growth and characterization of ga2o3 -...

Growth and characterization of Ga2O3

based wide bandgap semiconductor films

March 2016

Department of Science and Advanced Technology

Graduate School of Science and Engineering

Saga University

Fabi Zhang

I

Growth and characterization of Ga2O3 based wide

bandgap semiconductor films

Abstract

Wide bandgap semiconductor materials have become the hot spot of recent

research for the possible using in many fields such as light emitting devices, power

devices and flame detectors. Among all the wide bandgap materials, β-Ga2O3 film

with the monoclinic structure is considered as a promising candidate for its large

bandgap and chemical and physical stabilities. And it is also suitable for extreme

environment applications such as high temperatures, intense radiation and corrosive

environments. However, the bandgap should be tuned to realize high sensitive

wavelength tunable photodetectors, cutoff wavelength-tunable optical filters or to

introduce shallow impurity levels for good electronic properties.

In Chapter 1, the background including the properties and the review of studies

on Ga2O3, (Ga1-xInx)2O3, (AlxGa1-x)2O3 and Si doped Ga2O3 were described. The

purpose or the motivation of this study was presented.

In Chapter 2, the film growth and characterization methods were introduced.

In Chapter 3, we have investigated the influences of oxygen pressure, substrate

temperature and deposition time on the structure and optical properties of Ga2O3 films

grown by PLD. The influence of post annealing was also been discussed. (1) The

crystal quality and the thickness of films deposited at 600 oC increase with the

increasing of oxygen pressure. The growth mode of the films is island mode. (2) By

varying the substrate temperature, the evolutions of the structure, surface morphology

and bandgap have been clearly clarified. Films deposited at substrate temperature

below 400 oC show amorphous structure while those deposited at substrate

temperature higher than 500 oC are of high oriented monoclinic structure. (3) The

optimized growth substrate temperature and oxygen pressure for our experiment is

II

500 oC and 0.1 Pa. The growth relationship between the film and the substrate is:

sapphire (0001) // β-Ga2O3 (-201) and sapphire [11-20] // β-Ga2O3 [102]. The obtained

β-Ga2O3 film is of sixfold in-plane rotational symmetry. The hard X-ray

photoemission spectroscopy reveals that the valence band of the crystalline films is

mainly due to the hybridization of Ga 4sp. (4) By varying the growth time (film

thickness), the growth process has been investigated. (5) Post annealing (annealing

temperature from 700 to 900 oC) cannot be used to obtain films with better crystal

quality than the film deposited under the optimized growth conditions. The films with

post annealing show smaller blue/UV emission ratio.

In Chapter 4, we have investigated the Si doping influence on the structure and

properties of Ga2O3 films. (1) Ga2O3 films with different Si content were grown on

sapphire substrate at 500 oC by PLD. All of the films exhibit smooth surfaces and high

transmittances. The films of Si content lower than 4.1 at. % show high (-201) oriented

monoclinic structure. The carrier density of Ga2O3 film has been increased to 9.1×1019

cm-3 with conductivity of 2.0 S cm-1 by 1.1 at. % Si doping. Further increase of Si

content leads to the decrease of carrier density. (2) By varing the substrate

temperature, it is found that film deposited at 500 oC (1 wt.% Si doped) shows lowest

conductivity and highest carrier density while possesses best crystallinity. (3) Oxygen

pressure has no obviously influence on the electrical properties of Si-doped Ga2O3,

indicating the oxygen deficiency is not the main origin of the conductive carrier in our

study.

In Chapter 5, we showed the growth of crystalline and bandgap tunable

(Ga1-xInx)2O3 films on sapphire (0001) substrate. The elimination of the phase

separation was discussed in detail. (1) Optical analysis indicates that the bandgap of

the (GaIn)2O3 films grown by PLD can be tailored from 3.8 eV to 5.1 eV by

controlling the indium content. Single phase (GaIn)2O3 films were obtained although

films with nominal In content between 0.2 and 0.5 exhibit phases separation. (2)

(GaIn)2O3 films with nominal In content of 0.3 were deposited on sapphire substrate

by PLD at substrate temperatures from RT to 500 oC. The phase separation were

observed for the films grown at substrate temperature higher than 300 oC while the

III

films grown at substrate temperature lower than 200 oC revealed homogenous element

distributions with amorphous structures. Thermal annealing had no obvious effects on

(GaIn)2O3 films grown at substrate temperature higher than 300 oC. The clusters

remained on the surface of the films after thermal annealing treatment. On the other

hand, however, by thermal annealing the film deposited at RT in atmosphere,

(GaIn)2O3 film with smooth surface, homogenous element distribution, high

orientation crystal and high optical transmittance was successfully obtained. (3) In

order to understand the annealing effects, (GaIn)2O3 films with nominal In content of

0.3 as-deposited in room temperature have been annealed under different gas

ambients (N2, vacuum, Ar, O2) or at different temperatures (700~1000 oC). It is found

that gas ambient and temperature have important influence on crystal quality of

annealed (GaIn)2O3 films. Only oxygen ambient can crystallize (GaIn)2O3 film and

film annealed in 800 oC appears best crystal quality. X-ray photoelectron

spectroscopy analysis indicated that oxygen ambient annealing has greatly helped on

decreasing the oxygen vacancy. (4) (GaIn)2O3 films with different nominal In contents

from 0.2 to 0.7 annealed at 800 oC under O2 ambient also showed high crystal quality,

improved optical transmittance, and smooth surface. Thus, high oriented films with

nominal In content from 0.2 to 0.7 without phase separation can be obtained through

annealing process. Complementally, high oriented films without phase separation can

be obtained in the nominal indium content regions of 0 to 0.1 and 0.9 to 1.0 for the

film deposited at 500 oC. By combing the two processes, bandgap tunable high quality

(GaIn)2O3 films throughout the whole indium content range from 0 to 1 can be

successfully obtained.

In Chapter 6, bandgap tunable (AlGa)2O3 films were deposited on sapphire

substrates by PLD. The deposited films are of high transmittance as measured by

spectrophotometer. The Al contents in films increase linearly with that of the targets.

The measurement of bandgap energies by examining the onset of inelastic energy loss

in core-level atomic spectra using X-ray photoelectron spectroscopy is proved to be

valid for determining the bandgap of (AlGa)2O3 films as it is in good agreement with

the bandgap values from transmittance spectra. The measured bandgap of (AlGa)2O3

IV

films increases continuously with the Al content covering the whole Al content range

from about 5 to 7 eV, indicating PLD is a promising growth technology for growing

bandgap tunable (AlGa)2O3 films.

V

Contents

1 Introduction …………………………………………………….………………….1

1.1 Background………………………………………………………..……………1

1.2 Review of studies on Ga2O3, (Ga1-xInx)2O3 and (AlxGa1-x)2O3 films…….…..….2

1.2.1 Ga2O3……..………………………………………….……………….….…2

1.2.2 Si doped Ga2O3………….……………………………………………….…6

1.2.3 (Ga1-xInx)2O3……………………..…………………………………..……..8

1.2.4 (AlxGa1-x)2O3……………………..…………………………..…….……..11

1.3 Purpose and Outline…………………………………………………………...12

References……………………………………………………………………..15

2 Film growth and characterization methods……………………...……………...18

2.1 Film deposition techniques…………………………………………………….18

2.2 Pulsed laser deposition………………………………………………………...20

2.2.1 Basic of pulsed laser deposition………………………………………..…20

2.2.2 The deposition process……………………………………………………22

2.2.3 The pulsed laser deposition equipment used in this research……………..24

2.2.4 The film growth procedures………………………………………………26

2.3 Characterization methods……………………………………………………...28

References……………………………………………………………………..34

3 Growth and characterization of Ga2O3 films……………………………………35

3.1 Introduction……………………………………………………………………35

3.2 Oxygen pressure influence………………………………………………..…...36

3.2.1 Growth rate………………………………………………………………..36

3.2.2 Crystal structure…………………………………………………………..38

3.2.3 Transmittance and surface morphology………………………………......39

VI

3.2.4 Discussions………………………………………………………….…….41

3.3 Substrate temperature influence………….……………………………………42

3.3.1 Crystal structure……….………………………………………………….43

3.3.2 Optical properties…....................................................................................48

3.3.3 Surface morphology....................................................................................49

3.3.4 Valence band structure…………………..…………………………..…...53

3.4 Growth time influence…………………………………………………………54

3.5 Annealing effects…………………………………………………….…….…..58

3.5.1 Annealing effect on films deposited at RT………………………….…….59

3.5.2 Annealing effect on films deposited at 500 oC…………………………....63

3.5.3 Annealing effect on CL spectra…………………………………………...66

3.6 Conclusions……………………………………………………………………67

References……………………………………………………………………..69

4 Effect of Si doping on properties of Ga2O3 films…………………………..…....71

4.1 Introductions…………………………………………………………………...71

4.2 Si content influence…………………………………………………………....72

4.3 Substrate temperature influence…………………………………………….....81

4.4 Oxygen pressure influence………….................................................................86

4.5 Conclusions………………………………………………………………..…..89

References……………………………………………………………………..91

5 Growth and characterization of (Ga1-xInx)2O3 films……………...…………….92

5.1 Introduction…………………………………………………………….…..….92

5.2 Bandgap engineering of (Ga1-xInx)2O3 films…………………………………..93

5.2.1 Growth parameters………………………………………………….…….93

5.2.2 Optical properties…...…………………………………………………….94

5.2.3 Structure and surface morphologies……………………………..……..…98

5.3 Thermal annealing impact on crystal quality of (GaIn)2O3 alloys……....…...101

5.4 Toward the understanding of annealing effects on (GaIn)2O3 films…………108

VII

5.4.1 Influence of annealing gas ambient…………………………………...…109

5.4.2 Influence of annealing temperature…………………………………...…114

5.5 Annealing effect on films with different indium content…………………….118

5.6 Conclusions…………………………………………………………………..121

References…………………………………………………………….……...124

6 Growth and characterization of (AlGa)2O3 films………………………….…..126

6.1 Introduction…………………………………………………………………..126

6.2 The Al content in the film……………………………………………………127

6.3 Structure of the (AlGa)2O3 films……………………………………………..129

6.4 Transmittance and bandgap of the (AlGa)2O3 films…………………………130

6.5 Conclusions…………………………………………………………………..135

References……………………………………………………………………136

7 Summary……………………………………...………………………………….137

Acknowledgments……………………………………………….…….……………141

List of publications………………………………………………………………….142

Abbreviations……………………………………………………………………….144

1

Chapter 1

Introduction

1.1 Background

Oxide semiconductive novel thin films of homo and heterostructures are

technologically attractive for future devices because of their exciting fundamental

intrinsic and extrinsic optical, electrical, magneto-optical, and piezoelectric properties.

They have potential applications such as power device, piezoelectric sensors, thin film

transistors and ultraviolet detectors. Such materials are recognized as being suitable

for extreme environment applications such as high temperatures, intense radiation and

corrosive environments. They can be deposited using conventional process methods

such as sputting, pulsed laser deposition (PLD), metal-organic chemical vapor

deposition (MOCVD) and molecular beam epitaxy (MBE).

Of the oxide semiconductor, Ga2O3 is a potential candidate due to its direct and

wide bandgap and physical and chemical stabilities. For a semiconductor material, it

is important to tail its electronic, magnetic, optical properties through doping, alloying,

quantum wells and heterostructures and nano-engineering, thus one can realize many

technologically important devices like laser diodes, visible and solar blind detectors,

transparent electronic devices, thin film transistors, and spintronic devices. To realize

the tailing of those properties, one of the cores is the bandgap tailing engineering.

2

1.2 Review of studies on Ga2O3, (Ga1-xInx)2O3 and (AlxGa1-x)2O3 films

1.2.1 Ga2O3

(1) Properties

It is known that properties of Ga2O3 films depend highly on its structure. For the

phase of Ga2O3, five polymorphs 1: α, β, γ, δ and ε phases, are known so far as shown

in table 1.1. The first α phase with a hexagonal structure is obtained on heating

GaO2H in air about 500 oC. Bohm et al. 2 reported the third form, γ-Ga2O3, similar to

γ-Al2O3, having a unit cell edge of 8.22 Å. δ-Ga2O3, which has analogous structurally

to Mn2O3, In2O3,Ti2O3, etc, corresponds to the body centered, cubic structure of the

rare earths 1. ε-Ga2O3 has been considered to have a crystal structure similar to

κ-Al2O3 in Pna21 space group 3. β-Ga2O3 with a monoclinic structure is commonly

formed under ordinary conditions and is thought to be the most thermodynamically

stable phase. Heat treatment of the other metastable transition phases including α-, γ-,

δ-, and ε-Ga2O3 converts them to β-Ga2O3. The conversion relationships among the

forms are indicated in figure 1.1 1.

Table 1.1 Crystal structures, lattice parameters of Ga2O3 polymorphs.

Phase Space group

Crystal structure

Lattice parameters

Density

(g cm-3)

PDF card

number

α R3c Rhombohedral a=4.98 Å c=13.43 Å α=120o

6.47 PDF#43-1013

β C2/m Monoclinic a=12.23 Å b=3.04 Å c=5.8 Å β=103.7o

5.94 PDF#43-1012

γ Fd-3m Cubic a=8.22 Å 6.05 PDF#20-0426

δ Ia-3 Cubic a=10 Å 5.18 PDF#06-0529

ε Pna21 Orthorhombic a=5.12 Å b=8.79 Å c=9.41 Å

5.88 PDF#06-0509

3

Figure 1.1 Conversion relationships among the five polymorphs of Ga2O3.

Yoshioka et al. 3 have investigated the structures and energetics of five Ga2O3

polymorphs using first-principles calculations. The expansivity increases in the order

β, ε, α and δ, and the bulk modulus increases in the order β, ε, δ and α at low

temperatures. The formation energies are in the order β < ε < α < δ < γ at low

temperatures. With the increase of temperature, the difference in free energy between

the ε-phase and β-phase becomes smaller, and vanishes at around 1600 K 3 .

Monoclinic β-Ga2O3, with a melting point of 1740 oC, is the most stable

crystalline structure. The gallium ions are in distorted octahedral and tetrahedral sites,

with Ga-O bond distances of 2.00 and 1.83 Å, respectively, and the oxide ions are in a

distorted cubic closest packing arrangement. These distortions are in fact the reasons

for the great level of stability of β-Ga2O3 4. The octahedral coordination gallium ions

named Ga(VI) and the tetrahedral coordination Ga(IV) are shown in figure 2. As

shown in the figure, the doubly connected straight chains of GaO6 edge shared

octahedra run along b, and the chains are connected by GaO4 tetrahedra chains. On

another aspect, (-201) oriented films in which the b axis is parallel to the substrate is

suggested to enhance the conductivity of β-Ga2O3 film 5-7, because octahedral Ga(VI)

chains which are considered to constitute the paths followed by carrier electrons are

present in the lattice along the b axis, as shown in figure 1.2.

4

Figure 1.2 Arrangements of the atoms in β-Ga2O3.

The stability of a material basically includes the ability against high temperatures

and corrosive environments. β-Ga2O3 substrate is reported to be chemically stable

against both acids and alkalis except HF and NaOH. Aqueous HF solution can be used

to etch β-Ga2O3 uniformly at room temperature 8.The β-Ga2O3 surface was found to

decompose above 1150 oC in N2, while at only 350 °C in the presence of H2 9. These

properties promise the application of β-Ga2O3 based devices under most of the

extreme environments, while reserving the door for the fabricating of those devices

utilizing existing technology such as etching, lithography etc.

(2) Growth method of Ga2O3 thin film

One of the merit of β-Ga2O3 compared with other wide bandgap materials such

as SiC, GaN, AlN and diamond is that large-diameter, high-quality single-crystal

substrates can be fabricated by melt growth methods, such as edge-defined film-fed

growth (EFG) 10, Czochralski 11 and floating zone 12-14 methods. Two inch diameter

β-Ga2O3 wafers with low dislocation densities on the order of 104 cm-2 fabricated by

EFG are now available 15 .

One of the key to the fabrication of β-Ga2O3 based devices such as chips can be

Ga (IV)

Ga (VI)

b

5

regarded as the deposition technology, since the microelectronic solid-state devices

are almost based on material structures created by thin film deposition. Several

physical and chemical techniques have been employed to deposit β-Ga2O3 films such

as sputtering 16-18, chemical vapor deposition 19, 20, spray pyrolysis 21, sol-gel method

22, thermal evaporation 23, 24, molecular beam epitaxy 25-27, and pulsed laser deposition

7, 28-30. Recently, thick β-Ga2O3 films grown by halide vapor phase epitaxy (HVPE)

were also reported 31, 32. Because of the different growth mechanism of each method,

the substrate temperature for the growth of crystalline β-Ga2O3 is also different. Table

1.2 lists some growth methods of β-Ga2O3 films and the corresponding crystallization

temperature. As shown in that table, the deposition methods of MOCVD, MBE and

PLD show the lowest crystallization temperature around 400~500 oC.

Table 1.2 Growth methods of β-Ga2O3 films and the crystallization temperatures.

Growth techniques Substrates Crystalline temperature

HVPE31 β-Ga2O3 800-1050 oC

HVPE32 sapphire 1050 oC

thermal evaporation23 glass post annealing at 600 oC

evaporation in oxygen plasma24 sapphire 800 oC

Sol-Gel22 sapphire 600 oC

MOCVD19 sapphire 500 oC

sputtering18 Si (111) post annealing at 700 oC

MBE27 β-Ga2O3 560 oC

plasma-assisted MBE25 sapphire 500 oC

PLD29 sapphire 500 oC

PLD7 sapphire 380 oC

PLD30 sapphire 700~ 800 oC

(3) Research progress of Ga2O3 thin film grown by PLD

Some growth of Ga2O3 films by PLD has been reported recently. Orita et al. 33

6

have investigated the conductivity of Ga2O3 films by Sn doping using pulsed-laser

deposition method. They at first tried to deposit Sn-doped β-Ga2O3 on silica glass.

N-type conductivity up to 1 S cm-1 was obtained under low O2 partial pressure (10-5

Pa) at substrate temperatures above 800 oC 33. Two years later, they 7 reported an

increased conductivity of 8.2 S cm-1, which was obtained at a low deposition

temperature of 380 oC using sapphire substrate. However, the crystal structures of the

films easily changed from β to ε phase when the substrate temperature is above 435

oC 7. Sn doping effect on Ga2O3 films grown by PLD was also reported by Matsuzaki

et al. 29, 34. They found that Sn-doped Ga2O3 grown on yttria-stabilized zirconia (111)

with post-annealing at 1400 oC did not exhibit detectable electrical conduction. Films

with electrical conductivities applicable to FETs were obtained at substrate

temperature between 500 to 550 oC and oxygen pressure from 5×10-4 Pa to 1×10-3 Pa

on sapphire substrates. However, the structure was determined to be orthorhombic

with large possibility of a higher-symmetry hexagonal or rhombohedral system 29, 34.

Both researches by Orita and Matsuzaki indicated that Sn doping can increase the

conductivity of Ga2O3 films but easily leads to a non-stable phase.

Hayashi et al. 35 have investigated the Mn doping effect on Ga2O3 films. They

found that Mn-doped film grown on sapphire substrate at 500 oC shows γ phase with

spinel structure. Mn atoms are located at tetrahedrally coordinated Ga sites with a

valence of +2. The doped sample shows ferromagnetism up to 350 K. They also

demonstrated the fabrication of single domain Mn-doped γ-Ga2O3 films with a

defective spinel structure on spinel substrates by PLD at 500 oC 28.

Lee et al. 36 have utilized another kind of substrate, GaN, to deposit Ga2O3 films

by PLD. (-201) oriented β-Ga2O3 thin films were obtained at substrate temperature of

800 °C and the dielectric constant were about 13.9. The β-Ga2O3/GaN structure

showed sharp interface comparing with thermally oxidized β-Ga2O3/GaN film.

1.2.2 Si doped Ga2O3

It is well known that the pure Ga2O3 is an insulator 12. As for semiconductor

devices, suitable amount of carrier density is indispensable. For example, the typical

7

carrier density (Si based devices) is 1016 cm-3 for the channel of

metal-oxide-semiconductor field-effect transistor (MOSFET); 1019, 1017, and 1015

cm-3 for the emitter, base and collector of bipolar transistor, respectively 37. Thus

selecting the appropriate doping elements to improve the carrier density of Ga2O3 is

very important. It is well known that Sn and Si are efficient n type dopants for bulk

β-Ga2O3. Orita et al. 33 have obtained Ga2O3 films with n type conductivity up to 1

Scm-1 by Sn doping at substrate temperatures above 800 oC. And the conductivity can

be improved to 8.2 Scm-1 by optimizing the deposition conditions 7.On the other hand,

Varley et al. 38 have theoretically studied donor impurities in β-Ga2O3 based on the

density functional theory and suggested that Si is an efficient n type dopant. Actually,

the effect of Si doping on the carrier density for β-Ga2O3 bulk single crystals was

proved experimentally. Villora et al. 39 showed that the carrier density of bulk

β-Ga2O3 can be intentionally controlled over three orders of magnitude from 1016 to

1018 cm-3 by Si doping. Sasaki et al. 40 developed a donor doping technique for

β-Ga2O3 substrate by using Si-ion (Si+) implantation with followed high temperature

thermal annealing, the carrier density is reported to be 1019 cm-3. Thus, the

conductivity control of β-Ga2O3 bulk single crystal has been established well.

However, for β-Ga2O3 thin films, the intentional control of the electrical properties

is still a remaining issue. Gogova et al. 41 have grown Si-doped β-Ga2O3 films by

metal organic vapor phase epitaxy (MOVPE). Although secondary ion mass

spectrometry showed that Si was incorporated in the films, Hall effect measurements

demonstrated that the resulting material was not electrically conductive. Takakura et

al. 42 also found that the conductivity does not increase by Si doping for β-Ga2O3

films grown by RF magnetron sputtering. Müller et al. 43 have tried to grow Si-doped

β-Ga2O3 thin films by PLD and obtained Ga2O3 film with conductivity of 0.2 S cm-1.

Unfortunately, they did not report on carrier density and optical properties of these

films. 4

8

1.2.3 (Ga1-xInx)2O3

(1) Properties

In-Ga-Al-O system is a promising system for semiconductor device applications

in the ultraviolet (UV) wavelength region as its possible tunable bandgap 3.5 (In2O3)

~ 4.9 (Ga2O3) ~ 8.6 eV (Al2O3) which covers the wavelength from 144 to 354 nm. As

the UV spectrum is usually divided into four regions which includes vacuum UV

(10-200 nm, UV-V), short wave UV (200-280 nm, UV-C), middle wave UV (280-315

nm, UV-B) and long wave UV (315-400 nm, UV-A), it is clear that the In-Ga-Al-O

system almost covers the whole range of the UV spectrum, thus it is a very promising

candidate for the UV device applications. It has promising usages such as (1)

destruction of microorganisms for air/water purification (240-280 nm), (2) UV-

identification , label tracking, bar coding (230-365 nm), (3) optical sensors and

analytical instrumentation (230-400 nm), (4) forensic analysis, drug detection

(250-300 nm), (5) protein analysis, DNA sequencing, drug discovery (270-300 nm),

(6) medical imaging of cells (280-400 nm), (7) solid-state lighting (300-400 nm) , (8)

curing of polymers and printer inks (300-365 nm) and (9) light therapy in medicine

(300-320 nm).

