MEE growth and characterization of I-III-VI2
chalcopyrite ternary compound thin films
(Applicant Name)
Sathiabama
T THIRUGNANA
( Department of Electrical Engineering and Bioscience, Research on
Semiconductor Engineering )
April , 2015
MEE growth and characterization of I-III-VI2
chalcopyrite ternary compound thin films
DOCTORAL THESIS
Department of Electrical Engineering and Bioscience
Graduate School of Advanced Science and Engineering
Waseda University
Sathiabama
T THIRUGNANA
( Department of Electrical Engineering and Bioscience, Research on
Semiconductor Engineering )
April , 2015
Contents
1. Introduction
1
2. CGS, CIS chalcopyrite materials
2.1
Crystal structure
17
2.2
Phase diagram
19
2.3
Band structure
21
2.4
Thermal expansion coefficient
23
3. MBE, MEE growth of CGS and CIS
3.1
MBE growth method
26
3.2
MEE growth method
30
3.3
Review of CGS and CIS growth by MBE
32
4. CGS, CIS growth and characterization
4.1
Motivation of CGS growth on GaAs (001)
44
4.2
Growth of CGS thin films
46
4.3
Characterization of CGS thin films
51
4.4
Conclusion/summary of CGS thin films
59
4.5
Motivation of CIS growth on GaAs (001)
60
4.6
Growth of CIS thin films
60
4.7
Characterization of CIS thin films
61
4.8
Conclusion/summary of CIS thin films
66
5. CGS/CIS double heterostructures
5.1
Motivation of CGS/CIS Double Heterostructures
i
69
5.2
Preparation of CGS/CIS Double Heterostructures
73
5.3
Properties of CGS/CIS Double Heterostructures
74
5.4
Summary of CGS/CIS Double Heterostructures
83
6. Doped CGS analysis
6.1
Motivation of Doped CGS
86
6.2
Preparation of Doped CGS
88
6.3
Optical and electrical properties of Doped CGS
89
6.4
Conclusion/Summary of Doped CGS
99
7. CGS/CIS single quantum wells
7.1
Motivation of CGS/CIS Single Quantum Wells
102
7.2
Preparation of CGS/CIS Single Quantum Wells
104
7.3
Properties of CGS/CIS Single Quantum Wells
105
8. Summary
115
Acknowledgments
119
List of publications and conference presentations
121
ii
Chapter 1
Research Overview
This chapter gives a thorough research introduction, background and motivation
regarding chalcopyrite material system. The CuIn 1-xGaxSe2 (CIGS) material system
is widely known for the solar cell application; however its application can be
broaden and utilized for other active opto-electronic devices.
1.1 Research Introduction
Research on I-III-VI2 chalcopyrite ternary and quaternary compound materials
had started since 1960s [1, 2]. They were all mainly computational studies about
photovoltaic effect that predicts the characteristics of these materials which would
operate with an optimum efficiency as solar cell devices [3]. Eventually,
polycrystalline thin films of CIGS became popular as solar cell absorber layer. There
are several growing techniques such as a chemical vapor transport method using
iodine as a transport agent [4], vacuum evaporation [5], non-vacuum deposition [6],
molecular beam epitaxy [7] and three-stage growth method [8]. As it is defined by
solar cell industries and/or National Renewable Energy Laboratory (NREL), solar
cells are classified into four main generations. The first generation includes Si and
Ge crystalline cells. The second generation is thin-film technology which includes
CIGS solar cells. The third generation is multi-junction cells and concentrated
photovoltaic (PV). And the fourth generation consists of emerging PV which
includes dye-sensitized cells. Nevertheless, the existence of these chalcopyrite
materials, CuInSe2 (CIS) and CuGaSe2 (CGS) gave hope for the second generation
1
thin film solar cells industry. However, nothing has been changed to date where
these materials are mostly, still in use only for the solar cell application.
Hence, in this thesis, we will mainly discuss about CGS and CIS structural,
optical properties and other possible applications for future opto-electronic devices.
In most cases, band gap energy (Eg) of the semiconductor plays an important role on
the device efficiency. By tuning the elemental compositions of CGS and CIS, its Eg
can be varied between 1.04 and 1.68 eV at room temperature. CGS and CIS have
high-density defects, where it is classified as ordered-defect-compound (ODC). Even
though chalcopyrite thin film solar cell materials have high defect densities, they
are able to produce an energy conversion efficiency rate (
) of about 19% [9] due to
its high absorption coefficient near band-edge (
104-105 cm-1). We have
attempted the growth of CGS/CIS heterostructures and a single quantum well
structure for the first time to increase the solar cell efficiency as well as to seek for
other device applications. Thus, we have confirmed that these chalcopyrite
materials have other possible applications in future optical and electronic devices.
1.2 Research Background
Conventional natural energy resources such as petroleum, natural gas, coal, etc.,
are exhausting in a fast pace. Energy crisis since early of 1970s has boosted up
research and development on solar cells or PV which is a form of renewable energy,
all around the world. This global energy crisis is due to rapid industrialization and
population growth. One of the solutions for this crisis is the renewable energies like
solar energy and wind energy. As these energies are ample on earth, researchers
and engineers are focusing on the fabrication of devices that could extract these
2
free of charge energies into usable and storable form of energy. The radiation
which travels from the Sun to the Earth is inexhaustible and less pollution. The sun
radiation is about 20,000 times more than the current energy that we require. If we
could convert even a fraction of this radiation into usable energy, our world would
become a better place to live on. The spectral distribution of the solar radiation with
different air masses (AM) [10, 11] is an important factor for the
. In general AM is
defined as the atmosphere which contains dust particles, air molecules, CO 2 etc.;
where solar radiation passes through this atmosphere before reaching earth. AM1
(nearer to equator) and AM1.5 are the standard solar spectra that are commonly
used to specify the energy conversion efficiency of terrestrial solar cells. Solar
spectral irradiance of AM1.5 is equivalent to 908 Wm-2.
Currently, the major stakeholder in the realm of photovoltaic industry is Si, with
amorphous, polycrystalline and single crystal Si based solar cells. Hence, in general
a Si solar cell is made out of about 300
m thickness of Si which is required due to
its low absorption coefficient. As the first generation solar cell candidate for
terrestrial applications, Si solar cell has reached its maximum efficiency of about
26 %. On the other hand, this indirect semiconductor Si is also in high demand for
other electronic devices. These factors have motivated researchers and engineers to
explore other new materials and processes. Thus, there are many other types of
solar cells; plastic/polymer based solar cells, multi-junction GaAs based thin film
solar cells and also chalcopyrite CIGS based thin film solar cells. This alternative,
ternary semiconductor compound CGS, CIS and CIGS with chalcopyrite structures
are of technological interest for solar cell absorber layer as a second generation thin
film solar cell. Here, we will focus on properties of chalcopyrite materials CGS and
3
CIS. In reality, solar cells of chalcopyrite material are often heterostructured with
CdS and ZnO. This is because by nature these chalcopyrite materials have p-type
conductivity due to high concentration of copper vacancy (VCu) [12].
1.3 Research Motivation
These chalcopyrite materials (CIS and CGS) possess great attractive
characteristics for opto-electronic devices such as:
High electron effective mass 0.14 m 0 and low refractive index 2.14 for CGS
High absorption coefficient
High hole mobility
105 cm-1
200 cm2/Vs (GaAs
400 cm2/Vs)
Less temperature (between 10 and 300 K) dependent of band gap energy Eg
(
Eg 10-50 meV [17]) compared to GaAs (
However,
these
materials
also
exhibit
Eg 100 meV [18])
unfavorable
characteristics
for
homojunction devices such as:
High concentration of acceptor type defects like VCu, interstitial and
anti-site defects which makes it strong p-type semiconductor
Critical Doping Limit Rules especially in equilibrium state [12]. Properties
of the host material, mainly controls its complex cation-vacancy and
anion-vacancy
One of our objectives in this dissertation is to grow n-type chalcopyrite material
with CGS and CIS. However, doping in these materials showed a peculiar
characteristic. Our experiment results show that up to a certain level of dopant
quantity, the hole concentration decreases. However, when the doping level is
increased further, the hole concentration increases. Thus, this behavior has
4
motivated us to introduce a new approach which is the modulation doping . While
this modulation doping was implemented in the GaAs/AlGaAs material system, it
has never been applied to the chalcopyrite material system. Furthermore,
theoretically, at the CGS/CIS interface, almost all the band gap energy difference
distributed to the conduction band-
[13]. The value of
-
100 meV. This is a beneficial characteristic for electron confinement and
quantization. Such a large difference in band-edge discontinuity provides
interesting applications of quantum well system. Furthermore, CGS has high
excitonic binding energy of approximately 20 meV because of its relatively large
electron effective mass of 0.14 m0 and low refractive index of 2.14. In addition, the
grown CGS and CIS thin films possess high hole mobility about 220 cm2/Vs and 200
cm2/Vs, respectively. These notable distinctions of CGS and CIS encourage and
motivate our research for the realization of new opto-electronic devices.
1.4 Characterization methods
This section consists of a brief introduction of all characterization techniques
that are mainly used in this research dissertation. Their basic principles and setup
diagrams are presented here.
1.4.1
Reflection High-Energy Electron Diffraction (RHEED)
RHEED is an in-situ surface analysis probing tool during epitaxial growths
using non-equilibrium MBE. RHEED observation can help us to understand the
basic mechanism of epitaxial growth in a qualitative manner and also to determine
5
the growth rate. Usually, the intensity of diffracted electron beams displays an
obvious damping oscillation during growth. In general, MBE systems today include
an electron gun and a phosphor screen for displaying RHEED patterns during
growth. Thus, electron beam of 20 keV strikes on crystal surface at a glancing
incidence of about 1 - 2
to produce a diffraction pattern on the phosphor screen as
shown in Fig. 1.4.1. Nevertheless, in this chalcopyrite material system, RHEED
observation analysis has never been reported. Hence, we have focused on the
amplitude change/difference during growth of CGS and CIS.
Figure 1.4.1 Schematic illustration of RHEED observation system
1.4.2
X-ray Diffraction (XRD)
XRD is a versatile technique which is generally used to identify or determine
crystal phases, thickness of thin films, phase composition, lattice constant and
growth orientation of semiconductors. In this thesis all samples were grown on
GaAs (001) and the crystal qualities of these thin films were determined by XRD
measurements. In this measurement system, a Ge two crystal monochromatic X-ray
beam impinges on the surface of sample. Depending on the incident angle, x-ray
6
wavelength and atomic spacing in the sample, constructive interface will take place
in the diffracted beam. This phenomenon is well known and explained by the
Bragg s relation as equation (1.4.2.1),
2d sin
where
n
(1.4.2.1)
is an integer ( th reflection condition),
is the wavelength of the
electromagnetic wave (Cu target), d is the crystal lattice spacing and
angle of incidence. Mainly 2
-
scan and
is the
scan were performed to characterize
the crystal quality. XRD measurements of all samples discussed in this thesis were
performed using a Rigaku diffractometer, model RINT (TTR III). The diffractometer
was equipped with 2 Germanium crystals and Copper cathodes emitting CuK
wavelength of 1.5405
1.4.3
and operated at 40kV
1
at a
30mA.
Photoluminescence (PL)
PL spectroscopy is defined as a spontaneous emission of light from a
semiconductor material that was absorbed initially and then excited optically. In
other words, we can measure physical properties by using photons to induce
electronic states in a material and then analyze the emitted light which corresponds
to its relaxed state. PL is a very efficient, contactless and non-destructive method to
probe the opto-electronic properties of semiconductor material. This measurement
only requires a small piece (< 5
5 mm) of sample. We have first measured CGS
thin films using Ar+ ion laser at a wavelength of 488 nm as excitation source. The
excitations of these CGS thin films were detected using Si CCD. On the other hand,
as shown in Fig. 1.4.3.1, PL measurements were performed using CW solid state
laser and InGaAs CCD detector for the double heterostructures and single quantum
7
wells (SQWs). As the energy band gap (Eg) of samples (in this thesis) ranges
between 1.04 - 1.68 eV, laser with
= 532 nm is sufficient for photo-excitement of
CGS and CIS. Cooled Andor InGaAs CCD detector (Quantum efficiency shown in
Fig. 1.4.3.2) is able to detect luminescence from both samples. We have mainly used
PL to analyze the recombination processes, impurity levels, band gap determination
and crystal quality.
Figure 1.4.3.1 Schematic illustration of PL measurement setup
Figure 1.4.3.2 Quantum efficiency of the Andor InGaAs CCD detector
8
1.4.4
Hall effect measurement
The fundamental physical principle in Hall measurement is the Lorentz force.
Lorentz force is explained by the equation (1.4.4.1).
F
q( E v B)
(1.4.4.1)
In electromagnetism, the Lorentz force
force on a charge
is the addition of electric
and magnetic
which has a velocity . Hall effect is the multiplication of a
voltage difference (Hall voltage,
) across a semiconductor material, transverse to
an electric current in the material and a magnetic field which is perpendicular to
the current. The
is the current,
is measured and its magnitude is as equation (1.4.4.2), where
is electron charge,
is electron/hole concentration and
is the
sample thickness. In the Hall effect measurement, formation of good ohmic contact
to semiconductors is very important. In the case of GaAs, Cr-Au or Au-Zn is used for
p-type material while Au-Ge-Ni is used for n-type material. We have investigated
various metals including those used for GaAs and determined that Au-Ge-Ni gives
lowest resistivity ohmic contact to CGS and CIS. In order to determine the mobility
and electron/hole concentration
, we have used the well-known symmetric
4-contact Van der Pauw technique for Hall measurements.
n
IB
qd | VH |
(1.4.4.2)
d | VH |
IB
(1.4.4.3)
9
1.4.5
Other characterization methods
We have implemented many characterization techniques other than those
explained previously. Structural properties and film thickness were analyzed with
Field Effect Scanning Electron Microscopy (FE-SEM), Transmission Electron
Microscopy (TEM). Material compositions were determined by Electron Probe
X-Ray Microanalysis (EPMA) and Energy-Dispersive X-ray Spectroscopy (EDS).
1.5 Outline of this dissertation
This research thesis presents a study on the growth and characteristics of
chalcopyrite materials such as CGS, CIS and their mixed compounds CIGS. The
dissertation consists of 8 chapters as described in the following. Chapter 1 gives a
thorough introduction to our research motivation and background of this material.
