. Here, we observe the same phenomenon with a phonon energy of 36
meV. The peaks at 1.484 eV and 1.47 eV are also impurity related (free-to-bound transition)[61].
In this case, the blueshift is not observed and the donor-band transition is not dominant. In Figure
3.1.2b, three main peaks from the moderately Si-doped are observed. The peaks at 1.469 eV and
1.49 eV correspond to a band-acceptor transition due to unintentional carbon impurities. The
peak at 1.53 eV indicates the donor-band transition and is blue-shifted by 20 meV. Figure 3.1.2c
shows the PL response of heavily Si-doped GaAs nanowires. Here, a blueshift of 25 meV is
observed (peak at 1.54 eV). The peak at 1.516 eV corresponds to an excitonic transition as
reported by Wood and Ploog[60, 61]. The peaks at 1.496 eV, 1.488 eV, 1.47 eV, and 1.46 eV
originate from acceptor impurity levels. It has been shown that, in the presence of high
concentrations of Si dopants, the behavior of some dopant impurities may switch from donor to
acceptor [62, 63]. The observed peaks between 1.48-1.49 eV are likely due to Si dopant
impurities. As observed in our lightly doped GaAs nanowires, the peak at approximately 1.45 eV
is the phonon sideband of the 1.49 eV peak. The PL spectra of these six samples were also
measured at 77K, as shown in Figure 3.1.3. While the features in the 4K spectra are a little
sharper than those taken at 77K, all of the same features can be resolved. Furthermore, the
overall FWHM which can be used to measure carriers concentration are the same. This indicates
that the spectra taken at 77K are sufficient to resolve the features required to establish doping
concentrations. Based on previous literature for impurity levels below 5x1017cm-3, the band-
acceptor transition is dominant[56]. However, as the minority carrier concentration increases,
the impurity level within the band gap (ET) extends into the band, due to inhomogeneous dopant
distributions and potential fluctuations (Vrms). The following formula is obtained for Vrms [62]:
41

where rs is the screening length

and for moderate to heavily doped cases,
we can assume

. ET can be calculated as:

As the concentration of the dopant increases, ET increases until it reaches the ionization energy
(EA ~ 26.4 meV) of the C dopants and creates localized acceptor states. At the same time, an
increase in n leads to the formation of a conducting impurity band in the energy gap and creation
of free electrons. Therefore, the dominant transition would occur as an indirect band-acceptor
transition between the free electrons in the conduction band and localized acceptor states in the
impurity level within the band gap near the valance band (Figure 3.1.4).
In Figure 3.1.5, we use the empirical relation reported by Queisser et al. in order to estimate the
carrier concentrations from the energy half-width ( E) of the dominant emission line in PL
spectra[56]:

Where E is measured in meV and n in cm-3.We can also estimate the band gap variation as a
function of dopant concentration using the indirect band-localized acceptor transition model;
where Eg is:

and EF can be derived from the work of Raymond et al.[64]:

42

Using the above formula, we have calculated the expected band shift from the photoluminscent
FWHM and compared that to the experimental values. The results are plotted in Figure 4b;
where the experimental values and calculated values are in relatively close agreement. However,
the experimental results show band gap values of about 10 meV below the calculated values, and
the blueshift due to Pauli blocking in the band gap is evident.

Figure 3.1.1 SEM image of n-type doped GaAs nanowires a) arrays of nanowires, b) individual
nanowire on Si/SiO
2
substrate

43

Figure 3.1.2 Photoluminescence response of n-type doped GaAs nanowires excited by a CW
532 nm laser at 4K, a) lightly doped, b) moderately doped and c) heavily doped.
44

Figure 3.1.3 PL response of n-type doped GaAs nanowires at 4K and 77K, a) lightly doped
GaAs NW at 77K, b) moderately doped GaAs NW at 77K and c) heavily doped GaAs NW at
77K. d) lightly doped GaAs NW at 4K, e) moderately doped GaAs NW at 4K and g) heavily
doped GaAs NW at 4K.

Figure 3.1.4 Schematic for indirect band-acceptor transition.
45

20 30 40 50 60 70 80 90
5x10
17
1x10
18
2x10
18
2x10
18
3x10
18
3x10
18
4x10
18
Carrier Concenteration
Energy FWHM(eV)

20 30 40 50 60 70 80 90 100
1.49
1.50
1.51
1.52
1.53
1.54
1.55
1.56
Eg(eV)
Energy FWHM (eV)
Experimental
Model

Figure 3.1.5 a) Carrier concentrations estimated from photoluminescence FWHM, b) effective
band gap estimated from both experimental photoluminescence response and model [56]

(a)
(b)
46

3.2 Optical Responses from PIN (p-type, intrinsic, n-type) GaAs Nanowires
PIN GaAs nanowires pose three regions of p-type (one end), n-type (the other end) and
the intrinsic region in the middle. As it was mentioned in the previous sections, the type of the
dopant in GaAs nanostructures will affect the PL response. P-type (acceptors) will lead to band
gap reduction or red shift which is due to the formation in acceptor states within the gap. N-type
dopant will lead to band gap widening or blue shift which is perfectly described by Burstein-
Moss effect [65-67].According to Burstein-Moss, the shift results from the filling of conduction
band. We expect to be able to observe such shifts (red and blue) along the PIN nanowire using
PL. In Figure 3.3.1, a SEM micrograph of the PIN sample is demonstrated. Figure 3.3.2
demonstrates the PL spectra changes as we traverse the laser along the wire. From point 1 to 5
the red shift (p-type region) can be observed and from point 8 to 10 the blued shift is noticeable
(n-type region). In Figure 3.3.3, the peak position corresponding to each position point on the
wire is demonstrated. The maximum concentration of dopant for p-type regions is 2.5x10
19
/cm
3

(point 1) and that for n-type region is 5x10
18
/cm
3
(point 10).

Figure 3.2.1 SEM micrograph of the PIN GaAs nanowire.

1
1 0
47

Figure 3.2.2 The PL spectra of the PIN GaAs nanowire as we traverse from one end of the wire
to the other.

Figure 3.2.3 The peak position corresponding to each position point on the PIN GaAs nanowire.

0 1 2 3 4 5 6 7 8 9 10
848
850
852
854
856
858
860
862
864
866
Wavelength(nm)
Position

820 830 840 850 860 870 880 890 900
0.00
0.05
0.10
0.15
0.20
0.25
0.30
Intensity(a.u.)
Wavelength(nm)
Point (1)
Point (2)
Point (3)
Point (4)
Point (5)
Point (6)
Point (7)
Point (8)
Point (9)
Point (10)

48

Chapter 4

Observation of Fabry-Perot Resonance from GaAs Nanosheets

GaAs nanosheets with no twin defects, staking faults, or dislocations are excellent
candidates for optoelectrical applications. Their outstanding optical behavior and twin free
structure make them superior to traditionally studied GaAs nanowires. While many research
groups have reported optically resonant cavities (i.e., Fabry-Perot) in 1D nanowires, here, we
report an optical cavity resonance in GaAs nanosheets consisting of complex 2D asymmetric
modes, which are fundamentally different from one-dimensional cavities. These resonant modes
are detected experimentally using photoluminescence (PL) spectroscopy, which exhibits a series
of peaks or “fringes” superimposed on the bulk GaAs photoluminescence spectrum. Finite -
difference time-domain (FDTD) simulations confirm these experimental findings and provide a
detailed picture of these complex resonant modes. Here, the complex modes of this cavity are
formed by the three non-parallel edges of the GaAs nanosheets. Due to the asymmetrical nature
of the nanosheets, the mode profiles are largely unintuitive. We also find that by changing the
substrate from Si/SiO
2
to Au, we enhance the resonance fringes as well as the overall optical
emission by 5X at room temperature. Our FDTD simulation results confirm that this
enhancement is caused by the local field enhancement of the Au substrate and indicate that the
thickness of the nanosheets plays an important role in the formation and enhancement of fringes.