For the complex (Ga1-xInx)2O3 alloy, the component oxides Ga2O3 and In2O3 have

different crystal structure. The cubic In2O3 (bixbyite) is stable from room temperature

to the melting point. This body-centered cubic structure has 16 In2O3 per unit cell and

a cell dimension of 10.117 Å. The stable monoclinic β-Ga2O3 has four Ga2O3 per unit

cell. Half of the cation sites are octahedral coordinated and the other half are

tetrahedrally coordinated as already described. The alloyed (Ga1-xInx)2O3 can either be

of monoclinic or cubic structure or the mixture, depending on the cation composition

of the alloy.

A band structure calculation of β-Ga2O3 showed that the conduction band is

anisotropic and exclusively constituted by the antibonding states of the Ga–O bonds

in octahedral sites, with a major contribution arising from the 4s-AOs (atomic

orbitals ) of Ga ions 6. From X-ray diffraction 44 and the study of the hyperfine

9

interaction of Cd in In-doped β-Ga2O3 45, it was proved that In ions occupy

preferentially the octahedral sites. Thus, owing to the preference of In ions for

octahedral sites, a significant contribution of In-AOs to the conduction band edge can

be expected. And Binet et al. have demonstrated the participation of In orbitals in the

conduction band edge 46. And because of the above features, the bandgap engineering

of monoclinic structured (Ga1-xInx)2O3 solid state seems easy to realize. However, the

bandgap engineering of cubic structured (Ga1-xInx)2O3 solid solution becomes difficult

because very small amount of Ga can be incorporated into the cubic In2O3 44.

(2) Research progress

The research on (Ga1-xInx)2O3 alloy was began in 1997 by Doreen D. Edwards et al.

44. They reported that the solubility limit of In2O3 in the β-gallia structure decreases

with increasing temperature from 44.1 mol% at 1000 °C to 41.4 mol% at 1400 °C.

The solubility limit of Ga2O3 in cubic In2O3 increases with temperature from 4 mol%

at 1000 °C to 10.0 mol% at 1400 °C 44.

Before 2001, research of this system was focused on the bulk solid solution and

mainly the solubility of (Ga1-xInx)2O3 system. Binet et al. 46 reported that the samples

(Ga1-xInx)2O3 elaborated by coprecipitation and treated at 1400 oC show the β-gallia

phase for x<0.4. Optical spectroscopies show a decrease of the bandgap when x

increases. Patzke et al. 47 found β-(Ga1-xInx)2O3 with a maximum In2O3-content of 48

mol% and cubic (Ga1-xInx)2O3 with less than 10 mol% Ga2O3 can be obtained. Ratko

et al. 48 reported that the solubility of Ga2O3 in In2O3 cubic lattice has a limit up to

15-16.5 mol%. Vigreux et al. 49 investigated solid solutions of β-(Ga1-xInx)2O3 (x≤0.4)

with β-gallia structure by Raman spectroscopy. A continuous evolution with a linear

shift of the Raman lines with increasing x was observed in the existence range of the

solid solution (x≤0.4). For 0.4≤x≤0.5, strong alterations of the Raman spectra were

observed, corresponding either to the demixing of the solid solutions for the samples

elaborated at 1400 oC or to the occurrence of a new phase with γ-alumina structure for

samples elaborated at 1550 oC 49. From the above results, the solubility above 1000 oC

was proved to be less than 45 mol% for In2O3 in monoclinic Ga2O3 and 10 mol% for

10

Ga2O3 in cubic In2O3.

From 2002 to 2010, research on nanowires, films of (Ga1-xInx)2O3 appeared. But

the number of reports was also small and the reports were focused on electrical

properties.

Wang et al. 50 reported that (GaxIn1−x)2O3, where x=0.0~0.55, with a

homogeneously Ga-substituted, cubic In2O3 microstructure was grown by MOCVD.

The lowest electrical resistivities were found in films deposited on quartz at 500 °C

with Ga/(In + Ga) atomic ratios near 0.06. They concluded that the major determinant

of carrier scattering is not high-angle grain boundaries or lattice distortions but rather

the presence of neutral or ionized impurity scattering centers.

Lim et al. 51 reported that gallium-doped indium oxide use as highly transparent

and low resistance ohmic contacts to p-GaN. The contact annealed at 600 °C for 1

min in a nitrogen ambient showed a low specific contact resistance and high light

transmittance.

Oshima et al. 52 reported (Ga1-xInx)2O3 alloy thin films grown on c-plane sapphire

substrates with a thin Ga2O3 buffer layer by MBE. Apparent phase separation and In

segregation was observed even at x=0.08 when grown at high temperatures of 700 or

800 oC. In order to suppress the phase separation, low temperature growth at 600 ℃

or lower was found to be essential and effective.

Su et al. 53 investigated In-doped Ga2O3 zigzag-shaped nanowires and undoped

Ga2O3 nanowires synthesized by vapor–solid (VS) growth mechanism. Compared

with undoped Ga2O3 nanowires, the shapes of the as-synthesized In-doped Ga2O3

nanowires have changed obviously.

After 2011, the researches on this system become prosperous, especially for

nano-wires. Thin films growth by sol-gel method, CVD, and MBE have also been

reported.

Kokubun et al. 54 prepared (Ga1-xInx)2O3 thin films at 900 oC using a sol–gel

method, on (0001) sapphire substrates. For In content below 0.4, films are of single

phase with the same monoclinic structure as β-Ga2O3. For films with x>0.8 cubic

In2O3 solid solution structure exhibited. Between In content 0.4 and 0.8, both phases

11

exist but with poor transmittance. Knapp et al. 55 used Aerosol-Assisted Chemical

Vapor Deposition (AACVD) reaction and post annealing at 1000 oC to obtain a

Ga-substituted cubic In2O3 structure (Ga0.6In1.4O3.1). They also considered that for the

growth of (Ga1-xInx)2O3 films with high Ga content, the crystallization temperature is

high. Typically deposition of β-Ga2O3 below 650 oC often results in amorphous oxide

56.

Farvid et al. 57 investigated (Ga1-xInx)2O3 nano crystals at low temperature (<250

oC). For In content below 13%, the nano crystals exhibit γ-Ga2O3 structure while for

In content above 47%, exhibits corundum-type In2O3 structure. In3+ occupies only

octahedral, rather than tetrahedral, sites in the spinel-type γ-Ga2O3 nanocrystal host

lattice. Lin et al. 58 investigated tunable growth of In-doped Ga2O3 (Ga2O3:In) and

Ga-doped In2O3 (In2O3:Ga) nanowires (NWs) on Au-coated Si substrates by

modulating the amount of water vapor in flowing Ar at 700–750 oC. In Ar only the

Ga2O3:In NWs were grown, while in wet Ar the In2O3:Ga NWs were synthesized

instead. López et al. 59 found that indium plays a major role in the observed

morphologies of the nano structures.

1.2.4 (AlxGa1-x)2O3

Because of the similar electron structures of Al and Ga, Al can alloy with

Ga2O3 to form (AlxGa1-x)2O3 alloy. (AlxGa1-x)2O3 has a tunable larger bandgap than

Ga2O3 because Al2O3 has a big bandgap (~8.8 eV for bulk material, ~6.4 eV for

amorphous Al2O3 films).

The investigation of Ga-Al-O system is mainly focused on the formation of

γ-Al2O3-Ga2O3 by means of hydrothermal to serve as catalyst. Zahir et al. has

reported Ga2O3 as dopants to mesoporous γ-Al2O3 membrane 60. Watanabe et al.

synthesized γ-Ga2O3–Al2O3 solid solutions by several methods 61, 62.

The reports on the growth of (AlxGa1-x)2O3 films by vacuum techniques appeared

only in recent years and the number is still seldom. Oshima et al. 63 has successfully

grown β-(AlxGa1-x)2O3 on (100)-oriented β-Ga2O3 single-crystal substrates by plasma

assisted MBE. The films were grown almost coherently and maintained the β-phase

12

up to an Al content of x=0.61. Below an Al content of about x=0.4, step-flow growth

was achieved. Ito et al. 64 have realized corundum structured α-(AlxGa1-x)2O3 thin

films on sapphire by spray-assisted mist chemical vapor deposition. However, the

bandgap engineering on that system was not investigated in detail, which is believed

due to the limitation of the spectrometer wavelength range of the optical method

(commonly longer than 200 nm, which restrict the measurements for material with

bandgap higher than 6.2 eV).

Recent researches revealed that the measurement of bandgap energies by

examining the onset of inelastic energy loss in core-level atomic spectra using X-ray

photoelectron spectroscopy (XPS) is an efficient way to measure the bandgap of wide

bandgap materials 65-68. And that method has driven the bandgap engineering of

(AlxGa1-x)2O3 to go forward. The inelastic collisions include exciting the plasmon in

the bulk and at surface (a fast-moving charged particle can lose energy in the bulk

material to collective high-frequency plasma oscillations of electrons in the valence

band) and single-particle excitations due to band-to-band transitions 68. The onset of

the inelastic energy loss spectra equal to the bandgap energy because the fundamental

lower limit of inelastic loss is the bandgap energy 69.

1.3 Purpose and Outline

As has discussed above, Ga2O3 is a material that has a high dielectric constant,

good thermal stability, and superior performance as a wide-bandgap semiconductor

material. It has potential applications such as power device, piezoelectric sensors, thin

film transistors and ultraviolet detectors. And it is also suitable for extreme


environments. However, the growth of Ga2O3 by PLD is still in its infancy. The

researches of Ga2O3 films grown by PLD mainly aimed to change its electrical and

magnetic properties. Specifically, to reach those goals, impurities such as Sn and Mn

have to be incorporated into the films, which often result in a non-stable structure.

The detailed investigation of substrate temperature and oxygen influence on the

13

structure and properties of non-doped Ga2O3 deposited by PLD is thus highly desired.

On another aspect, for a semiconductor material, it is important to tail its

electronic, magnetic, and optical properties through doping, alloying, quantum wells

and heterostructures, and nano-engineering thus one can realize many technologically

important devices like laser diodes, visible and solar blind detectors, transparent

electronic devices, thin film transistors, and spintronic devices. To realize the tailing

of those properties, one of the cores is the bandgap tailing engineering. In and Al can

alloy with Ga2O3 to form (Ga1-xInx)2O3 and (AlxGa1-x)2O3 alloys to realize the bandgap

engineering because they all belong to the same element group and have similar

electron structures. However, for (Ga1-xInx)2O3 alloy, the bandgap reported is only

restricted in the region of indium content lower than 0.35. Bandgap engineering of

(Ga1-xInx)2O3 films with In content in the whole range is desired. And the bandgap

engineering on (AlxGa1-x)2O3 system is also not investigated in detail, which is

believed due to the limitation of the spectrometer wavelength range of the optical

method (commonly longer than 200 nm, which restricts the measurements for

material with bandgap higher than 6.2 eV).

The purpose of this dissertation mainly includes:

(1) Investigation of growth parameter influence on the structure and properties

of non-doped Ga2O3 deposited by PLD.

(2) Control of the conductivity of Ga2O3 film by Si doping.

(3) Bandgap engineering of (Ga1-xInx)2O3 films with In content in the whole

range.

(4) Bandgap engineering of (AlxGa1-x)2O3 films with Al content in the whole

range.

This dissertation is divided into seven chapters, and the outline is as follows.

In Chapter 1, the background and the purpose of this study are presented.

In Chapter 2, the film growth and characterization methods are introduced. In Chapter 3, the influence of oxygen pressure, substrate temperature, deposition

time on the structure and optical properties of Ga2O3 films grown by PLD is described.

The influence of post annealing is also been discussed.

14

In Chapter 4, the effect of Si doping on the properties of Ga2O3 films grown by

PLD is described. The properties of films deposited with different Si content and Si

doped films grown under different growth conditions are discussed in detail.

In Chapter 5, the growth of crystalline and bandgap tunable (Ga1-xInx)2O3 films

on sapphire (0001) substrate is described. The elimination of the phase separation was

discussed in detail.

In Chapter 6, bandgap tunable (AlGa)2O3 films are deposited on sapphire

substrates by PLD. The measurement of bandgap energies by examining the onset of

inelastic energy loss in core-level atomic spectra using X-ray photoelectron

spectroscopy is discussed.

In Chapter 7, the summary of this study is described.

15

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18

Chapter 2

Film growth and characterization methods

2.1 Film deposition techniques

One of the key to the fabrication of β-Ga2O3 based devices can be regarded as

the deposition technology, since the microelectronic solid-state devices are almost

based on material structures created by thin film deposition. Several physical and

chemical techniques have been employed to deposit β-Ga2O3 films such as sputtering

1-3, chemical vapor deposition 4, 5, spray pyrolysis 6, sol-gel method 7, thermal

evaporation 8, 9, MBE 10-12, and PLD 13-16.

In the following, the commonly used techniques are briefly described. Since in

this work, PLD was exclusively used for the deposition of oxide thin films, this

method is described in more detail in a separated section.

(1) Sol-gel method

The sol-gel process is a colloidal processing route used for the fabrication of

both glassy and ceramic materials. In this process, the sol (or solution) evolves

gradually towards the formation of a gel-like network which contains both liquid

phase and solid phase. Metal alkoxides and metal chlorides are typical precursors,

which undergo hydrolysis and poly condensation reactions to form a colloid.

Colloidal behavior was governed by the interactions between dispersed particles

within the suspending medium. The basic structure or morphology of the solid phase

can range anywhere from discrete colloidal particles to continuous chain-like polymer

networks 17.

The deposition of thin films on solid substrate by sol-gel coatings generally

includes the following steps: (a) preparation of the sol, (b) inducing the sol to change

19

into a gel by polycondensation, (c) deposit the gel on substrate, (d) sintering the

shaped gel to desired material by drying and heating the coated layer on substrate.

The advantages of the sol-gel method mainly include: It is a simple, economic

and effective method to produce high quality coatings; it can easily shape materials

into complex geometries in a gel state; composition can be highly controllable.

Despite its advantages, sol-gel technique has some limitations, e.g. weak bonding, low

wear-resistance, high permeability, and difficult controlling of porosity.

(2) MOCVD

MOCVD is a chemical vapour deposition method used to produce thin films in a

substrate. The gases are injected into a reactor. The reaction of organic compounds or

metal organics and hydrides to form crystalline films takes place at the surface of the

substrate.

For example, indium phosphide could be grown in a reactor on a heated substrate.

The first step is to introduce trimethyl indium ((CH3)3In) and phosphine (PH3) into the

reactor. The heated organic precursor molecules will be pyrolyzed. Pyrolysis leaves

the atoms on the substrate surface in the second step. These atoms will bond to the

substrate surface or act with each other thus a new crystalline layer will form in the

last step.

The advantage of this method includes the ability to grow uniform layers, high

purity and potential for large scale commercial applications. The major limitation of

this method is the availability of suitable precursor material. And another limitation is

the use of large quantities of poisonous gases.

(3) Sputtering

A well-known of the glow discharge processes is sputtering. It ejects surface

atoms from an electrode surface (target) by momentum transfer from bombarding ions

to surface atoms 18. Since sputtering produces a vapor of target material, it is similar

to evaporative deposition. The target is bombarded by ionized atoms (mostly argon

atoms). If the energy transferred to the surface atom is three times than the surface

20

binding energy, the surface atoms will be sputtered. And a secondary electron is

released from the target material. A magnetic field can be used in the sputtering

processes in order to trap the secondary electrons, thus the ionization of the plasma

near the target can be enhanced 19. The advantages of the sputtering include better step

coverage, easier to deposit alloys and etc. The disadvantage is that some plasma

damage including implanted argon is easy produced.

(4) MBE

MBE is one of several methods to grow single-crystal epitaxial films. It was

invented in the late 1960’s at Bell Telephone Laboratories by J. R. Arthur and Alfred Y.

Cho 20. MBE takes place in high vacuum (~10-8 Pa). The sources for evaporation are

placed in separate Knudsen effusion source cells (deep crucibles in furnaces with

cooled shrouds). The films are deposited on substrates by evaporating the elemental

or molecular constituents of the film evaporated from the sources. The substrates are

held at a temperature appropriate for chemical reaction, epitaxy, and re-evaporation of

excess reactants. Fast shutters are often interposed between the sources and the

substrates for the convenience of growing superlattices.

MBE exhibits many advantages over similar thin film deposition processes,

like vapour deposition such as: significantly improved purity, arbitrarily sharp

deposition resolution, and operation at low temperatures. The limitations of that

method mainly include the complex operation and the expensive equipment.

2.2 Pulsed laser deposition

2.2.1 Basic of pulsed laser deposition

PLD is a physical vapor deposition process which is carried out under a vacuum

system. It has some process characteristics similar with MBE and some with sputter

deposition. A schematical diagram of the PLD system is shown in figure 2.1. A

pulsed laser is excited from a KrF excimer laser source with wavelength 248 nm. The

pulsed laser is focused onto a target of the material to be deposited. The target

21

material will be vaporized or ablated by each of the laser pulse if it has sufficiently

high laser energy density, and thus create a plasma plume. The ablated material will

transfer forward until it reaches the surface of the substrate and provides materials

flux for the deposition of film.

Figure 2.1 Schematic diagram of the PLD equipment.

PLD process has some features which make it suitable for complex material film

deposition. The first one is that the stoichiometric transfer of material from the target

to film is possible, because that the high laser energy is absorbed by a small volume

of target material, which makes the ablation process a non-equilibrium one. The

second feature is that background gas can be used to react with the ablated captions.

Especially, the background pressure can be adjusted from ultrahigh vacuum to several

Pa. The third feature is that multicomponent films can be deposited with PLD using

single, stoichiometric targets of the material of interest, or with multiple targets for

each element. Especially, by using programmable multiple targets rotation system, the

growth of complex multi-layers such as quantum well becomes feasible.

22

2.2.2 The deposition process

The PLD deposition process can be roughly divided into three main stages. I) the

interaction of the pulsed laser with the target; II) the expansion and the propagation of

plume/ablated plasma; III) thin film growth.

(1) The interaction of the pulsed laser with the target

When the pulsed high-energetic laser beam is focused on a target, the material

will be ablated. At beginning a dense layer of vapor is formed in front of the target.

When the vapor absorbs the reminder energy of the laser, the pressure and

temperature of this vapor will increase and it will be ionized. Because the pressure of

this layer is high, it will expand from the surface of the target and form the so-called

plasma plume. In the duration of expansion, the thermal and the ionization energies

will convert into kinetic energy of the ablated plume with the order of hundreds eV.

Several parameters such as the absorption coefficient and reflectivity of the

target material, the duration of the pulse τ, wavelength λ and frequency of the laser

influence the interaction of the pulsed laser with the target

(2) The expansion and the propagation of plume

At early stage, the shape of the vapor plasma is known to be a cloud which is

strongly forwarded in the direction normal to the surface of the ablated target. Once

the plume expansion into the background gas, it looks like a piston and compresses

the background gas. During the expansion and interaction with background

atmosphere process, the plume expansion along the direction normal to the target

surface is braked to some extent, and the plume shape tends to be more and more

hemispherical.

(3) Thin film growth

At typical PLD conditions, the deposition of the ablated material on the substrate

surface can be regarded as instantaneous for every pulse. That is, each instantaneous

deposition is followed by a relative long time interval, where no deposition takes

place.

There commonly exist two steps for the growth of films including the nucleation

23

and the growth of islands. Both the nucleation and growth are far from

thermodynamic equilibrium condition and are determined by kinetic processes. As the

plume vapor possesses high supersaturating atoms, the nucleation rate is large. Once

the critical atom density reaches, nuclei will be formed. And since that, the nuclei will

grow up and crystallization process will begin.

Figure 2.2 Film possible growths: (a) Frank-Van der Merwe or layer-by-layer growth, (b)

Volmer–Weber, island (c) Stranski-Krastanov growth.

Several possible growth modes exist depending on the thermodynamic approach

of the balance between the free energies of film (γF ) , substrate (γS) and the interface

between film and substrate (γI) as shown in figure 2.2. If the total free energy of the

film surface and the interface is less than the free energy of the substrate surface

(γF+γI<γS), significant wetting is expected. Layer by layer growth mode as described

by Frank and Van der Merwe 21 will take place. However, when the bonding between

film and substrate is not strong, the interface free energy will be big (γF+γI >γS) and

three-dimensional (3D) islands will form to reduce the interface area between the film

and the substrate. This is the case of Volmer-Weber mechanism 22. In heteroepitaxial

growth, the growth mode changing from layer by layer to island growth can also been

24

observed. A layer-by-layer growth takes place in the first stage. With the deposition

process, the elastic strain induced by the mismatch between the film and substrate

becomes big. Such large strain energy can be lowered by forming islands in which

strain is relaxed. This mechanism is called as Stranski-Krastanov mode 23.

2.2.3 The pulsed laser deposition equipment used in this research

Figure 2.3 The appearance of the PLD equipment.

The appearance of the PLD equipment used in this research is shown in figure

2.3. It is composed of laser, gases, chambers and the control system.

(1) The laser

The excimer laser (KrF, 248 nm) of the equipment uses combination of a noble

krypton and a reactive fluorine gas. In an excited state (induced by an electrical

discharge or high-energy electron beams, which produce high energy pulses), noble

gas krypton can form temporarily bound molecules with themselves (dimers) or

with fluorine. The excited compound can give up its excess energy by

undergoing spontaneous or stimulated emission, resulting in a strongly repulsive

ground state molecule which very quickly (on the order of a picosecond) dissociates

Gases

Laser

Main chamber Load chamber

Control system

25

back into two unbound atoms. This forms a population inversion 24. The direction of

the laser is modulated towards the growth chamber by reflection mirrors. Before it

gets into the camber, it is focused by a lens.

(2) The chambers

The main structure of the system includes a main chamber and a load chamber.

The load chamber is used for the exchange of the substrate. The load chamber is

equipped with TMP (turbo molecular pump) and RP (rotary pump) for evacuation.

The exchange of the substrate is done after the load chamber is leaked with nitrogen.

There is valve between the load chamber and the main chamber, in order to keep the

main chamber in vacuum while the load chamber is leaked. The main chamber

(growth chamber) is also equipped with TMP and RP for evacuation. The base

pressure for the deposition is in the range of ~10-6 Pa. The heater for the samples uses

pyrolytic boron nitride (PBN) resistance heating elements. The maximum temperature

of the heater is 1200 oC. The heater has some advantages such as: superior

performance in ultra-high vacuum; chemically inert to most corrosive gases, liquids;

long life, dimensionally and electrically stable; high resistance for low cost power

supplies; mechanically durable; thermally shock resistant; unaffected by vibration;

tailored thermal gradients for specific requirements and ultra-fast response, low

thermal mass. The substrate then is heated through a surface contacted substrate

holder which is loaded on the heater. The targets are loaded on a target holder system.

There are six target holders on that system. The targets can revolue around the center

of that system. They can also rotate around themselves. The others of the targets are

shielded and only the one needed is exposed and faced to the laser and the substrate.

Through this, the cross contaminant among the targets is avoided.

(3) The gas supplying system

The gases include those used for the laser, for the growth chamber and for the

load chamber. The gases used for the laser include the krypton, hafnium, neon (buffer

gas) and helium (flush the pipe). High purity oxygen and nitrogen are used for the

26

environments needed during the film growth. Purity nitrogen is used for the leak of

the load chamber.

(4) The control system

Figure 2.4 The main view of the control panel.

The control system has many functions. It can be used to monitor the pressure

the load and main chamber, to monitor the temperature of the substrate as well as the

target. It is also used to control the actions of valves, pumps in the load chamber and

in the growth chamber. The motions of the substrate holder and the target holder are

controllable by this system. And also, the temperature of the substrate is

programmable by using that system. The main view of the control panel is shown in

figure 2.4.