The CIGS material system is of great interest for solar cell applications. There are
many techniques to grow CIGS thin films. Conventional CIGS solar cells are grown
on soda-lime glass (SLG) using the three-stage-process evaporation method.
Majority of these thin films are polycrystalline with a range of grain sizes
depending on the annealing temperature. Then it will be heterojunctioned using
CdS for solar cell application. In chapter 2, to further understand the growth recipes
and its challenges, it is important to provide some theoretical and fundamental
details about these ternary material systems which include their crystal structure
and thermal expansion coefficient. We will also show a review on its phase diagram
to identify other common secondary phases and their roles. Chapter 3 is devoted to
the growth methodology; molecular beam epitaxy (MBE) and migration enhanced
epitaxy (MEE). In this thesis we are probing the physics behind these material
10
systems and their possibility for new device applications with high-quality
chalcopyrite single crystals using MEE. In this chapter a detailed review is done to
explain the difference between MBE and MEE growth. Here we have surveyed
literatures of single crystal CGS and CIS grown by MBE, which shows that their
hole mobility are as low as about 5 cm2/Vs [14] and 10 cm 2/Vs [15], respectively.
Thus, in this research we have confirmed that MEE growth is essential to increase
the single crystal quality which contributes to higher hole mobility for these
chalcopyrite materials. Furthermore, we have realized that its application can be
broadened and utilized for other active optical devices.
To understand and characterize each material individually, we have separately
optimized the growth conditions of undoped CGS and CIS. Thus, we have grown
-insulating GaAs (001)
substrate. We have employed GaAs as a substrate because the lattice constants of
CIS and CGS (5.784
and 5.614 , respectively) are close to that of GaAs (5.653
).
CGS and CIS thin films are characterized by means of in-situ RHEED, FE-SEM,
XRD, EPMA, PL and Hall measurements. The Hall measurements on these thin
films are performed by Van der Pauw method. Hence, in chapter 4, we have
assessed the thin film CGS and CIS crystal qualities, optical properties and
electrical properties. PL analysis shows that the emission energies of CGS and CIS
have very small temperature dependence compared to that of GaAs. We have
confirmed that the optical and electrical characteristics of CGS and CIS are rather
unique and attractive. Therefore, this material system could be applied to variety of
semiconductor active devices other than solar cells. Characteristic features of this
system include a wide range of direct band gap energies in CIGS (1.04 ~1.68 eV).
11
However, to date there are no heterostructures grown with CGS/CIS material
systems.
Therefore, prior to the CGS/CIS SQWs growth and characterization, it was
necessary to perform a thorough study on CGS/CIS double heterostructure. Hence,
chapter 5 shows a summary of the growth and results of the structural and
interfacial properties of CGS/CIS double heterostuctures. RHEED, XRD, TEM,
secondary ion mass spectroscopy (SIMS) and PL are used to discuss the structural
and interfacial properties in further details. Here, we have emphasized that the
structural properties of CGS/CIS interface can be analyzed by means of RHEED.
The peak intensity of the RHEED specular beam decreases while preserving the
initial RHEED oscillation amplitude for several MLs. This decrease could be due to
the relaxation process of lattice strain when CIS is grown on CGS. However when
CGS is grown on CIS, the amplitude of the RHEED specular beam intensity
oscillation increases for more than 70 MLs. This increasing in RHEED amplitude
intensity oscillation is not ascribed to lattice strain but also to the intermixing of
Gallium (Ga) and Indium (In) atoms. As the bonding energies of Ga-Se and In-Se
are different and qualitatively In-Se bonding is weaker [16]. In other words, Ga
atoms replaces In atoms throughout the growth. For the first time in these material
systems, we have confirmed the atomic replacement of In atoms by Ga atoms using
RHEED specular beam intensity oscillation amplitude. Hence, we have concluded
that growth temperature and sequence (CGS on CIS or CIS on CGS) are critical
parameters when we hybrid CGS and CIS and that, using RHEED and XRD
analysis we confirm that low temperature growth is required.
In order to explore new applications of CIS, CGS and CIGS, we have
12
investigated
fundamental
characteristics
of
these
materials
and
their
heterostructures. Even though chalcopyrite material has high-density defects and
being classified as an ODC, as explained in chapter 3, MEE low temperature growth
method is implemented to control/eliminate the defect formation process. One of the
most serious problems of this material system to develop new devices such as those
done in III-V compound semiconductors is the difficulty in growing n-type materials.
This problem has been a major research theme for many research groups.
Another problem, which we have found in this dissertation, is the difficulty to
grow sharp CGS/CIS heterostructures mainly due to Ga and In atomic replacement.
This phenomenon is qualitatively determined for the first time in these material
systems.
Chapter 6 summarizes the growth and characterization of doped CGS. Optical
and electrical properties of Si, Zn and Ge doped CGS are discussed. Although we
have obtained highly resistive material with about 105 106
on doped samples, n-type conductivity has never been achieved. However, PL
results show several dominant donor acceptor pair emissions which are attributed
to impurities such as VCu (acceptor - 80 meV), Si (donor - 30 meV) and Zn (donor - 50
meV).
Here, we propose a new approach which is the modulation-doped structure
using CGS/CIS single quantum well (SQW) and superlattices (SLs). In the SL
structure, quantum well and barrier would be CIS and CGS, respectively.
c
in CGS/CIS quantum wells, electrons in the deep donors
of barriers are expected to be activated into the conduction band of CIS, which
would contribute to n-type conductivity. As summarized in chapter 6, our study of
13
doping in CGS with Si and Zn produce donor levels with depths of 30-50 meV. If the
modulation doped structure mentioned above works successfully, and n-type
conductivity is achieved, a new application field will be opened in the chalcopyrite
materials. It can be also applicable to solar cells which were our initial objective. In
this case, pn junction solar cells without lattice mismatch at the interface unlike,
the conventional CGS/CdS can be achieved. If interface issues are eliminated,
dramatic increase in the solar cell efficiency can be achieved.
In chapter 7, we have described the growth and characterization of CGS/CIS
SQW which is one step before the growth of CGS/CIS SLs. We have successfully
grown CGS/CIS SQW and first time ever reported. In order to prevent intermixing
between Ga and In atoms, a special temperature control process is required while
growing a CGS barrier on a CIS well. As predicted, the excitonic luminescence from
10nm CIS well at 1.11eV shows a quantization level of about 60 meV and its high
intensity emission compared to CGS emission, which confirms that almost all free
electrons are injected into the CIS well. This has been also confirmed by PL power
excitation measurements, where no significant energy shift was observed. Moreover,
PL results prove that there is no obvious temperature effect (only 10 meV) on the
CGS/CIS SQW which is an excellent character for optical devices.
Lastly, chapter 8 presents the summary of this research with results evaluation
and provides an outlook of future research and application possibilities.
14
References
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Schrott, Thin Solid Films 517 (2009) 2158
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16
Chapter 2
Basics of Chalcopyrite Materials
In this chapter, to further understand the chalcopyrite material growth
conditions and its challenges, it is important to comprehend some theoretical and
fundamental details about these ternary material systems which include their
crystal structure, band structure and thermal expansion coefficient. These
compounds have tetrahedral structures; which means each atom is surrounded by
four neighboring atoms. We will also show a review on its phase diagram to identify
other common secondary phases and their roles.
2.1 Chalcopyrite crystal structure
Among I-III-VI2 materials, CGS and CIS have been most intensively researched
ternary chalcopyrite materials. Figure 2.1.1 shows a schematic representation
zinc-blende structure and an ideal CGS or CIS chalcopyrite structure. As compared
in Fig. 2.1.1, chalcopyrite or also known as superlattice structure of zinc-blende is
constructed by two zinc-blende lattices. In other words, two Zn atoms are
transmuted into Cu and Ga as derived in equation 2.1.1.
30
ZnSe
30
ZnSe
30
Zn2 Se2
29
Cu 31GaSe2
(2.1.1)
Here, it is obviously shown that the presence of two different cations makes
chalcopyrite materials different from those of their isoelectric zinc-blende
compounds. Lattice parameters
and
are shown in Fig. 2.1.1; the ratio
of CGS
and CIS are 1.965 and 2.008, respectively [1]. These ternary compounds belong to
17
the space group
12
D 2d
with eight atoms per unit cell. However, zinc-blende
structure only has 4 atoms per unit cell.
Figure 2.1.1 Zinc-blende (left) and chalcopyrite (right) crystal structure
In epitaxial growth monolayers (MLs) is a unit that is often used. One unit cell
of CGS or CIS consists of 4 MLs. Each ML is constructed by cations (metal atoms,
Cu and Ga/In), anions (Se) and cations (metal atoms, Cu and Ga/In). Nevertheless,
deformations always take place in a unit cell scale because Cu-Se and Ga-Se bonds
have different strengths. In usual semiconductor materials, s and p electron
configuration is responsible for bonding. However, semiconductor hybridization
with transition metal elements such as, Cu occupies the d-electron shells for
bonding. This is the main reason for the different bonding strengths. Apparently, Se
atoms possess stronger bonding energy with the two nearest Cu atoms [2].
18
2.2 Chalcopyrite phase diagram
In general, a phase diagram which follows the Gibbs phase rule is in a stable
mode. Furthermore, there are possibilities of two phases of the component or
material system to coexist at equilibrium. A phase diagram displays regions of
potential space and various phases of the material system. Basic potential space
variables are pressure, temperature and composition. A single-component phase
diagram which involves a solid, liquid and gaseous phase is common for most
materials. Phase diagram can also be defined as a state where certain atomic
percentage or mole percentage of a respective material composition at a certain
temperature where it possesses a definite crystal structure [3].
Chemical thermodynamics of the material system plays an important role in
phase diagram as it is usually discussed at equilibrium state, where a phase can be
in solid, liquid or even gaseous form and the existence of one or more distinct crystal
phase is possible. Major potentials that define a system would be thermal,
mechanical and chemical. A simplified Gibbs-Duhem equation [4] is shown below for
ternary compound material system:
F Ph C 2 5
(2.2.1)
is the number of degrees of freedom,
and
is the number of phases at equilibrium
is the number of components in the system (for chalcopyrite
is 3).
The phase diagrams of CGS and CIS are rather complicated as they have many
different crystal structures, such as stannite and sphalerite structures. As shown in
Fig. 2.2.1 (temperature versus composition), we can understand that at equilibrium
CGS and CIS crystallizes as a chalcopyrite structure at different temperatures [5, 6].
In order to have a 1-1-2-stoichiometric CGS, the tolerance in growth temperature
19
and material compositions are very narrow. However, our MBE/MEE growth
method is a non-equilibrium one and beneficial for these type of materials.
Figure 2.2.1 CuGaSe2 CuInSe2 phase diagram for temperatures > 700
at 1 atm
CGS and CIS phase diagrams also confirm that there are binary phases such as
Cu2Se, Ga2Se3 and In2Se3 [5, 6] at equilibrium. In order to avoid these binary phases
and grow only 1-1-2-stoichiometric CGS and CIS, Cu/Ga and Cu/In ratio has to be
controlled between 1.00-0.75 and 1.00-0.80, respectively. Figure 2.2.1 shows that at
equilibrium CGS and CIS growth temperatures have to be also precisely controlled
between 800 - 1050
and 700 - 800
, respectively.
20
2.3 Chalcopyrite band structure
Similar to zinc-blende structure, chalcopyrite materials are chemically bonded
through hybrid sp3-orbitals. Conduction band of CGS is built from s-like state
having symmetry at
. On the other hand, valence band maximum is a p-like state.
However, the uppermost valence bands are influenced by the transition metal Cu
with a d-like state [7]. Figure 2.3.1 shows the electronic band structure of CGS
where transitions between three valence sub-bands maximum and conduction band
minimum are labeled as EA, EB and EC. This split off is due to the
which causes the chalcopyrite material to be under uniaxial stress.
Figure 2.3.1 Electronic band structure of CuGaSe2
21
hybridization
Figure 2.3.2 shows the electronic band structure of ternary compound CIS [7].
The indications
and
show that the optical transition is allowed for the light
polarization, in perpendicular and parallel to the c-axis, respectively. Similar to
CGS and other tetragonal structures, CIS valence band is split into three bands,
due to the influence of spin-orbit interaction
so and crystal field effect
cf.
Obvious electronic band structure difference between CGS and CIS is the crossing
of topmost valence band in spin orbit region which is due to crossing of
in the cf region. In most cases of CIS [6], topmost
6
and
7
5
and
4
valence bands are
nearly degenerate. Hence, the transition between the degenerate bands and
conduction band is often denoted as EAEB transition.
so
Figure 2.3.2 Electronic band structure of CuInSe2
22
2.4 Thermal expansion coefficient comparing with GaAs
The temperature dependence of the thermal expansion coefficient (
GaAs is shown in Fig. 2.4.1 [8].
Obvious decrease in
(T)) of
(T) between 60 and 300 K increases monotonically.
(T) is seen at temperatures below 50 K. Furthermore, the
major derivatives here are material elastic stiffness constants and phonon
frequencies. Similar phenomenon is observed even in InP compound semiconductor.
The phonon dispersion curves of both GaAs and InP at their covalent-metallic
transitions are confirmed in numerical values [8].
Figure 2.4.1 The thermal expansion coefficient (
23
(T)) of GaAs
It was reported that the dependence of tetragonal distortion (
ternary compounds including CGS [9], where
,
and
) and d /dT of
are vertical and
horizontal lattice constants, respectively. Here, the thermal expansion coefficients
of CGS are measured between 300 and 670 K. The thermal expansion coefficient in
a-axis (
and 5.2
a)
and thermal expansion coefficient in c-axis (
c)
of CGS are 13.1 10-6K-1
10-6K-1, respectively. Many measurement attempts have been made to
probe the root cause of this distortion, especially, regarding covalent and ionic
binding forces, the two kinds of bonds in these compound materials. These
measurement results confirmed that the tetragonal distortion increases with
increasing temperature. However, there are no literatures to confirm the
(T) of
CGS below 50 K.
P. Deus et al. have reported that the temperature dependence of the lattice
constant
and
of chalcopyrite compound CIS [10]. The measurements were
performed between 30 and 300 K. As the temperature increases till 60 K, lattice
constant
decreases. For temperatures between 60 and 300 K, the lattice constant
increases. Similar results are observed for the lattice constant
temperature increases till 80 K, the lattice constant
between 80 and 300 K, the lattice constant
of CIS; as the
decreases. For temperatures
increases. This phenomenon is similar
to GaAs compound semiconductor. However for a better understanding, it is
important to confirm the
(T) in the low temperature regions to probe the lattice
dynamics related to these chalcopyrite materials. We can confirm that the critical
temperatures (temperature when
constant
(T) becomes negative) are different for lattice
and . This difference between
and
clearly proves the root cause of tetragonal distortion,
24
at a respective temperature,
.