49

4.1 Observation of Fabry-Perot Resonance from Bare GaAs Nanosheets at
77K and Room
GaAs is used in different applications including solar cells, LEDs and lasers. However,
these structures, particularly nanowires, suffer from different defects like twin formation,
dislocations and stacking faults. While GaAs nanowires suffer from such defects, nanosheets
grow easily without twin defects or stacking faults; hence, overall they show higher quality and
optical properties in comparison to GaAs nanowires [68]. Among different applications,
nanoscale lasers have been studied in a variety of nanostructure geometries. So far, both
nanowire lasers and thin film semiconductors have demonstrated high intensity emission under
optical or electrical pumping [5, 69-72][6-7]. Nanowire lasers have also been studied broadly.
Most recently, Saxena et al. [73] reported passivated GaAs nanowire lasing at room temperature
by optimizing the size and surface passivation. Duan et al. reported electrically pumped lasing
from n-type CdS nanowires [5]. Fukui et al. reported observation of Fabry-Perrot resonances and
later lasing from GaAs nanowires [19, 70, 74]. In our previous work, we observed FP resonance
from AlGaAs-passivated, ionic liquid-passivated and non-passivated GaAs nanowires at room
temperature [75]. Considering mass interest in nano-cavities and nanolaser, formation of such
optical cavities in a defect-free nanostructure is the point of interest for many research groups.
This issue is interesting since defect-free cavities show superior electrical properties and can be
more efficiently electrically pumped.
Most previously reported cavities are cross-plane (1D) cavities formed between two
parallel mirrors. In the work presented here, we demonstrate an in-plane (2D) optical cavity
formed by the three non-parallel edges of the GaAs nanosheets. The nature of this resonance is
studied using photoluminescence spectroscopy as a function of GaAs nanosheet dimensions,
50

laser power, temperature, and underlying substrate. In order to obtain a fundamental
understanding of the resonances and underlying PL enhancement mechanism, finite-difference
time-domain (FDTD) simulations are carried out on an individual GaAs nanosheet. The
simulations are configured to separate various enhancement mechanisms of light absorption and
emission. Optical cavity modes are visualized by plotting the electric field intensity through a
cross-sectional slice of the nanosheet. Both the strength and quality factors of the resonances are
extracted from taken measurements and calculations. GaAs nanosheets are synthesized on
(111)B GaAs substrates by metal organic chemical vapor deposition (MOCVD) with selective
area growth (SAG). The detailed synthesis method has been described in previous publications
[76, 77]. Briefly, trimethylgallium (TMGa) and arsine are used as precursors for Ga and As
deposition, and nanosheets are n-type doped using disilane (Si
2
H
6
). A 28nm layer of SiN
x
is
deposited on the GaAs substrate using plasma enhanced chemical vapor deposition (PECVD),
which serves as the growth mask. Electron beam lithography is used to create nanoscale strips in
the SiN
x
which is selectively etched, leading to the formation of windows exposing the
underlying GaAs substrate. The nanosheets are grown in a hydrogen environment at 0.1atm
using a V/III ratio of 1.5, and Ga/Si ratio of about 16000 for 30 minutes at 790˚C. The partial
pressure of TMGa, AsH
3
, and Si
2
H
6
are 1.124x10
-6
atm, 1.63x10
-6
atm, and 2.857x10
-9
atm,
respectively. The as-grown nanosheets are sonicated for 5 sec in isopropanol and subsequently
transferred onto Si/SiO
2
(300nm SiO
2
) or Si/SiO
2
/Au (300nm- SiO
2
/200nm Au) substrates.
These GaAs nanosheets grow on (111)B surfaces through a self-terminating growth mechanism,
which results in their triangular shape with a base length of 6-10 µm, and height of 2-4 µm, as
shown in Figure 1. The thickness of these nanostructures is 200nm ± 25nm. Micro-Raman and
photoluminescence spectroscopy are collected using a 40X objective lens, an 1800 l/mm grating,
51

and a silicon CCD camera to detect the Photoluminescence (PL) in the range from 750nm to
950nm. Low to moderate power excitation (10
2
-10
5
mW/cm
2
) using a 532nm laser is used to
avoid optical heating. Figure 4.1.1a shows an SEM image of a GaAs nanosheet deposited on a
Si/SiO
2
substrate. The center of the sample is excited by a 532 nm CW laser focused to a spot
size of 1.5 µm with an excitation power of 10
5
mW/cm
2
. The thickness of the nanosheet is
measured using atomic force microscopy (AFM), which gives a measured value of 200nm ±
10nm. The corresponding PL spectrum at 77 K is shown in Figure 4.1.1b. The spacing between
the resonance peaks varies between 6.5nm to 10 nm, with an average spacing of 8.47 nm. In this
spectrum, two predominant peaks are observed at 1.52 eV and 1.44 eV. The peak at 1.52 eV
corresponds to the nominal band gap of GaAs (1.507eV at 77K), blueshifted by 12 meV due to
n-type doping by Si dopant impurities, which results in Pauli blocking of band-edge to band-
edge optical transitions. The 1.44 eV sub-band gap peak is attributed to an acceptor-band
transition arising from carbon impurities originating from the MOCVD growth, carbon-doping
of samples is expected from growth precursors[78, 79]. The quality factors of the resonance
peaks range from 88 to 125, indicating that trapped light is totally internally reflected roughly
100 times. These resonance fringes are only observed in nanosheets with thickness of 200-
220nm, whereas no such resonance is observed in thicker GaAs nanosheets ranging from of
250nm to 300nm. As mentioned, we have previously observed formation of the optical cavity in
the GaAs nanowires; therefore, comparing the behavior of such cavities is demonstrated in Table
4.1.1.

52

Table 4.1.1 Comparison between the properties of the optical cavity in the nanowires vs.
nanosheets
GaAs nanowires GaAs nanosheets
Peak-to-peak Spacing (Experimental) 8-10nm 6.5-10nm
Peak-to-peak Spacing (Theoretical) 11nm 12.3 nm
Quality Factor 181 125
Optimal Thickness 270nm±20nm 210nm±10nm

Previously, we have shown that using Au as substrate not only enhances the
photoluminescence intensity of the GaAs nanowires, but also leads to observation of FP
resonance in GaAs nanowires at room temperature. Here, we imply the same mechanism for
GaAs nanosheet and study the effects of using Au substrate. In Figure 4.1.2, the PL spectra of a
similar GaAs nanosheet deposited on a Si/SiO
2
/Au (200nm of Au deposited on Si/SiO
2
) substrate
is compared to that of one taken on Si/SiO
2
. For both datasets, the PL spectra are centered at 840
nm, corresponding to a band gap of 1.47eV. Here, we observe a blueshift of 50 meV compared
to the room temperature GaAs band gap (1.42 eV), again due to n-type doping of GaAs with Si.
A 5X enhancement in the PL intensity and emergence of fringes are observed as a result of Au
the substrate. The effective cavity length can be determined using the following formula and the
experimental value of the peak-to-peak spacing:
2
2
l
l
l
l








dn
Ln
d
.
Assuming n = 3.655, Δλ = 10.8 nm, and dn/dλ= -0.822/µm, the cavity length is approximately
7.45µm, consistent with the base length of the nanosheet. As a comparison in our previous work,
a pronounced resonance was observed in GaAs nanowires grown with the SAG-MOCVD
method[75]. The reported mode spacing for these GaAs nanowires was 11.5 nm (at 4K).
53

In order to understand the underlying optical resonances and PL enhancement
mechanism, finite-difference time-domain (FDTD) simulations are carried out on a GaAs
nanosheet on top of both Si/SiO
2
and Si/SiO
2/
Au substrates, as illustrated in Figures 4.1.3a and
3b. In each configuration, the geometry of the nanosheet resembles those grown experimentally,
with dimensions determined from SEM and AFM. The simulated nanosheet is an obtuse triangle
with a maximum base length of 6 m, an obtuse angle of 119.5°, minimum angle of 21.8°, and
thickness of 200 nm. In each system we account for the dispersive nature of the materials and
impose perfectly matched layers (PML) along the simulation boundaries. A 532 nm optical
pump, indicated by the green spot in each figure, is focused onto the GaAs nanosheet. We
simulate the pump excitation using a Gaussian beam with a 1.5 m spot size. We can quantify the
power absorbed in each system using the following formula:
2
"
ABS
12 PE w 

where w is the angular frequency,
2
E
is the electric-field intensity, and
"
 is the imaginary
component of the dielectric function. In order to calculate the absorptive enhancement in the
GaAs nanosheet due to the presence of an underlying Au film, we take a horizontal slice across
the triangular profile of the nanosheet (100 nm above the SiO
2
and Au substrates), integrate the
electric-field intensity, and take the corresponding ratio. The absorptive enhancement factor (
ABS
EF ) for a single wavelength is simply the ratio of integrated electric-field intensities:
2
2
ABS 2
Au
SiO
E dA
EF
E dA