2.2.4 The film growth procedures

(1) Clean of the substrate

The substrates are cleaned ultrasonically in methanol and acetone for 20 minutes,

followed each with a nitrogen gun to blow off the liquid drop. After rinsed in

27

deionized water for about 5 minutes, the substrates are chemically etched. The

sapphire substrate is chemically etched in a hot H3PO4:H2SO4 (1:3) solution, rinsed in

deionized water, and then blown dry with nitrogen gas before they were introduced

into the growth chamber.

(2) Setting of the substrate on the substrate holder

The operation is done as follows; Move the substrate holder in the growth

chamber to a proper position and angle which faces a valve between the load and

growth chamber by the control panel; open that valve and then transfer the substrate

holder to the load chamber; close that valve. Stop the TMP and RP pump of the load

chamber; open the valve of the purity nitrogen gas and leak the load chamber with the

nitrogen gas. After the pressure of the load chamber increasing to the atmospheric

pressure, open the door of the load chamber and take the substrate holder out. Fix the

substrate on the substrate holder and put it on the transfer tube which is located at the

load chamber; close the door of the load chamber and then turn on the RP pump and

TMP pump to evacuate the load chamber.

(3) Growth of film

After the pressure of the load chamber is lower than about 5×10-4 Pa, open the

valve between the load chamber and the growth chamber, transfer the substrate holder

to the growth chamber and then close that valve. Move the substrate holder to desired

position and adjust the angle of the holder to about 0o using the control panel. Move

the target holder to desired position. Set the substrate temperature and turn on the

power of the substrate heater. After the substrate temperature reaches the set point,

open the valve of the gas (oxygen or nitrogen) and adjust the chamber pressure. Turn

on the power of the excimer laser, adjust the parameters such as the laser power, laser

mode and frequency. Select the target and start the target rotation program. Start the

laser thus the film begins to deposit on the substrate surface.

(4) Cooling of the system

28

Stop the laser when deposition time reaches. Turn off the power of the laser.

Stop the target rotation program. Turn off the power of the heater. When the

temperature of the substrate is lower than 200 oC, turn off the valve of the gas.

2.3 Characterization methods

In this section, we presented the material characterization methods used to

analyze the thickness, surface morphology, chemical composition, crystal structure

and conductivity of the deposited films. Each technique is explained in the individual

subheading below.

(1) Scanning Electron Microscope (SEM)

Figure 2.5 The overview of the SEM equipment.

SEM with built-in energy-dispersive X-ray spectroscopy (EDS) analysis

capability was used for morphological characterization of the deposited layers. The

system used is Philips XL30 FEG SEM, a field emission high resolution scanning

electron microscope which can be operated both at low accelerating voltage (200V)

29

and high accelerating voltage (30kV). The SEM resolution is 2.0 nm at 30 kV and 5.0

nm at 1 kV. Oxford Instruments INCA X-sight EDS is equipped for characterize the

composition of the deposited films. The cathodoluminescence (CL) spectra were

measured in a SEM using an Oxford Instrument MonoCL system.

(2) Atomic Force Microscopy (AFM)

Figure 2.6 The outlook of the AFM equipment.

The equipment used in this experiment is a multi-mode scanning probe

microscope (MM-SPM). It is designed for imaging small samples (approx. 1.5cm dia.)

using a series of interchangeable scanners and is able to provide images from the

atomic scale to 175μm in size. In the measure process of the AFM, a sharp tip (probe)

is used to scan the surface of the sample by measuring forces between the tip and the

surface at a very short distance. The topographic image of the surface is produced by

recording the variations of the tip height above the surface while the tip is scanned

repeatedly across the surface. The AFM can be operated in a number of modes,

30

depending on the application. In general, possible imaging modes are divided into

static (also called contact) modes and a variety of dynamic (non-contact or "tapping")

modes where the cantilever is vibrated or oscillated at a given frequency.

(3) X-ray diffraction (XRD)

Figure 2.7 The outlook of the XRD equipment.

XRD is a rapid non-destructive analytical technique which primarily used for

phase identification of a crystalline material; it also can provide information on unit

cell dimensions. By observing the scattered intensity of an X-ray beam hitting a

sample as a function of incident angle, one can obtain information about the

crystallographic structure, chemical composition and physical properties.

The XRD measurement used was PANalytical X'Pert MRD (Materials Research

X-ray Diffraction) system. The system can be used for detailed structural analysis of

advanced semiconductors, thin film, and nanomaterials. It handles a wide variety of

X-ray scattering methods, including high resolution diffraction, in-plane diffraction,

reflectivity, thin film phase analysis, wafer mapping, grazing incidence small angle

X-ray scattering, stress, texture and nonambient analysis. Identification of the patterns

31

was achieved by comparison of the XRD diffractogram with the international centre

for diffraction data base. The crystalline quality of the films was evaluated by X-ray

rocking curve (XRC) of a double-crystal method using Cu Kα1 radiation

(4) Hall effect

Figure 2.8 The outlook of the Hall equipment.

The conduction type and the carrier concentration of the films were characterized

by Hall measurements (Resitest 8300) under the Van der Pauw configuration at room

temperature.

The Hall effect is due to the nature of the current in a conductor which is

composed of the movement of many small charge carriers, typically electrons

and holes. When a magnetic field is present, these charges experience a force, called

the Lorentz force. The electrons and holes will be deflected in opposite directions by

the magnetic field. The separation of charge establishes an electric field that opposes

Lorentz force, so a steady electrical potential is established as long as the charge is

flowing.

32

(5) Other characterization methods

X-ray photoelectron spectroscopy measurements were performed using Mg Kα

X-ray source. The X-ray excites the electrons of the sample atoms and if their binding

energy is lower than the X-ray energy, they will be emitted from the parent atom as a

photoelectron. Only the photoelectrons at the extreme outer surface (10-100 Å) can

escape the sample surface, making this a surface analysis technique. A typical XPS

spectrum is a plot of the number of electrons detected as a function of the binding

energy of the electrons. Each element produces a characteristic set of XPS peaks at

characteristic binding energy values. These characteristic spectral peaks correspond to

the electron configuration of the electrons within the atoms, e.g., 1s, 2s, 2p, 3s, etc.

The XPS can be used to measure the elemental composition, chemical

state and electronic state of the elements that exist within a material.

The XPS experiments were done in Kyusyu University, Japan, using MgKα line

in an ultrahigh vacuum system with a Perkin–Elmer 5600 X-ray photoelectron

spectrometer, and the spectra were collected at a takeoff angle of 15°. The surface of

the samples was Ar+ ion (3 keV) etched for 2 min before XPS analysis. The typical

ion beam current was about 1.8 μA. And the raster size of the irradiation area was

about 3 mm×3 mm.

The transmittance of films was measured using a Jasco V-570 spectrophotometer

which is shown in figure 2.9. It is a double beam system with single monochromatic.

The wavelength range is between 190~2500 nm with a wavelength accuracy of 1.5

nm.

33

Figure 2.9 The outlook of the spectrophotometer.

It is well known that for films with direct bandgap, the absorption follows a

power law of the form:

/ (2-1)

where is the energy of the incident photon, is the absorption coefficient, B is

the absorption edge width parameter, and is the bandgap. The optical absorption

coefficient of the films is evaluated using the standard relation taking the film

thickness into account 25. By extrapolating the linear part of ~ to the

horizontal axis, one can obtain the bandgap of the films.

34

References: 1. M. Fleischer, W. Hanrieder and H. Meixner, Thin Solid Films, 1990, 190, 93.


3. Y. Zhang, J. Yan, Q. Li, C. Qu, L. Zhang and T. Li, Physica B: Condensed Matter, 2011, 406,

3079.


011119.




8. A. A. Dakhel and W. E. Alnaser, Microelectron. Reliab, 2013, 53, 676.

9. S. Nakagomi and Y. Kokubun, J. Cryst. Growth, 2012, 349, 12.


11. T. Oshima, T. Okuno, N. Arai, N. Suzuki, S. Ohira and S. Fujita, Applied Physics Express,

2008, 1, 011202.






15. M. Orita, H. Hiramatsu, H. Ohta, M. Hirano and H. Hosono, Thin Solid Films, 2002, 411,

134.


2014, 583, 214.

17. L. C. Klein and G. J. Garvey, J. Non-cryst. Solids., 1980, 38, 45.

18. W. Kern and K. K. Schuegraf, Handbook of Thin Film Deposition Processes and Techniques

(Second Edition), ed. K. Seshan, William Andrew Publishing, Norwich, NY, 2001, DOI:

http://dx.doi.org/10.1016/B978-081551442-8.50006-7, pp. 11.

19. M. Hellwig, Doctor, Ruhr-University Bochum, 2011.

20. A. Y. Cho and J. R. Arthur, Prog. Solid State Ch., 1975, 10, 157.

21. F. C. Frank and J. H. van der Merwe, One-Dimensional Dislocations. II. Misfitting

Monolayers and Oriented Overgrowth, 1949.

22. M. Volmer and A. Weber, Z. phys. Chem, 1926, 119, 277.

23. I. N. Stranski and Krastanov, Acad. Wiss. Math.-Naturw.Klasse IIb, 1938, 146, 797

24. D. Basting, Excimer laser technology, Springer Science & Business Media, 2005.

25. E. J. Rubio and C. V. Ramana, Appl. Phys. Lett., 2013, 102, 191913.

35

Chapter 3

Growth and characterization of Ga2O3 films

3.1 Introduction

As we have stated that Ga2O3 is a material that has a high dielectric constant,

good thermal stability, but most obviously, it shows superior performance as a wide

bandgap semiconductor material 1. Thus Ga2O3 can be used as green light-emitting

diode 2, ultraviolet photodetector 3, 4, deep-ultraviolet transparent electrode 5, 6, metal

oxide semiconductor field-effect transistors 7, and high dielectric oxide 8 or active

material for FET device 9. Ga2O3 films have been prepared by various methods such

as sputtering 10, 11, chemical vapor deposition 1, 12, spray pyrolysis 13, sol-gel method 3,

MBE 4, 14, and PLD 15, 16. Among them, PLD has many advantages due to completely

compositional consistency between a target and a deposited film, and is especially

suitable for low temperature growth of thin films for the relative high kinetic energies

that the ablated species have 17. Moreover, PLD has another advantage for growing

oxide films such as Ga2O3 because the carrier concentration generated by oxygen

vacancies can be controlled by oxygen pressure in PLD 5.

As I have reviewed in Chapter 1, the reported Ga2O3 films by PLD often

contained impurities such as Sn and Mn, and often appeared in different phases 15, 18.

A recent first-principles study on the energetics of the Ga2O3 polymorphs suggested

that the differences in free energy between α, β and ε phases are small and, therefore,

the polymorph of Ga2O3 is sensitively selected depending upon the preparation

conditions 19. Thus systematic investigation of growth factors such as oxygen pressure

and substrate temperature influence on the structure and properties of pure Ga2O3

films is highly required.

36

3.2 Oxygen pressure influence

The background oxygen mainly plays two important roles in the PLD deposition

process. One is to supply oxygen atoms for the film; the other is influences on the

plume shape and propagation speed.

In order to investigate the oxygen pressure influence in present work, the oxygen

pressure is adjusted from 1×10-4 to 1×10-1 Pa by a valve between the growth chamber

and the oxygen gas line. The growth temperature was set at 600 oC during the growth

because the reported crystallization temperature of β-Ga2O3 films deposited on

sapphire substrates by PLD are mainly from 400 to 800 oC 5, 16, 18, 20. The growth

conditions are listed in table 3.1.

Table 3.1 Growth condition for oxygen pressure influence.

Targets Ga2O3

Substrate Sapphire (0001)

Oxygen pressure 1×10-4, 1×10-3, 1×10-2, 1×10-1Pa

Substrate temperature 600 oC

Laser frequency 1 Hz

Laser power 225 mJ

3.2.1 Growth rate

Figure 3.1 shows the oxygen pressure influence on growth rate of β-Ga2O3

deposited at substrate temperature 600 oC. The growth rate is obtained by the

thickness of the film considering the deposition time. The growth rate of Ga2O3 thin

film increases monotonically with the increase of oxygen pressure.

37

10-4 10-3 10-2 10-158

60

62

64

66

68

70

72

74

76

Gro

wth

rat

e (

nm

/h)

O2 pressure (Pa)

Figure 3.1 Oxygen pressure influences on growth rate of β-Ga2O3 deposited at substrate

temperature of 600 .

It is known that the film growth rate of PLD affects by many parameters such as

laser energy, substrate temperature, the distance between the target and substrate, and

the pressure of the ambient gas, etc. Tyunina et al. 21 have investigated the ambient

gas pressure influence on the growth rate of lead zirconate titanate films under

different laser fluency and found different results: the growth rate almost increased

with the ambient gas pressure at high laser fluence but the growth rate decreased with

the increase of gas pressure at low laser fluence. Thus they suggested the growth rate

is affected by the number of adatoms that can reach the substrate surface as well as

the desorption of these adatoms. The film growth rate per laser pulse can be expressed

as 21:

1 exp （3-1）

Where is the density of the film; is the flux of the species arriving on the

substrate; is the activation energy of the desorption; is the temperature relating

to the activated complex; kB is the Boltzmann constant; A is the coefficient.

The gas pressure can affect the growth rate by two aspects: (1) with decreasing

gas pressure, the energy of the ablated species increases, corresponding to an increase

in the number of adatoms; (2) the decreasing gas pressure also can led to an increase

38

in the rate of desorption. Both the rate of arrival and the rate of desorption of the

species in a gas pressure could finally determine the growth rate. In our case, the

desorption process is the main factor that influence the deposition rate. At high

oxygen pressure, the desorption of the adatoms decreases, thus result in the relative

high growth rate of the Ga2O3 films.

3.2.2 Crystal structure

Figure 3.2 XRD patterns of Ga2O3 films deposited on (0001) sapphire substrates with different

oxygen pressure.

Figure 3.2 shows XRD patterns of Ga2O3 films deposited on sapphire c plane

with different oxygen pressure. All the peaks were indexed as (-201) and higher order

diffractions of β-Ga2O3, indicating the orientation relationship β-Ga2O3

39

(-201)//sapphire (0001). It is obviously that with the decrease of oxygen pressure, the

peak intensity decreased, indicating the decreased crystal quality with the decrease of

oxygen pressure.

3.2.3 Transmittance and surface morphology

200 400 600 800 10000

20

40

60

80

100

T

rans

mitt

anc

e (%

)

Wavelenth (nm)

10-1Pa

10-2Pa

10-3Pa

10-4Pa

Figure 3.3 Optical transmission spectrum of the thin films deposited at different oxygen pressure.

All films have high transmittances of higher than 85%.

Figure 3.3 shows the optical transmission spectrum of the thin films deposited at

different oxygen pressure. The spectrum exhibited a high transmittance of over 85%

with clear fringes in the visible and UV regions. The spectrum decreased rapidly at

around 250 nm for all of the films. No obvious difference has been observed for films

deposited at different oxygen pressure.

Figure 3.4 shows the AFM morphology of films deposited at different oxygen

pressure, respectively. Island like structure appeared in all samples. Film deposited at

10-1 Pa shows lowest aspect ratio island which looks like a sharp mountain. Film

deposited at 10-4 Pa shows highest aspect ratio island which appears like ball. The

surfaces of all the films are very uniform. The roughness values of the films deposited

40

at oxygen pressure from 10-4 Pa to 10-1 Pa are 1.0, 1.6, 1.9 and 3.4 nm, respectively,

indicating very smooth surfaces of these films.

Figure 3.4 AFM (4 μm×4 μm) surface plots (right) and top views (left) of Ga2O3 films

deposited at different oxygen pressure. (a and a’)10-1 Pa , (b and b’) 10-2 Pa, (c and c’) 10-3 Pa and

(d and d’) 10-4 Pa, respectively.

41

3.2.4 Discussions

It has been stated that the background oxygen pressure mainly plays two

important roles in the deposition process. One is to supply oxygen atoms for the film

and to prevent the desorption of ad-atoms; the other is influences on the plume shape

and propagation speed.

It has been observed that crystal quality increase with the increasing of oxygen

pressure. This is because that oxygen pressure helps to preserve the Ga content in the

film through the formation of gallium oxide in the gas phase. The oxygen will interact

with the species in the plume like plasma. It is also possible that the oxidation of those

species will generate additional heat and energy thus assists the crystallization of the

Ga2O3 film.

Another mechanism is the oxygen pressure influence on the kinetic energy of the

ablated plumes. The oxygen pressure also influences the kinetic energy of the

impinging species as it is easy to imagine the role of the deposition pressure in

slowing down the plume. One of the kinetic parameters, the plume stopping distance

(that is the average path of the ablated species before they thermalize) can be

expressed as following 22:

√

(3-2)

where represents the mean diameter of species, while is the temperature of

plume-like plasma generated by PLD process and p is the pressure in the chamber. It

is clear that, when fixed the other parameters, the value of is inversely

proportional to that of and decrease the oxygen pressure will increase the value of

. Sambri et al. 23 have investigated the growth parameters influence on the

growth mode of SrTiO3 films. They found that when the plume stopping distance

is much larger than that of the distance between substrate and target ,

island mode dominates. When the has similar value with that of the , the

gowth mode changes to layer by layer. In our case, as the film deposited under the

highest oxygen pressure (10-1 Pa, which has the smallest value of ) showed

island growth mode, thus the plume stopping distance is much larger than that

42

of the distance between substrate and target in all of films deposited with the oxygen

pressure range from 10-1 Pa to 10-4 Pa. And they are all deposited with island mode. A

further increase of the kinetic energy (decrease the oxygen pressure) seems deteriorate

the crystallinity of the growing film, probably by creating small defects on the surface

of the film that cannot be recovered by the ad-atoms growth kinetics. From this view

of point, to improve the crystallinity of the Ga2O3 films, further increase of the

oxygen pressure is suggested.

3.3 Substrate temperature influence

The role of substrate temperature in the deposition process of PLD is similar

with in other film growth techniques. The atom mobility on the surface and within the

bulk will be increased at high temperature. And that will cause film structure to

progress from amorphous to polycrystalline and further to single crystal epitaxial

growth. In order to investigate the substrate temperature influence, the Ga2O3 films

were prepared by PLD using a KrF excimer laser source (λ=248nm) on (0001)

sapphire substrates. Facing the substrate, Ga2O3 (99.99%) was set as the target. The

pulsed laser with a frequency of 1 Hz was irradiated with a target-substrate distance of

about 40 mm. The oxygen pressure was set at 0.1 Pa while substrate temperature was

varied from 200 °C to 600 °C. The deposition time was 3 hours for all films. The

growth conditions are listed in table 3.2.

Table 3.2 Growth condition for substrate temperature influence.

Targets Ga2O3


Oxygen pressure 1×10-1Pa

Substrate temperature 200~600 oC


Laser power 225 mJ

43

3.3.1 Crystal structure

(1) XRD analysis


substrate temperatures.

Figure 3.5 shows the XRD patterns of Ga2O3 films deposited on (0001) sapphire

substrates with different temperature. When substrate temperature is below 400 °C, no

peaks can be observed, indicating the amorphous structure of these films. When the

substrate temperature is increased to 500 °C, three different diffraction peaks located

at 2θ value of 18.89°, 38.05° and 58.93° appear. They can be assigned as (-201), (-402)

and (-603) planes of β-Ga2O3, respectively. The result suggests that (-201) oriented

β-Ga2O3 can be epitaxially grown on (0001) sapphire substrates. The decreased peak

intensities at substrate temperature 600 oC are suggested due to the re-evaporation of

Ga2O3 at high temperature.

44

Figure 3.6 The distribution of oxygen atoms on the surface of the c-plane sapphire (a) and

β-Ga2O3 (-201) plane (b).

(-201) oriented β-Ga2O3 grown on (0001) sapphire substrates is due to the fact

that the arrangement of oxygen atom on c plane of sapphire is similar to that on the

(-201) plane of β-Ga2O3. Figure 3.6(a) reveals the distribution of oxygen atoms on the

surface of the c-plane sapphire. The spacing between the oxygen atoms affected by

aluminum atoms is in the range 0.251–0.287 nm. The oxygen atoms arrangement in

the (-201) plane of β-Ga2O3 as shown in figure 3.6 is similar to the oxygen atoms of

a

b

b

c

a

(a)

(b)

45

the (001) c-plane sapphire. In addition, the distance between the oxygen atoms in the

direction of the b-axis is 0.304 nm and the distances between the other oxygen atoms

are about 0.289 nm. Therefore, gallium atoms can bond to the oxygen atom layer

almost without feeling the difference between the (001) c-plane sapphire and the

(-201) plane of β-Ga2O3. This will lead to the formation of (-201) oriented β-Ga2O3

film on the (001) c-plane sapphire. Another advantage of the using sapphire as

substrate is that the Ga2O3/sapphire structure can be directly used in UV or

transparent devices as sapphire possesses big bandgap.

0 50 100 150 200 250 300 350

-Ga2O3 {-4 0 1}

Inte

nsity

(ar

b.un

it)

Figure 3.7 The in-plane XRD Φ scan result for the β-Ga2O3 (-4 0 1) diffraction of film grown at

substrate temperature of 500°C.

Figure 3.7 shows the in-plane XRD Φ scan result for the β-Ga2O3 (-4 0 1)

diffraction of film grown at substrate temperature of 500 °C. Six peaks that appear

every 60° indicate sixfold in-plane rotational symmetry. The monoclinic β-Ga2O3 (-4

0 1) planes originally have twofold in-plane rotational symmetry. Thus it can be

concluded that the grown film contains in-plane rotational domains. The threefold

rotational symmetry of the c-sapphire substrate surface is suggested to be the reason

of the in plane rotational domains of the β-Ga2O3 films; It is equivalent for the

originally twofold β-Ga2O3 to grow along the three equivalent directions of sapphire

46

substrate, resulting in the sixfold rotational symmetry. Similar result was also reported

by Oshima et al. using MBE method 14.

(2) Reflection High Energy Electron Diffraction (RHEED)

Figure 3.8 The RHEED patterns of Ga2O3 film deposited at 500 oC on sapphire substrate along

direction of (a) 0 o, (b) 90 o, (c) 180 o and (d) 270 o, respectively. The 90 o direction is parallel to

the [11-20] azimuth of the sapphire substrate.

RHEED is a surface sensitive technique. The electrons probe only topmost

atomic layers due to small incident angle (below 3 degs). Additionally the energy of

electrons is high (usually 10~20 keV). Therefore the Ewald's sphere radius is large. A

streaky RHEED pattern will appear in the case of smooth surface of a monocrystal.

As for the rough surface of a monocrystalline, the streaks become non-uniform to

form a spotty structure. This is a result of additional diffraction at parallel atomic

planes (in the case of rough surface the electrons can enter into "individual islands"

through their side surface). Spotty pattern can be observed as long as the lateral size

of the "islands" forming roughness is smaller than a characteristic coherence length.

When the surface of a polycrystalline film is probed, the pattern takes a form of rings.

The RHEED patterns of Ga2O3 film deposited at 500 oC on sapphire substrate

along different direction of 0 o (a), 90 o (b), 180 o (c) and 270 o (d) are shown in figure

3.8. All of the patterns showed streaky shapes, indicating the obtained film is

monocrystal along those directions. In order to verify the in plane relationship

(a) (b)

(c) (d)

47

between the films and the sapphire substrate, the RHEED patterns of the Ga2O3

substrate (from Tamura cooperation) is also measured as shown in figure 3.9.

Figure 3.9 The RHEED patterns of (-201) Ga2O3 substrate along directions of (a) 0 o, (b) 30 o, (c)

45 o, (d) 60 o, (e) 90 o, (f) 120 o and (g) 135 o, respectively. (h) The crystal orientation of the (-201)

β-Ga2O3 substrate. The 90 o direction is parallel to β-Ga2O3 [102].