References
[1] H. W. Spiess, U. Haeberlen, G. Brandt, A. Rauber, and J. Schneider, Phys.
Status Solidi B 62, (1974) 183
[2] M. Robbins, J. C. Phillips, V. G. Lambrecht Jr., J. Phys. Chem. Solids 34
(1973) 1205-1209
[3] Svein Stolen and Tor Grande, John Wiley & Sons, (2004) Ltd. ISBN:
0-471-49230-2
[4] Gibbs, J. W., Scientific Papers (Dover, New York, 1961)
[5] J. C. Mikkelsen, J. Electron. Mater. 10 (1981) pp. 541-558
[6] U. C.Boehnke and G. Kuhn : J. Material Science 22 (1987) 1641
[7] S. Chichibu, T. Mizutani, K. Murakami, T. Shioda, T. Kurafuji, H.
Nakanishi, S. Niki, P. J. Fons and A. Yamada, J. Appl. Phys., 83, 7 (1998)
3678-3689
[8] T. Soma, J. Satoh and H. Matsuo, Solid State Commun. 42, 12 (1982),
889-892
[9] H. -G. Bruhl and H. Neumann, Solid State Commun. 34 (1980) 225-227
[10] P. Deus, H. Neumann, G. Kuhn and B. Hinze, Phys. Stat. Sol. (a) 80,
(1983) 205
[11] J. E. Jaffe and Alex Zunger, Phys. Rev. B 28 (1983) pp. 5822-5847
[12] C. L. Burdick and J. H. Ellis, J. Am. Chem. Soc. 39 (1917) pp. 2518
[13] G. Donnay, L. M. Corliss, G. Donnay, J. D. H., N. Elliot and J. M.
Hastings, Phys. Rev 12 (1958) pp. 1917-1923
25
Chapter 3
Molecular Beam Epitaxy and Migration Enhanced Epitaxy
CGS and CIS thin films epitaxial growth details are discussed in this chapter to
comprehend the fundamental aspects of growth methodology. Here, we have
thoroughly reviewed the difference between MBE and MEE growth methodologies.
Joyce et al. started MBE growth (known as molecular beam techniques) of
semiconductors in 1966 [1]. Later, the investigation of compound semiconductors
using MBE was reported by Al Cho et al. [2] from Bell s Lab in the early 1970s. They
have successfully confirmed that MBE can produce high quality epitaxial films with
well-defined composition; doping and its thickness can be controlled at atomic level.
MBE growth became popular since then and there are many literatures on the
physics of this growth technique and their device applications [3, 4]. One of the
unique features of MBE is the main growth chamber at ultra-high vacuum (UHV)
which enables the in-situ RHEED observation.
3.1 MBE growth method
In this research work, all samples are grown epitaxially using solid sources in a
MBE machine (V80H). Historically, the Greek word epitaxy ,
above and
which means
means in an ordered manner . Epitaxial growth on a substrate
simply means that the grown material adjusts its lattice constant according to the
lattice constant of the respective substrate. Our MBE system consists of three
chambers, which are load-lock (LL) chamber, preparation chamber (PC) and the
main growth chamber (MC). All the samples in this thesis are grown using the
26
Vacuum Generator s MBE machine as shown in Fig. 3.1.1. The V80H is composed of
eight Knudsen-cells, which contain core materials such as Cu, Ga, In, Se and
dopants like Zn, Si and Ge. In short, MBE denotes the epitaxial growth of compound
semiconductor films by a process which involves reaction of one or more thermally
induced molecular beams with a crystalline surface in a UHV condition. MBE is
similar to vacuum evaporation; however, it provides much improved control over the
incident molecular beam fluxes where the sticking coefficient differences between
respective materials are taken into account which allows rapid change in beam
species. Comparing with other growth processes such as magnetron sputtering and
liquid phase epitaxy (LPE), the sensitivity or ability of MBE to precisely control the
epitaxial layers in an atomic scale plays a vital role in the development of
optoelectronic devices. Furthermore, the average MBE growth rate is as high as 1
m/h. Important breakthrough of MBE is the demonstration of quantum well and
superlattice structures. Detailed physical processes during MBE growth can be
found in the referred textbook [5]. In comparison to metalorganic vapor phase
epitaxy (MOVPE) and metalorganic chemical vapor deposition (MOCVD), MBE
growth method is unique because;
In-situ RHEED observation can be performed
By implementing MEE, low temperature growth can be performed.
MBE has confirmed the preparation of novel devices with multilayered epitaxial
film structures. The typical MBE machine that is used for this research has a MC
which is usually at a back pressure of
10-11 mbar, PC is with a pressure of
10-10
mbar and the LL is with a pressure of
10-8 mbar. The layout of the V80H machine
which is used in this research work is shown in Fig. 3.1.2, where this machine uses
diaphragm pump, turbo molecular pump, ion pump and cryo pump to preserve the
27
UHV growth condition. Furthermore, the growth chamber is provided with liquid
N2 cryoshrouds, which is kept cold during growth of CGS and CIS epitaxial layers.
On the other hand, the effusion Knudsen cells are cooled by external water jet-pipe.
The crucibles used in these Knudsen-cells include pyrolytic boron nitride (PBN) and
tungsten (W). The constituent materials of the semiconductor are evaporated from
effusion cells which are equipped with magnetic-mechanical shutters. For in-situ
growth analysis on planar surface, this MBE machine possesses RHEED system. As
explained in the referred textbook [5], MBE growth processes take place in three
steps;
chamber,
Molecular beam fluxes are evaporated from Knudsen-cells into growth
Evaporated beam fluxes in the form of atoms or/and molecules arrive
onto the heated substrate,
Crystallization processes take place on the substrate.
These growth processes are illustrated in Fig. 3.1.3 [5]. There are three growth
modes in MBE;
Layer-by-layer growth known as Frank-van-der-Merve (FM)
Layer-plus-island growth known as Stranski-Krastanov (SK)
known as Vollmer-Weber (VW).
28
Island growth
Figure 3.1.1 The V80H system used in this research work
Figure 3.1.2 Layout of the V80H machine used in this research work
29
Figure 3.1.3 Illustration of physical processes during MBE growth
3.2 MEE growth method
MEE growth method is different from MBE growth method, where a special
shutter sequence is implemented for MEE. In other words, cation molecular beams
are supplied at a time and then the shutters are closed. Next, the anion molecular
beams are supplied and this process is performed repetitively till the required
thickness. This process is unique in MEE and it enhances the migration of the
lateral adatoms on the growth surface compared to MBE growth methodology [7],
where the crystal quality of the grown layers is high. The most important
requirement for a step-flow growth of GaAs is the large migration distance or length,
denoted as (
), where
is the diffusion coefficient and
is the average
lifetime of the Ga atoms. Generally, the high migration velocity of Ga atoms has no
influence on the surface flatness; it is the large migration length that makes the flat
surface. In other words, MEE has a larger migration length compared to MBE. For
better understanding, the difference between MBE and MEE is shown in Fig. 3.2.1.
30
MEE growth is very suitable especially for growing heterostructures. In general,
when two different materials are grown in the same chamber, its interface
abruptness is an important parameter. This fact has been confirmed in the growth
of other III-V compound semiconductors [7]. Thus, MEE growth makes it possible
for the construction of atomic layer structured devices. Furthermore, since MEE has
a larger migration length compared to MBE, growth temperature can be decreased
in MEE. Both kinetic and potential energy that each atom possesses in MEE are
higher than MBE. Thus, growth temperatures can be reduced. At lower growth
temperatures, we are able to avoid or decrease the diffusion of impurity atoms. This
directly increases the crystal quality. In principle, the flatness of a growing surface
is driven by the balance between two-dimensional nucleation and surface migration.
Usually, GaAs layers are grown using MBE under As-stable condition [8, 9]. Ga
atoms which arrive at the surface form strong and stable bonds with As atoms.
These chemically bonded GaAs would then construct molecular-like small islands
on the growing surface. In the process of epitaxial growth of MBE, two small islands
would then join to form one large island and this process goes on to form monolayers
(MLs). If two-dimensional nucleation is strong and surface migration is weak,
eventually the growing surface would end in a rough state. The development and
implementation of MEE have solved these issues in III-V compound semiconductors
[7, 8, 9]. These evidently motivated us to implement MEE into chalcopyrite ternary
compound material.
31
Figure 3.2.1 Image to show the difference between MBE and MEE. Horizontal
arrows represent the lateral migration length and vertical arrows represent the
adatoms desorption value
3.3 CGS and CIS growth by MBE and other methods
CGS and CIS bulk, thin films and epi-layer grown by MBE and other epitaxial
growth methods are summarized in Table 3.3.1. In general, the grown epitaxial
CGS and CIS layers possess poor crystal quality due to microstructures/twinning,
stacking faults and secondary phases (i.e. Cu 2Se). Plausibly, these defects have
decreased the hole mobility of the grown samples. We have mainly compared the
growth method, growth parameters, grown layers hole mobility and its PL
excitation energy level.
32
Table 3.3.1. CGS and CIS growth and their characterization (crystal quality, transport and optical properties)
Growth
Growth condition details
Hole concentration (p) Other results
Method
[cm-3] and its mobility
(reference)
(
MBE (10)
MBE (11)
) [cm2/Vs] at RT
Epitaxial CGS grown on GaAs Not mentioned
Even near stoichiometric epi-layers contain twin crystals. XRD
(001) at 480
reveals compression in c-axis. PL shows donor acceptor pair (DAP) at
with different
Cu/Ga ratio
1.622 eV and its LO phonon.
Epitaxial CIS grown on InGaAs p=1.5 1017
XRD confirms reduction of residual strain when CIS is grown on
buffer
layers
=110
InGaAs. Sharp excitonic PL emissions also indicate improvement in
(pseudo-lattice-matched substrate)
crystal quality and substantial reduction in the point defect density.
and GaAs (001) at 550
33
Growth
Growth condition details
Hole concentration (p) Other results
Method
[cm-3] and its mobility
(reference)
(
MBE (12)
) [cm2/Vs] at RT
CGS and CIS are grown on GaAs Not mentioned
This group has mainly identified that in-situ annealing in the
(001) and these epitaxial films are
presence of Se is effective to remove antiphase domains or stacking
free from orientation and antiphase
faults. Discussions are supported by sharp PL band-to-band emission
domains. Growth temperature is
which confirms high crystal quality.
not mentioned but after MBE
growth, CGS was annealed at
450
MBE (13)
and CIS at 400
Cu-rich CGS films grown on GaAs
(001) at 490
Not mentioned
XRD results show that near stoichiometric sample (Sample No.1
. 2 samples with
with Cu/Ga ratio of 1.09) is of better crystal quality. Even the LT PL
different Cu/Ga ratio, 1.09 and 1.14
emissions confirm that sample No. 1 possesses sharp excitonic
transition.
34
Growth
Growth condition details
Hole concentration (p) Other results
Method
[cm-3] and its mobility
(reference)
(
MBE (14)
Epitaxial CIS grown on GaAs (001)
CIS on GaAs
InGaAs p= 3.4
and
pseudo-lattice-matched
350-600
at
) [cm2/Vs] at RT
at
XRD results show no noticeable improvement in the usage of InGaAs
1017,
buffer layer, however, optical and transport properties show a
significant improvement. Authors confirm that by controlling the
=38.6,
. Here, samples grown CIS on InGaAs
450
and
550
are p= 1.5
1017,
misfit strain between the substrate and the epitaxial film, they have
=110
discussed.
MBE (15)
Cu-rich CGS films grown on GaAs
(001) at 490
successfully reduced the defect density and increase the hole
mobility.
p=
1017 and
LT PL emission peaked at 1.62 eV is attributed to a DAP pair with its
=5
LO phonon at 1.587 eV.
35
Growth
Growth condition details
Hole concentration (p) Other results
Method
[cm-3] and its mobility
(reference)
(
) [cm2/Vs] at RT
Normal-
Bulk single crystal, polycrystalline Not mentioned
Band gap energies of CGS and CIS were determined using optical
freezing,
thin films and epitaxial layers of
absorption, photoreflectance and photoluminescence. Measurements
molecular
CGS and CIS
and calculations are performed to determine the CGS and CIS band
beam
gap (
), excitonic transition (
) and exciton binding energy (
).
deposition
Furthermore, strain effects on band gap structures are confirmed
(MBD), MBE,
mainly in the epitaxial layers.
MOVPE (16)
MOVPE (17)
Epitaxial CGS films grown on Not mentioned
TEM observation and first-principles calculation of phase stability in
GaAs (001) at 600
CGS were done. Only chalcopyrite ordering was found in CGS.
36
Growth
Growth condition details
Hole concentration (p) Other results
Method
[cm-3] and its mobility
(reference)
(
MBE (18)
Epitaxial CIS films on GaAs (001)
at 450
) [cm2/Vs] at RT
Net
with Cu/In ratio between concentrations
(
0.82-1.79.
)p=
2
MBE (19)
carrier
Epitaxial Cu-rich CGS films on p=
GaAs (001) at 490
with Cu/Ga
Piezoelectric photoacoustic (PPA) measurements were performed to
determine the nonradiative recombination in comparison to the PL
2
1016 (radiative recombination). Defect levels and band gap energy were
1020
confirmed. AB and C excitonic peaks at 1.04 and 1.50 eV.
1017 and
Piezoelectric photoacoustic (PPA) measurements were performed to
=5
determine the nonradiative recombination in comparison to the PL
ratio of 1.09.
(radiative recombination). A, B and C excitonic peaks at 1.73, 1.83
and 2.04 eV were confirmed.
37
Growth
Growth condition details
Hole concentration (p) Other results
Method
[cm-3] and its mobility
(reference)
(
MBE (20)
Epitaxial CIS films on GaAs (001)
) [cm2/Vs] at RT
Mobility
=
10
XRD analysis confirmed Ga from substrate interdiffused into CIS
and InGaAs (pseudosubstrates) at
layers. In the case of pseudosubstrate, no interdiffusions were
500
observed. Pole figures of {112} and {103} pole shows drastic twinning
with Cu/In ratio between
0.82 and 1.3.
concentration in In-rich. This anisotropy of twinning is confirmed in
all grown CIS layers.