54

With the pump excitation at 532 nm, the absorptive enhancement factor is nearly 11.5X,
indicating that the Au substrate leads to 11.5X more optical absorption within the nanosheetThe
metallic film appears to form a vertical cavity with the top of the nanosheet that traps light and
results in a large PL enhancement factor. We model PL emission within the GaAs nanosheet
using electric dipole sources. Each electric dipole emits between 750 – 1050 nm. A Gaussian
distribution is superimposed upon the emission spectra so that the results more closely resemble
the PL data collected experimentally. In order to examine the impact of emitter position on the
observed optical resonances, we consider four emitter positions: one position near each corner of
the nanosheet and one position near its center. At each emitter position, we consider three
orthogonal polarizations (i.e. x, y, and z). In order to avoid coherence effects, we include one
dipole per simulation and later perform an incoherent summation of the emission results. We
introduce a quantity, the extracted power   EP , which is defined as the total power collected
from the top of the GaAs nanosheet, measured using a flux-plane monitor. Figure 4.1.4a shows
the incoherently summed EP results for the cross-sectional position located near the smallest
corner of the nanosheet. Optical fringes are clearly visible in the simulated spectra for both the
SiO
2
and Au underlying substrate. Figure 4.1.4b shows the electric-field intensity patterns for the
wavelengths corresponding to the dashed line in Figure 4.1.4a, which coincides with a resonance
in each system. The solid white outline marks the triangular boundary of the nanosheet, whereas
the green dot indicates the position of the dipole emitters. The nanosheet forms an in-plane
cavity with clear interference fringes evident in the 2D field intensity plots, further supporting
the presence of resonant modes. The integrated extracted power enhancement  
ENH
EP is
defined as:
55

2
Au
ENH
SiO

EP d
EP
EP d
l
l




and represents the comparative enhancement in collected emission. We find that this
enhancement is independent of the emitter’s position with a value always near unity. The
broadband EP enhancement for the emission results shown in 4.1.4a is 1.2X. If we consider the
overall PL enhancement to be a product of absorptive enhancement and extracted power (EP)
enhancement, then we theoretically predict a PL enhancement of nearly 13.8X. The majority of
the enhancement contribution is owed to absorptive enhancement, resulting from the formation
of a vertical cavity due to the underlying metallic film. The observation of optical resonances,
however, is a consequence of the in-plane cavity formed by the cross-sectional profile of the
nanosheet. Both the number and quality-factor of observed resonances is strongly dependent on
the emitter’s position.
In conclusion, we report the observation of resonance in GaAs nanosheets resulting from
the formation of an optical cavity created within these nanostructures themselves. The in-plane
resonance modes are attributed to complex modes resulting from total internal reflections. A 5X
enhancement in PL intensity and resonance fringes are observed from GaAs nanosheets on
Si/SiO
2
/Au substrates with respect to those on Si/SiO
2
substrates. Our FDTD simulations
indicate that the enhancement is due primarily to absorptive enhancement rather than
photoemission enhancement, and confirms the formation of two cavities: one vertical and one in-
plane. The vertical cavity forms at the interface of Au and GaAs nanosheet, while an in-plane
cavity is formed between the three facets of the nanosheet. The theoretical enhancement due to
both absorptive and extracted power is 13.8X, and does not include any losses due to non-
56

radiative recombination effects. The low quality factors observed may be due to high optical loss
as the result of the high density of surface states, typical of GaAs nanostructures. Although the
results observed in this study is similar to that of GaAs nanowires, GaAs nanosheet are
electrically superior to GaAs nanowires due to the fact that they are defect free. Therefore,
possibility of efficient electrical pumping of GaAs nanosheets is expected. Passivation of these
surface states may decrease the non-radiative recombination and enhance the optical activity of
these nanostructures.

750 800 850 900 950
0.0
5.0x10
4
1.0x10
5
1.5x10
5
2.0x10
5
PL Intensity
Wavelength(nm)
GaAs NS on SiO
2

77K

Figure 4.1.1 (a) SEM image of a GaAs nanosheet and (b) Photoluminescence spectrum taken at
77K. The dominant peak at 816 nm is attributed to the band-to-band transition, and the peak at
860 nm corresponds to carbon acceptor impurities.

(a)
(b)
57

750 800 850 900 950 1000
0.0
2.0x10
3
4.0x10
3
6.0x10
3
8.0x10
3
1.0x10
4
300K
PL Intensity
Wavelength(nm)
GaAs NS on Si/SiO
2
GaAs NS on Au
5X

Figure 4.1.2. Photoluminescence spectra from GaAs nanosheets on Si/SiO
2
and Si/SiO
2
/Au.
Spectra were collected at room temperature under 532 nm CW laser excitation.

Figure 4.1.3. Schematic diagrams illustrating the two simulated pump geometries: GaAs
nanosheet on top of (a) 300 nm of SiO
2
and a semi-infinite Si substrate and (b) 200 nm of Au on
top of 300 nm of SiO
2
and a semi-infinite Si substrate. The nanosheet has a maximum side length
of 6 m, obtuse angle of 119.5°, minimum angle of 21.8°, and a thickness of 200 nm. The green
spot in each diagram represents a 532 nm optical pump with a beam waist of 1.5 m focused
onto the nanosheet.

(a) (b)
58

Figure 4.1.4. (a) Incoherently summed extracted power (EP) spectra calculated for the
topologies shown in Figure 3. (b) Electric-field patterns for resonances in each material system;
the wavelength of each resonance is indicated by the dashed lines in (a). The solid white lines
indicate the triangular outline of the nanosheet, whereas the green dots indicate the position of
the dipole emitters.

59

Figure 4.1.5. (a-c) Incoherently summed extracted power (EP) results for three different cross-
sectional emitter positions. Resonances are present in each material system with the number and
quality factor of the resonances varying significantly with emitter position. (d-f) Electric-field
patterns for resonances in each material system; the wavelength of each resonance is indicated by
the dashed lines in (a-c), respectively. The solid white lines indicate the triangular outline of the
nanosheet whereas the green dots indicate the position of the dipole emitters. Complicated
interference patterns are present in each figure and depend sensitively on emitter position. The
broadband extracted power (EP) enhancement for each dipole position in (a-c) is 1.0X, 1.0X, and
1.1X, respectively.

60

4.2 Effects of Ionic Liquid Passivation of Optical Response from GaAs
Nanosheets

Here, we demonstrated the PL emission of such nanosheets. As for nanowires,
nanosheets are sonicated and deposited on Si/SiO
2
substrate or Si/SiO
2
substrate with thin layer
of Au. As the passivation layer, ionic liquid (ethyl-methylimidazolium bis(trifluoromethyl-
sulfonyl) imide) is used. Here, biasing voltage is also applied and the increase in PL intensity is
studied. The PL spectra show about 41% enhancement after application of ionic liquid and bias
voltage application, the results are shown in Figure 4.2.1. We hypothesize that this form of
passivation may also improve and aid the formation of Ohmic contacts.
In Figure 4.2.2, we observe the effects of ionic liquid passivation without application of
bias voltage. An enhancement factor of 8X is observer.

Figure 4.2.1 PL spectra of GaAs nanosheet after application of ionic liquid and bias voltage
application
61

750 800 850 900 950
0.0
2.0x10
3
4.0x10
3
6.0x10
3
8.0x10
3
1.0x10
4
PL Intensity (Count)
Wavelength(nm)
GaAs NW without passivation at 300K
GaAs NW without ionic liquid passivation at 300K

Figure 4.2.2 PL spectra of non-passivated and passivated n-type GaAs nanowires with ionic
liquid at 300K

62

Chapter 5

Optical Enhancement of GaAs Nanosheets after AlGaAs
Passivation

Unlike nanowires, GaAs nanosheets exhibit no twin defects, stacking faults, or
dislocations even when grown on lattice mismatched substrates. As such, they are excellent
candidates for optoelectronic applications, including LEDs and solar cells. We report substantial
enhancements in the photoluminescence efficiency and the lifetime of passivated GaAs
nanosheets produced using the selected area growth (SAG) method with metal organic chemical
vapor deposition (MOCVD). Measurements are performed on individual GaAs nanosheets with
and without an AlGaAs passivation layer. Both steady state photoluminescence and time-
resolved photoluminescence spectroscopy are performed to study the optoelectronic performance
of these nanostructures. Our results show that AlGaAs passivation of GaAs nanosheets leads to a
30- to 40-fold enhancement in the photoluminescence intensity. The photoluminescence lifetime
increases from less than 30ps to 300ps with passivation, indicating an order of magnitude
improvement in the minority carrier lifetime. We attribute these enhancements to the reduction
of non-radiative recombination due to the compensation of surface states after passivation. The
surface recombination velocity decreases from an initial value of 2.5x10
5
cm/s to 2.7 x10
4
cm/s
with passivation.