The RHEED pattern of the Ga2O3 film along direction of 90 o (the electron beam

parallel to the [11-20] azimuth of the sapphire substrate) matches well with the pattern

along the 90 o azimuth of the Ga2O3 substrate (the electron beam parallel to the

orientation of β-Ga2O3 [102] substrate). Thus the growth relationship between the

Ga2O3 film and the sapphire substrate is that sapphire (0001)// β-Ga2O3 (-201) and

sapphire [11-20]// [102] β-Ga2O3. Hayashi et al. 24 have investigated the in plane

relationship between β-Ga2O3 and sapphire substrate using transmission electron

microscope, and have found similar results with ours.

(a) (b)

(c) (d)

(e) (f)

(g)

(h)

48

3.3.2 Optical properties

200 300 400 500 600 700 800 900 10000

20

40

60

80

100

Tra

nsm

itta

nce

(%

)

Wavelength (nm)

200C 300C 400C 500C 600C

Figure 3.10 Transmittance spectra of Ga2O3 films prepared at different substrate temperatures.

Transmittance spectra of the films prepared at various substrate temperatures are

plotted in figure 3.10. The transmittances of all the films in infrared region are above

80%, and exhibit clear fringes in the visible and UV regions. The films deposited at

substrate temperature higher than 400 oC show high transmittance for wavelength

lower than 350 nm and are transparent in appearances. Films deposited at substrate

temperature of 200 and 300 oC show decreased transmittance from wavelength 500

nm to 300 nm and their colors are brown in appearance.

β-Ga2O3 is a material that exhibits direct transition, although recent research also

showed the existing of a slightly smaller and weak indirect bandgap 25, 26. Analysis of

the dipole matrix elements by Varley et al. 27 based on density functional theory using

novel hybrid functional reveals that while the vertical transitions are dipole-allowed at

the Г point and at the valence band maximum (VBM), they are much weaker at the

VBM and decrease to 0 at the M-point rapidly. The similar energy of indirect and

direct gaps as well as the weakness of the indirect transitions effectively makes

β-Ga2O3 a direct-gap material, consistent with the experimentally observed sharp

absorption onset.

49

3.0 3.2 3.4 3.6 3.8 4.0 4.2 4.4 4.6 4.8 5.0 5.20

5

10

15

20

(ah)

2 (

10

10eV

2cm

-2)

h (eV)

200C 300C 400C 500C 600C

Figure 3.11 (αhν)2 vs. hν plot of Ga2O3 films deposited at different substrate temperatures.

It is well known that for films with direct bandgap, the absorption follows a

power law of the form:

/ (3-3)

where is the absorption coefficient, is the absorption edge width parameter,

is the energy of the incident photon, and is the bandgap. The optical absorption

coefficient αof the films is calculated using the standard relation taking the film

thickness into account 28. Figure 3.11 is the plot of as a function of photon

energy. The absorption coefficient increases rapidly at the photon energy range from

4.0 to 5.2 eV, depending on the substrate temperature. The plots of ~ fit the

straight line quite well. The result indicates that the Ga2O3 films are of direct

transition. The bandgap value about 4.9 eV at 500 and 600 °C, which agrees with that

of bulk Ga2O3 29. It is noticeable that the optical bandgap values of the films grown at

low substrate temperature such as 200, 300 and 400°C are smaller than that of the

bulk Ga2O3.

3.3.3 Surface morphology

The substrate temperature influences on the film thickness is shown in figure

3.12. The film thickness decreases almost linearly with the increase of substrate

temperature above 300 °C, which is due to the re-evaporation of the adsorbed spices

50

on the surface of the substrate.

200 300 400 500 600200

300

400

500

600

700

Thi

ckne

ss (

nm)

Substrate temperature (C)

Figure 3.12 Substrate temperature influences on the film thickness.

The surface AFM morphologies and roughness of films deposited at different

substrate temperatures are shown in figure 3.13 and figure 3.14. The morphologies are

different at different substrate temperatures as shown in figure 3.13. When the

substrate temperature is lower than 300 °C, very flat surfaces appear. The surface of

the film deposited at substrate temperature of 400 oC is composed of slim needle like

structures as shown in figure 3.13 (c). The surface morphologies change to islands

when substrate temperature is higher than 500 °C as shown in figure 3.13 (d) and (e).

The roughness of the films increases with substrate temperature when the substrate

temperature is not higher than 500 °C as shown in figure 3.14. The roughness values

of all the films are below 6 nm, indicating smooth surfaces.

It has been reported that thicker film often possess rougher surface 30. However,

the thickness of films decreases almost linearly with the increase of substrate

temperature when it is higher than 300 °C, as shown in figure 3.12. That indicates the

film thickness different is not the main reason of the roughness difference in present

experiments. For example, the film deposited at 200 °C shows smallest roughness but

a larger thickness. Thus, we ascribe the roughness difference to the growth

mechanism at different substrate temperatures.

51

Figure 3.13 AFM (4 μm×4 μm) surface plots (left) and top views (right) of Ga2O3 films deposited

at different substrate temperatures. (a and a’) 200 °C, (b and b’) 300 °C, (c and c’) 400 °C, (d and

d’) 500 °C, and (e and e’) 600 °C, respectively.

52

200 300 400 500 600-1

0

1

2

3

4

5

6

7

RM

S (

nm

)

Substrate temperature (C)

Figure 3.14 Dependence of roughness of Ga2O3 films on substrate temperature.

The changing of surface morphologies with temperature is supposed due to the

change of energy/mobility of ad-atoms. When the substrate temperature is low, the

atomic jump process period of ad-atoms on the substrate surface is very long. Thus

the species may stay stuck to the regions where they are landing 31. This process can

be observed in the films deposited at temperature lower than 300 °C. And because the

energy of ad-atoms is not sufficient for crystallization, the films are amorphous. The

mobility of ad-atoms on the surface increases with the substrate temperature increased

to 400 °C. Assemble of ad-atoms began. But at this temperature, the mobility of those

ad-atoms is not higher enough, only the ad-atoms near the nuclei can move to the

nuclei and thus contribute to crystallization. And thus slim needle like structure

appears as has been observed in the AFM morphology of figure 3.13(c). It is worth

noting that if the substrate structure is not similar to that of the film, the appeared

morphology should be spherical shaped grains in order to decrease the surface free

energy 31. With further increasing of the substrate temperature, the mobility of

ad-atoms increases, that will cause the nearby structures to joint together thus islands

like structure appears as shown in the AFM images figure 3.13(d) and (e). The film

growth mode is island as judging from the AFM morphologies.

53

3.3.4 Valence band structure

The real understanding on electronic states of Ga2O3 is very important for its

industrial applications. However, the mostly utilized methods for the investigation of

electronic structure of Ga2O3 are photoemission spectroscopy and X-ray absorption

spectroscopy. Unfortunately, both of them are much surface sensitivity and cannot

directly observe the bulk electronic states. On the other hand, Hard X-ray

Photoemission Spectroscopy (HAXPES) has a large detection depth that is sufficient

for the observation of bulk sensitive electronic states in a precise and non-destructive

manner. The HAXPES measurements were carried out at with a VG Scienta R4000

hemispherical electron analyzer at undulator beamline BL46XU of SPring-8. The

excitation X-ray beam was monochromatized with Si (111) double crystal and Si (444)

channel-cut monochromators.

The HAXPES valence band spectra are plotted in figure 3.15. The films

deposited at 500 oC and 600 oC which are of monoclinic crystalline structure have

similar valence band spectrum. The spectra of amorphous films (deposited at 300 and

400 oC) are somewhat different from that of the crystalline film.

For both of the crystalline and amorphous films, the peaks around 23 eV and 20

eV are mainly due to the contributions of O 2s and Ga 3d bands, respectively. The

peak and shoulder structures from 2.7 to 8 eV of VB structures of the crystalline films

(500 oC and 600 oC) mainly due to the hybridization of Ga 4sp orbits. The valence

band structures of these films are in good agreement with the reported structure of

β-Ga2O3 single crystals by HAXPS by Li et al.32. And, the valence band structures of

these films are of much different from the reported results of 150 eV (the voltage used

in laboratory XPS experiments), where three prominence maxima are assigned as O

2p 33. Li et al.32 contributed this difference to the different PICSs and DADs of Ga

and O elements at the different photon energies of 8 keV (HAXPS) and 150 eV. The

PICSs of Ga 4s and 4p of 150 eV are nearly 0.232 and 0.134 times of that of O 2p,

respectively. While these values of 8 keV are 170 and 35 times, respectively. It means

54

that at 150 eV the O 2 p is predominated, while at 8 keV the Ga 4s and 4p govern the

VB. Therefore, the VB by HAXPES can be assigned to the main contributions of Ga

4sp 32. The tailed valence structure of films deposited at 300 and 400 oC is contributed

to the amorphous nature of these films which associated with the disorder and defects

in the amorphous films.

35 30 25 20 15 10 5 0 -5

300 oC

Ga

3d

Binding energy (eV)

O 2

s

Inte

nsity

(ar

b.un

its)

400 oC

500 oC

Ga

4sp

Ga

4s600 oC

Figure 3.15 The HAXPES valence band spectra of films deposited at different temperature.

3.4 Growth time influence

By the above experiments, we have optimized the growth temperature and

oxygen pressure for high quality β-Ga2O3 film by PLD at present work as 500 oC and

10-1 Pa. In order to make the growth mechanism more clear, films with different

growth time were prepared at the optimized substrate temperature and oxygen

55

pressure of 500 oC and 10-1 Pa. The growth condition is listed in table 3.3. The

thickness of the films increased lineally with the deposition time as shown in figure

3.16.

Table 3.3 Growth condition for growth time influence.

Targets Ga2O3


Oxygen pressure 1×10-1Pa



Laser power 225 mJ

Growth time 10,30,60,180 min.

0 20 40 60 80 100 120 140 160 180 200

0

100

200

300

400

Thi

ckn

ess

(nm

)

Growth time (min)

Figure 3.16 The thickness dependence on growth time of Ga2O3 films by PLD.

56

Figure 3.17 AFM surface (4 μm×4 μm) plots (left) and top views (right) of Ga2O3 films

deposited with different time. (a and a’) 180, (b and b’) 60, (c and c’) 30 and (d and d’) 10 min,

respectively. (E and f) are the surface plots of films grown with 30 and 10 min with smaller scale.

57

Figure 3.17 shows the surface plots and top views of Ga2O3 films deposited

with different time. The films exhibit granular top view morphologies as observed

from figure 3.17 (a´) to (d´). The grains of the film deposited with 10 min are difficult

to recognize. As growth time increased to 30 min, the grains become apparent as

shown in figure 3.17 (c´). With the increase of growth time, the grain size becomes

bigger by the coalescence of the gains (figure 3.17 (b´) and (a´). The surface plots of

these films on the left side reveal that all of these grains are of conical shapes which

are vertical to the substrate as shown from figure 3.17 (a) to (f). With the increase of

growth time, both the length and the diameter of the cones increase.

The surface depth evaluated from the AFM morphologies together with the

thickness of the films are listed in table 3.4. It is ready observed that the surface depth

part relative to the whole film thickness decreases with the deposition time, from

about 33% of film with thickness of 12 nm to 14% of film with thickness of 387 nm.

Table 3.4 Thickness of surface of films deposited with different time.

Deposition

time (min)

Total thickness

(nm)

Surface depth

(nm)

Surface depth/film

thickness

10 12 3~5 0.33

30 53 10~12 0.21

60 126 25~30 0.21

180 387 50~60 0.14

58

3.5 Annealing effects

Annealing is a heat treatment that alters the physical and sometimes chemical

properties of a material. It involves heating a material to above its recrystallization

temperature, maintaining a suitable temperature, and then cooling. Annealing occurs

by the diffusion of atoms within a solid material, so that the material progresses

towards its equilibrium state. Heat increases the rate of diffusion by providing the

energy needed to break bonds. The movement of atoms has the effect of redistributing

and eradicating (to a lesser extent) the defects in the material. Annealing can also

release the internal stress of a material. The relief of internal stresses is a

thermodynamically spontaneous process; however, at room temperature, it is a very

slow process. The high temperature at which annealing occurs serve to accelerate this

process.

The three stages of the annealing process that proceed as the temperature of the

material are: recovery, recrystallization, and grain growth. The first stage is recovery,

and it can remove the defects such as dislocations and the internal stresses they cause.

It occurs at the lower temperature stage of all annealing processes and before the

appearance of new strain-free grains. At this stage, the grain size and shape do not

change 34. The second stage is recrystallization, where new strain-free grains nucleate

and grow to replace those deformed by internal stresses. If annealing is allowed to

continue once recrystallization has completed, then the third stage of grain

growth begins 34.

Table 3.5 Annealing conditions.

Films for annealing Ga2O3 as-grown at room temperature

Ga2O3 as-grown at 500 oC

Annealing temperature 750, 825, 900 oC

Annealing atmosphere air

Annealing time 1 h

59

In order to release the probably existing strain and improve the crystal quality of

the Ga2O3 films, annealing was processed at different temperature in atmosphere with

an electric furnace. The samples were put on a quartz boat which was placed into the

center of the electric furnace. The samples were held at the annealing temperature for

l h in an air atmosphere and then furnace cooled to room temperature. The annealing

conditions are shown in table 3.5.

3.5.1 Annealing effect on films deposited at RT

Figure 3.18 XRD patterns of Ga2O3 films annealed at different temperature. The film before

annealing (As-grown) is deposited at RT by PLD. Peaks not assigned belong to the sapphire

substrate.

We at first investigated the annealing effects on the Ga2O3 films deposited at RT.

The thickness of the films are about 270~350 nm. It is clear that only one peak which

60

located at the 2 theta value about 38.3 o appears after annealing as shown in figure

3.18. This peak can be attributed to the (-402) face of the monoclinic β-Ga2O3.

However, the intensity of the only peak is so weak that it is comparable to the

intensity of the noise. Modreanu et al.35 called such phase as “a small amount of

crystallization phase with an amorphous phase”. They suggested it as evidence of a

transition from an amorphous to a crystallization phase.

200 400 600 800 10000

20

40

60

80

100

Wavelength (nm)

Tra

nsm

ittace

(%

)

As-grown 750C 825C 900C

Figure 3.19 Transmittance of Ga2O3 films deposited at RT annealed at different temperature.

Transmittance spectra of the films annealed at different temperatures are shown

in figure 3.19. It is clear that the Ga2O3 films show sharp absorption edge. After

annealing, the absorption edge shifts to shorter wavelength. The film without

annealing shows high transmittance. The transmittance values of films annealed at

temperature of 700, 800 and 900 oC in visible region are about 50%, 80% and 90%,

which increase with the increase of the annealing temperature. The color of the films

showed in table 3.6 matches the transmittance spectra quite well. The film without

annealing is of light yellow and transparent. The film annealed at 750 oC shows black

color of half transparent. That black color is lightened when annealing at a higher

temperature 825 oC and the film become transparent at 900 oC.

61

Table 3.6 Color of films annealed at different temperature.

As-grown Anneal at 750oC Anneal at 825oC Anneal at 900oC

Light

yellow/transparent

Black/ half

transparent

Light

black/transparent

Transparent

The appeared black color during the annealing process is suggested due to the

accumulation of Ga2O. The plume of the ablated spices during the PLD process is

composed of many types of material such as monoatomic atoms and ions, molecules,

clusters, particles, etc. The low valence species, presumably Ga2O should exists in the

film before annealing because the substrate temperature is so low thus cannot make

the low valence species oxidized or vaporized. Orita et al. 18 also contributed the

black color of Ga2O3 films grown by PLD at relative low temperature to the existence

of Ga2O. Thus, it is not difficult to image that when annealing the film at 750 oC, the

Ga2O began to evaporate from the inside of the film and accumulated at the near

surface region and caused the black color and the low transmittance of the film. At

higher temperature of 825 oC, the amount of Ga2O decreased. When temperature was

raised to 900 oC, there should be no existence of the low valence species Ga2O as

judging from the transparent film.

It could be conclude from the above results that the annealing temperature to

remove the low valence species and form the stoichiometric Ga2O3 film in present

work is about 900 oC. However, that temperature seems not high enough for fully

crystallize the Ga2O3 film deposited at RT by PLD because the XRD peak of the film

annealed is too weak.

Figure 3.20 is the AFM surface plots (right) and top views (left) of Ga2O3 films

deposited at RT annealed at different temperatures. The film without annealing shows

very smooth surface. The surfaces of films annealed are composed of rectangle blocks

with some grains on their tops. Those grains are suggested to be crystal while most of

the rectangle blocks are still of amorphous. The grains only occupied parts of the film

surface, indicating not good crystallinity of those films. Most of those grains appeared

at the interface of two rectangle blocks or the interface of two groups of rectangle

62

block arrays with different orientation as shown in figure 3.21. This is suggested due

to the relative high energy of the interface places which favorites the nucleation

process.

Figure 3.20 AFM (4 μm×4 μm) surface plots (right) and top views (left) of Ga2O3 films

deposited at RT annealed at different temperatures. (a and a’)as-grown, (b and b’)750 oC, (c and

c’)825 oC and (d and d’)900 oC, respectively.

63

Figure 3.21 Typical enlarged AFM top graphic of film annealed at 750 oC.

The above results reveal that there are several steps existing in the annealing

process of the films deposited at RT. The first is to evaporate the low valence species

while the second to crystalline the stoichiometric film. The utilized annealing

temperature of 900 oC is enough for the first step but not for the second one. Higher

annealing temperature is necessary to further improve the crystalline quality of those

films.

3.5.2 Annealing effect on films deposited at 500 oC

The results of the upper section suggest that the annealing temperature from 700

to 900 oC is insufficient to crystallize the Ga2O3 film deposited at RT by PLD. So can

it be used to further improve the crystallinity of the film deposited at 500 oC? Similar

annealing experiments were carried out for the films deposited at 500 oC.

The XRD patterns of Ga2O3 films annealed at different temperature are shown in

figure 3.22. All of the films show similar XRD patterns. All of the films are of high

(-201) oriented monoclinic structure. All of the films show same degree of diffraction

peak intensity, indicating that the crystallinity of the film deposited at 500 oC has not

by greatly changed by thermal annealing.

64

Figure 3.22 XRD patterns of Ga2O3 films annealed at different temperature. The film before

annealing is deposited at 500 oC by PLD.

The AFM morphologies of films annealed at different temperatures are shown in

figure 3.23. There is no obvious difference between the surface morphology of Ga2O3

film without annealing and the annealing ones, indicating that the surface morphology

of the film deposited at 500 oC has not by greatly changed by thermal annealing.

65

Figure 3.23 AFM surface plots (right) and top views (left) of Ga2O3 films deposited at 500 oC

annealed at different temperatures. (a and a’)as grown, (b and b’)750 oC, (c and c’)825 oC and (d

and d’)900 oC, respectively.

66

3.5.3 Annealing effect on CL spectra

200 300 400 500 600 700

10

10

CL

inte

nsity

(ar

b.un

its)

Wavelength (nm)

900oC

825 oC

750 oC no anneal

385 nm

Film deposited at RT and then annealled

Figure 3.24 CL spectra at RT from Ga2O3 deposited at RT and then annealed at different

temperatures.

200 300 400 500 600 700

Film deposited at 500 oC and then annealled

Wavelength (nm)

900 oC

825 oC

750 oC as-grown

CL In

tens

ity (

arb.

units

)

395nm

Figure 3.25 CL spectra at RT from Ga2O3 deposited at 500 oC and then annealed at different

temperatures.

CL spectra at RT from Ga2O3 deposited at RT and then annealed at different

temperatures are shown in figure 3.24. The film without annealing shows no obvious

peaks while other films show emission peak located at about 385 nm. Those peaks

67

covered a range from about 320 to 500 nm. CL spectra at RT from Ga2O3 film

deposited at 500 oC and then annealed at different temperatures are shown in figure

3.25. The film without annealing shows a broader peak located at about 415 nm while

the other films show emission peaks located at about 395 nm. Those peaks covered a

range from about 320 to 600 nm.

The luminescence of β-Ga2O3 has been discussed in the past. Broad emissions

have been observed in the UV and blue regions in intrinsic β-Ga2O3.2, 36, 37 The above

CL results suggested that: (1) film deposited at RT shows no emission peak; (2) film

deposited at 500 oC has relatively big blue/UV emission ratio; (2) annealing the film

deposited at 500 oC has decreased the blue/UV emission ratio; (3) films deposited at

RT and then annealed show smallest blue/UV emission ratio.

The UV emission was independent of growth conditions and dopant and has

been attributed to self-trapped excitons 36. Blue luminescence has been associated

with the recombination of a trapped electron in a donor with a trapped hole in an

acceptor. Shimamura et al. have investigated the photoluminescence of pure and

Si-doped β-Ga2O3 37

. They attributed the systematically diminishes of the blue

emission with Si concentration to a decrease in the oxygen vacancy VO (charge

compensation to the Si) because the VO can act as a donor which will trap the

electrons. The decreased blue/UV emission ratio for films after annealing in this study

is also contributed to the decrease of the oxygen vacancy.

3.6 Conclusions

In this chapter, we have investigated the oxygen pressure, substrate temperature,

deposition time influences on the structure and optical properties of Ga2O3 films

grown by PLD. The influence of post annealing was also been discussed.

(1) The crystal quality and the thickness of films deposited at 600 oC increase

with the increasing of oxygen pressure. The growth mode of the films is island mode.

It is suggested that at high oxygen pressure, the desorption of the ad-atoms decreases,

thus result in the relative high growth rate and high crystal quality of the Ga2O3 films.

68

(2) By varying the substrate temperature, the evolutions of the structure, surface

morphology and bandgap have been clearly observed. Films deposited at substrate

temperature below 400 oC show amorphous structure while those deposited at

substrate temperature higher than 500 oC are of high oriented monoclinic structure.

The changing of surface morphologies with temperature is supposed due to the

change of energy/mobility of ad-atoms. The bandgap of the crystalline films are about

4.9~5.0 eV while the amorphous films have much lower values. The hard X-ray

photoelectron spectroscopy reveals that the valence band of the crystalline films (500

oC and 600 oC) are mainly due to the hybridization of Ga 4sp orbits and O 2p orbit.

Additional peak has appeared in the valence band of the amorphous Ga2O3 films

deposited at 300 and 400 oC.

(3) The optimized growth substrate temperature and oxygen pressure for our

experiment is 500 oC and 0.1 Pa. The growth relationship between the Ga2O3 film and

the sapphire substrate is: sapphire (0001)// β-Ga2O3 (-201) and sapphire [11-20]// [102]

β-Ga2O3. The obtained β-Ga2O3 film is of sixfold in-plane rotational symmetry.

(4) By varying the deposit time (film thickness), the growth process has been

observed.

(5) Post annealing (annealing temperature from 700 to 900 oC) cannot be used to

obtain films with better crystal quality than the film deposited at the optimized growth

condition. The films with post annealing show smaller blue/UV emission ratio.