MBE (21)
Epitaxial
CGS
LT PL emissions of Cu-rich layers peaked at 1.71, 1.67 (CB
films Not mentioned
CuGa),
near-stoichiometric grown on GaAs
1.63 (Cui CuGa) and 1.59 eV (Cui VGa). And Ga-rich layers peaked
(001) at 500
at 1.62 eV (DAP) (GaCu VCu). 1.71 eV is attributed to bound exciton
with Cu/Ga ratio
between 0.9 and 1.4.
while 1.67 eV is free-to-bound transition. 1.63 and 1.59 eV peaks are
confirmed to be DAPs.
38
3.4 Chalcopyrite thin films growth by MEE
There are limited attempts to grow epitaxial thin films of chalcopyrite materials.
However, the benefits of MEE are very clear and persistent as explained below:
Higher crystal quality, less defects
High hole mobility
200 cm2/Vs
These beneficial parameters have made us to attempt and continue using MEE
growth method to grow n-type chalcopyrite material.
Stanbery et al. [24], implemented MEE to grow epitaxial CIS thin films on GaAs
2 off (001) substrate at 550
50
with a growth rate of 0.3
concrete reason mentioned for using the special 2
m/h. There are no
off (001) substrate. However,
XRD analysis confirmed that under certain growth conditions, not only chalcopyrite
ordered (CH-ordered) but also CuAu ordered (CA-ordered) cation sublattices exist in
the grown films. The transmission electron diffraction and Raman scattering
results are consistent with this structural polytype.
Previously, in our group, CGS was grown with MBE method. However, the
in-situ RHEED pattern disappeared after a few MLs of growth. Hence, we have
implemented MEE [25] to grow epitaxial CGS thin films on GaAs (001) substrate at
580
with a growth rate of 1ML/2s. Hole concentration varied between 1
and 5 1018 cm-3 and hole mobility was
1017
200 cm2/Vs. There were some RHEED
pattern changes observed when Cu2Se segregation occurred (Cu rich condition).
XRD analysis results were also consistent with RHEED observation, where Cu 2Se
peaks were confirmed. Thus, after optimizing the Cu BEP, we have eliminated the
Cu2Se (secondary phase). Then we have successively grown the stoichiometric CGS
39
thin film layers. Electrical characteristics of the grown CGS with MEE sequence of
(Cu+Ga) for 2s, followed by Se for 2s and (Cu+Ga) for 2s, followed by Se for 1s of
deposition time shows a drastic difference in hole concentration (latter one shows an
increase). This has been referred to as an increase in the copper vacancy.
As we have confirmed the benefits of MEE, in my research work I have taken to
challenge to grow n-type conductive chalcopyrite material. Initially, I have
attempted to grow CGS with uniform doping. I have used Si, Zn and Ge as dopants.
However, uniform doping in CGS failed to produce n-type CGS. Thus, we have
introduced modulation doping into CGS/CIS material system.
3.5 Review on chalcopyrite material - Calculation/simulation
Zhao et al. [22], have used the first-principles total energy calculations to
determine the microscopic mechanism limiting n-type doping in CGS/CIS material
system. They have considered both cation-site donors Cd Cu and anion-site donors
ClSe. Theoretically, they have confirmed that:
(i) Bulk CIS can be doped n-type (equilibrium) with Cd or Cl but CGS cannot
(ii) Doping-limit rule VCu pins the Fermi level of CGS farther below the conduction
band compared to CIS
(iii) In CIS, Cd doping is more effective compared to Cl, where Cd Cu yields a higher
net donor concentration compared to ClSe
(iv) CIS/CGS system shows massive compensation of acceptors and donors.
In conclusion, they have confirmed that in CGS, n-type conductivity is hampered by
VCu (low formation energy, electron-killer)
40
Zunger et al. [23], further introduces some new insights by performing
calculations based on quantum-mechanical electronic structure on the properties of
CIS and CGS. Main calculations are regarding dopability, metastability and carrier
reflection at grain boundaries of these chalcopyrite materials.
(i) Dopability of CIS and CGS mostly depend on the growth conditions (i.e.
contamination of O2- or H-)
(ii) Anion vacancy (i.e. VSe) cause persistent hole photoconductivity in CIS and CGS.
Here, it is predicted that VSe is amphoteric with 2 transitions (2+ charge or 0
charge). (VSe-VCu) complex are types metastable defect. It is a concern for the solar
cell efficiency.
(iii) Existence of grain boundaries (GB) is not beneficial because in contrast to GaAs,
CIS is more stable in the polarized surface compared to non-polar surface. The
surface VCu is charged neutral because its negative charge has been used to cancel
the electrostatic dipole. This causes photogenerated holes to repel between GB and
the Grain Interior.
41
References
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157-191
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Phys. Rev. Lett., 48, (1982) 170
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[5] M. A. Herman, H. Sitter, Molecular Beam Epitaxy - Fundamentals and
Current Status (1996) Springer
[6] T. Nishinaga and K. -I. Cho, Jpn. J. Appl. Phys. 27 (1988) L12
[7] H. Yamaguchi, M. Kawashima, Y. Horikoshi, Appl. Surface Science, 33/34
(1988) 406
412
[8] J. H. Neave, P. J. Dobson and B. A. Joyce, Appl. Phys. Letters, 47 (1985)
100
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Microelectronics Journal 27 (1) (1996) 53
[11] S. Niki, P. J. Fons, A. Yamada, T. Kurafuji, S. Chichibu, H. Nakanishi, W.
G. Bi and C. W. Tu, Appl. Phys. Lett. 69 (5) (1996) 647
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1391
[13] A. Yamada, Y. Makita, S. Niki, A. Obara, P. Fons, H. Shibata, M. Kawai,
S. Chichibu and H. Nakanishi, J. Appl. Phys. 79 (8) (1996) 4318
[14] S. Niki, P. J. Fons, H. Shibata, T. Kurafuji, A. Yamada, Y. Okada, H.
42
Oyanagi, W. G. Bi, C. W. Tu, J. Crys. Growth 175/176 (1997) 1051
[15] A. Yamada, P. Fons, S. Niki, H. Shibata, A. Obara, Y. Makita and H.
Oyanagi, J. Appl. Phys. 81 (6) (1997) 2794
[16] S. Chichibu, T. Mizutani, K. Murakami, T. Shioda, T. Kurafuji, H.
Nakanishi, S. Niki, P. J. Fons and A. Yamada, J. Appl. Phys. 83 (7) (1998)
3678
[17] D. S. Su and Su-Huai Wei, Appl. Phys. Lett. 74 (17) (1999) 2483
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Yamada and S. Niki, J. Appl. Phys. 86 (8) (1999) 4354
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Phys. Lett. 77 (2) (2000) 259
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JCPDS-International Centre for Diffraction Data 2000, Advances in X-ray
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43
Chapter 4
CGS and CIS growth and characterization
Growth of undoped CGS and CIS epitaxial layers and their characterization
results are discussed and summarized in this chapter to have an advanced
understanding of these chalcopyrite materials for further new applications. All
samples in this research work are grown by MEE. Here, we will explain and show
the challenges of growing CGS and CIS thin films, mainly the optimization of beam
equivalent pressure (BEP) of Cu, Ga, In and Se. In-situ surface morphology has
been monitored using RHEED to confirm the optimized BEP. The growth
temperature is another important growth parameter. We have calibrated the
growth temperature using IR-APOCSC (CHINO), which is a Silicon pyrometer (it
possesses sensitivity in the temperatures between 500 to 1000
) to monitor and
confirm the optimized growth temperature.
4.1 Motivation of CGS growth on GaAs (001)
Chalcopyrite semiconductors such as CIS, CGS and CIGS are promising
materials for efficient thin film solar cells [1]. Among these materials, CGS has
superior optical properties for solar cell application. Although the band gap energy
of CGS (1.68 eV) is slightly high compared with the predicted optimal band gap [2]
of 1.4 eV, its absorption coefficient is considerably high even compared with those of
CIS and CIGS. Because of the large electron mass of 0.14m0, the joint density of
states of CGS becomes larger which directly results in a higher absorption
coefficient. In addition, the larger electron mass and lower dielectric constants
44
make the excitonic binding energy as high as 20 meV. Such a higher excitonic
edge, but also above the band gap energy through the Sommerfeld effect [3].
On the other hand, there exist some difficulties in the electrical transport
characteristics of CGS. One of the main drawbacks is the effect of grain boundaries.
It has been shown that the grain boundaries of CIS and CIGS produce no fatal
leakage paths in the solar cell device structures [4]. This beneficial effect cannot be
expected in CGS because of the energetic structure of cation terminated grain
boundaries of CGS [5]. Another difficulty of CGS is the formation of n-type material.
The electrical transport of undoped CGS is dominated by high hole concentration
caused by Cu-vacancy (VCu) [6] and it is very challenging to achieve n-type material.
This is ascribed to the fact that the formation energy of Cu-vacancy decreases with
increase in Fermi energy. Even when the donor impurity concentration is increased,
the acceptor concentration due to Cu-vacancy also increases with the donor
concentration. A similar mechanism is operative in CIS but for narrow band gap
materials such as CIS, the n-type conductivity could be achieved because n-type
conversion needs only a little rising up of Fermi energy to approach the conduction
band edge [6].
To overcome above difficulties, we have focused on and have grown high quality
CGS single crystal thin layers. We have succeeded in lowering the residual hole
concentration by optimizing the MEE growth condition. Then the impurity doping is
performed to achieve n-type conductivity which will be discussed in the following
chapter. So far there are no reliable reports available for n-type doping in CGS
using epitaxial growth method. However, Ge-doped n-type CGS has been reported
45
using ion implantation [7] on growing bulk CGS. Also group-IV elements such as C,
Si, Ge and Sn may produce efficient donor centers in ternary CGS which will be
discussed in the following chapter. Here, we will first discuss electrical and
photoluminescence (PL) studies of undoped CGS epitaxial layers.
4.2 Growth of CGS thin films
High-quality CGS single crystal layers are grown on (001) GaAs substrate
between 570 - 600
using the MEE [8] deposition sequence. Since the lattice
constants of GaAs and CGS are 5.653 Å and 5.614 Å, respectively, the lattice
mismatch of this system is less than 0.7 %. Therefore, high quality CGS epitaxial
layers are expected to be grown on a GaAs (001) substrate. The details of the high
quality CGS growth process by Fujita et al. were reported elsewhere [9]. During the
MEE growth, the RHEED pattern of the films remains streaky till the end of the
growth process (see Fig. 4.2.2 (RHEED pattern while thermal etching of GaAs
substrate) Fig. 4.2.3 and Fig. 4.2.4). The BEP ratio of Cu and Ga during growth has
been chosen at the value where no surface segregation of other binary phases such
as, CuSe or Cu2Se is observed in RHEED pattern. When the Cu flux is large,
electrically conductive CuSe or Cu2Se layers are segregated on the surface. However,
by reducing the Cu flux, the segregation of CuSe or Cu2Se decreases gradually and
finally disappears completely. This BEP ratio is confirmed to be the optimized
growth condition and is used in this experiment. Typical values of Cu-BEP and
Ga-BEP are 1.1 10
7
mbar and 3.0 10
7
mbar, respectively. These BEP values
provide a growth
Since CGS consists of alternate piling up of layers composed of Cu and Ga atoms,
46
and those composed of Se atoms in the [001] direction, we have adopted an MEE
deposition sequence composed of simultaneous deposition of Cu and Ga (metal
layer) for 2 sec, followed by 2 sec of Se (non-metal layer) deposition for undoped
layers growth as shown in Fig. 4.2.1.
In the MEE growth, Cu+In(Cu+Ga) deposition period, the RHEED specular
beam intensity increased and reached its maximum at the end of this deposition
period. In contrast, the specular beam intensity dropped in the Se deposition period.
These are shown in Fig. 4.2.5. Therefore, during MEE growth, RHEED specular
beam intensity shows periodic variation with constant amplitude according to the
deposition sequence (hereafter, this periodic RHEED intensity variation is called
RHEED oscillation). However, continuous RHEED intensity oscillation appeared
only when the deposition parameters, such as molecular beam intensity and
deposition duration for each component are optimized. Otherwise, RHEED
oscillation quickly disappears after the growth starts.
47
Fig. 4.2.1. MEE sequence during growth of CGS and CIS
Fig. 4.2.2. RHEED specular beam of GaAs substrate (thermal etching)
48
Fig. 4.2.3. RHEED specular beam during Cu+Ga deposition for 2 sec
Fig. 4.2.4. RHEED specular beam during Se deposition for 2 sec
49
Fig. 4.2.5. RHEED specular beam intensity oscillation in time space
50
4.3 Structural and optical characterization of CGS thin films
Undoped CGS thin films were grown at 600
by MEE. At this growth
temperature CGS samples showed hole concentration as high as 6.5×1018 cm-3. By
lowering the substrate temperature the residual hole concentration decreases
considerably. At 530
the typical hole concentration reaches approximately
2×1016 cm-3. Nevertheless, in this chapter we will discuss the characteristics of the
grown layers at 600
because the grown CGS layer exhibits a superior yet unique
PL characteristics compared with those grown at lower temperatures.
Figure 4.3.1 and 4.3.2 are the surface and cross sectional SEM images of CGS
epitaxial layers when the growth parameters are optimized. Those images show
that a two dimensional layers with high crystal quality and high interface
abruptness layers are achieved. According to the calibrated growth rate, at an
average the nominal thickness of CGS layers are
1.2
m.