63

5.1 Passivation Methodology and Measurement Setup

GaAs is an excellent candidate material for high efficiency solar cells[80, 81],
nanolasers[73, 74] and LED applications[82]. GaAs nanostructures (e.g., nanowires and
nanosheets) with high surface-to-volume ratios, however, suffer from high surface state densities
and high surface recombination velocities, which typically limit their optoelectronic device
performance[83]. Passivation of GaAs nanostructures has been widely studied in the literature,
including cladding of GaAs nanostructures with wide gap materials (e.g., AlGaAs, GaP and
GaAsP)[55, 74, 82, 84, 85]. A number of chemical routes to passivation of GaAs surfaces have
been studied, including sulfur solution passivation[86, 87], thiol passivation[88], and ionic liquid
passivation[89, 90]. Recently, selective area growth (SAG) of GaAs nanosheets has been
demonstrated using a slit instead of a hole geometry in the SiN growth masking layer[76]. GaAs
nanosheets have the distinct advantage over GaAs nanowires in that they grow easily without the
formation of twin defects and stacking faults, thus improving the overall sample quality and
optoelectronic properties in comparison to GaAs nanowires[68]. Longer minority carrier
diffusion lengths make GaAs nanosheets particularly well-suited for solar cell and nanolaser
applications. Previous literature provides extensive studies on AlGaAs passivation of GaAs
nanowires[91, 92]. However, no such studies have been reported for GaAs nanosheets.
In the work presented here, the effects of passivation of GaAs nanosheet structures are
explored using heteroepitaxial growth of an AlGaAs passivation layer on the GaAs nanosheet.
The effects of passivation on the photoluminescence efficiency and minority carrier lifetimes are
studied systematically using the techniques of room temperature and low temperature
photoluminescence (PL) spectroscopy and time-resolved photoluminescence (TRPL)
64

spectroscopy. All measurements are performed on individual GaAs nanosheets that have been
removed from the growth substrate and transferred to Si/SiO
2
substrates in order to eliminate any
contribution to the PL from the underlying substrate.
GaAs nanosheets are synthesized in a vertical showerhead, low-pressure metal organic
chemical vapor deposition (MOCVD) reactor with selective area growth (SAG)[68, 76].
Trimethylgallium (TMGa), trimethylaluminum (TMAl), and arsine are used as precursors for Ga,
Al, and As deposition. Nanosheets are doped n-type using disilane (Si
2
H
6
). High density arrays
of GaAs nanosheets are grown on (111)B GaAs substrates. The substrate preparation for
nanosheet growth is the same as that previously reported for nanowires[55]. First, a thin layer of
SiN
x
(approximately 28nm thick) is deposited onto the (111)B GaAs substrate by plasma
enhanced chemical vapor deposition (PECVD). 10 µm long, 100 nm wide openings are
patterned in the SiN
x
layer by electron-beam lithography and reactive ion etching. These
nanoscale stripes are patterned along the   2 11 direction of the GaAs growth substrate. The
nanosheets are grown in a hydrogen environment at 0.1atm using a V/III ratio of 1.5 and a Ga/Si
ratio of 16000 for 75 minutes at 790˚C. The partial pressure of TMGa, AsH
3
, and Si
2
H
6
were
1.12x10
-6
atm, 1.63x10
-6
atm, and 7.14x10
-11
atm, respectively. The AlGaAs passivation layer is
grown at the same temperature for 150 seconds with TMGa, AsH
3
, and TMAl at partial pressures
of 3.74x10
-7
atm, 1.85x10
-4
atm and 3.82x10
-7
atm. The approximate thickness of the AlGaAs
passivation layer is 20nm while that of the nanosheets is 300nm, as measured by scanning
electron microscopy (SEM) and atomic force microscopy (AFM). Micro-Raman spectroscopy of
the nanosheets was performed using a 100X objective lens with a numerical aperture of 0.85 and
a 40X objective lens with a numerical aperture of 0.6. A Si CCD camera was used to collect PL
spectra in the 750 to 1000 nm wavelength range. A continuous wave, 514 nm or 532 nm laser
65

was used for excitation. For the time-resolved PL measurements, a 76 MHz pulsed Ti-Sapphire
laser (Mira 900) was used to pump a supercontinuum fiber and provide optical excitation at 590
nm. The 590 nm excitation light was focused onto the nanosheet using a 50X/0.5NA objective
lens, and the emitted PL was dispersed by a MS260i (Newport) imaging spectrograph and
detected by a MCP-PMT (Hamamatsu) phototube. Time-correlated single photon counting was
used to obtain time-resolved PL data. Spatial imaging of single nanosheets was obtained using
slit-confocal microscopy[93]. The 514 nm laser was defocused to a ~30 um spot (encompassing
the full nanosheet) while the nanosheet was imaged with the short axis of the nanosheet oriented
parallel to the entrance slit of the spectrometer. As the nanosheet was scanned across the
entrance slit along the long axis, full CCD images of the PL were obtained, Figure 5.1.1a shows
an SEM image of the as-grown nanosheets on the underlying GaAs/SiN growth substrate. The
nanosheets are then removed by sonication in IPA and deposited on a Si/SiO
2
substrate, as
shown in Figure 5.1.1b. Nanosheets are grown along the (111) direction through a self-
terminating process. The base triangle is referred to as the initial growth region of the nanosheet.
Once the initial triangle self-terminates, additional triangles grow from the apex of the initial
triangle, which is referred to as the over-grown region [68, 76]. In these nanosheet structures, the
exposed surface of the nanosheet is the (011)A direction. In the subsequent passivation growth
step, AlGaAs is grown on this surface, which is not the usual surface for epitaxial growth. It is,
therefore, not known a priori how epitaxial the GaAs/AlGaAs interface will be in these novel
nanostructures, thus providing additional motivation for this study.
66

Figure 5.1.1 SEM images of non-passivated GaAs nanosheet: a) array of as-grown
nanosheets, b) individual nanosheet deposited on Si/SiO
2
substrate

Nanosheet
(a)
5 µm
Nanosheet
(b)
67

5.2 Effects of Passivation on the Steady State and Time-resolved PL
Responses of GaAs Nanosheets

Figure 5.2.1 shows a comparison of the PL spectra taken at a single point from an
individual nanosheet with and without AlGaAs passivation. Nanosheets are excited with 532 nm
wavelength (continuous wave) laser light at room temperature with a 100X objective lens. Here,
the PL intensity is enhanced by a factor of 42X after AlGaAs passivation. To assess how the
passivation of the nanosheet might affect the spatial distribution of PL, spatial maps of the PL
intensity were obtained using slit-confocal microscopy. Figure 5.2.2a shows such a spatial map
where the integrated intensity of the band-edge emission is displayed as a function of position for
an as-grown and, thus, unpassivated GaAs nanosheet. One can see points of relatively bright
emission at the two ends of the nanosheet, while the interior of the nanosheet emits at
approximately 30 – 50% the maximum intensity. Figure 5.2.2b shows an optical image of the
nanosheet passivated with a 20 nm AlGaAs layer, where the arrows show the direction of the
line scans used for spatially-resolved micro-Raman measurements shown later. Figure 5.2.2c
shows a PL image of the passivated GaAs/AlGaAs nanosheet. By comparing the two PL images
in Figures 5.2.2a and 5.2.2c we find that the maximum intensity integrated over the whole area
of the passivated GaAs nanosheet is nearly 30 times that of the unpassivated nanosheet The fact
that the PL intensity maximum occurs along the intermediate edge may indicate that the
AlGaAs-on-GaAs growth is considerably better and perhaps more epitaxial at this edge. This is
not surprising considering that this is the last surface deposited during the initial nanosheet
growth. The weak PL intensity along the long edge indicates that the long edge is not passivated
since it is cleaved from the growth substrate.
68

To assess whether the reduced intensity at a position in the passivated nanosheet is due to
an imperfection or inhomogeneity in the AlGaAs layer, we perform spatially-resolved micro-
Raman spectroscopy on the same nanosheet. The 514.5nm line from an Argon-ion laser is used
for excitation with 0.5mW power and the scattered light is collected by a 100X 0.7NA objective
lens. The scattered light is dispersed by an xy-Dilor spectrometer and detected by a liquid-
nitrogen cooled Si-CCD. The sample is mounted on a piezoelectric translation stage so that the
nanosheet can be sampled along two orthogonal lines, longitudinal and transverse, as shown in
the optical image of the nanosheet (Figure 5.2.2a). The aman spectra (see Figure 5.2.3) reveal
GaAs-like TO, AlAs-like TO and LO features, along with a weak GaAs-like LO response. These
phonon energies are obtained by fitting the Raman spectra to Lorentzian lineshapes for both
scans, as displayed in Figures 5.2.4a and 5.2.4b for the longitudinal and transverse line scans,
respectively. Since the position of the AlAs-like LO phonon is a sensitive measure of Al
concentration[94], we obtained the Al% based on the observed phonon energies using the
equation:

361 643 62 888x 27 91x
2
.
Here, x is the percentage of Al concentration and

is the AlAs-like LO phonon. We find the
Al% is consistently 55% across the entire area of the nanosheet, with only a small deviation of
1%, indicating good growth quality and control across the sample. As such, there is no obvious
correlation of the reduced region of PL intensity with the concentration of the AlGaAs layer.
Time-resolved PL spectroscopy was also performed on these nanosheets in order to
determine the carrier lifetimes before and after passivation, as shown in Figure 5.2.5. We find
that passivation of n-type GaAs nanosheets with an AlGaAs layer yields carrier lifetimes of up to
300 ps, which is an order of magnitude longer than the non-passivated nanosheets with a lifetime
69

of less than 30 ps, which is our system response. The increase in carrier lifetime is due to
passivation of the surface states, which greatly reduces the non-radiative recombination of photo-
excited carriers. Given the 30-fold increase in PL intensity seen in the PL maps of Figure 5.2.2,
we infer that the unresolved recombination lifetime in the unpassivated nanosheet may be only
~10 ps and is limited by the non-radiative recombination at the surfaces. This result is consistent
with previous reports of GaAs nanowires passivated with AlGaAs[55], which showed a lifetime
of 1.3 ns for AlGaAs-passivated nanowires. The surface recombination velocity (SRV) can be
approximated using the following formula[83, 95]:

Here, τ is the minority carrier lifetime, τ
b
is the bulk GaAs lifetime (1.3 µsec), d is the thickness
of the nanosheets, and S is the surface recombination velocity. Considering a GaAs nanosheet
with an average thickness of 300nm, the surface recombination velocity measured for an
AlGaAs-passivated sample is 2.7x10
4
cm/s, which is about an order of magnitude smaller than
the value measured for the non-passivated case of greater than 2.5x10
5
cm/s). Lower SRV values
of 500 cm/s have been reported for passivated bulk GaAs in the literature[96]. The slightly
higher recombination velocity may result from the long unpassivated cleaved edge in the
nanosheets measured here.

