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이학박사학위논문

Dislocations and high electron mobility in

epitaxial Ba1-xLaxSnO3 and SnO2-x

결정 성장된 Ba1-xLaxSnO3 와 SnO2-x의 전위결함과

높은 전하 이동도 연구

2017 년 2 월

서울대학교 대학원

물리천문학부

문 효 식

Dislocations and high electron mobility in

epitaxial Ba1-xLaxSnO3 and SnO2-x

by

Hyosik Mun

Supervised by

Professor Kookrin Char

A Dissertation in physics

A dissertation submitted to the Faculty of

Seoul National University

In Partial Fulfillment of the Requirements

For the Degree of

Doctor of Philosophy

February 2017

Department of Physics and Astronomy

Graduate School

Seoul National University

i

SEOUL NATIONAL UNIVERSITY

Abstract

Department of Physics and Astronomy

Doctor of Philosophy

Dislocations and high electron mobility in

epitaxial Ba1-xLaxSnO3 and SnO2-x

By Hyosik Mun

Perovskite-type oxides have an advantage of a simple and flexible structure that

enables ionic substitution on both A and B sites which can form a vast set of

important materials for a wide variety of academic and industrial research. Our

research group have recently published about high mobility in Ba0.96La0.04SnO3

(BLSO) single crystal: as high as 320 cm2/Vs at the carrier concentration 8.0 ×

1019 cm−3 in single crystals. On the other hand, epitaxial BLSO films were

found to exhibit lower mobility, 27 cm2/Vs at the carrier concentration 6.0 ×

1019 cm−3. The large difference in mobility between single crystals and epitaxial

films was attributed to the existence of the dislocations in the epitaxial films from

the mobility’s dependence on the carrier density.

In this dissertation, we have investigated the influence of dislocation for

electrical properties of BLSO films on SrTiO3 (STO) substrates in detail. In our

research, we studied the relations between the mobility and the dislocation density

of the recently discovered high mobility Ba0.96La0.04SnO3 thin films. The effect

of dislocations on the electrical transport properties have hardly been studied for

ii

perovskite films. We have found that the carrier density and the mobility, as high

as 4.0 × 1020 cm−3 and 70 cm2/Vs, decrease as the dislocation density increases.

We provide the values for the density of dislocations by TEM micrographs as well

as AFM images after surface etching. Furthermore, we found the effect of

dislocations on the mobility to be large, when compared with GaN with similar

density of dislocations. The importance of dislocation scattering in the perovskite

structure is emphasized for the first time, especially in the low carrier density

regime.

We, also, investigated the electronic transport properties of epitaxial SnO2-x thin

films on r-plane sapphire substrates. The films were grown by pulsed laser

deposition technique and its epitaxial growth direction was [101] and the in-plane

alignment was of SnO2-x [010] // Al2O3 [1210]. The mobility of SnO2 films has

delicate sensitivity of thickness due to the high density of the dislocations between

the tin dioxide film and the sapphire substrate. Therefore, we studied the effect

of buffer layer on the mobility of SnO2, intensively.

We have searched the various buffer layers to reduce the interfacial defects and

it is found that fully oxidized SnO2 at grown 800 °C and Ruthenium-doped SnO2

(Ru-doped SnO2) buffer layer demonstrates good insulating and crystalline

properties. From the X-ray diffraction analysis, phase pure SnO2 film is grown

on Ru-doped SnO2 on r-plane sapphire substrate. Through the omega scan, SnO2

and Ru-doped SnO2 shows good crystallinity. As a result of using 1um Ru-doped

SnO2 buffer layer, we can get a mobility of almost about 80 cm2/Vs at carrier

density of 3.7×1018 cm-3. On the other hand, the SnO2-x films were grown with

fully oxidized SnO2 buffer layer, we have found the mobility of the 30 nm thick

SnO2-x thin films strongly dependent on the thicknesses of the fully oxidized

insulating SnO2 buffer layer. When the buffer layer thickness increased from 100

nm to 700 nm, the mobility values increased from 23 to 106 cm2/Vs and the carrier

density increased from 9×1017 cm-3 to 3×1018 cm-3, which we attribute to reduction

of large density of dislocations as the buffer layer thickness increases. In addition,