Figure 4.3.3 shows the PL spectrum of undoped CGS at room temperature. Two
distinct peaks appear at 1.763 and 1.679 eV. Well-structured CGS/GaAs (001) has
its c axis perpendicular to the substrate plane [10]. Figure 4.3.4 shows the XRD
results of
2
-
appears at 68.38
scan in the vicinity of GaAs (004), where CGS (008) peak
. Here, we are able to confirm that CGS grown on GaAs
substrate has smaller lattice (compressive strain along c-axis) in comparison to the
strain-free CGS (dashed vertical line [16]). Thus, our PL measurement always
satisfies E
c polarization configuration. Hence, these peaks observed (at room
temperature) at 1.679 eV and 1.763 eV are ascribed to free exciton emissions;
A-exciton and B-exciton, respectively, according to the literature [10]. Peak 1.763 eV
is considered to be a radiative transition associated with the
51
energy region at room temperature indicating that the purity of the undoped CGS
layer is fairly high with high crystal quality. The PL spectra measured at various
temperatures (4.2
200 K) are shown in Fig. 4.3.5. The PL emission peaks at 4.2 K
of the undoped CGS layer exhibit three major PL bands; (a), (b) and (c). The peak (a),
a faint shoulder at 1.730 eV, is probably related to the free exciton (A-exciton), and
the peak (b), the dominant peak at 1.715 eV, is related to the bound exciton, because
these energies agree well with the reported values [11]. As the measurement
temperature increases, the intensity of the peak (a) increases and becomes
dominant above approximately 150 K, while the bound exciton peak (b) becomes
less intense and the peak disappears above 125 K as shown in the inset of Fig. 4.3.5.
This phenomenon is explained by considering the excitation of the bound exciton
[11] to the free exciton at elevated temperatures.
The peak (c), observed at 1.680 eV (c) can be attributed to the band-to-acceptor
[11] radiative recombination involving a shallow acceptor level like VCu, according to
the literature [12]. Figure 4.3.6 shows the temperature dependence of the PL peak
energies of peaks (a), (b) and (c).
52
Fig. 4.3.1. Surface SEM image of undoped CGS/GaAs epitaxial layer
Fig. 4.3.2. Cross section SEM image of undoped CGS/GaAs epitaxial layer
53
Fig. 4.3.3. PL spectrum of undoped CGS/GaAs (001) at room temperature
54
[16]
Fig. 4.3.4. 2 -
scan of CGS/GaAs (001) where [16] (dashed line) shows the CGS
strain-free value
55
Fig. 4.3.5. Photoluminescence spectra of undoped CGS/GaAs at various
temperatures (4.2
56
200 K)
Fig. 4.3.6. Temperature dependence of PL peak energy of near band edge emissions
of undoped CGS/GaAs
57
4.4 Conclusion/summary of CGS thin films
High-quality undoped CGS single crystal thin films were successfully grown on a
GaAs (001) substrate using MEE. Undoped CGS showed clear free exciton emission
lines (A-exciton and B exciton) at room temperature. At 4.2 K these samples exhibit
intense near band edge emissions composed of A-exciton, bound exciton and
band-to-acceptor related emissions. As the measurement temperature increases,
the intensities of the bound exciton and band-to-acceptor emission decrease and
above 125 K they disappear. At higher temperatures, the free exciton emission
dominates.
58
4.5 Motivation of CIS growth on GaAs (001)
Similar to CGS, CIS layers have been grown on SLG substrates by variety of
deposition methods for solar cell application. The grown layers usually have been
polycrystalline films. In these structures the interface flatness and thin films
crystal quality may not be very good. Heterointerface with large lattice
mismatching, in particular, may introduce high density defects. However, solar cells
using n-type CdS/ZnO [13, 14] show high performances as used for practical solar
cells. Our work is focused on the growth of single crystal thin films of CIS and CGS
to obtain high quality layers with flat heterointerfaces. We also predict that low
defect density single crystal layers would be beneficial to achieve materials with
n-type conduction. If the heterojunction interface issues are eliminated, dramatic
increase in efficiency of solar cells can be achieved. Prior to the fabrication of
CGS/CIS heterostructures, we have first attempted to optimize the growth
condition of CIS thin films on a GaAs substrate.
4.6 Growth of CIS thin films
Single crystal CIS thin films were grown on a GaAs (001) substrate using MEE
method. Similar to the CGS growth, the RHEED pattern of the CIS films remains
streaky patterns till the end of the growth process. In the CIS growth, MEE
deposition sequence is composed of alternate deposition of Cu+In and Se. The
duration of the deposition cycle is 2 sec for each of Cu+In and Se deposition which
provides about 2.0MLs/cycle (2 cycles makes 1 complete unit cell of chalcopyrite
crystal). Optimized beam equivalent pressures of Cu, In and Se used in these
experiments are 1 10
7
mbar, 3 10
7
59
mbar and 1 10
5
mbar, respectively.
Background pressure of the growth chamber is as low as 5.0×10 -10 mbar. After
removing the native oxide from the GaAs (001) substrate at 590
, we have
confirmed (2×4) reconstruction streaky patterns by means of RHEED observations.
We have successively grown single crystal CIS thin films at temperatures of 580
and 530
.
4.7 Structural and optical characterization of CIS thin films
Structural studies using XRD measurements are performed for the grown CIS
thin films. Figure 4.7.1 shows the results of double-crystal X-r
thin films grown on GaAs (001) at temperatures of 580 and 530
(004) GaAs Bragg s angle. Each CIS layer is about
grown at 530
1.2
in the vicinity of
thick. The CIS layer
exhibits well-distinguished diffraction peaks closer to its
strain-free angle of CIS (008) as shown by a dashed vertical line [16] in Fig. 4.7.1. In
the CIS layer grown at 580
, however, the diffraction peak appears at 64.74 °,
which corresponds to a smaller vertical lattice constant
compared to CIS (11.616
Å). Plausible reason for this obvious shift is due to the Ga incorporation into CIS
layer from the GaAs substrate during the epitaxial growth. Even in the CIS layers
grown at 530
, diffraction peak shifts occur slightly to the higher angle direction,
suggesting Ga atoms from the GaAs substrate are incorporated into the CIS layer.
The reason for this process is probably caused by the fact that the growth
temperature of 530
which is high enough to induce intermixing of Ga from the
substrate with the In during growth process.
Assuming th
580
valid in this system, the CIS layers grown at
contains 16 % of Ga in average, while 2 % of Ga contamination is expected
60
in the CIS layer grown at 530
temperature of 530
. This implies that even at a low growth
, Ga atoms diffuse from the GaAs substrate into growing CIS
prominently. Rather broad CIS related diffraction peaks are probably caused by the
gradual variation of Ga/In ratio in epitaxial layers. Hence, these results imply that
the low temperature growth is vital to achieve an abrupt interface between CIS and
GaAs substrate.
PL measurements were carried out for the CIS layer grown at 530
between
10 and 300 K using 532 nm excitation source with 2~30 mW using Andor model
InGaAs CCD detector. The PL spectra measured at different temperatures are
shown in Fig. 4.7.2. At 10 K, the dominant peak appears at 1.043 eV, which is
probably caused by the CIS excitonic transition. The intense and sharp spectra of
near band edge emissions are obtained which confirms high crystal quality. Using
results from Fig. 4.7.2, a temperature dependence of PL peak energy is analyzed
using the dominant emission peaks. The PL result is used to discuss the
temperature effect of CIS thin films. As shown in Fig. 4.7.3, the PL peak energy
difference between temperatures of 300 and 10 K, is only
phenomenon is caused by the small band gap energy (
10 meV [17]. This
) variation on temperature.
Furthermore, this value is an order of magnitude smaller than that of GaAs (
110
meV). This is one of the unique characteristics of CIS which will be useful for optical
devices. In addition, as shown in Fig. 4.7.3, these emission peaks have a slight
increase in PL peak energy between 10 and 50 K. In temperatures above 50 K, these
emission peaks show a decrease in PL peak energy till 300 K. One of the possible
reasons is that this phenomenon is related to the thermal expansion characteristics
of the grown CIS thin film layers where it exhibits a U shape between 4 and 50 K
61
which corresponds to the n shape in the PL peak energy.
Furthermore, CIS thin films grown at 530
resistivity of 2 10
3
(p-type) exhibit a moderate
cm , mobility of 86 cm2/Vs and hole concentration of 3~4
1017 / cm 3 .
[16]
Fig. 4.7.1. XRD measurement results of CIS grown on GaAs (001) at 580 and 530
where [16] (dashed line) shows the CIS strain-free value
62
,
Fig. 4.7.2. PL emissions of CIS thin film grown on GaAs (001) at 530
63
Fig. 4.7.3. Temperature dependence of excitonic emissions of CIS thin film grown on
GaAs at 530
64
4.8 Conclusion/summary of CIS thin films
We have successfully grown high-quality CIS single crystal thin films on GaAs
(001) substrates using the MEE growth sequence. The epitaxial layers were
characterized by X-ray diffraction and PL measurements. We have confirmed that
low temperature growth is essential for CIS growth on GaAs substrate in order to
reduce Ga incorporation into CIS layers. PL results from CIS thin films prove that
there is no obvious temperature dependence (10
300 K) of band gap energy (only
10 meV) which could be an excellent characteristic for opto-electronic devices.
65
References
[1] K. Ramanathan, M. A. Contreras, C. L. Perkins, S. Asher, F. S. Hasoon, J.
Keane, D. Young, M. Romero, W. Metzger, R. Noufi,
J. Ward and A. Duda,
Prog. Photovolt: Res. Appl. 11 (2003) 225
[2] W. Shockley and H. J. Queisser, J. Appl. Phys. 32 (1961) 510
[3] T. Ogawa, T. Takagahara, Phys. Rev. B 44 (1991) 8138
[4] A. L. Fahrenbruch and R. H. Bube, Fundamentals of Solar Cells:
Photovoltaic Solar Energy Conversion (Academic, New York, 1983)
[5] I. Repins, M. A. Contreras, B. Egaas, C. DeHart, J. Scharf, C. L. Perkins,
B. To, R. Noufi, Prog. Photovolt: Res. Appl. 16 (2008) 235
[6] S. B. Zhang, S. H. Wei, A. Zunger and H. Katayama-Yoshida, Phys. Rev. B
57 (1998) 9642
[7] J. Krustok, J. Raudoja and J. H. Schon, Phys. Stat. Sol (a) 178 (2000) 805
[8] Y. Horikoshi, M. Kawashima and H. Yamaguchi, Jpn. J. Appl. Phys. 27
(1988) 169
[9] M. Fujita, T. Sato, T. Kitada, A. Kawaharazuka, Y. Horikoshi, J. Vac. Sci.
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Shirakata, S. Isomura, H. Higuchi, J. Appl. Phys. 76 (5) (1994) 3009
[11] S. Shirakata, S. Chichibu, H. Miyake and K. Sugiyama, J. Appl. Phys. 87
(2000) 7294
[12] M. P Vecchi, J. Ramos and W. Giriat, Solid-State Electronics 21 (1978)
1609
66
[13] I. Repins, M. A. Contreras, B. Egaas, C. DeHart, J. Scharf, C. L. Perkins,
B. To and R. Noufi, Progress in Photovoltaics: Research and Applications
16(3), 235-239 (2008)
[14] M. A. Green, K. Emery, Y. Hishikawa, W. Warta and E. D. Dunlop,
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Makimoto and Y. Horikoshi, Appl. Phys. Lett., to be published
[16] H. W. Spiess, U. Haeberlen, G. Brandt, A. Rauber, and J. Schneider,
Phys. Stat. Sol. B 62, 183 (1974)
[17] S. Niki, P. J. Fons, A. Yamada, T. Kurafuji, S. Chichibu, H. Nakanishi, W.
G. Bi and C. W. Tu, Appl. Phys. Lett. 69 (5) (1996) 647
67
Chapter 5
Structural
and interfacial
properties
of
CGS/CIS
Double
Heterostructures grown on a GaAs substrate
In the previous chapter, we have gained a vast amount of knowledge about
optical properties of CGS and CIS thin films grown on a GaAs (001) substrate. Now
in this chapter, we shall further explore the uniqueness of these chalcopyrite
materials when they are grown as CGS/CIS double heterostructures on a GaAs
substrate. The realization of heterostructures is vital in the process of fabricating
semiconductor devices. Moreover, the high crystal quality heterointerface is
essential to improve device characteristics. Major challenges for growing CGS/CIS
structures are
intermixing and
the optimization of growth parameters,
the reduction of In-Ga
the optimization of growth sequence. Here, we will mainly
analyze results of RHEED, XRD, TEM, PL and Hall effect measurements.
5.1 Motivation of CGS/CIS Double Heterostructures
In general, growth of abrupt heterostructures is vital for device application.
Prior to the growth of CGS/CIS SLs, we have first investigated the growth
mechanism of CGS/CIS double heterostructures. CGS and CIS are interesting
materials due to their characteristic features which include a wide range of band
gap energies between 1.04 and 1.68 eV [1-3]. As mentioned earlier, these materials
have direct band gap, they can be applied to light-emitting devices and
photodetectors. At the CGS/CIS interface, almost all the band gap energy difference
Eg is distributed to the conduction band-edge discontinuity
68
EC [4]. This large
EC difference in band-edge discontinuity provides interesting applications of
quantum wells, such as resonant tunneling devices and quantum cascade lasers of
these chalcopyrite materials. An attempt by Afshar et al. to grow CGS/CIS
multi-quantum well on GaAs (001) has been reported [5]. They have grown the
CGS/CIS multi-quantum well using metal-organic vapor phase epitaxy (MOVPE),
where an appreciable interdiffusion between In and Ga atoms at the CGS/CIS
interface were detected. The CGS/CIS multi-quantum well structures were grown
at 500
(CIS) and 570
(CGS), which consist of 50 nm layers of CGS and CIS of
10 cycles. Because of the interdiffusion, the band gap energy of the quantum well
increased to over 1.2 eV. Furthermore, there are no results that confirm if those
CGS/CIS structures are really quantum well structures.
In order to explore new applications of CIS, CGS and CIGS, we have
investigated
fundamental
characteristics
of
these
materials
and
their
heterostructures. Even though chalcopyrite material has a high density of defects
and being classified as ODC, we employed MEE method to control the defect
formation process. The most serious problem of this material system in developing
new devices is the difficulty in growing n-type materials. To date there are no n-type
CGS/CIS single crystal thin films.
One of our research objectives was to fabricate n-type chalcopyrite materials.
Although we have obtained a highly resistive CGS thin film layers (2-3
104
cm),
n-type conductivity has never been achieved. Thus, we propose a new approach
which is modulation-doped structure using CGS/CI(G)S heterojunctions or
superlattices (SLs), grown using non-equilibrium method MBE/MEE. In the SL
structure, quantum well and barrier layers would be CIS and CGS, respectively.
69
Considering the large
EC [4] and the length of well(barrier) of CGS/CIS quantum
wells as shown in Fig. 5.1.1, electrons in the deep donors of barriers are expected to
be activated into the conduction band of CIS (confinement effect), which would
result in n-type conductivity. The quantization level is explained by equation 5.1.1.
is the length of the well structure.