In summary, our steady state and time-resolved photoluminescence measurements of
individual GaAs nanosheets indicates that AlGaAs passivation improves the photoluminescence
efficiency by 30 to 40 times and increases the minority carrier lifetime by more than an order of
magnitude. We attribute these enhancements to the reduction of non-radiative recombination due
70

to the compensation of surface states after passivation. The surface recombination velocity
shows a substantial decrease from 2.5x10
5
cm/s to 2.7x10
4
cm/s. These results indicate that
passivated GaAs nanosheets

750 800 850 900 950 1000
0
1x10
5
2x10
5
3x10
5
4x10
5
5x10
5
42X
200 250 300 350 400 450
AlAs
Intensity (a.u.)
Raman Shift (cm-1)
GaAs
AlAs

PL Intensity
Wavelength (nm)
Passivated
Non-passivated

Figure 5.2.1 Photoluminescence spectra of an individual GaAs nanosheet with and without
AlGaAs passivation. Inset graph: Raman spectra of passivated GaAs nanosheet [97].

71

Figure 5.2.2 Spatial maps of the PL intensity a) PL spatial map for an as-grown and thus
unpassivated GaAs nanosheet, b) Optical image of the nanosheet passivated with a 20 nm
AlGaAs layer, where the arrows show the direction of the line scans used for spatially-resolved
micro-Raman measurements shown later, c) PL spatial map for the passivated GaAs/AlGaAs
nanosheet.

72

Figure 5.2.3 Selected Raman spectrum near the center of the GaAs/AlGaAs sample. GaAs TO
(267cm
-1
) and AlAs-like TO (360cm
-1
) and LO (388cm
-1
) phonon structures are clearly
identified, along with a weak GaAs LO response (~290cm
-1
). The weak structure at the lower
energy side of the GaAs TO is presumably GaAs-like TO response from AlGaAs.

73

0.0 0.5 1.0 1.5 2.0 2.5 3.0
264
266
268
356
358
360
386
388
390

Longitudinal Position ( m)
GaAs TO

AlAs-like TO
Phonon Energy (cm
-1
)

AlAs-like LO
0.50
0.55
0.60
Al%
0.0 0.5 1.0 1.5 2.0 2.5
265
266
267
268
360
361
362
363
387
388
389
390

Transverse Position ( m)
GaAs TO

AlAs-like TO

Phonon Energy (cm
-1
)
AlAs-like LO
0.52
0.56
0.60
Al%

Figure 5.2.4 Longitudinal and transverse Raman scattering scan analysis. The Al-concentration
is calculated based on AlAs-like LO phonon energy due to its strong Al-concentration dependent
nature[98]. From both of the scans, we observe a consistent Al-concentration 55% across the
sample with a deviation of ±1%.

74

0.0 0.5 1.0 1.5 2.0 2.5
Passivated GaAs nanosheet
Non-passivated GaAs nanosheet
non-passivated NS=24ps
Passivated NS=300ps
Photoluminescence Intensity
Time (nsec)

Figure 5.2.5 Time resolved PL data from passivated and unpassivated GaAs nanosheets at 300
K. The non-passivated GaAs nanosheet PL lifetime is less than the system response (30 ps). The
passivated nanosheet PL exhibits a lifetime of 300 ps.

75

Chapter 6

Electronic Characterization

In this chapter we present an overview of the different electrical characterization methods
and measurements which have been performed on GaAs nanowires and nanosheets.

6.1 Challenges of Making Ohmic Contacts to GaAs Nanowires
A possible material system offering the best photovoltaic performance is GaAs. GaAs
optoelectric application requires Ohmic contacts to both n and p regions. In Ohmic contact, the
carriers should tunnel through the electronic barrier. The work function difference between the
metal and the semiconductor defines the height of the barrier. The work function difference
thickness is inversely proportional to the square root of dopant concentration. Fabrication of
Ohmic contact to n-type GaAs is particularly the matter of interest and yet quite difficult.
Although there are several reports on fabrication of Ohmic contacts to bulk GaAs, the
formation of Ohmic contact to GaAs nanowires is not widely investigated. In the literature, L.C.
Wang et al. used Pd/Ge (50nm/136nm) and 30 min of rapid thermal annealing (RTA) at 325˚C for
bulk n-type GaAs [99]. AuGe/Ni/Au (100nm/20nm/50nm) is another recipe used by T.S. Kuan et
al for n-type GaAs. Here, AuGe referrers to Au and Ge alloy with about 10% Au content, this
alloy is commercially available. In their method, after liftoff, RTA of 90 sec, 115 sec or 200 sec at
450 ˚C is used [100]. Y.C. Shih et al. used the same recipe with minor alteration of using
additional Ni layer as the first layer, again for n-type bulk GaAs. Deposition of Ni/AuGe/Ni/Au
76

(0.5-10nm/100nm/35nm/50nm) is then followed by 350-600 ˚C annealing for 2 min [101].
Among the very first works for n-type GaAs nanowire we can mention the work of C. Gutsche et
al. where two separate methods are used, method one Ge/Ni/Ge/Au (5nm/10nm/15nm/300nm)
followed by 30 sec RTA at 320 ˚C, 340 ˚C and 360 ˚C. In the second recipe, Pd/Ge/Au
(50nm/170nm/80nm) is used followed by 30 sec of RTA at 280 ˚C, 300 ˚C and 320 ˚C. Both of
which show Ohmic contact but degradation of nanowires near the metal contact [102]. C. Gutsche
et al. also reported formation of Ohmic contact to n-type GaAs nanowires using Pd/Ge/Au
(50nm/170nm/80nm) recipe followed by 30 sec of RTA at 280 ˚C [103]. In our case, we have
tried three of the above recipes for bulk n-type GaAs. These recipes are demonstrated in Table.
6.1.1. We first tried our contact recipes for n-type GaAs wafers. For metal evaporation we used an
ebeam metal evaporator and before evaporation we dip the samples (here n-type GaAs wafer) into
HCl: H
2
O (1:1) for 15 sec. This step is critical to remove the native oxide layer formed over the
samples. Samples are directly loaded for metal evaporation afterwards.
Table 6.1.1: Ohmic contact recipes for n-type GaAs
Metals Thicknesses (nm) RTA time and temperature
Recipe A AuGe/Ni/Au 100/30/100 30sec at 375˚C (forming gas)
Recipe B Ge/Ni/Ge/Au 5/10/15/250 30sec at 320˚C (forming gas)
Recipe C Pd/Ge/Au 50/170/80 30sec at 280˚C (forming gas)

The I-V results for each recipe are shown in Figure 6.1.1. In order to make metal contacts
to the GaAs nanowires we need ebeam lithography. The thickness of these electrodes varies
between 0.6 um to 6 um. For E-beam lithography we need to spin coat the sample with a layer of
MMA and a layer of PMMA (both are ebeam resists). The total thicknesses of the resists should
be more than 300nm to be able to do the liftoff. After ebeam lithography (Raith) samples are
77

developed in IPA:MIBK (3:1) and ready for evaporation. Before evaporation samples are dipped
in HCl:H
2
O (1:1) for 15 sec to remove the native oxide layer. We have tried all the mentioned
recipes (A,B and C) for different batches of n-type GaAs nanowires. However, we only were able
to make Ohmic contact to one heavily doped batch of wires. The results are demonstrated in
Figures 6.1.2 and 6.1.3. Our conclusion is that high dopant concentration (high 10
18
/cm
3
) is
required for acceptable Ohmic contact. Also, recipe A and C with even higher annealing
temperatures are effective for n-type GaAs nanowires.