69

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71

Chapter 4

Effect of Si doping on properties of Ga2O3 films

4.1 Introduction

Wide bandgap materials offer the possibility of fabricating high power density

electronic devices, deep-ultraviolet transparent electrode, short wavelength optical

emitter and deep-ultraviolet transparent photodetector 1. Ga2O3 is a promising wide

bandgap material for its large bandgap and chemical and physical stabilities. It is well

known that the pure Ga2O3 is an insulator 2. As for semiconductor devices, suitable

amount of carrier density is indispensable. For example, the typical carrier density (Si

based devices) is 1016 cm-3 for the channel of metal-oxide-semiconductor field-effect

transistor (MOSFET); 1019, 1017, and 1015 cm-3 for the emitter, base and collector of

bipolar transistor, respectively 3. Thus selecting the appropriate doping elements to

improve the carrier density of Ga2O3 is very important. It is well known that Sn and Si

are efficient n-type dopants for bulk β-Ga2O3. Orita et al. 4 have obtained Ga2O3 films

with n type conductivity up to 1 Scm-1 by Sn doping at substrate temperatures above

800 oC. And the conductivity can be improved to 8.2 Scm-1 by optimizing the

deposition conditions 5. On the other hand, Varley et al. 6 have theoretically studied

donor impurities in β-Ga2O3 based on density functional theory and suggested that Si

is an efficient n type dopant. Actually, the effect of Si doping on the carrier density for

β-Ga2O3 bulk single crystals was proved experimentally. Villora et al. 7 showed that

the carrier density of bulk β-Ga2O3 can be intentionally controlled over three orders of

magnitude from 1016 to 1018 cm-3 by Si doping. Sasaki et al. 8 developed a donor

doping technique for β-Ga2O3 substrate by using Si-ion (Si+) implantation with

72

followed high thermal temperature annealing, the carrier density is reported to be 1019

cm-3.

However, for β-Ga2O3 thin films, the intentional control of the electrical properties

is still a remaining issue. Gogova et al. 9 have grown Si-doped β-Ga2O3 films by

metal organic vapor phase epitaxy (MOVPE). Although secondary ion mass

spectrometry showed that Si was incorporated in the films, Hall effect measurements

demonstrated that the resulting material was not electrically conductive. Takakura et

al. 10 also found that the conductivity does not increase by Si doping for β-Ga2O3

films grown by RF magnetron sputtering. Müller et al. 11 have tried to grow Si-doped

β-Ga2O3 thin films by PLD and obtained Ga2O3 film with conductivity of 0.2 S cm-1.

Unfortunately, they did not report on carrier density and optical properties of these

films.

In this chapter, we present on the Si doping influence on the properties of Ga2O3

films grown by PLD. The properties of Si doped films deposited with different Si

content under different growth conditions are discussed in detail. The growth


Table 4.1 Growth condition for Si doped Ga2O3

Targets Ga2O3 disks with different Si content (0~6 wt.%)


Oxygen pressure 1×10-4~1×10-1 Pa

Substrate temperature 200~600 oC


Laser power 225 mJ

4.2 Si content influence

Because controlling the carrier density of a film by its doping level is a common

idea, Ga2O3 disks with different silicon content (0, 1, 3 and 6 wt.%) were used as

targets in this study. The pulsed laser with a frequency of 2 Hz was irradiated with

73

energy of 225 mJ. High purity oxygen gas (99.999%) was introduced through a mass

flow controller and the oxygen pressure was set as 10-1 Pa. The target was rotated

during the growth to avoid crater formation. The substrate temperature was 500 °C.

The deposition time was 3 hours for all films. The thicknesses of non-doped, 1 wt.%

Si-doped, 3 wt.% Si-doped and 6 wt.% Si-doped films were 211, 252, 249 and 143

nm, respectively.

1 2 3 4 5

2 3 4 5

0wt%

6wt%

1wt%

Inte

nsity

(ar

b. u

nits

)

Energy (keV)

Si

3wt%

0wt.%

1wt.%

3wt.%

Inte

nsity

(a

rb. u

nits

)

Energy (keV)

6wt.%

O GaAl

Si

(a)

1 2 3 4 5 60

2

4

6

8

10

12

Si content in targets (at.%)7.86.23.1 9.34.7

Si content in targets (wt.%)

S

i cont

ent in

film

s (a

t. %

)

(b)

1.6

Figure 4.1 (a) Typical EDS of the films with various Si contents in the targets. Insert is the

enlarged figure. (b) The Si content in films obtained from the EDS measurement as a function of

Si content in the targets.

Figure 4.1(a) shows the typical EDS of the films with various Si content in the

74

targets. From the spectra, elements of oxygen, silicon, gallium and aluminum are

observed. The appearance of aluminum is ascribed to the sapphire substrate. From

figure 4.1(a) and the enlarged insert figure, it is clear that the intensity of Si peak

increases with Si content of the targets. The Si content in films obtained from the EDS

measurement is shown in figure 4.1(b) as a function of Si content in the targets. The

Si content in the films increases almost linearly with the increase of Si content in the

targets, indicating the composition of films can be tailored by changing the Si content

of the targets. The average Si contents of films deposited with 1 wt.% (1.6 at. %), 3

wt.% (4.7 at. %) and 6 wt.% (9.3 at. %) Si-doped targets are 1.1 at. %, 4.1 at. % and

10.4 at. %, respectively. The increase in Si concentration of film deposited from 6 wt.%

Si-doped target may be attributed to volatile Ga species generated by laser-ablation

during growth.

0 1 2 3 4 5 6 7 8 9 10 11

1010

1012

1014

1016

1018

1020

Si content (at. %)

C

arrie

r de

nsity

(cm

-3)

10-7

10-5

10-3

10-1

101

103

Con

duct

ivity

(S

cm

-1)

Figure 4.2 Carrier density and conductivity of the Ga2O3 films with different Si contents.

Figure 4.2 shows the Si content influence on the carrier density of the Ga2O3

films deposited on (0001) sapphire plane at substrate temperature of 500 oC. Hall

measurements show that the films are of n type. The carrier density of Ga2O3 films

has been greatly increased by 4 orders of magnitude: from an average value of

75

8.9×1015 cm-3 for non-doped Ga2O3 film to 9.1×1019 cm-3 for 1.1 at. % Si-doped film.

However, further increasing the Si content decreases the carrier density: the carrier

density of 4.1 at. % Si-doped and 10.4 at. % Si-doped films are 5.2×1017 cm-3 and

3.9×1011 cm-3, respectively. The conductivity of the films with different Si content is

shown on the right side logarithmic plot of figure 4.2. The variation tendency of the

conductivity is quite similar with that of the carrier density: the conductivity of Ga2O3

film has also been increased by 4 orders of magnitude, from an average value as low

as 2.2×10-4 S cm-1 for non-doped film to 2.0 S cm-1 for 1.1 at. % Si-doped film. We

note that this value is higher than the value (0.2 S cm-1) of Si-doped Ga2O3 film

reported by Müller et al. 11. Further increasing the Si content decreases the

conductivity: the conductivity is 8.9×10-3 S cm-1 for 4.1 at. % Si-doped and 2.4×10-7 S

cm-1 for 10.4 at. % Si-doped film. The donor activation ratio of 1.1 at. % Si-doped

film (with a Si concentration about 4.2×1020 cm-3) is about 22%, which is similar with

the ratio of the reported Si-ion implantation doping at the same impurity level 8.

The mobility is obtained to be 0.1, 0.1, 0.5 and 5.5 cm2V-1s-1 for non-doped, 1.1,

4.1 and 10.4 at. % Si-doped Ga2O3 films, respectively. The relatively high mobility

for 10.4 at. % Si-doped Ga2O3 film is suggested due to the amorphous like structure

of the film. Since electron transport paths of the amorphous oxides composed of

heavy-metal cations are made of spherically spreading metals orbitals, the amorphous

oxides often possess relatively high mobility 12, 13.

Figure 4.3 shows the typical XRD patterns of Ga2O3 films deposited on (0001)

sapphire substrates with different Si contents. The sharp peak at 2θ=41.6o can be

identified as the (0006) Bragg reflection of the sapphire substrate. Additional three

different diffraction peaks located at 2θ value of 18.9°, 38.2° and 58.8° appear at the

XRD patterns of films with Si content lower than 4.1 at. %. They are assigned to the

(-201), (-402) and (-603) planes of β-Ga2O3, respectively, based on the PDF card (No:

43-1012). The result suggests that highly (-201) oriented β-Ga2O3 films are grown on

(0001) sapphire substrates when Si content is lower than 4.1 at. %. The peak intensity

of the 10.4 at. % Si-doped film decreases dramatically. The corresponding RHEED

76

images for films with different Si contents are shown in the insert of figure 4.3. The

XRD as well as RHEED images indicate that the film with carrier concentration of

～1020 cm-3 (1.1 at. % Si-doped) has same grade of crystallinity with the non-doped

one and the crystallinity is degraded for higher Si-doped films. Degraded crystallinity

of Ga2O3 films with high Si-doping were also reported by other groups 9, 10.

Figure 4.3 Typical XRD patterns of Ga2O3 films deposited on (0001) sapphire substrates with

different Si contents. Insert is the RHEED images for films with different Si contents. The

electron beam is parallel to the [11-20] azimuth of the sapphire substrate.

Figure 4.4 shows the typical in-plane XRD Φ scan for the β-Ga2O3 (-401)

diffraction of the films. Six peaks appear every 60° at the diffraction of films with Si

content lower than 4.1 at. %, indicating that those films contain in-plane rotational

77

domains, which is the same as that of the β-Ga2O3 film deposited on sapphire by other

growth techniques 14, 15. No obvious peak has been observed for 10.4 at. % Si-doped

film.

0 50 100 150 200 250 300 350

non-doped

-Ga2O3 {-4 0 1}

10.4 at.%

4.1 at. %

Inte

nsi

ty (

arb

. units

)

1.1 at. %

Figure 4.4 Typical in-plane XRD Φ scan for the β-Ga2O3 (-401) diffraction of the films with

different Si contents.

Typical transmittance spectra of films with different Si content are plotted in

figure 4.5. All of the films exhibit high transmittance of over 90% in the visible region

and show clear fringes in the visible and UV regions. The optical absorption edge of

films with Si content lower than 4.1 at. % is almost the same with each other while

that of the 10.4 at. % film shifts towards longer wavelength. The result indicates that

the optical performance of the films with Si content lower than 4.1 at. % has no

obvious degradation compared with the un-doped sample.

The spectra of transmittance are converted into (αhν)2~hν plot (α denotes

absorption coefficient) as shown in the insert of figure 4.5. The absorption coefficient

increases rapidly at the photon energy range around 4.5-5.4 eV depending on the Si

content, and (αhν)2 as a function of hν fits the straight line quite well. By extrapolating

78

the linear part of (αhν)2~ hν to the horizontal axis, one can obtain the bandgap of the

films. The derived typical bandgap values of non-doped, 1.1, 4.1 and 10.4 at. %

Si-doped films are 5.11, 5.13, 5.18, and 4.75 eV, respectively. The bandgap of films

with Si content lower than 4.1 at % is almost the same with each other while that of

the 10.4 at % film decreases. The decreased bandgap of the 10.4 at. % Si-doped film

is ascribed to the amorphous structure of the film.

200 400 600 8000

20

40

60

80

100

4.4 4.8 5.2 5.60

20

40(a

hv)

2 (10

10 eV

2 cm-2)

hv (eV)

T

ransm

ittance

(%

)

non-doped 1.1 at. % 4.1 at. % 10.4 at. %

Wavelength (nm)

Figure 4.5 Typical transmittance spectra of Ga2O3 films with different Si contents. Insert shows

the (αhν)2 vs. hν plot of Si doped Ga2O3 films.

Figure 4.6 shows the typical surface morphologies of films with different Si

content. For films with Si content lower than 4.1 at. %, grain like structure appears as

shown in figure 4.6(a), (b) and (c). Only few gains appeared on the surface of 10.4 at. %

Si-doped film shown in figure 4.6(d). The maximum roughness of the films is 7.3 nm

as shown in figure 4.6(e).

79

0 2 4 6 8 100

2

4

6

8

10

RM

S (

nm)

Si content (at. %)

Figure 4.6 Typical AFM morphology of Ga2O3 films deposited on (0001) sapphire substrates with

different Si contents of (a) non-doped, (b) 1.1 at. % Si-doped, (c) 4.1 at. % Si-doped and (d) 10.4

at. % Si-doped, respectively. (e) Dependence of roughness of Ga2O3 films on Si content.

The carrier density of Ga2O3 film has been greatly increased by the 1.1 at. %

Si-doping, which suggest that Si can be act as effective dopant. Varley et al. 6 reported

the formation energies for donor impurities based on density functional theory and

suggested that Si is readily incorporated. The Si substituting on the Ga site is shallow

donor that may contribute to the n type conductivity of Ga2O3. Irmscher et al. 16 have

derived the donor ionization energy of β-Ga2O3 bulk crystals based on temperature

2μm

(a)

50nm 0 50nm 0

2μm

(b)

50nm 0

2μm

(c)

20nm 0

(d)

2μm

80

dependent conductivity and Hall effect measurements as well as deep level transient

spectroscopy. They reported that the ionization energy of the isolated donor is about

36 meV, and at room temperature nearly all donors should be ionized 16. Both the

above theoretical and experimental data together with the high carrier density

9.1×1019 cm-3 obtained by PLD in present work proved that Si is an efficient n type

dopant. It is noticeable that the formation energy of Si_Ga donors is higher in the

O-rich rather than the O-poor limit 6. That is a possible reason of the reported

nonconductive Si-doped Ga2O3 films grown by MOVPE 9 and sputtering 10 because

their films were both grown under O-rich conditions. On the other hand, the

instantaneous growth rate of PLD is very high because of the large flux of species

(~1019 species /cm2 s) incident onto the surface during each pulse, thus the oxidation

during growth can be kinetically limited by the availability of sufficient oxygen to

oxidize the high flux of cation species 17. Moreover, PLD is reported to be a growth

technical lacking in oxidizing activity 18, which benefits the Si substituting on the Ga

site in Ga2O3 films. Additionally, as PLD offers high controllability of chemical

composition 18, it is suggested that by using target with Si content lower than 1 wt%,

controllable carrier density is possible.

The decreased carrier densities of films with Si content higher than 1.1 at. % are

suggested mainly due to the decreased crystallinity as identified from the results of

XRD, RHEED and AFM. It is suggested that much of the Ga is not incorporating in

the desired substitutional site 19. In this way the Si is passivated and carrier density in

Ga2O3 film decreased 9. Another possible reason is the Ga vacancies because the Ga

vacancies can also be favorable compensating acceptors in the n type limit 20, 21. Very

recently, Korhonen et al. have studied the vacancy defects in undoped and Si-doped

Ga2O3 thin films by positron annihilation spectroscopy. The results show that Ga

vacancies are formed efficiently during growth of Ga2O3 thin films by MOVPE. Their

densities are high enough to fully account for the electrical compensation of Si doping

22.

81

4.3 Substrate temperature influence

It is necessary to investigate the substrate temperature influence on the electrical

properties of Si doped Ga2O3 films because the composition and microstructure of

as-deposited thin films are influenced by the substrate temperature during deposition.

The 1 wt.% Si-doped Ga2O3 bulk was selected as the target. The substrate

temperatures were varied from 200 to 600 oC while the oxygen pressure was 0.1Pa.

10 20 30 40 50 60 70 80

2 ()

200oC

300oC

Inte

nsity

(ar

b. u

nits

)

400oC

(-6

03

)

(-4

02

)

500oC

(-20

1)

600oC

Figure 4.7 XRD patterns of Si doped Ga2O3 films deposited on (0001) sapphire substrates with

different substrate temperatures.

Figure 4.7 shows the XRD patterns of Ga2O3 films deposited on (0001) sapphire

substrates with different temperatures. When the substrate temperature is lower than

300 °C, there is no peak except the (0006) reflection from sapphire substrate. When

substrate temperature increasing to 400 °C, four weak peaks have appeared, indicating

the beginning of the crystallization process. For films deposited at substrate

82

temperature of 500 °C, three strong diffraction peaks which can be attributed to the

(-201), (-402) and (-603) planes have been observed, indicating high orientation of the

obtained film. When the temperature reaches to 600 oC, the diffraction peaks become

weaker, indicating the deterioration of the crystallinity.

200 400 600 8000

20

40

60

80

100

Tra

nsm

ittan

ce (

%)

Wavelength (nm)

200oC

300oC

400oC

500oC

600oC

Figure 4.8 Transmittance spectra of Si doped Ga2O3 films prepared at different substrate

temperatures.

3.2 3.6 4.0 4.4 4.8 5.20

2

4

6

8

10

hv (eV)

(hv)

2 (10

10eV

2 cm-2)

200oC

300oC

400oC

500oC

600oC

5.01 eV

4.94

4.683.883.81

Figure 4.9 (αhν)2 vs. hν plot of Sn doped Ga2O3 films deposited at different substrate temperatures.

83

The transmittance spectra of films prepared at various substrate temperatures are

plotted in figure 4.8. All of the films grown above 400 oC exhibited high transmittance

of over 80% in the visible and the UV regions.

The spectra of transmittance were converted into (αhν)2~hν plot (α denotes

absorption coefficient) as shown in figure 4.9. The absorption coefficient increases

rapidly at the photon energy range around 3.8-5.0 eV depending on the substrate

temperature, and (αhν)2 as a function of hν fits the straight line quite well. The result

indicates that the obtained Ga2O3 films are of direct transition. By extrapolating the

linear part of (αhν)2~ hν to the horizontal axis, one can obtain the bandgap of the films.

The derived bandgap values of films deposited at substrate temperature of 200, 300,

400, 500, 600 oC are 3.81, 3.88, 4.68, 5.01 and 4.94 eV, respectively. The variation of

the bandgap is similar with the non-doped Ga2O3 films as reported in chapter 3.

It is known that the surface morphology of films is also important for

device fabrication. We investigated the surface morphologies of Si doped Ga2O3 films

by AFM as shown in figure 4.10. All of the film surfaces are very uniform in

appearance. For films deposited with substrate temperature higher than 400 oC, grain

like structure appears.

84

Figure 4.10 AFM top-view of Si doped Ga2O3 films deposited at different substrate temperatures.

(a) 200 °C, (b) 300 °C, (c) 400 °C, (d) 500 °C, (e) 600 °C, respectively.

The electron carrier density and the conductivity of Ga2O3 films obtained by Hall

measurement as a function of substrate temperature are shown in the logarithmic plots

in figure 4.11. Due to the insulating properties of sapphire, the conducting features of

the films are not affected by the sapphire substrate and are completely attributed to

Ga2O3 films. The electron carrier density continuously increases by over 10 orders of

magnitude, from a value as low as 2.0×1010 cm-3 for substrate temperature of 200 oC

to 3.0×1021 cm-3 for substrate temperature of 500 oC. It is noticeable that the carrier

0 10nm 0 10nm

0 100nm 0 100nm

0 100nm

(a) (b)

(d)

(e)

(c)

1μm 1μm

1μm 1μm

1μm

85

density increases abruptly for substrate temperature from 300 to 500 oC. Further

increasing the substrate temperature to 600 oC decreased the carrier density.

The conductivity of the films exhibits similar tendency as that of the carrier

density. Films deposited at substrate temperature lower than 300 oC show very low

conductivities of 10-8 Scm-1. The conductivity increases linearly with the substrate

temperature from 300 to 500 oC and reaches a top value about 1 Scm-1 at 500 oC.

Further increase of the substrate temperature to 600 oC decreased the conductivity.

200 300 400 500 600

1011

1013

1015

1017

1019

1021

Carrier density

Con

duct

ivity

(S

cm-1)

Car

rier

dens

ity (

cm-3)

Substrate temperature (oC)

10-9

10-7

10-5

10-3

10-1

101

Conductivity

Figure 4.11 Carrier density and conductivity as a function of substrate temperatures.

The variation tendency of the carrier density agrees well with that of the

crystallinity of the films. The films deposited at the substrate temperature lower than

300 oC have worst crystallinity and exhibit the lowest conductivity. The film

deposited 500 oC have best crystallinity and exhibits the highest conductivity. The

degraded crystallinity of films deposited at 400 and 600 oC also correspond to lower

conductivity. Thus it is easily concluded that the crystallinity is the main factor that

influences the conductivity of the films in present study. Xiu et al. also found that the

high crystallinity Mo-doped ZnO films have higher carrier density 23. Rao et al. 24

investigated the effect of substrate temperature on electrical properties of vacuum

deposited zinc telluride thin films. They also suggested that the decrease in the

86

resistivity with substrate temperature can be seen as a consequence of improved

crystallinity and stoichiometry of the films deposited at higher substrate temperatures.

The films deposited at lower substrate temperatures should have large number of

defect states due to incomplete atomic bonding or lack of stoichiometry, which results

in an increase of trapping states. These trapping states could trap carriers and then

reduce the number of free carriers. The films deposited at higher temperatures are

nearly stoichiometric and have better crystallinity. This would result in a decrease in

the defect states and hence an increase in the carrier concentration of the films.

4.4 Oxygen pressure influence

Because the oxygen pressure can influence the ablated plume as well as growth

process of the films, it also has the possibility to affect the conductivity of the Si

doped Ga2O3 films. The 1 wt.% Si-doped Ga2O3 bulk was selected as the target. The

oxygen pressure was varied from 1.0×10-4 to 1.0×10-1 Pa while the substrate

temperature was set as 500 oC. The thickness of the films deposited at oxygen

pressure of 1.0×10-1, 1.0×10-2, 1.0×10-3 and 1.0×10-4 Pa are 439, 384, 403 and 256 nm,

respectively. The decreased thickness with decreasing of oxygen pressure is attributed

to the evaporation of Ga species, which is similar with that of the pure Ga2O3

described in chapter 3.2.

Figure 4.12 shows XRD patterns of Ga2O3 films deposited on sapphire c plane

with different oxygen pressure. All the peaks were indexed as (-201) and higher order

diffractions of β-Ga2O3, indicating the orientation relationship β-Ga2O3

(-201)//sapphire (0001). There is no obvious difference on the XRD patterns of films

with different oxygen pressures, indicating the films have similar level of crystallinity.

It is noticeable that we have observed obviously degraded crystallinity with the

decreasing of oxygen pressure in non-doped Ga2O3 films grown at a higher

temperature of 600 oC. Thus it is concluded that the crystallinity dependence on

oxygen pressure is more sensitive at higher temperature. We suggest that the greatly

87

enhanced evaporation at higher temperature (600 oC) should correspond for that

variation.


oxygen pressure.

Figure 4.13 shows the optical transmission spectra of the thin films deposited at

different oxygen pressure. The spectrum exhibited clear fringes in the visible and UV

regions and decreased rapidly at around 250 nm for all of the films. The obtained

bandgaps are shown in the insert of the transmission spectra. All of the films have

similar bandgap values around 5.1 eV. However, it is noticeable that the

transmittances of film deposited at oxygen pressure of 1×10-3 and 1×10-4 Pa are lower

than that of the films deposited at higher oxygen pressure. On another hand, we have

88

observed “black lines” in films deposited at oxygen pressure of 1×10-3 and 1×10-4 Pa.

These “black lines” should correspond for the decreased transmittances. These “black

lines” are currently ascribed to the low valence species such as Ga2O, which exiting in

the ablated plumes but has not been fully oxidized due to the insufficient oxygen

pressure.

200 400 600 8000

20

40

60

80

100

3.6 4.0 4.4 4.8 5.20

20

40

hv (eV)

(ahv

)2 (

10

10e

V2cm

-2)

5.16

10-1 Pa

10-2 Pa

10-3 Pa

10-4 Pa

5.12 eV

Tra

nsm

ittan

ce (

%)

Wavelength (nm)

Figure 4.13 Optical transmission spectra of the thin films deposited at different oxygen pressure.

Insert shows the (αhν)2 vs. hν plot.

10-4 10-3 10-2 10-11015

1017

1019

1021

Con

duct

ivity

(S

cm-1)

Carr

ier

densi

ty (

cm-3)

Oxygen pressure (Pa)

-5

0

5

10

15

20

Figure 4.14 Carrier density and conductivity as a function of substrate temperatures.