(5.1.1)
In our previous study of doping in CGS, we have confirmed that Si and Zn would
produce donor levels with activation energies of 30 - 50 meV [9] as shown in Fig.
5.1.1. If the modulation doped structure mentioned above works successfully, then
n-type conductivity will be achieved. Furthermore, the n-type conductivity will open
a new application field in the chalcopyrite materials; it can be also applicable to
solar cells. In this case, pn junction solar cells without large lattice mismatch at the
interface unlike CGS/CdS can be achieved. If interface issues are eliminated,
dramatic increase in the solar cell efficiency can be achieved.
70
Ec
Donor
(30-50 meV)
0.54eV
1.68eV
810nm
nm
1.04eV
0.10eV
CGS
CIS
Ev
CGS
Fig. 5.1.1. The ideal band alignment of CGS/CIS double heterostructure at room
temperature
71
5.2 Preparation of CGS/CIS Double Heterostructures
We have grown CIS, CGS and CIS/CGS, CGS/CIS heterostructures using MEE
with a solid source MBE machine. Since the lattice constant of CGS (5.614
very close to that of GaAs (5.653
) is
), we have first grown CGS on GaAs (001)
substrate, then followed by CIS. Furthermore, CGS barrier layers are selectively
(modulated) doped with Ge (1050
). When only cations (Cu+Ga/In) are supplied,
group I and III adatoms have sufficient time to find their preferable/suitable
nucleation sites and able to produce well-ordered and smooth surface layers [10].
This directly suppresses the defect formation.
As mentioned in the previous chapter, the optimized BEPs values of Cu, In, Ga
and Se were used to grow these double heterostructures. All other growth
parameters are confirmed to be the same. First, 500nm-thick CGS layer is grown at
580
. Then, the subsequent CIS and CGS layers are successively grown at 580,
500 and 400
.
For structural studies, XRD measurements and TEM observations are
performed for the grown CGS/CIS/CGS heterostructures. We have measured and
analyzed the PL of this double heterostructures, together with undoped CGS and
CIS thin films (bulk layers) as references. Hall measurements were also performed
at room temperature.
72
5.3 Structural and optical properties of CGS/CIS Double Heterostructures
The RHEED specular beam intensity oscillation during the MEE deposition
sequence of CIS on CGS at 580
is already shown in Fig. 4.2.1. Two cycles of
RHEED intensity oscillation correspond to the growth of about 1 unit cell of
chalcopyrite structure. During the CGS growth, the highest intensities are obtained
at the end of (Cu+Ga/In), metallic components deposition period, while the lowest
intensities are obtained after the Se deposition. The constant amplitude oscillation
implies that the deposition condition is optimized (smooth surface) as shown in
previous chapters.
In the CIS growth on CGS as shown in Fig. 5.3.1, persistent oscillation with
constant amplitude is observed. This amplitude in the CGS growth, however, is a
little higher than that in the CIS growth. In addition, the peak intensity level
decreases after several MLs from the interface (shown as
), while the amplitude
remains constant. This transient is probably caused by the strain relaxation [12] at
the CGS/CIS interface (partial relaxation). The thickness of this transient is
approximately equal to the critical thickness ( ) predicted by using Mathew
estimation [13]. Similar behavior of the RHEED oscillation has been reported in
heterostructures of III-V compound semiconductors [14].
The growth of CGS on CIS exhibits a very different feature compared with the
deposition of CIS on CGS described above. Figure 5.3.2 shows the RHEED specular
beam intensity trace during MEE deposition process of CGS on CIS at 580
. In
the growth of CIS on CGS, RHEED oscillation with constant amplitude is observed.
On the other hand, when the growth is switched from CIS to CGS, higher oscillation
73
amplitude is observed as expected. The amplitude increased gradually for first 70
MLs from the heterointerface. This phenomenon is probably caused not only by the
strain relaxation but also by the intermixing between the constituent elements of
CIS and CGS. Since In and Ga atoms have different bonding energies (Ga>In) in
CIS and CGS, respectively, In atoms in CIS can be easily replaced by Ga atoms
during CGS growth. Not limited to the growth sequence, the substrate/growth
temperature is also a vital parameter of this atom-exchange process, the result in
Fig. 5.3.2 clearly indicates that the substrate temperature of 580
is too high to
grow abrupt CGS/CIS interfaces. Thus, low-temperature growth is essential for the
heterostructures in this material system.
Fig. 5.3.1. RHEED oscillation during CIS growth on CGS
74
Fig. 5.3.2. RHEED oscillation during CGS growth on CIS
75
Figure 5.3.3 shows double-crystal X-
-
CGS/CIS/CGS
heterostructures in the vicinity of (004) GaAs Bragg s angle. Each layer has a
thickness of 500 nm. The first CGS layer in each sample is grown at 580
the other layers are grown at 400, 500 and 580
, while
. The sample grown at 400
exhibits well-distinguished diffraction peaks close to the angle [15] for CIS (008),
CGS (008) and GaAs (004) indicating that high crystalline and interfacial quality of
the heterostructure. In epitaxial growth the relaxation of lattice distortion takes
place vertically (c-axis) and horizontally (a-axis). The lattice mismatch (a-axis)
between GaAs and CGS is only 0.7 %, while it is 2.3 % between GaAs and CIS.
Hence, we have chosen to grow CGS on GaAs, subsequently followed by CIS.
According to Matthews and Blakeslee equation, the l
proportional to the reciprocal of critical thickness
. It was revealed [16] that hc is
-0.08%. Plausible reason for this contradiction is
due to the difference of thermal expansion coefficient between the GaAs substrate
and coherently grown epitaxial layer. As shown in Fig. 5.3.3, the sample grown at
580
has no CIS peak. However, obvious CIGS peak with 40% of Ga incorporation
is detected, instead. On the other hand, the sample grown at 500
confirms a
mixed phase, CIGS peak appears at 15% of Ga incorporation. This implies that the
In atoms are replaced by the Ga atoms even at the low growth temperature of
500
.
76
Fig. 5.3.3 XRD measurement results of CGS/CIS grown at 580, 500 and 400
77
TEM measurement is performed by using Hitachi HF2200 with 200 kV
acceleration voltage to analyze the heterostructure in details. Cross-sectional
scanning TEM images of CGS/CIS heterointerfaces grown at 400 and 500
shown in Fig.5.3.4. The sample grown at 400
while the sample grown at 500
are
exhibits flat and abrupt interfaces,
shows rough and diffused interfaces. The TEM
images also confirm the nominal thickness of the grown layers which is
approximately equal as calibrated monolayer/cycle.
PL measurements are carried out on the CGS/CIS/CGS double heterostructures
to identify the energy levels associated with various compositions of the samples. PL
measurements are performed between 10 and 300 K using 532 nm excitation source
with 2~30 mW and InGaAs CCD detector. The PL spectra measured at 300 K
observed from CIS related emission peaks of the heterostructures grown at 400 and
500
are shown in Fig. 5.3.5. The peaks of these spectra appear at 1.01 eV for
heterostructures grown at 400
and 1.06 eV for heterostructures grown at 500
respectively, which are probably caused by the transition between bands or band to
impurity levels. 1.06 eV transition of heterostructures grown at 500
is caused by
Ga intermixing. On the other hand, intense spectra are obtained when the structure
is grown at 400
grown at 400
. This implies that there are less impurity in the heterostructures
.
Figure 5.3.6 demonstrates the PL result at 10 K for CIS related emission.
Although the intense spectra are observed in both 400 and 500
grown samples,
they exhibit dominant emission peaks at 0.96 and 1.01 eV, respectively. These
energies are in the lower energy side of the 300 K spectra. Hence, the detected
peaks at 10K are attributed to radiative recombination of CIS defect levels, such as
78
,
VCu, VIn, VGa and VSe [17].
Electrical properties are also investigated for these three CGS/CIS double
heterostructure samples by Hall effect measurements. Although the samples grown
at 400 and 500
, exhibit high sheet resistivity and low hole concentration (i.e.
104 ~ 105 / sq and 4~5 1016 / cm 3 ), no n-type conductivity has been obtained.
The hole mobility of both the samples were about 200 cm2/Vs at room temperature
and >1500 cm2/Vs at 100 K. To achieve efficient modulation doping, much shorter
periods, such as SL structures may be needed.
Fig. 5.3.4. Cross sectional S-TEM images of CGS/CIS grown at Ts 500 and 400
79
1.2E+3
1.0E+3
8.0E+2
6.0E+2
4.0E+2
2.0E+2
Fig. 5.3.5. PL spectra of CIS related emissions at 300K, which show a shift of 50
meV due to Ga incorporation
80
Fig. 5.3.6. PL spectra of CIS related emissions at 10K, which show a shift of 50 meV
due to Ga incorporation
81
5.4 Conclusion/Summary of CGS/CIS Double Heterostructures
In conclusion, we have successfully grown CGS/CIS double heterostructures on
GaAs (001) substrates using MEE growth sequence. Study on RHEED specular
beam intensity oscillation during MEE growth at 580
revealed that the
interface of CGS on CIS shows diffused structure compared with CIS on CGS
interface because of the In-Ga intermixing. Indeed XRD results confirmed that CGS
growth on CIS produces no signal of CIS but only mixed CIGS phase. Although
intermixing of CGS on CIS occurred to some extent at 500
, the effect is not so
serious and double heterojunction of CGS/CIGS/CGS can be grown.
When the growth temperature is reduced to 400
, flat and abrupt interfaces
are obtained as confirmed by XRD and TEM image observation. Hence, the
essential growth temperature for CGS/CIS heterostructures should be below 400
This low temperature growth condition has also been proved beneficial using PL
spectra analysis. The CGS/CIS double heterostructure grown at 400
showed an
efficient PL emission spectrum. The PL emission can be attributed to transition
between bands or band to impurity levels. Even though n-type CGS/CIS was not
achieved through modulation doping with Ge as a dopant, high sheet resistivity and
low hole concentration of the samples motivates this research.
82
.
References
[1] J. Stankiewicz, W. Giriat, J. Ramos and M. P. Vecchi, Solar Energy
Materials 1, 369-377 (1979)
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B. To and R. Noufi, Progress in Photovoltaics: Research and Applications
16(3), 235-239 (2008)
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Progress in Photovoltaics: Research and Applications 19, 565-572 (2011)
[4] A. Kawaharazuka, M. Fujita and Y. Horikoshi, WeA-1-4, The 17th
International Conference on Molecular Beam Epitaxy, Nara, Japan, 23-28
Sept. 2012
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Marron, A. A. Rockett, E. Rasanen and M. C. Lux-Steiner, Adv. Energy Mater.
1, 1109-1115 (2011)
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Shirakata, S. Isomura, H. Higuchi, J. Appl. Phys. 76 (5), 3009 (1994)
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3032 (1991)
[13] J. E. Matthews, A. E. Blakeslee, J. Crystal Growth, 27 (1974), 118
[14] H. Yamaguchi and Y. Horikoshi, J. Appl. Phys. 68 (4), 1610 (1990)
[15] H. W. Spiess, U. Haeberlen, G. Brandt, A. Rauber, and J. Schneider,
Phys. Stat. Sol. B 62, 183 (1974)
[16] K. Nakajima, S. Komiya, K. Akita, T. Yamaoka and O. Ryuzan, J.
Electrochem. Soc. 127, 1568-1572 (1980)
[17] C. Parlak and R. Eryigit, Phys. Rev. B 73, (2006) 245217
[18] R. People and J. C. Bean, Appl. Phys. Lett. 49, 229 (1986)
84
Chapter 6
Electrical and optical properties of Doped CGS grown on a GaAs
substrate
To realize modulation doping in CGS/CIS material system, first we have to
identify the possibility to dope a CGS single layer. In the previous chapter, we have
confirmed the growth of CGS and CIS layers and characterized the double
heterostructure of CGS/CIS. In this chapter, we will determine the doping levels of
Si (30 meV) and Zn (50 meV) impurities in CGS thin films. Impurity doping of a
material (to become n-type and p-type) is vital in the process of fabricating
semiconductor devices. Moreover, high crystal quality with homojunction is
essential to produce a high performance device. In general, major challenges of
doping in CGS are
cation site doping,
high density of intrinsic VCu and
amphoteric dopant. Here, we will mainly analyze results of PL and Hall effect
measurements.
6.1 Motivation of Doped CGS
So far there are no reliable reports on n-type chalcopyrite material. One of our
objectives was to determine whether CGS can be doped n-type with group IV
materials, such as Si and Ge using MEE. CGS has superior optical properties for
solar cell application. Although the band gap energy of CGS (1.68 eV) is slightly
higher compared to the predicted optimal band gap of 1.4 eV, its absorption
coefficient is considerably higher compared to those of CIS and CIGS. Due to the
large electron mass of 0.14m0, the joint density of states of CGS becomes larger
85
which directly results in a higher absorption coefficient. There are several groups
working on the modeling of optical constants (photon + phonon) in CGS and CIS
[1-4]. They have implemented the commonly used complex dielectric function as
below:
where, the dielectric function consists of real and imaginary values. The fitting
parameters are tuned to a respective model according to the experimental values.
Usually, the discrete and continuum excitonic effects are taken into account to avoid
the overestimation of excitonic binding energy.
The electrical transport of undoped CGS is dominated by high hole concentration
caused by Cu-vacancy [6] and it is very challenging to pursue n-type CGS.
To overcome the above difficulties, we have grown high quality CGS single
crystal thin layers on closely lattice-matched GaAs (001) substrate by using MEE.
We have succeeded in lowering the residual hole concentration by optimizing the
deposition parameters, mainly the growth temperature. Then the dopants (Si, Zn
and Ge) are incorporated into cation sites, to achieve n-type conductivity. No report
is available for n-type doping in CGS using epitaxial growth method. However,
n-type CGS co-doped with Ge and Zn has been reported using ion implantation [7]
on bulk CGS. Group-IV elements such as C, Si, Ge and Sn may produce efficient
donor centers in ternary CGS. Among these group-IV elements, we have first
employed Si as the donor impurity. Although, no n-type thin film has been achieved
yet, the formation of donors is confirmed through PL measurement. The
donor-acceptor pair (DAP) emission intensity increases considerably by Si-doping in
86
CGS in comparison to the undoped CGS.