we studied the doping dependence of the mobility of SnO2-x thin films grown on

iii

top of 500 nm thick insulating SnO2 buffer layers. The oxygen vacancy doping

level was controlled by the oxygen pressure during deposition. As the oxygen

pressure increased to 45 mTorr, the carrier density were found to decrease to

1.8×1017 cm-3 and the mobility values to 22 cm2/Vs, which is consistent with the

dislocation limited transport properties. We also checked the conductance change

of the SnO2-x during thermal annealing cycles, demonstrating unusual stability of

its oxygen. The correlation between the electronic transport properties and

microstructural defects investigated by the transmission electron microscopy were

drawn. The excellent oxygen stability and high mobility of low carrier density

SnO2-x films demonstrates its potential as a transparent oxide semiconductor.

In applicative works, the author demonstrated a field effect transistor made with

an un-doped layer of SnO2 on r-plane Al2O3 substrate, with the gate dielectrics as

HfO2. The field effect mobility, the Ion/Ioff ratio, and the subthreshold swing of

the device are 72.1 cm2/Vs, 6.0×106, and 0.48 Vdec-1, respectively.

Keywords: Transparent conducting oxide, Transparent oxide semiconductor,

Perovskite oxide, BaSnO3, Wide band gap, Oxygen stability, High electrical

mobility, SnO2, Field effect transistor

Student number: 2012 -30102

iv

v

Contents

Abstract………………………………………..………………i

Contents………………………………………….……….......v

List of Figures…………………………………..……………ix

List of Tables ……………………………..………...……...xiv

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

1.1 Perovskite oxide semiconductor……..…………………….....1

1.2 High electron mobility of BaSnO3…..……………………..…2

1.3 Thermal stability ...………..………..……………………..…4

1.4 Transparent conducting oxide (TCO) and .……………………

transparent oxide semiconductor (TOS)..………..…5

2 Epitaxial growth of BaSnO3 thin films ..…..….………..7

2.1 High quality BaSnO3 sample growth using ..............................

pulsed laser deposition.............7

vi

2.2 Structure properties of La-doped BaSnO3 thin films ............12

2.2.1 X-ray diffraction analysis (XRD) ......………..……..12

2.2.2 Atomic force microscope (AFM) ......………..……..17

2.2.3 Transmission electron microscopy (TEM) .....……..18

2.2.4 Etch-pit development (EPD) .....……………..….…..20

3 Electrical transport properties of La-doped BaSnO3 .....21

3.1 Scattering mechanisms and carrier mobility in BaSnO3 ...….21

3.1.1 Electrical mobility and effective mass ……....……..22

3.1.2 Total scattering rate ……………………….....……..22

3.1.3 Acoustic phonon scattering ……………….....……..25

3.1.4 Optical phonon scattering ……...………….....……..27

3.1.5 Ionized impurity scattering ……...…………...……..29

3.1.6 Neutral impurity scattering ……...…………...……..33

3.1.7 Threading dislocation scattering ……..……………..34

3.2 Electrical transport of La-doped BaSnO3 ……………...………

single crystal and thin film. ……………36

3.2.1 BLSO single crystal ……................................……..36

3.2.2 BLSO thin film ……...…................................……..40

vii

4 High mobility SnO2 thin film with various buffer layers 50

4.1 Motivation of SnO2 research for transparent ………………….

oxide semiconductor …..…..50

4.2 Structural properties of SnO2 on r-plane Al2O3 substrate ….52

4.2.1 Epitaxial SnO2 film growth condition …….....……..53

4.2.2 X-ray diffraction analysis (XRD) …………………..54

4.2.3 Thickness dependence of electrical transport properties

…………..55

4.2.4 Transmittance electron microscopy (TEM) ….……..56

4.3 Ruthenium-doped SnO2 buffer layer …………………….….59

4.4 Rutile TiO2 buffer layer ……...…………………………..….62

4.5 Fully oxidized SnO2 buffer layer at 700 °C ….…………..….64