89

The electron carrier density and the conductivity of Ga2O3 films obtained by Hall

measurement as a function of oxygen pressure are shown in the logarithmic plot of

figure 4.14. The films deposited at different oxygen pressure show similar levels of

conductivity (about 4 Scm-1) and carrier density (about 1020 cm-3). That means oxygen

pressure has no obvious influence on the electrical properties of the 1 wt.% Si doped

Ga2O3 films

For a long time, the conductivity of Ga2O3 was attributed to oxygen vacancies

(Vo). In present work, we have found that the oxygen pressure has little influence on

the conductivity of the 1 wt.% Si doped Ga2O3. The results suggested that the oxygen

vacancies are not the main contribution to the conductivity of the 1 wt.% Si doped

Ga2O3. Varley et al.6 have investigated the role of oxygen vacancies in the electrical

and optical properties of the transparent conducting oxide Ga2O3 and found that

oxygen vacancies are deep donors, and thus cannot explain the unintentional n-type

conductivity. Our results also proved that Si is an efficient n type dopant for the

Ga2O3 films.

4.5 Conclusions

We have investigated the Si doping influence on the structure and properties of

Ga2O3 films.

(1) Ga2O3 films with different Si content were grown on sapphire substrate at

500 oC by PLD. All of the films exhibit smooth surfaces and high transmittances. The

films of Si content lower than 4.1 at. % show high (-201) oriented monoclinic

structure. The carrier density of Ga2O3 film has been increased to 9.1×1019 cm-3 with

conductivity of 2.0 S cm-1 by 1.1 at. % Si doping. Further increase of Si content leads

to the decrease of carrier density.

(2) The film deposited at substrate temperature of 500 oC (1 wt.% Si doped)

shows lowest conductivity and highest carrier density while possesses best

crystallinity.

90

(3) Oxygen pressure has no obviously influence on the electrical properties of

Si-doped Ga2O3, indicating the oxygen deficiency is not the main origin of the carrier

in our study.

91

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M. Albrecht and R. Fornari, J. Cryst. Growth, 2014, 401, 665.

10. K. Takakura, S. Funasaki, I. Tsunoda, H. Ohyama, D. Takeuchi, T. Nakashima, M. Shibuya,

K. Murakami, E. Simoen and C. Claeys, Physica B: Condensed Matter, 2012, 407, 2900.

11. S. Müller, H. von Wenckstern, D. Splith, F. Schmidt and M. Grundmann, physica status solidi

(a), 2014, 211, 34.

12. H. Hosono, M. Yasukawa and H. Kawazoe, J. Non-cryst. Solids., 1996, 203, 334.

13. H. Q. Chiang, J. F. Wager, R. L. Hoffman, J. Jeong and D. A. Keszler, Appl. Phys. Lett., 2005,

86, 013503.

14. Y. Oshima, E. G. Vίllora and K. Shimamura, J. Cryst. Growth, 2015, 410, 53.


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17. A. Gupta, J. Appl. Phys., 1993, 73, 7877.

18. M. Hiratani, K. Imagawa and K. Takagi, J. Appl. Phys., 1995, 78, 4258.

19. P. Kozodoy, H. Xing, S. P. DenBaars, U. K. Mishra, A. Saxler, R. Perrin, S. Elhamri and W.

C. Mitchel, J. Appl. Phys., 2000, 87, 1832.

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92

Chapter 5

Growth and characterization of (Ga1-xInx)2O3 films

5.1 Introduction

As have previously stated, wide bandgap semiconductor materials have become

the hot spot of recent research for the possible applications in many fields such as

light emitting devices, power devices and flame detectors 1-4. Among all the wide

bandgap materials, β-Ga2O3 film with the monoclinic structure is considered as a

promising candidate for its large bandgap and chemical and physical stabilities 3, 5-10.

However, the bandgap should be tuned to realize high sensitive wavelength-tunable

photodetectors, cutoff wavelength-tunable optical filters or to introduce shallow

impurity levels for good electronic properties 11. In2O3 is a candidate to realize the

bandgap engineering of Ga2O3 since both indium and gallium belonging to the same

elements group have similar electronic structures.

In chapter 1, we have reviewed that (Ga1-xInx)2O3 films have been prepared by

various method such as plasma-assisted molecular beam epitaxy 12 and sol-gel method

4. Oshima et al. 12 have grown (Ga1-xInx)2O3 films on sapphire substrates by MBE.

However, they only studied (Ga1-xInx)2O3 films with indium content lower than 0.35

and have not reported the bandgap value. Kokubun et al. 4 have tried to prepare

(Ga1-xInx)2O3 thin films using a sol-gel method on sapphire substrates and reported the

bandgap engineering. Unfortunately, the indium content is also limited within 0.35.

To the best of knowledge, there is no reports on bandgap engineering of (Ga1-xInx)2O3

films with indium content in the whole range.

In this chapter, we reported the bandgap engineering of (Ga1-xInx)2O3 films

with indium content in the whole range. The solution of the phase separation which

93

occurs due to the different nature structure of Ga2O3 and In2O3 was stated in detail.

5.2 Bandgap engineering of (Ga1-xInx)2O3 films

5.2.1 Growth parameters

In2O3 is a candidate to realize the bandgap engineering of Ga2O3 since both

indium and gallium belonging to the same elements group have similar electron

structures. The films of (Ga1-xInx)2O3 were deposited by PLD on (0001) sapphire

substrate at temperature of 500 oC. The targets are mixture of Ga2O3 and In2O3, with

In2O3 content (weight ratio) of 0, 0.1, 0.2, 0.3, 0.5, 0.7, 0.9 and 1.0. The growth


Table 5.1 Growth condition for (Ga1-xInx)2O3 films.

Targets Mixture of Ga2O3 and In2O3, with In2O3 content

(weight ratio) of 0, 0.1,0.2,0.3,0.5,0.7,0.9 and 1.0.


Oxygen pressure 1×10-1 Pa



Laser power 225 mJ

94

5.2.2 Optical properties

0.8 1.6 2.4 3.2 4.0 4.8

0

0.1

0.2

0.3

0.4

0.7

0.9

In

Inte

nsity

(ar

b. u

nits

)

Energy (keV)

Ga

O

Al

1.0

Indium content of the target

Figure 5.1 EDS of (GaIn)2O3 films deposited from targets with various In content.

Figure 5.2 In content (x) in the (Ga1-xInx)2O3 films as a function of the In content of the targets.

95

The thickness of the films determined by a surface step profile analyzer were

about 200~250 nm. The element composition in the prepared films was measured by

EDS. The EDS of (Ga1-xInx)2O3 films deposited from targets with various In content

are shown in figure 5.1. Elements of oxygen, indium, gallium and aluminum are

observed from the spectra. The observed aluminum element spectrum is coming from

the sapphire substrate as the thickness of the film was only about 200 nm. It is

obvious that the intensity of indium increases with the increase of indium content of

the targets while the gallium content decreases. The indium content (x) in

(Ga1-xInx)2O3 films calculated from the EDS spectra is shown in figure 5.2 as a

function of indium content of the targets. The indium content (x) in the films increases

almost linearly with the increase of indium content in the targets, indicating the

composition of (Ga1-xInx)2O3 films can be tailored by changing the indium content of

the targets. The obtained indium atomic content x in the films are 0.04, 0.16, 0.18,

0.33, 0.56 and 0.83, corresponding to the targets with indium content of 0.1, 0.2, 0.3,

0.5, 0.7 and 0.9, respectively.

200 300 400 500 600 700 800 900 10000

20

40

60

80

100

Tra

nsm

ittan

ce (

%)

Wavelength (nm)

x=0 x=0.04 x=0.16 x=0.18 x=0.33 x=0.56 x=0.83 x=1.0

Figure 5.3 Transmittance spectra of (Ga1-xInx)2O3 films prepared at substrate temperature of 500oC

on (0001) sapphire substrates.

Figure 5.3 shows the transmittance spectra of the films with various In contents.

96

A sharp absorption edge which is caused by the fundamental absorption of light for

(Ga1–xInx)2O3 films is observed. The absorption edges shift towards longer wavelength

with the increasing of In content (x). Films with In content x of 0, 0.56, 0.83 and 1.0

show high transmittances of higher than 80% in visible and infrared region. Films

with In content x from 0.04 to 0.33 show degraded transmittances. Film with In

content x of 0.18 shows lowest transmittance.

The electronic structure of (Ga1–xInx)2O3 has not been reported either by

calculation or experiment up to now. In fact, the research on the electronic structure of

both Ga2O3 and In2O3 is far from accomplished. Early experimental works on In2O3

single crystals showed that the onset of strong optical absorption is found to be 3.75

eV, but with indications of a much weaker absorption onset at 2.62 eV tentatively

assigned to indirect transitions 13. However, the studies published so far have not

provided clear evidence for the presence of an indirect bandgap 14. Recently, the

investigation on the electronic structure of In2O3 using a combination of valence band

X-ray photoemission spectroscopy, oxygen K-edge resonant X-ray emission

spectroscopy, and oxygen K-edge X-ray absorption spectroscopy by Piper et al. 15 has

experimentally rule out the possibility of an indirect bandgap. Thus, the separation

between weak and strong optical onsets is due to the fact that transitions of valence

bands into the conduction band are either symmetry forbidden or have very low

dipole intensity, and the electron structure of In2O3 is of direct bandgap 15, 16. On the

other hand, β-Ga2O3 is a material that exhibits direct transition, although recent

research also showed the existing of a slightly smaller and weak indirect bandgap 17, 18.

Analysis of the dipole matrix elements by Varley et al. 19 based on density functional

theory using novel hybrid functional reveals that while the vertical transitions are

dipole-allowed at the Г point and at the valence band maximum (VBM), they are

much weaker at the VBM and decrease to 0 at the M-point rapidly. The similar energy

of indirect and direct gaps as well as the weakness of the indirect transitions

effectively make β-Ga2O3 a direct-gap material, consistent with the experimentally

observed sharp absorption onset. Thus we believe that the transitions for both In2O3

97

and β-Ga2O3 are of direct bandgap and expect that the (Ga1–xInx)2O3 is also of direct

transition since indium occupies only octahedral site of Ga2O3 20 which has

exclusively constitution to the conduction band 21, yet up till now there is no

calculation reports on the electronic structure of (Ga1–xInx)2O3.

3.0 3.2 3.4 3.6 3.8 4.0 4.2 4.4 4.6 4.8 5.0 5.2 5.40

10

20

30

40

50

(hv

)2 (10

10eV

2 cm-2)

hv (eV)

x=0 x=0.04 x=0.16 x=0.18 x=0.33 x=0.56 x=0.83 x=1

Figure 5.4 (αhν)2 vs. hν plot of (GaxIn1-x)2O3 films with different In content (x).

0.0 0.2 0.4 0.6 0.8 1.03

4

5

Ban

dgap

(eV

)

Indium content x

This work Ref. 4 Ref. 22 Ref. 23

Figure 5.5 Dependence of bandgap of (GaxIn1-x)2O3 films on In content (x).

It is well known that by extrapolating the linear part of (αhν)2~hν to horizontal

axis, one can obtain the bandgap of a direct bandgap material. Figure 5.4 is the plot of

98

(αhν)2 as a function of photon energy of the (Ga1–xInx)2O3 films. The optical

absorption coefficient αof the films is evaluated using the standard relation taking

the film thickness into account 22. The absorption coefficient increases rapidly at the

photon energy range around 3.7-5.4 eV depending on the indium content x. The steep

parts are represented fairly well in linear form, indicating the allowed direction

transition in (Ga1–xInx)2O3. Therefore, the value of the bandgap can be determined by

the extrapolation of the linear part to the horizontal axis in figure 5.4. The derived

bandgap value as a function of indium content (x) is shown in figure 5.5. For

comparison, the reported bandgap in references 4, 23, 24 of (Ga1–xInx)2O3 are also

plottedin the figure. Both the bandgaps reported by other groups and our work and

decrease almost lineally with indium content (x). The results indicate that the bandgap

of (Ga1–xInx)2O3 can be controlled by changing the indium content (x), paving a way

to design optoelectronic and photonic devices based on this material.

5.2.3 Structure and surface morphologies

The XRD patterns of (Ga1-xInx)2O3 films deposited on (0001) sapphire substrates

with different In content (x) at substrate temperature of 500 °C are shown in figure 5.6.

Three different diffraction peaks located at 2θ value about 18.9°, 38° and 58.9° which

can be ascribed to the (-201), (-402) and (-603) faces of monoclinic β-(Ga1-xInx)2O3

appear when In content is lower than 0.04. With increasing the In content x to 0.16, a

peak attributed to cubic structured (Ga1-xInx)2O3 begins to appear at 2θ value about

30°, indicating the co-exist of two phases in the films. On the other hand, only the

peaks of cubic (Ga1-xInx)2O3 are observed when In content is higher than 0.56. The

simultaneously existing of both cubic and monoclinic structured (Ga1-xInx)2O3 is

called phase separation. The phase separation occurs mainly in the In content x range

from 0.16 to 0.33 as observing from figure 5.6.

99

Figure 5.6 XRD patterns of (GaxIn1-x)2O3 films with different In content (x) deposited on (0001)

sapphire substrates at substrate temperature of 500oC. Peaks marked by triangle (Δ) belong to

monoclinic structure while that marked by circle (Ο) belong to cubic structure.

The enlarged XRD patterns for (-402) face of monoclinic structure and (222)

faces of cubic structure are shown in figure 5.7 (a) and (b), respectively. The 2θ value

for the (-402) face of monoclinic phase decreases with the increase of In content (x).

Same tendency has been observed for the (222) face of cubic phase. These 2θ value

variations with In content reflect the increase of lattice constant with the increase of

indium content (x). This is an evidence that the indium /gallium has occupied the

lattice position of gallium/indium in the monoclinic/cubic structured (Ga1-xInx)2O3

films.

100

36 37 38 39 40

Inte

nsity

(ar

b. u

nits

)

2 ()

0

Ga1-x

Inx)

2O

3 (-402)

0.04

0.16

x=0.18 (a)

28 30 32 34

2 ()

0.33

0.56

0.83

(b)

Inte

nsi

ty (

arb

. un

its)

cubic Ga1-x

Inx)

2O

3 (222) x=1

Figure 5.7 Enlarged XRD patterns of cubic (222) (a), and monoclinic (-402) (b) faces of the

(GaxIn1-x)2O3 films, respectively.

Figure 5.8 shows the AFM surface plots of (Ga1-xInx)2O3 with different In content

x. The surfaces of the (Ga1-xInx)2O3 films with In content x of 0, 0.83 and 1.0 are very

uniform. These are projections on the surfaces of the (Ga1-xInx)2O3 films with In

content x from 0.04 to 0.33. These projections are very obvious for (Ga1-xInx)2O3 films

with In content x from 0.04 to 0.33, of which In content the transmittances have been

degraded as shown in figure 5.5. The roughness of the films with In content x of 0,

0.04, 0.16, 0.18, 0.33, 0.56, 0.83 and 1.0 are 2.1, 6.4, 72, 99, 37, 23, 6 and 2.1 nm,

respectively. The existing of those big projections has also greatly degraded the

surface quality of the (Ga1-xInx)2O3 films. The appearance of those projections is

suggested due to the phase separation.

101

Figure 5.8 AFM (4 μm×4 μm) surface plots of (Ga1-xInx)2O3 with different In content x of (a)0,

(b)0.04, (c)0.16, (d)0.18, (e)0.33 , (f)0.56, (g)0.83 and (h) 1.0, respectively.

5.3 Thermal annealing impact on crystal quality of (GaIn)2O3 alloys

In the section before, we have demonstrated the growth and bandgap engineering

of (Ga1-xInx)2O3 films by PLD. The bandgap of (Ga1-xInx)2O3 films can be tailored

(a) (b)

(c) (d)

(e) (f)

(g) (h)

40 nm

600 nm

1000 nm

1000 nm

600 nm

300 nm

100 nm

60 nm

3

2

1 1

2

3μm

3

2

1 1

2

3μm

3

2

1 1

2

3μm 3

2

1 1

23

μm

3

2

1 1

2

3μm

3

2

1 1

2

3μm

3

2

1 1

2

3μm

3

2

1 1

2

3μm

102

between 3.8 eV and 5.1 eV by controlling the In content in the target. Unfortunately

phase separation (the simultaneously existing of both monoclinic and cubic phases)

was observed in some of the films. Suzuki et al. 25 have observed similar phase

separation phenomenon in (Ga1-xInx)2O3 films and have tried to solve it by the

fabrication of rhombohedral corundum-structured α-(Ga1-xInx)2O3 films. However

phase separation still appeared when In content was between 0.08 and 0.67. In order

to meet the requirements of device applications, it is necessary to suppress phase

separation.

In this section, we report the thermal annealing effects on the (GaIn)2O3 films

(We will use (GaIn)2O3 to represent the films for convenience unless intentionally

mentioned.) with nominal In content of 0.3 (the indium content in the target). This

composition was selected because it showed obviously phase separation as observed

from both the XRD and AFM shown in figure 5.6 and 5.8, respectively.

It is necessary to find the way to obtain (GaIn)2O3 films with desired chemical

composition, but without phase separation. Annealing is a common way to recover the

defects and improve the crystal quality. However, additional experiments showed that

once the phase separation formed, it cannot be removed by annealing process. Thus,

we suggest a two-step method to achieve the goal. The first step of which is to obtain

the films without phase separation. The second step is to improve the crystal quality

of the film by annealing. In order to achieve the first object, (GaIn)2O3 films with

different substrate temperatures were grown on sapphire substrate by PLD. The films

were deposited on sapphire substrate at substrate temperature from RT to 500 oC. The

oxygen pressure was set at 0.1Pa. The target with In content of 0.3 (weight ratio) was

used.

103

Figure 5.9 XRD patterns of (GaIn)2O3 films deposited on (0001) sapphire substrates with

different substrate temperatures. Peaks marked by triangle (Δ) belong to monoclinic structure

while that marked by circle (Ο) belong to cubic structure.

The XRD patterns of the as-deposited (GaIn)2O3 films with nominal In content

of 0.3 grown at different substrate temperatures are shown in figure 5.9. The

“as-deposited” means films deposited just by PLD and no annealing has been carried

out. No peaks except for those from sapphire substrate are observed when the

substrate temperature is lower than 300 oC, indicating the amorphous structure of

these films. With increasing the substrate temperature above 400 oC, both the

characteristic peaks of monoclinic and cubic structure have appeared, suggesting the

existing phase separation in this film. At a substrate temperature of 500 oC, an

additional peak belonging to cubic structure appears.

104

Figurer 5.10 SEM and EDS mapping of In of (GaIn)2O3 films deposited at different substrate

temperatures. (a) and (a’) room temperature, (b) and (b’)200 °C, (c) and (c’)300 °C, (d) and

(d’)400 °C, (e) and (e’)500 °C, respectively.

SEM EDS

(a)

(b)

(c)

(d)

(e)

(a’)

(b’)

(c’)

(d’)

(e’)

2μm

2μm 2μm

2μm

2μm 2μm

2μm 2μm

2μm 2μm

105

Figure 5.10 shows the surface morphology together with the indium element

distribution of the (GaIn)2O3 films measured by means of SEM and EDS. The surface

of the films is very smooth and indium element distribution is homogenous when

substrate temperature is lower than 200 oC as shown in figure 5.10(a), (a’) and (b),

(b’). However it is clear that clusters have formed on the film surface when the

substrate temperature rises to 300 oC as observed in figure 5.10(c).

The morphology of the clusters changes with the substrate temperatures as

shown in figure 5.10(d) [400 oC] and (e) [500 oC]. The shapes of In element

distributions detected by EDS shown on the right side of figure 5.10 match well with

the results of SEM on the left side, indicating that the clusters are of In rich. These

clusters are ascribed to cubic structure while the back-body can be assigned as

monoclinic structure. The appeared phase separation at high substrate temperatures is

suggested due to the none-equilibrium process of PLD 26 because the In content in the

films doesn’t exceed the equilibrium solubility of indium in Ga2O3 (about 0.45) 21,

27-29. It is known that PLD is a process including laser ablation of the target material

(creation of a plasma), plasma plume expansion (plasma propagation) and deposition

of the ablation material on the substrate. Deposition of ablated species on the substrate

surface can be regarded as instantaneous for every pulse in PLD process. The ablated

particles exhibit shallow implantation into the upper monolayer of the surface

material. During the time interval of the laser pulse, the adatoms rearrange on the

surface by migration and subsequent incorporation through nucleation and growth.

The surface diffusion coefficient is many orders of magnitude bigger than the volume

diffusion 26. Thus, we believe that the phase separation at high substrate temperatures

is due to the different surface diffusion coefficient between indium and gallium. On

the other hand, when the substrate temperature is lower than 200 oC, the adatoms

cannot travel far for the low diffusion coefficient which decreases with an exponential

function of temperature 30, resulting in the homogenous structure of the films.

Next, thermal annealing was carried out on these as-deposited (GaIn)2O3 samples

in order to obtain crystalline film without phase separation. However, no obvious

106

effects on phase separation were observed on (GaIn)2O3 films grown at substrate

temperature higher than 300 oC. The observed clusters remained on the surface of the

films after thermal annealing treatment. Thus, the film deposited at RT was selected

for thermal annealing because it exhibits homogenous element distribution as shown

in figure 5.10(a) and (a’). Here, the RT deposited film is of amorphous, thus the

annealing process was expected to improve the crystallinity of the film because the

post thermal annealing has been reported to be an effective method to change the

properties of semiconductor films 31, 32. The samples were put on a quartz boat which

was placed into the center of an electric furnace. The annealing temperature was 800

oC. The samples were held at that temperature for l hour in an air atmosphere and then

furnace cooled to room temperature.

10 20 30 40 50 60 70

(-6

03

)

(-4

02

)

(-2

01

)

2()

Inte

nsity

(ar

b. u

nit)

Sa

pp

hir

e (

00

6)

Figure 5. 11 XRD pattern of (GaIn)2O3 film annealed at 800 oC in air. Film for annealing was

prepared at room temperature by PLD.

Figure 5.11 shows the XRD pattern of the (GaIn)2O3 film annealed at 800 oC in

air. It is clear that only peaks belonging to (-201) face of β-(GaIn)2O3 and its higher

order can be observed, indicating that oriented monoclinic structured film without

107

phase separation has been successfully obtained by post thermal annealing. The

orientation relationship between the substrate and film is: β-(GaIn)2O3 (-201) //

sapphire (0001), which is the same orientation as the as-deposited β-Ga2O3 on

sapphire substrate at 500 oC.

Figure 5.12 SEM (a), indium element distribution (b), and AFM (4 μm×4 μm) (c) morphology of

(GaIn)2O3 film annealed at 800 oC.

Figure 5.12 shows the SEM image (a), indium element distribution (b), and AFM

image (c) of the (GaIn)2O3 film annealed at 800 oC, respectively. The surface of the

film is very smooth with the roughness as low as 1.94 nm. The In element is

uniformly distributed as shown in figure 5.12 (b). Both the results from XRD in figure

5.11 and the surface morphology in figure 5.12 indicate that crystalline (GaIn)2O3

film with nominal In content of 0.3 without phase separation has been successfully

(a)

(b)

(c) 50nm

0 nm

2μm

2μm

108

obtained.

200 400 600 800

20

40

60

80

100

Tra

nsm

itta

nce

(%

)

Wavelength (nm)

As-deposited Annealing

Figure 5.13 Transmittance spectra of the annealing film and the one as-deposited at substrate

temperature of 500 oC.