6.2 Preparation of Doped CGS
Similar to undoped CGS single crystal layers, Si, Zn and Ge doped CGS are also
grown on closely lattice-matched (001) GaAs substrate by MEE deposition sequence
at 580 - 600
. Since the lattice mismatch of this system is less than 0.7 %, high
quality CGS thin films are grown. During MEE growth, the RHEED pattern
remains streaky till the end of the growth. The BEP ratio of Cu and Ga during
growth has been chosen at the value where no surface segregation of Cu2Se
(metallic) is observed in RHEED pattern. This BEP ratio is used in this experiment.
Since CGS consists of alternate piling up of layers composed of Cu and Ga atoms,
and those composed of Se atoms in the [001] direction, we have adopted MEE
deposition sequence composed of simultaneous deposition of Cu and Ga followed by
Se deposition for doped layers growth. Si doping has been performed during the
Cu+Ga deposition sequence of MEE, because Si may be incorporated into Ga lattice
sites. Si doped CGS thin films have the Si atom concentration of about 2×1018 cm-3.
PL and Hall effect measurements are carried out for Si doped CGS between 4.2 K
and room temperature. 30 mW Ar+ ion laser at a wavelength of 488 nm was used as
excitation source for PL measurement. The laser beam was focused onto the sample
and detected using silicon CCD. PL emissions were dispersed by Nanofinder model
with a 0.5 m grating monochromator.
87
6.3 Optical and electrical properties of Doped CGS
In this chapter, we will discuss the PL characteristics of Si-doped CGS thin films.
As described previously, Si doping has been performed during the Cu+Ga deposition
sequence of MEE process at 600
with an atomic Si concentration of 2×1018 cm-3.
The resulting thin films showed p-type conductivity with a residual hole
concentrations of approximately 4.1×1018 cm-3. Although this value is still high, the
hole concentration of Si-doped CGS decreases compared with that of the undoped
CGS (6.5×1018 cm-3). However, further increase of the Si doping concentration
slightly increases the hole concentration. This phenomenon may be explained by
considering the reduction of acceptor formation energy when the Fermi level is
raised.
Figure 6.3.2 demonstrates the PL characteristics of Si-doped CGS at low
temperatures. Among these emission bands, those denoted as (a), (b) and (c) are also
observed in the undoped thin films. However, Si doped CGS thin films exhibit
additional two PL bands in the lower energy region as shown in Fig. 6.3.2. These
new peaks appear at 1.64 eV and 1.61 eV and denoted as (d) and (e), respectively.
Since these peaks are detected only in the Si-doped sample, they are probably
caused by the doped Si impurity. No significant emissions are observed in the lower
energy region. The peak energies of peaks (d) and (e) exhibit positive temperature
dependence as demonstrated in Fig. 6.3.3. This anomalous characteristic suggests
that these emission bands are caused by donor-acceptor pair (DAP) recombination.
88
The blue shift at elevated temperatures can be explained using equation above,
the PL emission peak energy, Eg the band gap energy, E D the
donor level, EA the acceptor
r
the dielectric constant
and r the DAP separation. At low temperatures, both distant pairs and closer pairs
are operative. However, distant pairs tend to disappear at elevated temperatures
because of the weak coulomb interaction of distant pairs. As a result, contribution of
the closer pairs dominates the PL emission resulting in a distinct blue shift as
demonstrated in Fig. 6.3.3. By assuming the donor level of approximately 80 meV [9,
10, 11] and acceptor level of 70 - 80 meV (Hall effect measurement) as discussed in
connection with Fig. 6.3.3, the DAP emission spectrum at 4.2 K and 96 K suggests
an average DAP separations are about 2.1 nm and 1.8 nm, respectively. These
values are tentative but favorably compared with the Bohr radii of donor (4.9 nm)
and acceptor (2.0 nm). Since the emission peaks (d) and (e) show similar
temperature dependence and the energetic separation between peaks (d) and (e) is
close to the longitudinal optical (LO) phonon energy which is approximately 34 meV
as shown in Fig. 6.3.3. This value coincides well with the reported LO phonon
energy [12]. DAP recombination of these PL bands are also been confirmed by the
laser power excitation intensity dependence of the emission energy as shown in Figs.
6.3.4 and 6.3.5. By increasing the excitation intensity, blue shift occurs in the
emission energies of peaks (d) and (e). Since the distant pairs have low
89
recombination probabilities, the excitation of distant pairs tend to be saturated at
increased excitation intensity [13, 14, 15].
Figure 6.3.6 shows PL measurement results of undoped, Si doped and Zn doped
CGS at 4.2K. These PL spectra exhibit a similar characteristic near the band-edge.
However, deep level emissions are different. Thus, we are able to identify 3
dominant peaks at 1.62, 1.58 and 1.55 eV in the Zn doped CGS which are different
from the Si doped CGS.
As shown in Fig. 6.3.7, the excitation power dependence of these peak energies
exhibit a linear relationship with a positive slope. Thus, we have attributed peak
emission at 1.62 eV as a DAP and peak emissions at 1.58 and 1.55 eV as LO1 and
LO2, respectively. A simple Zn doped CGS band diagram is shown in Fig. 6.3.8.
90
Fig. 6.3.1. Growth temperature versus residual hole concentration of CGS layers
91
Fig. 6.3.2. PL spectra of Si doped CGS at low temperatures
92
Fig. 6.3.3. Temperature dependence of PL peaks energy of emission bands (d) & (e)
93
Fig. 6.3.4. Power excitation dependence of Si doped CGS at 6K
94
Fig. 6.3.5. Power excitation dependence of band energy of Si doped CGS
95
Fig. 6.3.6. PL spectra of undoped, Si doped and Zn doped CGS at low temperatures
96
Fig. 6.3.7. Power excitation dependence of band energy of Zn doped CGS
Fig. 6.3.8. Band diagram of Zn doped (donor level of 50 meV) CGS at 10K
97
6.4 Conclusion/Summary of Doped CGS
We have successfully grown undoped and doped (Si, Zn and Ge) CGS thin films
on a GaAs (001) substrate using MEE. Both undoped and Si doped CGS shows clear
free exciton emission lines (A-exciton and B exciton) at room temperature. At 4.2 K
these thin films exhibit intense near band edge emissions composed of A-exciton,
bound exciton and band-to-acceptor related emissions. With increasing temperature,
the intensity of the bound exciton and band-to-acceptor emission decreases and
above 125 K, they almost disappear. Then the free exciton emission dominates until
above room temperature. This phenomenon is even confirmed by other groups [13].
In addition to these near band edge emissions, new emission bands appear in Si
doped CGS in the lower energy region. They are attributed to the donor-acceptor
pair recombination emission and its phonon replica, because of their characteristics
depending on temperature and excitation power intensity. Although no n-type CGS
is obtained, doped Si atoms act as donors and produce DAP centers in CGS.
98
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B. To, R. Noufi, Prog. Photovolt: Res. Appl. 16 (2008) 235
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Albert, S. Rushworth, M. Ch. Lux-Steiner, Thin Solid Films 361-362 (2000)
426-431
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[8] S. Chichibu, Y. Harada, M. Uchida, T. Wakiyama, S. Matsumoto, S.
Shirakata, S. Isomura, H. Higuchi, J. Appl. Phys. 76 (5) (1994) 3009
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Umeda, Physica B 302-303 (2001) 357
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229 (1997) 199
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Oyanagi, J. Appl. Phys. 81 (1997) 2794
[14] J. Soon Il, Y. Kyung Hoon, A. Sejin, G. Jihye, Y. Jae Ho, Current Applied
99
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87 (2000) 7294
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035211
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100
Chapter 7
CuGaSe2/CuInSe2 single quantum wells growth on GaAs and its
characterization
High quality CuGaSe2 and CuInSe2 single crystalline layers are grown on (001)
GaAs by employing the MEE deposition sequence. When CuGaSe2 is grown on
CuInSe2 at moderate temperatures (
580
), severe interdiffusion of Ga and In
atoms take place at the CuGaSe2/CuInSe2 heterojunction. This problem has been
solved by optimizing the growth temperature. Thus, we have successfully grown a
CuGaSe2/CuInSe2 single quantum well (SQW) with sharp interfaces on (001) GaAs
for the first time. The intense PL emission from the 10 nm well SQW has been
observed.
7.1 Motivation of CGS/CIS SQW
As one of our objectives is to grow n-type chalcopyrite material and for the first
time we have challenged through the implementation of modulation doping in
CGS/CIS SLs structure. Prior to the growth of CGS/CIS superlattices (SLs), we
have first performed the growth of CGS/CIS SQWs. Well/barrier interface
abruptness is an important parameter for the efficiency of all device applications.
Particularly, we expect that the solar cell performances could be much improved
using high quality single crystal thin films of CIS and CGS instead of
polycrystalline layers [1, 2]. If the high quality single crystal layers are available,
flat interfaces are easily achieved. In addition, we expect that such low defect
density single crystal layers could be useful to achieve materials with n-type
101
conduction. If the heterojunction interface issues are eliminated, dramatic increase
in solar cell efficiency can be expected. In addition, if the flat interfaces are obtained
in CIS/CGS and CGS/CIS heterostructures, high-quality CGS/CIS SQWs and SLs
can be grown. In such cases, this material system can be applied to variety of
electro-optic devices in addition to solar cells.
Growth and characterization of single crystal CIS and CGS layers have so far
been grown by using MBE [3-6]. They have determined that Cu/Ga ratio and types
of substrate (GaAs (001) and InGaAs (pseudo-substrate)) are very important
parameters to improve the crystal quality which directly contributes to a higher
hole mobility of CGS and CIS. However, the values were still not very high
compared with those of the CGS bulk grown by chemical vapor transport method [7,
8].
When the MEE deposition sequence is applied instead of simultaneous
deposition of conventional MBE, the intense RHEED patterns are conserved
throughout the growth resulting high-quality epitaxial layers. Therefore, we have
adopted the MEE deposition sequence where cations (Cu and Ga) and anions (Se)
are alternately deposited. Similar improvement is observed for CIS growth on GaAs
(001). Thus, we have successfully grown high crystal quality of SQW on GaAs (001)
substrates for the first time. The grown structures were evaluated by means of low
temperature PL measurement.
102
7.2 Preparation of CGS/CIS Single Quantum Wells
Using the same machine with the same optimized BEPs of Cu, In, Ga and Se, we
have grown CGS/CIS SQWs. The duration of the deposition cycle is also same, 2 sec
for each of Cu+In(Ga) and Se deposition which provides about 2.0 MLs/cycle, where
2 MEE cycles complete one unit cell of chalcopyrite crystal. All other atmospheric
and growth parameters are the same, where after removing the native oxide from
the GaAs (001) substrate at 590
, we have confirmed (2×4) reconstruction. After
that, we have grown CGS/CIS SQWs. They were grown at temperatures between
580
and 350
.
For the growth of CIS/CGS quantum wells, we have first deposited
500nm-thickness of CGS layer at 580
because no discernible intermixing occurs
at the interface between the CGS layer and the GaAs substrate at 580
substrate temperature is lowered to 350
. Then the
, where a 10nm-thickness of CIS well is
grown. Next, at the same temperature of 350
, a 30 nm-thickness of CGS growth
is continued. Then the growth is suspended and the substrate temperature is
increased to 580
. Then the 470-nm thick CGS is grown until the total CGS
thickness reaches 500 nm. This is a special growth process that we have developed
to grow the CGS/CIS SQW structures.
We have measured and analyzed low temperature PL of CGS/CIS SQWs. Same
PL measurement systems were used to confirm the temperature and excitation
power dependence. Hall measurements were also performed at room temperature.
103
7.3 Electrical and optical properties of CGS/CIS SQWs
As described above, the MEE deposition sequence keeps the RHEED pattern
bright throughout the epitaxial growth. This is probably caused by the fact that, by
separating the cation deposition with the anion deposition which is MEE, we are
able to produce well-arranged lattices and smooth surface layers [11-13]. Indeed,
MEE has been proved useful for producing high mobility CGS and CIS compared
with those grown by conventional MBE [9]. Undoped CGS with hole concentration
17
of 1 10 / cm
3
shows a hole mobility higher than 200 cm2/Vs at room temperature.
As discussed in the earlier chapter, another difficulty while growth of this
system is severe intermixing of Ga and In atoms at the CGS/CIS interface. It is
found that less intermixing occurs when CIS is grown on CGS, while drastic
intermixing is observed when CGS is grown on CIS at temperatures higher than
580
. This intermixing problem is circumvented by lowering the growth
temperature and optimizing the deposition beam fluxes. The intermixing is directly
detected through in-situ RHEED specular beam intensity observation.
Such intermixing occurs even between a CIS thin films and the GaAs substrate.
CIS thin films grown at 580
, exhibits a XRD diffraction peak at 64.74 °, which
corresponds to a smaller vertical,
-lattice constant compared to that of CIS.
Furthermore, as discussed in the previous chapter, low temperature PL
measurements are carried out on the CIS thin films grown on GaAs at 530
. The
PL peak energy difference between temperatures of 300 and 10 K, is only 10 meV,
indicating that the band gap variation on temperature is very small. This is one of
the superior characteristic of CIS which will be useful for optical devices.
The PL spectra from a CGS/CIS SQW composed of 500 nm of CGS barrier and 10
104
nm of CIS well are shown in Fig. 7.3.1. The excitonic emission from the SQW
appears at 1.11 eV, while the excitonic transition feature from CGS barrier layer is
seen at 1.71 eV. PL emission peak of pure CIS should be around 1.05 eV at 10 K. The
calculated energy of the transition between the lowest quantized levels of 10 nm
quantum well is about 30 meV higher than the bulk CIS band-to-band transition.
Therefore, the SQW emission occurs at 30 meV higher than the calculated value (30
meV). This discrepancy is probably caused by the possible small amount of Ga
diffusion into the CIS well during the growth of upper CGS barrier layer.
Furthermore, the full-width-at-half-maximum (FWHM) of the PL spectra (Fig.
7.3.1) of SQW ( 140 meV) is much broader than that of the excitonic emission (Fig.
7.3.2) of CGS (
10 meV). In addition, the FWHM is almost independent on the
temperature. Therefore, this broadness of SQW emission is probably caused by the
Ga and In interdiffusion at the interface between CGS and CIS. On the other hand,
there is no obvious temperature variation (10
(only
300 K) in the PL emission energies
10 meV) in CIS SQW. This is quite an interesting feature of this material
system, which makes it attractive for future device application. Figure 7.3.2 shows
PL emission spectra from the CGS barrier layer in the CGS/CIS SQW structure.