Figure 6.1.1 I-V characteristics of n-type GaAs wafer using, recipe A, the resistance is 1.75
ohms.
78

Figure 6.1.2 I-V characteristics of n-type GaAs wafer using, recipe B, the resistance is 1.42
ohms.

Figure 6.1.3 I-V characteristics of n-type GaAs wafer using, recipe C, the resistance is 1.67
ohms.

79

Figure 6.1.4 I-V characteristics of n-type GaAs nanowires (data retrieved for three nanowires
from the same batch).

-2.0 -1.5 -1.0 -0.5 0.0 0.5 1.0 1.5 2.0
-1.2x10
-5
-1.0x10
-5
-8.0x10
-6
-6.0x10
-6
-4.0x10
-6
-2.0x10
-6
0.0
2.0x10
-6
4.0x10
-6
6.0x10
-6
8.0x10
-6
I(A)
Vbias(V)
NW1_1
-2.0 -1.5 -1.0 -0.5 0.0 0.5 1.0 1.5 2.0
0.0
1.0x10
-8
2.0x10
-8
3.0x10
-8
4.0x10
-8
5.0x10
-8
I(A)
Vbias(V)
NW1_2
-2.0 -1.5 -1.0 -0.5 0.0 0.5 1.0 1.5 2.0
-1.0x10
-8
0.0
1.0x10
-8
2.0x10
-8
3.0x10
-8
4.0x10
-8
5.0x10
-8
I(A)
Vbias(V)
NW1_6
80

6.2 I-V Characterization of GaAs Nanosheets
In order to perform electronic measurement on GaAs nanosheets, we need to make low
resistance Ohmic contact to the nanosheets. GaAs nanosheets pose different regions with
opposite doping profile. In Figure 6.2.1 an SEM micrograph of nanosheet is shown. The dark
triangle in the middle of the nanosheet is the base (initial growth) and based on the SEM contrast
(and later on EBIC) we claim that this area is n-type doped. The small triangle on the top of the
based (dark region, distinct by dash lines) is also n-type doped. However, the two side areas
which are lighter (overgrown regions) are p-type. Based on this structure we claim to have a p-n
junction within the nanosheet. Our goal is to make four contacts, two in n-type region and two in
p-type region to be able to study all forms of I-V responses across these regions; four contacts
are also required for Hall measurement. Here, vertically grown GaAs nanosheets are sonicated
and deposited on Si/SiO
2
substrate. Samples are spin coated with two layers of MMA (1 min,
4000 rpm) followed by 5 min bake at 180˚C and one layer of PMMA (1 min 3500 rpm) followed
by 5 min bake at 180C. Then ebeam lithography is used in order to make 2, 3 or 4 contact
channels. After, ebeam lithography samples are developed in IPA: MIBK (3:1) for 60 sec and
rinsed with IPA. Before evaporation samples are dipped in HCl: H
2
O (1:1) for 15 sec. Recipe A
(AuGe/Ni/Au) 100nm/30nm/100nm is used. Considering the fact that the base of the nanosheet
is at most 8µm and the height is 4µm placing four contacts in the appropriate locations is non-
trivial. In Figure 6.2.2, an SEM micrograph of a nanosheet with contacts is shown. The
corresponding I-Vs across different electrodes are shown in Figure 6.2.3. In this case we were
not able to achieve Ohmic contact even after annealing at 375 ˚C for 30 sec. We tried to optimize
our annealing recipe to get Ohmic contact. In Figure 6.2.4.a to 6.2.4.e SEM images of nanosheet
81

under different RTA processes are demonstrated. In Figure 6.2.5 the I-Vs corresponding to each
annealing stage is shown and in Figure 3.2.6 the variation of resistance versus the applied RTA
temperature is shown.

The same experiment has been done on another GaAs nanosheet in Figures 6.2.7.a to
6.2.7.e, 6.2.8 and 3.2.9. Both cases show that the best RTA temperature, which leads to low
Ohmic contact resistance, is about 380 ˚C to 385˚C.

Figure 6.2.1 SEM micrograph of nanosheet with possible electric contacts.

82

Figure 6.2.2 SEM micrograph of nanosheet and corresponding I-V characteristics.

83

Figure 6.2.3 SEM micrograph of GaAs nanosheet after RTA annealing at 375˚C, 380˚C, 385˚C,
390˚C and 395˚C.

Figure 6.2.4 I-V characteristics of GaAs nanosheet after RTA annealing at 375˚C, 380˚C, 385˚C,
390˚C and 395˚C.

-2.0 -1.5 -1.0 -0.5 0.0 0.5 1.0 1.5 2.0
-400
-300
-200
-100
0
100
200
300
400
500
600
I(uA)
Vbias(V)
After 375C RTA
After 380C RTA
After 385C RTA
After 390C RTA
After 400C RTA
84

Figure 6.2.5 RTA temperature dependence of contact resistivity for GaAs nanosheet.
Table 6.2.1: Resistance vs. RTA temperature (nanosheet sample1)

Temp (˚C) R(Ω)
375 5.9 E6
380 1.18 E6
385 4.6 E5
390 4.39E5
400 2.6 E5
375 380 385 390 395 400
0
1000
2000
3000
4000
5000
6000
Resistance (K ohm)
Temperature ('C)
85

Figure 6.2.6 SEM micrograph of GaAs nanosheet after RTA annealing at 375˚C, 380˚C, 385˚C,
390˚C and 395˚C.

Figure 6.2.7 SEM micrograph of GaAs nanosheet after RTA annealing at 375˚C, 380˚C, 385˚C,
390˚C and 395˚C.

-2.0 -1.5 -1.0 -0.5 0.0 0.5 1.0 1.5 2.0
-200
-100
0
100
200
300
400
500
600
700
I(uA)
Vbias(V)
After 375C RTA
After 380C RTA
After 385C RTA
After 390C RTA
After 400C RTA
86

Figure 6.2.8 RTA temperature dependence of contact resistivity for GaAs nanosheet.
Table 6.2.2: Resistance vs. RTA temperature (nanosheet sample 2)

Temp (˚C) R(Ω)
375 E7
380 3.9 E5
385 3.1 E5
390 3.5 E5
400 5.7 E5
375 380 385 390 395 400
0
2000
4000
6000
8000
10000
Resistance (K ohm)
Temperature (C)
87

6.3 Photocurrent and Photo-spectroscopy Measurement of GaAs Nanosheets
Photocurrent and photocurrent spectroscopy are two other important electronic
measurements. In photocurrent measurement, the sample is excited by laser and the produced
current (as the applied voltage is varied) is measured. In Figure 6.3.1.a and b the photocurrent of
GaAs nanosheet is shown. This diagram shows that GaAs nanosheets (without passivation) are
not efficient.

88

Figure 6.3.1 Photocurrent measured for GaAs nanosheets, a) sample 1, b) sample 2.
As it was mentioned in the first chapter, photocurrent spectroscopy is measured when light
at different wavelengths are radiated at the sample with p-n junction or Schottky barrier
(with/without applied voltage) and the generated currents are measured and shown as the function
of wavelengths. Photocurrent spectroscoy measurement was performed on bare GaAs nanosheets.
Figure 6.3.2 shows the corresponding dataset. The photocurrent spectroscopy was measured in the
range of 400-1700 nm using a Fianium laser. The blue line shows the generated current at zero
applied voltage and the red line shows the laser power. From 1.5 eV to 3 eV the laser power is
decreasing. Hence, the only acceptable peak observed here is at 1.5 eV. This peak corresponds to
89

the band gap of GaAs at room temperature. No sub-band gap state was observed within this
range.