4.5 Fully oxidized SnO2 buffer layer at 800 °C ….…………..….65

4.6 Oxygen stability of SnO2-x …………………...…………..….71

5 Transparent thin film transistor based on SnO2

with amorphous gate oxides……. 73

5.1 Transparent thin film transistors (TTFT)……….….....……. 74

5.2 Investigation of HfO2 gate dielectrics ………........................79

viii

5.3 Devices utilizing HfO2 gate dielectric ...................................81

5.3.1 Device fabrication process ……………..............…..81

5.3.2 Output characteristics of HfO2/SnO2 FET ................83

5.3.3 Transfer characteristics of HfO2/SnO2 FET ..............84

6 Summary…………………….……………...…………...86

Bibliography…..…………………...……..…..……………..89

Abstract (In Korean) .……………..……………………….96

List of publications and Presentations ..………...….…......98

ix

List of Figures Figure 1.1: Electron mobilities of (Ba,La)SnO3 single crystal and other

semiconductors plotted against the carrier density.………...............………...… 4

Figure 1.2: Resistance change of BLSO epitaxial films at different gases. (a) Time

dependence of the resistance at T = 530 C. (b) Temperature dependence of the

resistance with the environments of three different gases…………………….... 5

Figure 1. 3: Flexible TTFTs. (a), Structure of TTFT fabricated on a plastic sheet.

(b), A photograph of the flexible TTFT sheet bent at R = 30 mm. The TTFT sheet

is fully transparent in the visible light region. (c), A photograph of the flexible

TTFT sheet ...……………………..…………………………………………….. 6

Figure 2.1: Pulsed laser deposition (PLD) system …………….……....…….…. 9

Figure 2.2: A schematic of pulsed laser deposition system ……….………….… 9

Figure 2.3: (a) theta-2theta diffraction pattern and (b) omega-scan of a 4% BLSO

thin film at various growth temperature of 600, 650, 700, and 750 °C ……...… 12

Figure 2.4: Theta-2theta diffraction patterns of BLSO thin films at various La-

doping level. …………………………………………………………..……… 13

Figure 2.5: Omega-rocking curves of the BLSO thin films at various La-doping

level. …………………………………………………………...……………… 14

Figure 2.6: Phi-scans of the (110) reflection for the 4% BLSO (001) film and the

(110) reflection for the SrTiO3 (001) substrate .………..…………..……..…… 15

Figure 2.7: Reciprocal space map (RSM) around the (103) asymmetric Bragg

diffraction point for BLSO film on STO.……………………………………….16

x

Figure 2.8: (a) Surface morphologies of BLSO films grown at various

temperatures. (b) Grain size and surface roughness as a function of growth

temperature ...…………………………………………………………………...17

Figure 2.9: The origin of (a) misfit dislocations and (b) threading dislocations 19

Figure 2.10: Cross-sectional TEM bright-field images of BLSO films grown at (a)

600 C and (b) 750 C …………………………………………………………….19

Figure 2.11: Surface morphologies of BLSO films by AFM (a) before and (b) after

etching with 0.4% nitric acid for 10 seconds ………………………………….20

Figure 3.1: The change of conduction band edge and the deformation potential

due to thermal expansion or contraction of the lattice spacing ..……..……..… 26

Figure 3.2: (Color online) Calculated mobility versus temperature for LO-phonon

scattering (solid red circles) for n=1020 cm-3. The calculated mobility due to

scattering from the individual phonon modes is shown:LO1 (purple open squares),

LO2 (green open circles), and LO3 (orange triangles) with energies 18, 51 and 88

meV, respectively…………………………………………..…………………. 29

Figure 3.3: (Color online) (a) Calculated drift mobility versus electron density

(cm-3) at 300 K (RT) for LO-phonon scattering, LO (orange dotted line), and

ionized impurity scattering, impurity (blue dashed line), as well as the total drift

mobility, total (black solid line). (b) Comparison of the screened (orange dotted