Here, we are able to confirm much larger temperature dependence than those of
CIS. The peak energy variation between 10 K and 300 K is as large as 25 meV
which is attributed to the CGS band gap characteristics [3]. Figures 7.3.3 and 7.3.4
show the PL excitation power dependence of CGS/CIS SQW measured at 10 K. We
have identified that there are no obvious PL peak energy shifts. This confirms that
the peaks at 1.11 and 1.71 eV are CIS SQW and CGS excitonic related emissions,
respectively.
105
Electrical properties are also investigated for undoped CIS and CGS single layer
films. CIS thin films grown at 530
17
and 580
exhibit similar hole
3
concentrations between 3~4 10 / cm . The hole mobility of these CIS samples are
86 and 220 cm2/Vs, respectively. CGS thin films also show similar mobility
characteristics as CIS. The CGS samples grown at 580
approximately 200 cm 2/Vs, while those grown at 550
exhibit a hole mobility of
give only about 70 cm 2/Vs.
However, in the latter, the resulting hole concentration is very different; the
samples grown at 580
sample grown at 550
17
3
show a hole concentration of ~1 10 / cm , while the
16
3
gives 4 10 / cm . In general, these results suggest that
higher temperature growth is preferable for obtaining high mobility material for
both CIS and CGS. However, in our SQW samples, low-temperature grown CIS well
and a part of CGS are annealed during the growth of upper CGS barrier. Thus, the
quality of the well region is predicted to be good. This may be the major reason of
the intense PL emission from the well.
106
Fig. 7.3.1. PL emissions of CGS/CIS SQW grown at 350
107
Fig. 7.3.2. PL emissions of CGS barrier layer in the CGS/CIS SQW structure
108
Fig. 7.3.3. Excitation power dependence of CGS/CIS SQW measured at 10 K
109
Fig. 7.3.4. Excitation power dependence of CGS barrier layer in the CGS/CIS SQW
structure measured at 10 K
110
7.4 Conclusion/Summary of CGS/CIS SQWs
In conclusion, we have successfully grown CGS/CIS SQWs on GaAs (001)
substrates using MEE growth sequence. In order to prevent intermixing between
CIS and CGS, a special temperature control process is developed in the growth of
CGS barrier on CIS well, in particular. Grown structures are characterized by low
temperatures PL. The PL measurement of 10 nm thick CIS SQW shows an intense
peak at 1.11 eV which is higher than the predicted value by about 30 meV. This
discrepancy is mainly due to intermixing of Ga and In atoms at the interface which
has made the well-length much narrower. Since the intensity of SQW PL emission is
much higher than those observed in CGS barrier layer, the photo-excited carriers
are efficiently injected into the SQW (also confirmed with the PL excitation power
dependence). Moreover, PL emission from CIS SQW shows only 10 meV difference
with temperature variation (10
300 K) which could be a useful characteristic for
optical devices.
111
References
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Phys. Stat. Sol. B 62, 183 (1974)
113
Chapter 8
Summary
Energy crisis since early of 1970s has boosted up research and development on
solar cell or photovoltaic which is a form of renewable energy, all around the world.
The radiation which travels from the Sun to the Earth is inexhaustible and
current energy that we require. If we could convert a fraction of this radiation into
usable energy, our world would become a better place to live on. As the first
generation solar cell candidate for terrestrial applications, a silicon solar cell has
reached its maximum efficiency of about 26%. Alternatively, ternary semiconductor
compound CuGaSe2 (CGS), CuInSe2 (CIS) and quaternary CuIn1-xGaxSe2 (CIGS)
with chalcopyrite structure are of technological interest as a second generation thin
film solar cell absorber layer. Nevertheless, in this research we have attempted the
growth of CGS/CIS heterostructures and SQWs. The research result has shown that
these chalcopyrite materials have other possible applications in the future optical
devices. For the first time, we have demonstrated the CGS/CIS SQWs with superior
photoluminescence. This implies that there are other possible applications for these
chalcopyrite materials, such as quantum cascade lasers.
In this thesis, we have first confirmed the high-quality undoped CGS single
crystal thin films which were successfully grown on a GaAs (001) substrate using
MEE. Undoped CGS showed clear free exciton emission lines (A-exciton and B
exciton) even at room temperature. At 4.2 K these samples exhibit intense near
band edge emissions composed of A-exciton, bound exciton and band-to-acceptor
114
related emissions. With increasing sample temperature, the intensities of the bound
exciton and band-to-acceptor emission decrease and above 125 K the emissions
disappear. At higher temperatures, the free exciton emission is stable and
dominates till room temperature.
Next, we have successfully grown single crystal CIS thin films on (001) GaAs
substrates using the similar MEE growth sequence. These thin films were
characterized by X-ray diffraction and PL measurements. CIS grown at 530
was
confirmed to have a high quality crystal structure. PL measurements for CIS thin
films prove that there is only 10 meV energy difference with temperature variation
(10
300 K) which could be an excellent characteristic for opto-electronic devices.
One of our objectives was to grow n-type chalcopyrite CGS. Hence, we have tried
to dope CGS with several dopants, such as Si, Zn and Ge. First, we have
successfully grown Si doped CGS thin films on GaAs (001) substrates using MEE.
Similar to undoped CGS, Si doped CGS shows clear free exciton emission lines
(A-exciton and B exciton) at room temperature. At 4.2 K these CGS thin films
exhibit intense near band edge emissions composed of A-exciton, bound exciton and
band-to-acceptor related emissions. In addition to these near band edge emissions,
new emission peaks appear in Si doped CGS in the lower energy region. They are
attributed to the donor-acceptor pair recombination emission and its LO phonon
replica, considering their temperature and excitation power intensity dependences.
Although no n-type CGS is obtained, doped Si atoms act as donors and produce DAP
centers in CGS.
These results on Si and Zn doped CGS has triggered us to challenge in the
modulation doped CGS/CIS structure. For the realization of modulation doping in
115
CGS/CIS material system, first we should identify the dopability in a CGS single
layer. Then, we have confirmed the growth and characterized the double
heterostructure of CGS/CIS. N-type and p-type doping are vital in the processes to
fabricate semiconductor devices. Moreover, high crystal quality is essential to
produce a highly efficient device.
Prior to the growth of CGS/CIS SLs, we have first successfully grown CGS/CIS
double heterostructures on GaAs (001) substrates using MEE growth sequence.
Study on RHEED specular beam intensity oscillation during MEE growth at 580
revealed that the interface of CGS on CIS shows drastic diffusion of Ga and In
atoms compared with that of CIS on CGS. Indeed XRD results confirmed that CGS
growth on CIS produces no signal of CIS but only mixed CIGS phase. Although
intermixing of Ga and In atoms at the interface of CGS on CIS occurred to some
extent at 500
, the effect is not so serious compared with those grown at 580
When the growth temperature is reduced to 400
.
, flat and abrupt interfaces
are obtained as confirmed by XRD and TEM image observation. Hence, the
essential growth temperature for CGS/CIS heterostructures should be below 400
This low temperature growth condition has been proved beneficial also by PL
characteristics. The CGS/CIS double heterostructure grown at 400
showed an
intense and sharp PL emission peak. This PL emission peak can be attributed to
transition between bands or band to impurity levels. Even though n-type CGS/CIS
was not achieved through modulation doping with Ge as a dopant, this modulation
doped structure showed high sheet resistivity and low hole concentration instead of
p-type conduction which was observed for almost all of the other CGS/CIS system.
In this thesis, we have also successfully grown CGS/CIS SQWs on GaAs (001)
116
.
substrates using MEE growth sequence. In order to prevent intermixing between
CIS and CGS, a special temperature control process is developed in the growth of
CGS barrier on CIS well, in particular. Grown structures are characterized by low
temperature PL. The PL peak corresponding to the 10 nm thick CIS well appears at
1.11 eV. Since the intensity of quantum well PL emission is much higher than that
of CGS barrier layer, the photo-excited carriers are efficiently injected into the
quantum well. Moreover, PL from SQW shows no obvious temperature variation (10
300 K) which could be a superior characteristic for optical devices. In general, by
implementing these CGS/CIS material systems, we can expect for an enhancement
in the device performance.
Further growth and analysis on the CGS/CIS SLs with much narrow well and
barrier is necessary. Doping optimization with different dopants is also required to
further comprehend the doping rules in this material system.
117
Acknowledgements
There are many people who have supported me in this journey of PhD and I
sincerely would like to thank them all.
First and foremost, my deepest gratitude to Prof. Yoshiji Horikoshi and Prof.
Toshiki Makimoto, who gave me the opportunity to be part of their research groups
at Waseda University and who have patiently guided me throughout this research
work. Their expertise in semiconductor physics, MBE and great dedication for this
research were always the motivating key element. I am always very grateful for
their valuable guidance and suggestions helping me to get a better understanding of
the complex chalcopyrite compound semiconductor system.
Next, I would like to convey my gratitude to the co-referees/committee members
Prof. Masakazu Kobayashi and Prof. Kyozaburo Takeda for spending their valuable
time to evaluate this research work. I would like to also thank Prof. Atsushi
Tackeuchi of Applied Physics department, Waseda University for his excellent
support in PL measurements.
Then, my gratitude to Dr. Atsushi Kawaharazuka, Dr. Miki Fujita and Dr. Jiro
Nishinaga who have introduced and trained me to operate the V80H MBE machine
which was used in this research work.
Here, I would like to acknowledge all the staffs at Kagami Memorial Material
Research Center (Zaiken) of Waseda University for their continuous support.
My thanks to the staffs of the Materials Characterization Central Laboratory
(Bussei Lab) at Waseda University, who have answered numerous questions and
provided operation trainings of the SEM, TEM and XRD measurements that were
118
performed for this research work.
Throughout my three years of this research work, I had the chance to meet many
students in the lab of Prof. Horikoshi and then Prof. Makimoto, here, I would like
thank all of them for their assistance.
I am greatly indebted to my father, husband, Razak School of Engineering and
Advanced Technology, University Technology Malaysia, Marubun Research
Promotion Foundation and Mitsubishi Materials Foundation for their financial
support.
Last but not least, I am happy to thank my father, mother, brother, sister-in-law,
husband, daughter and son for their support and trust in me. They were always
motivating me, whenever I was in my PhD dilemma .
119
List of publications and conference presentations
Publications
1. Sathiabama Thiru, Miki Fujita, Atsushi Kawaharazuka and Yoshiji Horikoshi
Photoluminescence study of Si doped and undoped Chalcopyrite CuGaSe2 thin films
Applied Physics A, V113 (2), 257-261 (2013)
2. Sathiabama Thiru, Masaki Asakawa, Kazuki Honda, Atsushi Kawaharazuka, Atsushi
Tackeuchi, Toshiki Makimoto and Yoshiji Horikoshi
Investigation of CuGaSe2/CuInSe2 double heterojunction interfaces grown by molecular
beam epitaxy
AIP Advances 5, 027120-1 (2015)
3. Sathiabama Thiru, Masaki Asakawa, Kazuki Honda, Atsushi Kawaharazuka, Atsushi
Tackeuchi, Toshiki Makimoto and Yoshiji Horikoshi
Study of single crystal CuInSe2 thin films and CuGaSe2/CuInSe2 single quantum well grown
by molecular beam epitaxy
Journal of Crystal Growth, http://dx.doi.org/10.1016/j.jcrysgro.2015.02.059 (Accepted)
4. Sathiabama Thiru, Miki Fujita, Atsushi Kawaharazuka and Yoshiji Horikoshi
Electrical and Photoluminescence study of Undoped CuGaSe2 Single crystal thin film
The Malaysia-Japan Model on Technology Partnership: International Proceedings 2013 of
Malaysia-Japan Academic Scholar Conference, Springer, 2013, Part IV, 265p
120
International conference presentations
1. Sathiabama Thiru, Miki Fujita, Atsushi Kawaharazuka, Koji Onomitsu and Yoshiji
Horikoshi
2
The 40th International Symposium on Compound Semiconductor, Kobe, Japan (May, 2013)
2. Sathiabama Thiru, Miki Fujita, Atsushi Kawaharazuka and Yoshiji Horikoshi
2:Zn
grown on GaAs(00
Fall Meeting of Materials Research Society, T7.10, Boston (December, 2013)
3. Sathiabama Thiru, Atsushi Kawaharazuka, Yoshiji Horikoshi
2/(CuGaSe2:Ge)
Superlattice grown on
onal Symposium on Compound Semiconductor, Montpellier,
France (May, 2014)
4. Sathiabama Thiru, Atsushi Kawaharazuka, Yoshiji Horikoshi
-ray Diffraction of CuInSe2/(CuGaSe2:Ge) Hetero-structure
nal Conference on Ternary and Multinary
Compounds, Niigata, Japan (September, 2014)
5. Sathiabama Thiru, Atsushi Kawaharazuka, Yoshiji Horikoshi
-ray Diffraction of CuInSe2/(CuGaSe2:Ge) Single quantum well
h International Conference on Molecular Beam Epitaxy,
Flagstaff, Arizona (September, 2014)
121
Domestic conference presentations
1. Sathiabama Thiru, Tomohiro Sato, Miki Fujita, Yoshiji Horikoshi
2
Autumn Meeting of The Jpn. Soc. Of Appl. Phys., 11a-J-11, Ehime (September,
2012).
2. Sathiabama Thiru, Tomohiro Sato, Miki Fujita, Yoshiji Horikoshi
2
thin films grown by migration-enhanced
Japan Academic Scholar Symposium, Waseda University (November, 2012).
3. Sathiabama Thiru, Tomohiro Sato, Tomotaka Sato, Kouki Toyoda, Miki Fujita,
Yoshiji Horikoshi
2
grown on
-G4-1,
Kanagawa (March, 2013).
4. Sathiabama Thiru, Tomotaka Sato, Kouki Toyoda, Miki Fujita, Yoshiji
Horikoshi
2
74th Autumn Meeting of The Jpn. Soc. Of Appl. Phys., 17a-D6-2, Kyoto (September,
2013).
122
5. Sathiabama Thiru, Miki Fujita, Atsushi Kawaharazuka, Yoshiji Horikoshi
2
single crystal thin
6. Sathiabama Thiru, Tomotaka Sato, Kouki Toyoda, Atsushi Kawaharazuka,
Yoshiji Horikoshi
-situ RHEED observation and optical properties of thin film CuGaSe2 grown on
-D7-2,
Sagamihara, (March, 2014).
123