Figure 6.3.2 Photocurrent Spectroscopy measured for GaAs nanosheets

1.5 2.0 2.5 3.0
0
5
10
15
20
25
30
Photocurrent (nA)
Energy (eV)
0.00
0.07
0.14
Fianium power (W)

90

6.4 Electron Beam Induced Current (EBIC) Measurement of GaAs Nanosheets
Electron Beam Induced Current (EBIC) is one of the methods which have been used for
determination of the minority carrier’s diffusion length of the semiconductors. EBIC has been
also used for visualization of dislocation and stacking faults[104]; here, the focused electron beam
creates electron−hole pairs in a p -type nanowire. Electron−hole pairs generated within one
minority carrier diffusion length of the metal−semiconductor Schottky junction will, on average,
result in a measurable current. Minority carriers generated far away from the Schottky junction
will recombine and, therefore, not contribute to the measured (EBIC) current. By spatially
mapping the EBIC current, we can determine the minority carrier diff usion length directly; figure
6.4.1 shows the EBIC setup. In Figure 6.4.2.a and b the SEM image and the corresponding EBIC
image is demosntrated for a GaAs nanowire. Here, an acceleration voltage of 5 kV was used to
create electron-hole pairs in this measurement. The EBIC signal is strongest near the Schottky
interface and gradually diminishes away from the contact. The minority carrier diff usion length,
L
diff
, is extracted by fitting with an exponenti ally decaying function (Figure 6.4.2.c). Here, we
compare the diffusion length of passivated nanowire versus a non-passivated nanowire. For the
passivated case the L
diff
is 180nm and for the non-passivated case it’s 30nm. This shows that
passiavtion layer helps the separated electrons and holes to be separated for a longer distance
before they recombine.
In Figure 6.4.3.a the SEM image of a GaAs nanosheet is demonstrated and Figures 6.4.3.b
and c show the EBIC resposnes when the polarity is changed. Fitting the exponential decay of the
EBIC diagram, we get the L
diff
to be 170nm.
91

Figure 6.4.1 EBIC setup located at CEMMA, USC
Schottky
Voltage source Current amplifier
Keithley 2400 Ithaco
Voltage output
back to machine
and form in-situ
EBIC image.
92

Figure 6.4.2 EBIC measurement on passivated vs. non-passsivated GaAs nanowire: a) EBIC
micrograph, b) SEM micrograph c) extracted and exponentially fitted signal, d) EBIC concept

p n
l
min_diff
e
+
e
-
ħ w ε
metal
e-beam
e
-
e
+
n
l
min_diff

(b) (a)
4
3
93

Figure 6.4.3 EBIC measurement of non-passsivated GaAs nanosheet: a) SEM micrograph of
GaAs nanosheet, b) EBIC micrograph of GaAs nanosheet, c) ) EBIC micrograph of GaAs
nanosheet with opposite polarity

(a)
(c)
(b)
94

Chapter 7
Enhancement of III-V Semiconductor in Photo-Catalysis

Here, we study the surface effects of TiO
2
passivation of several major III-V
semiconductors, including GaP, GaAs, and InP. Passivation of these III-V semiconductors with
TiO
2
is shown to enhance their photochemical stability and photo-conversion efficiency for solar
fuel production. We study the TiO
2
thickness dependence of the photocatalytic of these III-V
semiconductors, as well as surface composition and photoluminescence similarity and
differences of the TiO
2
-semiconductor interfaces among these species. These passivated
semiconductors are capable of sustaining continuous photo production of hydrogen from 0.5
molar aqueous sulfuric acid for over 48-hours while their non-passivated counterparts only show
substantial photo-corrosion within minutes of illumination. The thickness of TiO
2
passivation
layer has an important role in such enhancement. Although the photocatalysis performance after
passivation varies among different semiconductor, it strongly depends on the TiO
2
thickness. X-
ray photoelectron spectroscopy is done on all samples to provide further details on interface of
III-V semiconductor and TiO
2
.

95

7.1 Introduction
Artificial photosynthesis was introduced four decades ago [105]. Since then, a wide
range of materials and methods has been proposed towards facilitating and enhancing
photosynthesis in the visible wavelength range. Light absorbers, electrocatalysts, electrolytes and
separators are main factors for oxidation of water into O
2
and H
2
.[106, 107] Various material
classes have been investigated for photocatalysis including: oxides, III-V semiconductors and
chalcogenides. TiO
2
is one of the well-studied catalysts for water splitting applications due to its
chemical stability, being abundant and efficient catalysis behavior. While TiO
2
seems to be one
of the best catalysts for water splitting applications, its wide band gap (3.2eV) limits its
application to ultra violet wavelength range. On the other hand, III-V semiconductors (with
optimum band gap range of 1.1-1.7 eV) are among the most efficient materials for solar-based
water splitting; however, they are not stable once they are in contact with aqueous
electrolyte[108, 109]. Several research groups has proposed passivation schemes to protect such
III-V semiconductors. Heller et al. have studied p-type InP, where a transparent catalyst layer
was used to provide high hydrogen evolution photocurrent[110]. Shu et al. has investigated the
efficiency of GaAs nanowires for water splitting applications[111]. GaP has been studied by
many research groups for water splitting applications[112]; for instance, passivation with thick
film of metal was investigated by different groups [108, 113, 114]. Although the mentioned
passivation methods have helped improving the chemical stability, other groups did not report
any enhancement in the photo-conversion efficiencies. In our work, we observe both
enhancements in chemical stability and photo-conversion efficiencies due to the atomic layer
deposition (ALD) precursor, which we use. Shu et al. uses Tetrakis(dimethylamido)titanium(IV)
(i.e., [(CH
3
)
2
N]
4
Ti), which results in a high concentration of C and N impurities that form an
96

electrically conducting band[111]. We believe that understanding the underlying mechanisms of
photocatalytic enhancement at the interface of these III-V semiconductors and TiO
2
is important
but unanswered.
In this paper, we study the photcatalytic behavior and stability of TiO
2
-passivated p-type
InP, GaAs and GaP in order to achieve direct solar driven water splitting without corroding the
electrodes. The photo I-V characteristics are studied as a function of TiO
2
thickness. Thin films
of TiO
2
(1-5 nm) are deposited using atomic layer deposition (ALD) on p-type doped III-V
semiconductors. We study the surface chemistry of passivation III-V semiconductors using X-
ray photoelectron spectroscopy (XPS) and photoluminescence for these semiconductors
systematically as a function of TiO
2
thickness.

7.2 Photocatalysis Enhancement in III-V Semiconductor after TiO
2

Passivaiton
A very thin layer of TiO
2
(1-5 nm) is deposited on p-type doped GaAs, InP and GaP
wafers (University Wafer), using ALD (Cambridge NanoTech, Savannah) at 200 °C, using
alternating pulses of H
2
O and titanium tetrachloride (TiCl
4
). The anatase crystalline phase of the
TiO
2
passivation layer is verified using Raman spectroscopy. X-ray photoemission spectroscopy
(XPS) of these samples is taken using M-probe instrument interfaced with ESCA control
software (Service Physics, Inc.). Incident Al Kα X -ray (1486.6 eV) were directed at 35° with
respect to the sample and so is situated the analyzer. The hydrogen-evolution reduction (HER) is
measured in an aqueous solution of 0.5 molar sulfuric acid with pH of 0. The photocatalytic
97

reaction rates were measured using a three-terminal potentiostat with the prepared samples, a
Ag/AgCl electrode, and a graphite electrode functioning as the working, reference, and counter
electrodes, respectively, as shown in Figure 7.2.1.
Ohmic contacts to p-type GaP and InP is formed using a Ga-In eutectic paste (Sigma
aldrich). For p-type GaAs, Ti/Au (5/60nm) contacts are deposited on the backside of GaAs
samples using an electron-beam metal evaporator. In all three cases, the metal contact (Ga-In
paste or Ti/Au) was then connected to the external circuitry with a copper wire and coated with
epoxy cement to insulate it from the electrolytic solution, as illustrated in Figure 7.2.1.

Figure 7.2.1. Photo I-V measurement setup, 3-terminal potentiostat is used with Pt counter
electrodes and Ag/AgCl reference electrode. The illumination source is a 532 nm CW laser.

The photoelectrochemical behavior of p-type GaAs with and without TiO
2
passivation is
plotted in Figure 7.2.2a as a function of reference potential. Here, the TiO
2
-coated p-type GaAs
shows substantial enhancements for all three semiconductors. In particular, we see a clear shift of
98

these photo-I-V curves to lower overpotentials, as compared to the bare III-V sample. For GaAs,
1 nm TiO
2
is the most efficient case with over potential of 0.48 V (with respect to RHE).
Figure 7.2.2b shows the photo I-V curves of p-type GaP samples with and without TiO
2

passivation. Although GaP samples with TiO
2
passivation require lower overpotentials and
provide higher currents in comparison to the bare GaP, the enhancement is not as pronounced as
the case for GaAs samples. For the case of TiO
2
passivated GaP, 1, 3, and 5 nm coating leads to
similar required overpotential comparing to the bare case. Finally, in Figure 7.2.2c, the photo I-V
curves of the bare and TiO
2
passivated InP samples are shown. Here, the TiO
2
coated samples
require an overpotential two times lower than the corresponding the bare sample. From the I-V
curves, it is clear that thickness of the passivation does affect the performance and 1-3 nm thick
passivation layer lead to lower overpotentials. Summary of the overpotentials required for
different TiO
2
-passivated samples relative to their bare form is shown in Table 7.2.1.
The mechanism of photocatalytic enhancement created by the TiO2 passivation layer is
3-fold: Firstly, the TiO
2
deposited by ALD is n-type due to oxygen vacancies and forms a pn-
junction with the underlying p-type InP photocathode resulting in a built-in potential, which
reduces electron-hole recombination through charge separation. Secondly, the Ti
3+
active sites on
the TiO
2
lower the energy of the CO
2
-
intermediate through the formation of an intermediate
complex. Thirdly, the first few cycles of the ALD process strip off the native oxide layer, which
impedes charge transfer.