line) and unscreened (green open circles on dotted line) values for LO versus

electron density ……………………………………...…………..……………. 37

Figure 3.4: Electron mobility of the BLSO single crystals. The black dashed

curves are electron mobilities limited by each single scattering source: phonons

and ionized impurities. The red curve is the combination of the two black dashed

ones by the Matthiessen's rule .....………………….………………………….. 39

Figure 3.5: Electron mobility of the BLSO single crystals and epitaxial films on

STO substrate. …………………………………………………..…………….. 40

Figure 3.6: Formation energy of intrinsic and extrinsic defects in BSO and BLSO

xi

systems. (a) Intrinsic defects in Sn-rich environment. (b) Intrinsic defects in

O-rich environment. (c) Extrinsic defect in Sn-rich environment. (d) Extrinsic

defect in O-rich environment ....………………………………….…………… 41

Figure 3.7: Comparison of carrier concentration obtained by Hall measurements

and La impurity concentration. The dashed line indicates 100% doping efficiency

………………………………………………………………………………… 42

Figure 3.8: Electron mobility of the BLSO thin film. The black dashed curves are

electron mobilities limited by each single scattering source: phonons, ionized, and

threading dislocation impurities. The black curve is the combination of the three

black dashed ones by the Matthiessen's rule ……………………………......… 44

Figure 3.9: Electron mobility of the BLSO thin film. The black dashed curves are

electron mobilities limited by each single scattering source: phonons, ionized,

threading dislocation, and neutral impurity impurities. The black curve is the

combination of the three black dashed ones by the Matthiessen's

rule.…………………………………………………………………………..... 45

Figure 3.10: (a) The resistivity-temperature curves and (b) Carrier concentration

and Hall mobility measured from BLSO/STO(001) films grown at various

temperatures …………………....…………………………………….………. 46

Figure 3.11: Mobility of BLSO films are plotted as a function of carrier density n

at room temperature. GaN data are shown for comparison. {A, B} films and

{C, D} films have same carrier density but different dislocation density .……. 49

Figure 4.1: Interface structure of (110) SnO2 and r-plane Al2O3 …………..… 53

Figure 4.2: (a) XRD theta-2theta pattern of SnO2 films on r-Al2O3 substrate. (b)

XRD phi scan of 1um-thick SnO2 film on r-Al2O3 substrate ….…….……….. 54

Figure 4.3: Hall mobility μ, carrier concentrations n, resistivity ρ of SnO2 films on

r-plane sapphire substrate as a function of film thickness at oxygen pressure of 10

mTorr ……………..…………………………….……………………..….…. 56

Figure 4.4: Dark field image of SnO2/r-Al2O3 interface viewed in the SnO2[010]

xii

direction. Dark field image of SnO2/r-Al2O3 interface viewed in the in the SnO2[-

101] direction. High-resolution TEM of SnO2/r-Al2O3 interface viewed in the

SnO2[010] direction and in the in the SnO2[-101] direction .…………………. 58

Figure 4.5: XRD theta-2theta pattern of SnO2 and Ru-doped SnO2 films on r-

Al2O3 substrate ……………………………………………………………….. 60

Figure 4.6: (a) Hall mobility μ, carrier concentrations n, resistivity ρ of 30 nm

SnO2 films with and without 1 um Ru-doped SnO2 buffer layer on r-plane sapphire

substrate. (b) Hall mobility μ of 30 nm SnO2 films on 1 um Ru-doped SnO2 buffer

layer on r-plane sapphire substrate plotted as a function of carrier concentrations

n ..……………………………………………….…………………..…..……. 61

Figure 4.7: (a) XRD theta-2theta pattern of SnO2 and TiO2 films on r-Al2O3

substrate. (b) XRD omega scan of TiO2 and SnO2 film on r-Al2O3 substrate,

respectively …………………………………………………...……………… 63

Figure 4.8: XRD θ-2θ pattern of 1 um reflection of (101) SnO2-x in 10, 30, 50, and

100 mTorr oxygen pressure. (b) XRD φ scan of 500 nm SnO2 film on r-plane

Al2O3 substrate in an oxygen pressure of 100 mTorr ………..……...……….. 64

Figure 4.9: (a) Hall mobility μ, carrier concentrations n of 30 nm SnO2-x films on

r-plane Al2O3 substrate on SnO2 buffer layer of various thicknesses ……..….. 66

Figure 4.10: Hall mobility μ of 30 nm SnO2-x films on 500 nm fully oxidized

insulating SnO2 buffer layer on r-plane Al2O3 substrate plotted as a function of

carrier concentrations n ………………………………………………….….... 67

Figure 4.11: (a) XRD θ-2θ pattern of 500 nm SnO2-x on r-plane Al2O3 substrate

and (b) blow-up of (101) reflection of SnO2-x in 10, 30, 50, and 100 mTorr oxygen