Figure 5.13 shows the transmittance spectrum of the film annealed at 800 oC

together with the transmittance of the as-deposited film grown at 500 oC for

comparison. The transmittance of the annealed sample in infrared region is over 85%,

and exhibits clear fringes in the visible and UV regions, which is much higher than

the transmittance of the as-deposited film grown at 500 oC. On the other hand, two

absorption bands were observed in the transmission spectra for the as-deposited film

grown at 500 oC, which are attributed to phase separation of the film. Similar

phenomenon has been reported by other groups 12, 33, 34. The improvement of optical

properties of the annealed films verifies that thermal annealing is an effect way for

obtaining high quality (GaIn)2O3 film without phase separation.

5.4 Toward the understanding of annealing effects on (GaIn)2O3 films

In the above researches, we have observed that the bandgap of the (GaIn)2O3

films can be tailored between 3.8 eV and 5.1 eV by controlling the element

109

composition in the targets, however phase separation was observed in some of the

as-deposited films. Fortunately, further research found that by annealing the film

(nominal indium content of 0.3) deposited at room temperature in air ambient, high

quality β-(GaIn)2O3 film without phase separation can be obtained.

However, the mechanism of the annealing effect is still obscure because the

ambient air is a mixture of gases mainly composed of N2 and O2. And some

researchers found that inert gas annealing is effective for crystallizing the amorphous

Ga2O3 films 35-37. Huang et al. annealed the films at 800 oC under different ambient to

improve the crystallinity and found that the XRD of β-Ga2O3 annealed in a nitrogen

environment has the strongest peak while annealing in an oxygen environment shows

diminished performance due to the oxygen content reaching a saturation state during

annealing. In air ambient annealing which used both oxygen and nitrogen, the

observed crystalline properties were better than those obtained from purely oxygen

annealing 38. Thus, in this section, in order to understand the annealing effect, we

carry out further research on the (GaIn)2O3 film with same nominal indium content of

0.3 by investigating the annealing gas ambient and temperature influences.

The (GaIn)2O3 films for annealing were prepared by pulsed laser deposition on

(0001) sapphire substrates at room temperature. Post-annealing was carried out with a

quartz tube furnace with samples on a quartz boat placed into the center of it for one

hour. By changing the ambient (N2, vacuum, Ar, O2) and temperature (700~1000 oC),

the annealing parameters were optimized. The annealing effect on films with different

nominal indium content was investigated at the optimized temperature and gas

ambient. The total gas flow rate was 200 sccm for annealing.

5.4.1 Influence of annealing gas ambient

Figure 5.14 shows the X-ray diffraction patterns of (GaIn)2O3 film with nominal

In content of 0.3 annealed at 800 oC using different gas ambient. No obvious peaks

have been observed for films annealed in N2, vacuum and Ar ambient, indicating the

amorphous structure of these films. Film annealed at oxygen shows three diffraction

110

peaks which can be assigned as (-201) and its higher orders of monoclinic (GaIn)2O3

according to the JCPDS card (No: 43-1012). This reveals that the film becomes

textured with (-201) crystallographic planes oriented preferentially parallel to the

substrate surface after post annealing only at oxygen atmosphere.

10 20 30 40 50 60 70

10 20 30 40 50 60 70

2 ()

Inte

nsity

(a

rb. u

nits

)

N2

Ar

Vacuum(-

603)

(-4

02

)

(-2

01

)O

2

Figure 5.14 XRD patterns of (GaIn)2O3 films with nominal indium content of 0.3 annealed at 800

oC using different gas ambient. Peaks not assigned belong to the sapphire substrate.

The photos and AFM top morphologies of (GaIn)2O3 films as-deposited at RT

and annealed with different gas ambient are shown in figure 5.15. The as-deposited

film before annealing appears brown in color and becomes transparent after annealing

at 800 oC in all of the gas ambient as shown in figure 5.15(a). The surface

morphologies of these films are not same with each other. The surface of film without

111

annealing is very smooth and no grain has appeared as observed from figure 5.15(b).

The surface of film annealed in N2 is porous and grainy as shown in figure 5.15(c).

The surface of film annealed in vacuum appears somewhat inhomogeneous with

unclear grain boundary as shown in figure 5.15(d). And the surface of film annealed

in Ar as shown in figure 5.15(e) is also porous and grainy. The surface of the film

annealed in pure O2 shows uniformly distributed and tightly packed grains with

triangle shape as shown in figure 5.15(f). The surface roughness of the films annealed

in N2, vacuum, Ar and O2 is 14.2, 7.8, 3.7 and 1.2 nm, respectively.

Figure 5.15 Photos (a), AFM morphologies of (GaIn)2O3 films with nominal indium content

of 0.3 as-deposited at RT (b), and annealed at 800 oC under gas ambient of N2 (c), vacuum (d), Ar

(e) and O2 (f), respectively.

0 20nm

1μm

0

0 0 100nm

0 50nm

1μm

(e) (f)

(a)

1μm

(b)

(c)

1μm

(d)

1μm

20nm

150nm

112

1200 1000 800 600 400 200 0

Ga

LMM

C 1

sG

a LM

M

Ga

LMM

In M

NN

In 4

dG

a 3

dG

a 3

pG

a LM

M

In 3

d3/2

In 3

d5/2

In 3

p1/2

In 3

p3/2 O

1s

O K

LL

C K

VV

Ga

2p3/

2G

a 2

p1/2

Inte

nsity

(ar

b. u

nit)

Binding energy (eV)

Figure 5.16 The XPS wide scan spectrum of (GaIn)2O3 film with nominal indium content of 0.3

annealed at 800 oC under oxygen ambient.

The result of the XRD has showed that only oxygen ambient annealing can

crystallize the (GaIn)2O3 film while non-oxygen ambient (N2, vacuum and Ar)

annealing show amorphous structure. Thus it is suggested that oxygen vacancies

should exist in the (GaIn)2O3 films annealed in non-oxygen ambient as well as in the

film before annealing. In order to verify it, we carried out XPS measurement. The

typical XPS wide scan spectrum of the film annealed in oxygen is shown in figure

5.16. The peaks related to Ga, O and In elements are observed in the figure.

113

540 536 532 528 524

OIIIn

tens

ity (

arb.

uni

ts)

Binding energy (eV)

OI

(a)

540 536 532 528 524

OIIIn

tens

ity (

arb.

uni

ts)

Binding energy (eV)

OI

(b)

540 536 532 528 524

OIIIn

tens

ity (

arb.

uni

ts)

Binding energy (eV)

OI

(c)

540 536 532 528 524

OIIIn

tens

ity (

arb.

uni

ts)

Binding energy (eV)

OI

(d)

540 536 532 528 524

OIIIn

ten

sity

(a

rb. u

nits

)

Binding energy (eV)

OI

(e)

Figure 5.17 O 1s core level spectra with curve-fitting results obtained from (GaIn)2O3 films with

nominal indium content of 0.3 as-deposited (a), and annealed at 800 oC in N2 (b), vacuum (c), Ar

(d) and O2 (e) ambient.

Figure 5.17 shows the XPS core-level spectra of the O 1s peak for (GaIn)2O3

films annealed in different atmospheres. We used two components fitting to confirm

the oxygen vacancy because the oxygen vacancy often introduces another component

114

beyond the lattice oxygen component in the O1s peak 39. It is clear that the O 1s peak

can be consistently fitted by two near components, centered at 529.6 (OI) and 530.6

eV (OII) which is related to O2- on the lattice structure and oxygen vacancy,

respectively 39. The energy difference between OI and OII peaks for all films are about

1.0 eV. It is readily observed that the relative intensity ratios of OI /OII for (GaIn)2O3

film before annealing is the lowest, as shown in figure 5.17 (a). The OI /OII intensity

ratio has been increased by thermal annealing as shown from figure 5.17(b) to (e).

The (GaIn)2O3 film annealed in O2 shows the highest OI /OII ratio as shown in figure

5.17 (e). This suggests that thermal anneal has the effect to increase O2- on the lattice

structure and decrease oxygen vacancy. Similar results have been reported on Ga

doped ZnO films by Ahn et al. 40.

The oxygen vacancy in the (GaIn)2O3 film as deposited at RT can be contributed

to the low temperature deposition process of PLD. The film as-deposited at RT is of

great oxygen deficiency because low substrate temperature (RT) has restricted the

incorporation of oxygen into films 41. The decreased oxygen vacancy in films

annealed in non-oxygen ambient is suggested due to the disproportionation reaction

42-45 below:

GaIn O → GaIn O M （4-1）

where x is the oxygen composition in the films; M is the metallic gallium or indium or

the mixture. However, in our experiments, such disproportionation reaction did not

complete as none XRD peak of (GaIn)2O3 has been detected. On the other hand, the

oxygen ambient annealing has greatly helped on decreasing the oxygen vacancy thus

results in crystalline film.

5.4.2 Influence of annealing temperature

In order to optimize the annealing temperature, (GaIn)2O3 films with nominal

indium content of 0.3 were annealed at different temperatures in oxygen ambient.

115

10 20 30 40 50 60 70

10 20 30 40 50 60 70

Lo

g in

tens

ity (

arb.

uni

ts)

2()

700oC

800oC

900oC

(-6

03

)

(-4

02

)

1000oC

(-20

1)

Figure 5.18 XRD patterns of (GaIn)2O3 films with nominal indium content of 0.3 annealed at

different temperatures under oxygen ambient. Peaks not assigned belong to the sapphire substrate.

Figure 5.18 shows the XRD patterns of the films annealed at various

temperatures. It is clear that the film annealed at 700 oC shows amorphous structure as

no diffraction peaks can be observed. The films began to crystallization at temperature

of 800 oC, which agrees with the amorphous to crystalline transformation temperature

of Ga2O3 reported by other groups 46, 47. All films annealed at temperature higher than

800 oC exhibit only the (-201) peak and its higher orders, indicating that they

have preferred orientation due to self-texturing phenomenon. No secondary phases

and clusters can be observed from the XRD patterns. The rearrangement mechanism

may be responsible for the annealing effect. In general, crystallization is a process

116

strongly temperature dependent and is accommodated usually through a transfer of

matter by atomic diffusion. The temperature at which the atomic diffusion occurs

sufficiently for recrystallization is known to be about 0.44 of the melting temperature

for many cubic metals 48, 49. Thus the calculated crystallization temperature of our

system is about 606 oC (the melting point of β-Ga2O3 is 1725 oC 50), which is lower

than our experimental results. The increased crystallization temperature is perhaps due

to the influence of indium content or oxygen vacancy and needs further investigation.

-3 -2 -1 0 1 2 3

Inte

nsity

(ar

b. u

nits

)

()

800oC, FWHM=55.7'

1000 oC,65.4'

900oC, 57.8'

(-402)

Figure 5.19 XRC for the β-(GaIn)2O3 (-402) diffraction of the films with nominal indium

content of 0.3 annealed at different temperatures under oxygen ambient.

The crystallinity of the films is determined from the full-width at half-maximum

(FWHM) of the XRC measurement as shown in figure 5.19. From the figure, it is

readily observed that further increasing the annealing temperature above 800 oC does

not decrease the FWHM value of the peak, indicating that the film annealed at 800 oC

has the best crystallinity.

117

Figure 5.20 AFM morphologies of (GaIn)2O3 films with nominal In content 0.3 annealed under

oxygen ambient at temperatures of 700 oC (a), 800 oC (b), 900 oC (c) and 1000 oC (d), respectively.

The thickness of film before annealing is about 210 nm while the thicknesses of

films after annealing at temperature of 700, 800, 900 and 1000 oC are 201, 199, 179

and 168 nm, respectively. The decrease of film thickness with annealing temperature

can be ascribed to the re-evaporation at high temperature. However, no obvious

transmittance difference was observed for these films annealed from 700 to 1000 oC.

AFM images as given in figure 5.20 show the influence of post-annealing on the

microstructure of (GaIn)2O3 films. Seldom grains have appeared at the surface of

samples annealed at 700 oC as shown in figure 5.20, agreeing with the amorphous

structure of that film observing from the XRD result. The surface of the film annealed

at 800oC is comprised of uniformly distributed and tightly packed grains with triangle

shape as shown in figure 5.20(b). Further increasing the annealing temperature makes

the surface of the film porous, as shown in figure 5.20(c) and (d), indicating decreased

crystallinity which agrees with the results of XRC. Therefore, the optimized annealing

temperature is 800 oC both from the results of XRC and AFM. The degraded

0 200nm

500nm 0 0 100nm

(a)

(c) (d) 0 20nm

(b)

1μm 1μm

1μm 1μm

118

crystallinity for films annealed at 900 and 1000 oC is suggested due to the enhanced

evaporation process at high temperatures. As it is believed that, the evaporation rate is

significant at high temperature due to the fact that (1) nano sized material has a much

lower melting point than its bulk one and (2) the melting temperature of (GaIn)2O3-x is

generally much lower than that of the bulk 51. Similar phenomenon has also been

reported by Takakura et al. 35.

5.5 Annealing effect on films with different indium content

As observed from figure 5.6, phase separation occurs when the nominal In

content is between 0.2 and 0.5 (corresponding x from 0.16 to 0.33). Films with these

In contents were deposited at RT and then annealed under the optimized annealing

conditions (O2, 800 oC) on the purpose of obtaining high quality films without phase

separation. The film with nominal In content 0.7 was also included because the film

of that indium content as-deposited at 500 oC shows polycrystalline in figure 5.6.

Figure 5.21 shows the XRD patterns of 2θ/θ scanning measurements for films

with different nominal In content. The films maintain the (-201) oriented monoclinic

structure of (GaIn)2O3 up to an In composition of 0.3. On the other hand, films with In

content of 0.5 and 0.7 show cubic structured (222) oriented diffraction patterns.

Moreover, increasing the nominal indium content shifts both the (-402) peak position

of monoclinic and (222) peak position of cubic structures of (GaIn)2O3 films

gradually toward a lower angle. The co-existing of double structures has not been

observed clearly. That is, films with no phase separation have been obtained by using

the post annealing process.

119

10 20 30 40 50 60 70

10 20 30 40 50 60 70

0.2

(-60

3)

(-40

2)

(-20

1)

2(

Log

inte

nsity

(ar

b. u

nits

)

0.3

0.5

-A

l 2O3 (

0006

)

Nominal Incontent0.7

(444

) (222)

Figure 5.21 XRD patterns of (GaIn)2O3 films with different nominal indium content annealed at

800 oC under oxygen ambient.

AFM measurements were performed to study the differences on the surface

morphology between (GaIn)2O3 films. The images are shown in figure 5.22 over a

scale of 4 μm×4 μm. The films with indium content of 0.5 and 0.7 are very flat,

whereas those with nominal indium content 0.2 and 0.3 are grainy. The shapes of

these grains are ellipse and triangle with certain orientation, respectively. The

difference in morphology of the (GaIn)2O3 films indicates the different crystal

structure or element content of these films. The roughness of these films measured by

AFM are less than 3 nm, indicating the good surface quality of these films.

120

Figure 5.22 AFM morphologies of (GaIn)2O3 films annealed at 800 oC under oxygen ambient with

nominal indium content of 0.2 (a), 0.3 (b), 0.5 (c) and 0.7 (d), respectively.

The transmittance of annealed (GaIn)2O3 films with different nominal In content

is shown in figure 5.23. The transmission spectra of the annealed (GaIn)2O3 films

indicate: 1) transmittances of all samples in visible and infrared region above 85%, 2)

a sharp absorption edge for all (GaIn)2O3 alloys films caused by the fundamental

absorption of light, and 3) shift of the absorption edge to lower wavelength with

increase in nominal indium content. As has been stated, we demonstrated that the

bandgap of (GaIn)2O3 films can be tailored between 3.8 eV and 5.1 eV by controlling

the indium content. Unfortunately the transmittance of some of films grown at 500 oC

was not high due to phase separation as shown in figure 5.5. Thus transmittance

spectra of as-deposited films grown at 500 oC are also plotted for comparison as dash

line in figure 5.23. It is clear that the transmittance of annealed films is higher than the

as-grown films, especially for nominal indium content of 0.2 and 0.3, indicating that

thermal annealing is an effective way for obtaining high quality (GaIn)2O3 film.

0 20nm 0 20nm

1μm 1μm

0 20nm

(c) (d)

(a) 1μm

0 20nm

1μm (b)

121

200 400 600 8000

20

40

60

80

100

Inte

nsi

ty (

arb

. units

)

Wavelength (nm)

Annealled 0.2 0.3 0.5 0.7

As-grown 0.2 0.3 0.5 0.7

Figure 5.23 Transmittance spectra of (GaIn)2O3 films with different nominal indium content

annealed at 800 oC under oxygen ambient. Transmittance spectra of as-deposited films grown at

500 oC are also plotted for comparison as dash line.

5.6 Conclusions

In summary, we showed the growth of crystalline and bandgap tunable


separation was discussed in detail.

(1) Optical analysis indicates that the bandgap of the (GaIn)2O3 films grown by

PLD can be tailored from 3.8 eV to 5.1 eV by controlling the indium content (x),

indicating PLD is a promising growth technology for growing bandgap tunable

crystalline (GaIn)2O3 films. Single phase (GaIn)2O3 films were obtained although

films with nominal indium content between 0.2 and 0.5 exhibit phases separation

(2) (GaIn)2O3 films with nominal indium content of 0.3 were deposited on

sapphire substrate by PLD at substrate temperatures from RT to 500 oC. The phase

separation were observed for the films grown at substrate temperature higher than 300

oC while the films grown at substrate temperature lower than 200 oC revealed

homogenous element distributions with amorphous structures. Thermal annealing had

122

no obvious effects on (GaIn)2O3 films grown at substrate temperature higher than 300

oC. The clusters remained on the surface of the films after thermal annealing treatment.

On the other hand, however, by thermal annealing the film deposited at RT in

atmosphere, (GaIn)2O3 film with smooth surface, homogenous element distribution,

high orientation crystal and high optical transmittance was successfully obtained.

(3) In order to understand the annealing effects, (GaIn)2O3 films with nominal

indium content of 0.3 as-deposited in room temperature have been annealed under

different gas ambient (N2, vacuum, Ar, O2) or at different temperatures (700~1000 oC).

It is found that gas ambient and temperature have important influence on crystal

quality of annealed (GaIn)2O3 films. Only oxygen ambient can crystallize (GaIn)2O3

film and film annealed in 800 oC appears best crystal quality. X-ray photoelectron


decreasing the oxygen vacancy.

(4) (GaIn)2O3 films with different nominal indium content varying from 0.2 to

0.7 annealed at 800 oC under O2 ambient also showed high crystal quality, improved

optical transmittance, and smooth surface. Table 4.3 summarizes the crystal quality of

the as deposited films at 500 oC as well as the annealed films. It is obviously that high

oriented films without phase separation can be obtained in the nominal indium content

regions of 0 to 0.1 and 0.9 to 1.0 by deposited the film at 500 oC. Complementally,

high oriented films with nominal indium content from 0.2 to 0.7 without phase

separation can be obtained through annealing process. By combing the two processes,

bandgap tunable high quality (GaIn)2O3 films throughout the whole indium content

range from 0 to 1 can be successfully obtained.

123

Table 4.3 Crystal quality of films deposited by different methods

Indium content in target

Deposited at 500 oC

Suitable method or not

Deposited at RT then annealed at 800 oC

Suitable method or not

0 Highly oriented √ Amorphous ×

0.1 Highly oriented √ Amorphous ×

0.2 Phase separation × Highly oriented √



0.7 No orientation × Highly oriented √

0.9 Highly oriented √ No orientation ×

1.0 Highly oriented √ No orientation ×

124

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126

Chapter 6

Growth and characterization of (AlGa)2O3 films

6.1 Introduction

In the chapters before, we have reported the growth of high oriented β-Ga2O3

film on sapphire substrate. We have also obtained tunable bandgap from 3.8 eV to 5.1

eV by alloying Ga2O3 with indium (decreasing the bandgap). A wider bandgap range

is of great merit as it allows the design of devices such as high sensitive

wavelength-tunable photodetectors, cutoff wavelength-tunable optical filters in more

broad range. Al is a candidate to enlarge the bandgap of Ga2O3 because Al2O3 has a

bigger bandgap (~8.8 eV for bulk material, ~6.4 eV for amorphous Al2O3 films) and

the similar electron structures of Al and Ga makes the alloy (AlGa)2O3 possible. Until

now, most of studies on (AlGa)2O3 system were focused on γ-Ga2O3–Al2O3 solid

solution and its catalysts activities 1-3. Recently, Oshima et al. 4 have grown

β-(AlGa)2O3 thin films on (100) β-Ga2O3 single crystal substrates by plasma-assisted

molecular beam epitaxy. The films were grown almost coherently and maintained the

β-phase up to an Al content of 0.61. Ito et al. 5 tried to fabricate corundum-structured

α-(AlGa)2O3 thin films on sapphire by spray-assisted mist chemical vapor deposition.

However, they have not given details on the properties of these films. The

investigation of (AlGa)2O3 films is still in its infancy.

In this chapter, we present on the growth of (AlGa)2O3 films on sapphire

substrates by PLD in the whole Al content range. The Al contents in the films increase

lineally with that of the targets. The O 1s energy loss spectra obtained from X-ray

photoelectron spectroscopy are proved to be valid to determine the bandgap of

(AlGa)2O3 films. The bandgap of (AlGa)2O3 increases continuously with the Al

content according to the O 1s energy loss spectra.

127

6.2 The Al content in the film

We deposited (AlGa)2O3 films by PLD using a KrF excimer laser source with

wavelength 248 nm on (0001) sapphire substrates. The oxygen pressure was

controlled at 0.1 Pa while substrate temperature was set at 400 °C during the process

of deposition. Bulks (diameter of 20 mm) with different Al content (0.22, 0.53, 0.72,

0.86, 0.96 and 1.0) were used as targets. The thicknesses of the films were about

200~300 nm.

1200 1000 800 600 400 200 0

= x

1.0 0.96 0.86 0.72 0.53 0.22

Al 2

s

Ga

3d

Ga

3p

Ga

LMM

O 1

s

O K

LL

Ga

2p1

/2 Ga

2p3/

2

Inte

nsity

(a

rb. u

nits

)

B inding energy (eV)

Al 2

p

C 1

s(a)

t

1120 1116 1112 1108

In

tens

ity (

arb.

uni

ts) (b) Ga 2p

80 76 72 68 64

Al 2p(c)

Binding energy (eV)

Figure 6.1 XPS wide scan spectra (a), Ga 2p (b) and Al 2p (c) core level spectra of (AlGa)2O3

films with different Al content in targets.

128

Wide scan spectra in the binding energy range 0~1200 eV were obtained to

identify the elements in (AlGa)2O3 films as shown in figure 6.1(a). The determined

XPS peaks of Ga, O and Al are indicated in the figure. No other elements were

detected from wide scan spectra of the films, indicating the obtaining of high purity

(AlGa)2O3 films. It is obvious that the intensity ratio of Al/Ga increases with the

increase of aluminum content in the targets. Figure 6.1(b) and (c) shows the high

resolution XPS spectra of Ga 2p and Al 2p core levels for (AlGa)2O3 films. The Ga 2p

core levels of (AlGa)2O3 films with different Al content shown in figure 6.1(b) are

located at 1117.5 eV which is attributed to Ga-O bonding 6. The intensity of the peak

decreases with the increase of Al content and no obvious peak shift is observed. Al 2p

peaks of (AlGa)2O3 films are shown in figure 6.1(c). From figure 6.1(c), it is seen that

the intensity of the Al 2p peaks becomes more intense with increasing Al content. The

Al 2p peak in pure Al2O3 film is observed at 74.5 eV, which is considered as the Al-O

bonding and is in good agreement with the reported value of pure Al2O3 6. The peak

shifts almost uniformly towards lower binding energy with the decrease of Al content

and reaches 73.9 eV for (AlGa)2O3 films with a Al content of 0.22. This shift can be

contributed to the formation of Al-O-Ga bonds. Similar phenomenon has also been

observed by Andrulevičius et al. 7, who found that the Ti 2p peak positive shifts for

the formation of Ti-O-Si and Ti-O-Zr bonds. The shift is believed due to the fact that

Ga is a more ionic cation than Al in (AlGa)2O3 films, and thus the charge transfer

contribution changes with the increase of Al concentration 8.