99

Table.7.2.1. Shift of overpotential of TiO
2
-passivated samples with respect to their bare
counterparts.
GaAs (V) GaP (V) InP (V)
Bare -0.06 -0.63 -0.4
1 nm TiO
2
coating 0.55 -0.58 -0.20
3 nm TiO
2
coating 0.38 -0.5 -0.2
5 nm TiO
2
coating 0.25 -0.46 -0.35
10 nm TiO
2
coating 0.15 -0.23

The photoluminescence (PL) spectra of GaAs and InP provide further insight into the
mechanism of enhancement provided by TiO
2
passivation. Figure 7.2.3a shows GaAs with and
without TiO
2
passivation. We observe that TiO
2
passivation does not enhance the PL, indicating
that the TiO
2
passivation layer is introducing surface states (likely Ti
3+
). GaAs with a 3 nm TiO
2

passivation-layer shows the lowest intensity, indicating highest non-radiative recombination. In
Figure 7.2.3b, the PL spectra of InP samples are shown. Again, we find that the TiO
2
passivation
layer decreases the PL intensity, due to the introduction of catalytically active surface states that
act as recombination centers. For GaP samples due to their indirect band gap, no PL response
was observed. Both the PL spectra and the photo I-V measurements indicate/provide a consistent
picture in which higher non-radiative recombination leads to higher participation of carriers in
the photochemical process. However, two other processes, formation of p-n junction and surface
defects, also play important roles.
In order to further investigate the behavior of the TiO
2
passivation layer, we perform XPS
spectroscopy. As the thickness of the TiO
2
layer increases the intensity of both 458.9 eV (Ti
3p
3/2
) and 464 eV (Ti 2p
1/2
) increases, which is expected (Figures 7.2.4b, 4d and 4f).[115-117]
While the relative intensity of the peaks change with TiO
2
thickness, there is no shift in the peak
position and, hence, there is no change in composition. The oxygen peak however, does undergo
100

a shift in position. For the bare GaAs, the oxygen peak is centered at 532 eV, corresponding to
the native oxide.[118, 119] As the TiO
2
thickness in increased, a second peak appears at 530 eV
(O 1s due to TiO
2
).[120, 121] The evolution of the two oxygen peaks with TiO
2
thickness is
observed (Figure 4c) and as the thickness of the TiO
2
increase the oxygen peak due to TiO
2

increase and dominates the O peak at 532 eV which is due to native oxides or TiO
x
species. In
case of GaP, the oxygen peak at 531.5 eV (O 1s due to GaPO
4
) is dominant in the bare case
[122]. The oxygen peak due to TiO
2
(530 eV) shows up for samples with 1 nm TiO
2
on top and
become dominant for the case of 5 nm TiO
2
passivation. For InP samples, the oxygen peak due
to metal oxide formation is at 531.4 eV this peak eventually become less dominant when TiO
2

passivation is deposited [123-125] (Figure 7.2.4e). The behavior of the Ti peaks (Figures 7.2.4b,
4d, and 4f) is almost the same for GaP and InP samples (no position shift and only increase in
intensity).Our data shows that, for the samples with mixture of oxide peak and TiOx at the
interface more enhanced photo-I-V is obtained.
7.3 Conclusion
In conclusion, we have shown that passivation of III-V photocatalysts with TiO
2
not only
protects these materials against corrosion, but also enhances their photocatalytic properties. The
optimum thickness which would lead to lowest required over potential varies from one
semiconductor to another. However, the range is bellow 10nm in thickness. Our surface
measurements show that deposition of very thin layers of TiO
2
leads to formation of mixed TiO
2
,
TiO
x
which can be hold responsible to constructive trapping of charges and enhancement in
photocatalytic activities. The possible mechanism is formation of a pn junction at the
semiconductor-TiO
2
interface, which provides a photovoltage that acts like an externally applied
101

overpotential. We also believe that oxygen vacancies in the TiO
2
ALD layer leads to formation
of delocalized midgap states which help the charge transfer process.

102

-0.4 -0.2 0.0 0.2 0.4 0.6
-100
-80
-60
-40
-20
0
Current Density (mA/cm
2
)
Potential vs. Ag/AgCl (V)

Dark
Bare GaAs
GaAs w/1nm TiO
2
GaAs w/3nm TiO
2
GaAs w/5nm TiO
2
GaAs w/10nm TiO
2

-1.0 -0.5
-0.1
0.0
0.1
Dark
Bare GaP
GaP w/1nm TiO
2
GaP w/3nm TiO
2
GaP w/5nm TiO
2
GaP w/10nm TiO
2
Potential vs. Ag/AgCl (V)

Current Density (mA/cm
2
)

-0.8 -0.6 -0.4 -0.2 0.0
-60
-50
-40
-30
-20
-10
0
10
Potential vs. Ag/AgCl (V)
Current Density (mA/cm
2
)

Dark
Bare InP
InP w/1nm TiO
2
InP w/3nm TiO
2
InP w/5nm TiO
2

Figure.7.2.2. Photocurrent diagrams of III-V with and without TiO
2
passivation: a) GaAs, b) GaP and
c) InP
GaAs
GaP
InP
(a)
(b)
(c)
103

780 800 820 840 860 880 900 920
PL Intensity
Wavelength(nm)
Bare
GaAs with 1nm TiO2
GaAs with 3nm TiO2
GaAs with 5nm TiO2
GaAs with 10nm TiO2

800 850 900 950 1000
PL Intensity
Wavelength (nm)
Bare InP
InP with 1nm TiO2
InP with 3nm TiO2
InP with 5nm TiO2

Figure 7.2.3. Photoluminescence spectra of III-V with and without TiO
2
passivation: a) GaAs, b) InP

(a)
(b)
104

528 530 532 534 536
5nm TiO2

Binding Energy (eV)

Bare 1 nm TiO2
3 nm TiO2 10nm TiO2

GaAs with 10nm TiO2

GaAs with 5nm TiO2
GaAs with 3nm TiO2
GaAs with 1nm TiO2
Bare GaAs

456 460 464 468
Bare GaAs

GaAs with 1nm TiO
2

GaAs with 3nm TiO
2

GaAs with 5nm TiO
2
GaAs with 10nm TiO
2
Ti 2p1/2
Ti 2p3/2
Binding Energy (eV)

528 530 532 534 536
Bare GaP

GaP with 1nm TiO2

GaP with 3nm TiO2

GaP with 5nm TiO2

Bare
1nm TiO2 3nm TiO2 5nm TiO2 10nm TiO2

Binding Energy (eV)
GaP with 10nm TiO2
456 460 464 468
Binding Energy (eV)

Bare GaP
GaP with 1nm TiO
2
GaP with 3nm TiO
2
GaP with 5nm TiO
2

Ti 2p1/2
Ti 2p3/2
GaP with 10nm TiO
2

(c) (d)
(a) (b)
105

528 530 532 534 536

Bare
Bare
InP with 1nm TiO2

1nm TiO2
InP with 3nm TiO2

3nm TiO2 10nm TiO2

5nm TiO2
InP with 5nm TiO2
InP with 10nm TiO2

456 460 464 468

Bare InP

InP with 1nm TiO
2
Binding Energy (eV)

InP with 3nm TiO
2

InP with 5nm TiO
2

InP with 10nm TiO
2
Ti 2p3/2
Ti 2p1/2

Figure 7.2.4. XPS measurement results of III-V with and without TiO
2
passivation: a) GaAs, b)GaP, c)
InP

(e) (f)
106

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Abstract (if available)

Abstract In this thesis, we study the optical and electrical properties of one-dimensional and two-dimensional nanostructures used for different energy applications. These applications would include solar cells, LEDs, nanolasers and water splitting cells. Different characterization methodologies are used here

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University of Southern California Dissertations and Theses

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Creator Arab, Shermin (author)

Core Title Optical and electrical characterization of one-dimensional (1D) and two-dimensional (2D) nanostructures

School Viterbi School of Engineering

Degree Doctor of Philosophy

Degree Program Electrical Engineering

Publication Date 07/28/2015

Defense Date 03/02/2015

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

Tag contacts,doping,Fabry-Perot microcavity,GaAs,nanosheets,nanowires,OAI-PMH Harvest,optoelectronic,passivation,photoluminescence

Format application/pdf (imt)

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Cronin, Stephen B. (committee chair), Nakano, Aiichiro (committee member), Wu, Wei (committee member)

Creator Email arab@usc.edu,shermin.arab@gmail.com

Permanent Link (DOI) https://doi.org/10.25549/usctheses-c3-610306

Unique identifier UC11298793

Identifier etd-ArabShermi-3694.pdf (filename),usctheses-c3-610306 (legacy record id)

Legacy Identifier etd-ArabShermi-3694.pdf

Dmrecord 610306

Document Type Dissertation

Format application/pdf (imt)

Rights Arab, Shermin

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