Atomic concentrations were estimated from the XPS element peak area after

applying an atomic sensitivity factor. The aluminum content x in (AlGa)2O3 films

obtained from the XPS spectra is shown in figure 6.2 as a function of aluminum

content in targets. The elements contents in the films increase lineally with that in the

targets throughout the whole aluminum range. The results indicate desired

compositional film can be controlled by changing the aluminum content in target,

opening the possibility of designing (AlGa)2O3 based devices as the microstructures

and properties of semiconductor materials strongly depend on the composition. The Al

content in films deposited from targets with Al content of 0.22, 0.53, 0.72, 0.86, 0.96

129

and 1.0 are 0.24, 0.54, 0.76, 0.89, 0.98 and 1.0, respectively. From now on, we will

use the Al content in the film.

0.2 0.4 0.6 0.8 1.0

0.2

0.4

0.6

0.8

1.0

Al c

ont

ent

in fi

lms

x

Al content in targets x t Figure 6.2 Al content x in the (AlGa)2O3 films as a function of Al content in targets.

6.3 Structure of the (AlGa)2O3 films

Figure 6.3 shows the XRD 2θ/θ scanning for the films grown by changing the Al

content. For Al content of 0.24, three different diffraction peaks located at 18.8°, 38.0°

and 58.8° are observed. These peaks are ascribed to the patterns of monoclinic

β-(AlGa)2O3 and can be assigned as the (-201), (-402) and (-603) faces, respectively.

With increasing the Al content, the intensity of the diffraction peaks decreases. For Al

content higher than 0.89, almost no peak is observed. The degradation of crystallinity

with increasing Al content was also been observed by Ito et al., and was contribute to

the low substrate temperature for the higher inclusion level of Al, in spite of the

reduction in lattice mismatch to sapphire substrate 5. The (-603) peaks of the XRD

patterns display an apparent shift to high angle with increasing the Al content, which

is mainly due to the smaller atomic radius of Al compared with that of Ga. This also

verifies that the Al atoms enter into the crystal lattices of Ga2O3 to form ternary solid

solutions.

130

10 20 30 40 50 60 70

2(

(-6

03

)

(-4

02)

(-2

01

)

Inte

nsity

(ar

b. u

nits

)

0.24

0.54

0.76

0.89

0.98

x=1.0

Figure 6.3 XRD patterns of (AlGa)2O3 films with different Al content.

6.4 Transmittance and bandgap of the (AlGa)2O3 films

The transmittance spectra of (AlGa)2O3 films with different Al content are shown

in figure 6.4. The transmittances of all samples in visible and infrared region are

above 75%. A sharp absorption edge which is caused by the fundamental absorption

of light for (AlGa)2O3 film with Al content lower than 0.76 is observed. Those

absorption edges shift to lower wavelength with the increase of Al content. It is

noticeable that transmittance spectra of (AlGa)2O3 films with Al content higher than

0.76 is incomplete for the limitation of the measurement. By extrapolating the linear

part of (αhν)2~ hν to the horizontal axis, one can obtain the bandgap of (AlGa)2O3

films as shown in the insertion of figure 6.4. Here hν is the energy of the incident

photon, α the absorption coefficient which is evaluated using the standard relation

taking the film thickness into account 9. However, only bandgap of films with Al

131

contents lower than 0.54 can be obtained because the incomplete transmittance

spectra of other films. The bandgap values of films with Al content 0.24 and 0.54 are

5.34 and 5.74 eV, respectively, which increase with Al content.

200 400 600 8000

20

40

60

80

100

=

4.0 4.4 4.8 5.2 5.6 6.0 6.40

20

40

60

hv (eV)

5.34 eV5.74 eV(

hv)2

(10

10eV

2 cm-2)

Wavelength (nm)

Tra

nsm

ittan

ce (

%)

1.0 0.98 0.89 0.76 0.54 0.24

x

Figure 6.4 Transmittance of (AlGa)2O3 films with different Al content in the film. Insert is (αhν)2

vs. hν plot of (AlGa)2O3 films with different Al content.

XPS is generally used to analyze the binding energies of the atomic core-level

electrons in a material. It can also be used to analyze the inelastic collisions that occur

during photo excitation and photoemission of electrons from the sample. The inelastic

collisions mainly include exciting the plasmon in the bulk and at surface (a

fast-moving charged particle can lose energy in the bulk material to collective

high-frequency plasma oscillations of electrons in the valence band) and

single-particle excitations due to band-to-band transitions 10. The fundamental lower

limit of inelastic loss is equal to the bandgap energy, thus the onset of the inelastic

energy loss spectra corresponds directly to the bandgap energy 11. The determination

of the bandgap value of large bandgap materials using the energy-loss peak of the O

1s spectrum has been reported by many scientists 10, 12-14. A representative

high-resolution scan of the O 1s core level of (Al0.76Ga0.24)2O3 film obtained in this

work is shown in figure 6.5. The largest signal corresponds to the primary

132

photoelectron peak due to the O 1s core level electrons located at binding energy

530.6 eV. The bulk plasmon loss peak is observed at approximately 555.6 eV,

corresponding to bulk plasmon energy about 25 eV. Additionally, a small overlapping

peaks measured at 522 eV can be attributed to X-ray satellite peaks due to Kα

transitions from the non-monochromatic Mg X-ray source 10, 15.

550 540 530 520

Inte

nsity

(a

rb.

un

it)

Binding energy (eV)

(Al0.76

Ga0.24

)2O

3

X-ray satellite peak

Bulk plasmon loss peak

Eg

O 1s peak

Figure 6.5 High-resolution XPS scan of the O 1s core level for (Al0.76Ga0.24)2O3 film.

To find the bandgap energy, a linear fit is made to the measured loss spectra

curve near the approximate location of onset of inelastic losses. The bandgap energy

is equal to the difference between the core-level peak energy and the onset of inelastic

losses. This procedure is shown in figure 6.6 for (AlGa)2O3 films with different Al

contents. By this method, the energy gap of Al2O3 is measured as 7.1 eV as shown in

figure 6.6(f), which almost agrees with the reported 6.9 eV for amorphous Al2O3 14.

The energy gaps for (AlGa)2O3 films with Al content of 0.24, 0.54, 0.76, 0.89 and

0.98 are 5.2, 5.6, 6.1, 6.3, and 6.8 eV, respectively, as shown from figure 6.6(a) to

figure 6.6(e).

133

550 545 540 535 530 525

Inte

nsity

(ar

b. u

nit)

Binding energy (eV)

(a)

O 1s 530.4 eV

Eg=5.2 eV

535.6 eV

550 545 540 535 530

Eg=5.6 eV

O 1s 530.7eV

(b)

Inte

nsity

(ar

b. u

nit)

Binding energy (eV)

536.3 eV

550 545 540 535 530

O 1s 530.6 eVIn

ten

sity

(ar

b. u

nit)

Binding energy (eV)

(c)

536.7 eV

Eg=6.1 eV

550 545 540 535 530

O 1s

530.8eVInte

nsity

(ar

b. u

nit)

Binding energy (eV)

(d)

537.1 eV

Eg=6.3 eV

550 545 540 535 530

Eg=6.8 eV

O 1s

531 eVInte

nsi

ty (

arb

. un

it)

Binding energy (eV)

(e)

537.8 eV

550 545 540 535 530

Eg=7.1 eV

Inte

nsi

ty (

arb

. un

it)

Binding energy (eV)

(f)

O 1s

531.1 eV538.2 eV

Figure 6.6 The O 1s peak and inelastic scattering loss for (AlGa)2O3 films with Al content x of

0.24 (a), 0.54 (b), 0.76 (c), 0.89 (d), 0.98 (e) and 1.0 (f), respectively.

134

0.2 0.4 0.6 0.8 1.05.0

5.5

6.0

6.5

7.0

7.5

Al content in films x

Ban

dgap

(eV

)

By XPSBy transmittance

Figure 6.7 Dependence of bandgap of (AlGa)2O3 films on Al content x.

The derived bandgap value from figure 6.6 is shown in figure 6.7 as a function of

Al content x. The obtained bandgap from transmittance is also plotted for comparison.

The bandgap values by the two methods agree well with each other. It is well known

that the transmittance derived bandgap value is obtained from a Tauc plot, pioneered

by Jan Tauc, who proved that momentum is not conserved even in a direct optical

transition 16. However, the wavelengths range of the spectrometer for this method is

commonly higher than 200 nm, which restrict the measurements for material with

bandgap higher than 6.2 eV. On the other aspect, XPS core level energy loss spectra

are often applied to obtain the bandgap of dielectrics. The discrepancies in the

bandgap energies caused by the two different measurement techniques have not been

found in present work. This result indicates that the XPS is a valid way to determine

the bandgap of (AlGa)2O3 which beyond the ability of the common spectrophotometer.

It is also clear from the figure 6.7 that the bandgap of (AlGa)2O3 increases

continuously with the Al content covering the whole Al content range. The results

indicate that the bandgap of (AlGa)2O3 can be controlled by changing the Al content

x, paving a way to design optoelectronic and photonic devices based on this material.

135

6.5 Conclusions

Bandgap tunable (AlGa)2O3 films were deposited on sapphire substrates by

pulsed laser deposition. The deposited films are of high transmittance as measured by









136

References: 1. T. Watanabe, T. Masuda, Y. Miki, Y. Miyahara, H.-J. Jeon, S. Hosokawa, H. Kanai, H.

Deguchi and M. Inoue, J. Am. Ceram. Soc., 2010, 93, 3908.

2. T. Watanabe, Y. Miki, T. Masuda, H. Deguchi, H. Kanai, S. Hosokawa, K. Wada and M. Inoue,

Ceram. Int., 2011, 37, 3183.

3. M. H. Zahir, K. Sato, H. Mori, Y. Iwamoto, M. Nomura and S.-i. Nakao, J. Am. Ceram. Soc.,

2006, 89, 2874.

4. T. Oshima, T. Okuno, N. Arai, Y. Kobayashi and S. Fujita, Jpn. J. Appl. Phys., 2009, 48,

070202.

5. H. Ito, K. Kaneko and S. Fujita, Jpn. J. Appl. Phys., 2012, 51, 100207.

6. T. Mathew, Y. Yamada, A. Ueda, H. Shioyama and T. Kobayashi, Applied Catalysis A: General,

2005, 286, 11.

7. M. Andrulevičius, S. Tamulevičius, Y. Gnatyuk, N. Vityuk, N. Smirnova and A. Eremenko,

Materials Science-Medziagotyra, 2008, 14, 8.

8. H. Y. Yu, M. F. Li, B. J. Cho, C. C. Yeo, M. S. Joo, D. L. Kwong, J. S. Pan, C. H. Ang, J. Z.

Zheng and S. Ramanathan, Appl. Phys. Lett., 2002, 81, 376.


10. M. T. Nichols, W. Li, D. Pei, G. A. Antonelli, Q. Lin, S. Banna, Y. Nishi and J. L. Shohet, J.

Appl. Phys., 2014, 115, 094105.

11. F. G. Bell and L. Ley, Phys. Rev. B, 1988, 37, 8383.

12. Y. Hori, C. Mizue and T. Hashizume, Jpn. J. Appl. Phys., 2010, 49, 080201.

13. T. Kamimura, K. Sasaki, M. Hoi Wong, D. Krishnamurthy, A. Kuramata, T. Masui, S.

Yamakoshi and M. Higashiwaki, Appl. Phys. Lett., 2014, 104.

14. C. M. Tanner, Y.-C. Perng, C. Frewin, S. E. Saddow and J. P. Chang, Appl. Phys. Lett., 2007,

91, 203510.

15. C. Klauber, Surf. Interface. Anal., 1993, 20, 703.

16. J. Tauc, Mater. Res. Bull., 1968, 3, 37.

137

Chapter 7

Summary

Wide bandgap semiconductor materials have become the hot spot of recent

research for the possible using in many fields such as light emitting devices, power

devices and flame detectors. Among all the wide bandgap materials, β-Ga2O3 film

with the monoclinic structure is considered as a promising candidate for its large

bandgap and chemical and physical stabilities. And it is also suitable for extreme


environments. However, the bandgap should be tuned to realize high sensitive

wavelength tunable photodetectors, cutoff wavelength-tunable optical filters or to

introduce shallow impurity levels for good electronic properties.

In Chapter 1, the background including the properties and the review of studies

on Ga2O3, (Ga1-xInx)2O3, (AlxGa1-x)2O3 and Si doped Ga2O3 were described. The

purpose or the motivation of this study was presented.

In Chapter 2, the film growth and characterization methods were introduced.

In Chapter 3, we have investigated the influences of oxygen pressure, substrate

temperature and deposition time on the structure and optical properties of Ga2O3 films

grown by PLD. The influence of post annealing was also been discussed. (1) The

crystal quality and the thickness of films deposited at 600 oC increase with the

increasing of oxygen pressure. The growth mode of the films is island mode. (2) By

varying the substrate temperature, the evolutions of the structure, surface morphology

and bandgap have been clearly clarified. Films deposited at substrate temperature

below 400 oC show amorphous structure while those deposited at substrate

temperature higher than 500 oC are of high oriented monoclinic structure. (3) The

optimized growth substrate temperature and oxygen pressure for our experiment is

138

500 oC and 0.1 Pa. The growth relationship between the film and the substrate is:

sapphire (0001) // β-Ga2O3 (-201) and sapphire [11-20] // β-Ga2O3 [102]. The obtained

β-Ga2O3 film is of sixfold in-plane rotational symmetry. The hard X-ray

photoemission spectroscopy reveals that the valence band of the crystalline films is

mainly due to the hybridization of Ga 4sp. (4) By varying the growth time (film

thickness), the growth process has been investigated. (5) Post annealing (annealing

temperature from 700 to 900 oC) cannot be used to obtain films with better crystal

quality than the film deposited under the optimized growth conditions. The films with

post annealing show smaller blue/UV emission ratio.

In Chapter 4, we have investigated the Si doping influence on the structure and

properties of Ga2O3 films. (1) Ga2O3 films with different Si content were grown on

sapphire substrate at 500 oC by PLD. All of the films exhibit smooth surfaces and high

transmittances. The films of Si content lower than 4.1 at. % show high (-201) oriented

monoclinic structure. The carrier density of Ga2O3 film has been increased to 9.1×1019

cm-3 with conductivity of 2.0 S cm-1 by 1.1 at. % Si doping. Further increase of Si

content leads to the decrease of carrier density. (2) By varing the substrate

temperature, it is found that film deposited at 500 oC (1 wt.% Si doped) shows lowest

conductivity and highest carrier density while possesses best crystallinity. (3) Oxygen

pressure has no obviously influence on the electrical properties of Si-doped Ga2O3,

indicating the oxygen deficiency is not the main origin of the conductive carrier in our

study.

In Chapter 5, we showed the growth of crystalline and bandgap tunable


separation was discussed in detail. (1) Optical analysis indicates that the bandgap of

the (GaIn)2O3 films grown by PLD can be tailored from 3.8 eV to 5.1 eV by

controlling the indium content. Single phase (GaIn)2O3 films were obtained although

films with nominal In content between 0.2 and 0.5 exhibit phases separation. (2)

(GaIn)2O3 films with nominal In content of 0.3 were deposited on sapphire substrate

by PLD at substrate temperatures from RT to 500 oC. The phase separation were

observed for the films grown at substrate temperature higher than 300 oC while the

139

films grown at substrate temperature lower than 200 oC revealed homogenous element

distributions with amorphous structures. Thermal annealing had no obvious effects on

(GaIn)2O3 films grown at substrate temperature higher than 300 oC. The clusters

remained on the surface of the films after thermal annealing treatment. On the other

hand, however, by thermal annealing the film deposited at RT in atmosphere,

(GaIn)2O3 film with smooth surface, homogenous element distribution, high

orientation crystal and high optical transmittance was successfully obtained. (3) In

order to understand the annealing effects, (GaIn)2O3 films with nominal In content of

0.3 as-deposited in room temperature have been annealed under different gas

ambients (N2, vacuum, Ar, O2) or at different temperatures (700~1000 oC). It is found

that gas ambient and temperature have important influence on crystal quality of

annealed (GaIn)2O3 films. Only oxygen ambient can crystallize (GaIn)2O3 film and

film annealed in 800 oC appears best crystal quality. X-ray photoelectron


decreasing the oxygen vacancy. (4) (GaIn)2O3 films with different nominal In contents

from 0.2 to 0.7 annealed at 800 oC under O2 ambient also showed high crystal quality,

improved optical transmittance, and smooth surface. Thus, high oriented films with

nominal In content from 0.2 to 0.7 without phase separation can be obtained through

annealing process. Complementally, high oriented films without phase separation can

be obtained in the nominal indium content regions of 0 to 0.1 and 0.9 to 1.0 for the

film deposited at 500 oC. By combing the two processes, bandgap tunable high quality

(GaIn)2O3 films throughout the whole indium content range from 0 to 1 can be

successfully obtained.

In Chapter 6, bandgap tunable (AlGa)2O3 films were deposited on sapphire

substrates by PLD. The deposited films are of high transmittance as measured by






140




141

Acknowledgments

My deepest gratitude goes first and foremost to Professor Qixin Guo, my

supervisor, for his constant encouragement and guidance. He has walked me through

all the stages of the researches and preparation manuscripts for publications. Without

his consistent and illuminating instruction, this thesis could not have reached its

present form.

I also would like to express deepest gratitude to Professor Tooru Tanaka,

Professor Mitsuhiro Nishio, and Doctor Katsuhiko Saito for their help on the daily

experiments, instructions and discussions. I would also like to express my thanks to

my colleagues: Doctor Zhenwei Chen and Xu Wang, for their invaluable helps.

I also want to show my deepest gratitude to Doctor Makoto Arita in Kyushu

University, Doctor Yitao Cui in Spring 8 for the help on experiments and analysis.

I also want to show my deepest gratitude to Professor Tooru Tanaka, Professor

Mitsuhiro Nishio and Professor Kazutoshi Takahashi for their help on the review and

revise of this dissertation.

Finally, I would like to express my deep appreciation to my farther, Zhonghai

Zhang, mother, Yuyin Yuan, wife, Juan Zhou and daughter Yishan Zhang. Their

supports are my driving forces to go bravely forward.

142

List of publications:

A Original Paper Related To This Dissertation

1. Fabi Zhang, Katsuhiko Saito, Tooru Tanaka, Mitsuhiro Nishio, Makoto Arita, and

Qixin Guo, Wide bandgap engineering of (AlGa)2O3 films , Applied Physics Letters,

105（2014）162107

2. F. Zhang, H. Jan, K. Saito, T. Tanaka, M. Nishio, T. Nagaoka, M. Arita, Q. Guo,

Toward the understanding of annealing effects on (GaIn)2O3 films, Thin Solid Films,

578 (2015) 1-6.

3. Fabi Zhang, Katsuhiko Saito, Tooru Tanaka, Mitsuhiro Nishio, and Qixin Guo,

Thermal annealing impact on crystal quality of (GaIn)2O3 alloys, Journal of Alloys

and Compounds, 614, (2014) 173-176.

4. F. Zhang, K. Saito, T. Tanaka, M. Nishio, Q. Guo, Wide bandgap engineering of

(GaIn)2O3 films, Solid State Commun. 186 (2014) 28-31

5. F.B. Zhang, K. Saito, T. Tanaka, M. Nishio, Q.X. Guo, Structural and optical

properties of Ga2O3 films on sapphire substrates by pulsed laser deposition, J. Cryst.

Growth, 387 (2014) 96-100.

6. F. Zhang, K. Saito, T. Tanaka, M. Nishio and Q. Guo, Electrical properties of Si

doped Ga2O3 films grown by pulsed laser deposition, Journal of Materials Science:

Materials in Electronics, (2015), DOI: 10.1007/s10854-015-3627-6, 1-6.

143

B Original Papers of Other Subjects

1. Guo-Ling Li, Fabi Zhang, Yi-Tao Cui, Hiroshi Oji, Jin-Young Son, and Qixin

Guo, Electronic structure of β-Ga2O3 single crystals investigated by hard X-ray

photoelectron spectroscopy, Applied Physics Letters, 107, (2015) 022109.

2. X.H. Wang, F.B. Zhang, K. Saito, T. Tanaka, M. Nishio and Q.X. Guo, Electrical

properties and emission mechanisms of Zn-doped β-Ga2O3 films, Journal of Physics

and Chemistry of Solids, 75 (2014) 1201-1204.

3. X.H. Wang, L.Q. Huang , L.J. Niu , R.B. Li , D.H. Fan , F.B. Zhang , Z.W. Chen ,

X. Wan, Q.X. Guo, The impacts of growth temperature on morphologies,

compositions and optical properties of Mg-doped ZnO nanomaterials by chemical

vapor deposition, Journal of Alloys and Compounds 622 (2015) 440–445.

4. Wenwen Zhao, Akinobu Tanaka, Kyoko Momosaki, Shinji Yamamoto, Fabi Zhang,

Qixin Guo, Hideyuki Noguchi, Enhanced electrochemical performance of Ti

substituted P2-Na2/3Ni1/4Mn3/4O2 cathode material for sodium ion batteries,

Electrochim. Acta, 170 (2015) 171-181.

5. Wei Jiang, Qing-jun Chen, Jun Shen, Fa-bi Zhang, Xian-liang Zhou and Xiao-zhen

Hua, Salt Spray Corrosion Performance Associated with the Glass Forming Ability of

the FeCo-based Bulk Metallic Glasses, Applied Mechanics and Materials Vol. 618

(2014) 109-113.

144

Abbreviations:

α Absorption coefficient

at.% Atomic percent

AFM Atomic Force Microscopy

AACVD Aerosol-Assisted Chemical Vapor Deposition

CL Cathodoluminescence

dstop the plume stopping distance

EFG Edge-defined film-fed growth

EDS Energy-dispersive X-ray spectroscopy

Eg Bandgap

FWHM Full width at half maximum

HAXPES Hard X-ray photoelectron spectroscopy

HVPE Halide vapor phase epitaxy

PLD Pulsed laser deposition

wt.% Weight percent

MBE Molecular beam epitaxy

MOSFET Metal-oxide-semiconductor field-effect transistor

MOCVD Metal-organic chemical vapor deposition

MOVPE Metal organic vapor phase epitaxy nanowires

MM-SPM multi-mode scanning probe microscope

NWs Nanowires

PICS Photoionization cross section

RHEED Reflection High Energy Electron Diffraction

RT Room temperature

RF Radio Frequency

RP Rotary Pump

SEM Scanning Electron Microscopy

TMP Turbo molecular pump

145

3D Three-dimensional

UV Ultraviolet

VS Vapor–solid

VBM Valence band maximum

VB Valence band

XPS X-ray photoelectron spectroscopy

XRD X-ray diffraction

XRC X-ray rocking curve

xtar element content in the target

growth and characterization of ga2o3 -...

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