pressure. (c) XRD φ scan of 500 nm SnO2 film on r-plane Al2O3 substrate in an

oxygen pressure of 100 mTorr ..………………………………………...…….. 69

Figure 4.12: (a) XRD omega scan of 500 nm (101) SnO2-x on r-plane Al2O3

substrate and (b) full with half maximum value of 500 nm (101) SnO2-x in 10, 30,

50, and 100 mTorr oxygen pressure, respectively ……………………….….. 70

xiii

Figure 4.13: Conductance changes of the 100 nm SnO2-x during a thermal

annealing cycles under argon, oxygen and air atmosphere ..…………….….. 72

Figure 5.1. Schematics of conventional FETs .………………………………. 75

Figure 5.2: (a) Plot of 𝐼𝐷𝑆1/2- VGS , (b) IDS - VGS output, and (c) IDS-VGS transfer

characteristic of HfO2/SnO2 field effect transistor ............................................. 77

Figure 5.3: Optical microscope image and measurement schematic of ITO/oxide

dielectric/ITO capacitor ...……………………………………………………. 79

Figure 5.4: C-F characteristic of an ITO/HfO2/ITO capacitor are shown. Filled

circle and unfilled square markers denote the capacitance and dissipation factor,

respectively ...………………………………………..…..………………….… 80

Figure 5.5: The fabrication process of HfO2/BLSO devices ..…………...….… 82

Figure 5.6: optical microscope image of an Al2O3/BLSO FET device …….… 82

Figure 5.7: The output characteristics, IDS-VDS curves, of (a) 15 nm and (b) 150

nm SnO2 TFTs for various gate bias VGS at room temperature. Gate voltages of

TFTs are varied from 8 V to -2 V in 1 V steps …………..…………………… 83

Figure 5.8: The transfer characteristics, IDS-VGS and 𝜇𝐹𝐸-VGS, of (a) 15 nm and (b)

150 nm SnO2 TFTs are represented as the solid line and the symbols, respectively.

The Ion/Ioff ratio at VGS of 15 nm and 150 nm SnO2 shows 105 and 106, respectively.

..……………………………………………………………..………………… 85

xiv

List of Tables

Table 1: Frequency of experimental IR and Raman peaks in comparison with the

calculated modes of cubic BSO at the zone center (C-point) and at the M-point of

the Brillouin zone .……………….……….……………………………….…... 28

Table 2: Electrical properties of SnO2-x on 10 mTorr, 50 mTorr, and 100 mTorr

oxygen growth pressure of TiO2 on r-plane sapphire substrate …….……...…. 62

1

Chapter 1

Introduction

1.1 Perovskite oxide semiconductor

Perovskite-type ABO3 oxides are an important class of metal oxides since they

exhibit variety of physical properties such as high dielectric constant,

superconductivity, colossal magnetoresistance, and multiferroicity [1-3].

Recently, the interest in perovskite-type transparent conducting oxides (TCOs)

and transparent oxide semiconductors (TOSs) with large band gap and high

mobility has been increasing due to huge potential for integrating display and logic

devices. For example, perovskite-type ternary compounds such as La-doped

SrSnO3, Sb-doped ZnSnO3 and SrSnO3, all promising as material for TCO films,

have been reported [4-6]. Perovskite oxides are an extremely important class of

ternary oxide materials. The ideal perovskite structure is cubic (space group

Pm3m) with the larger A cation having 12-fold coordination and the smaller B

cation in 6-fold coordination with an octahedron of O anions. Simple substitution

2

of the A or B cation induces new properties in the perovskite structure partial

substitution of the A and B lattice sites is also an effective way to enhance

desirable properties. For example, partial substitution of Ti atoms in lead titanate,

PbTiO3, with Zr atoms to form Pb(Zr,Ti)O3 is common to manipulate the

ferroelectric and piezoelectric properties. [7-11] In addition, development of

growth technology of complex structure may have enlarged the application to

perovskite oxide heterostructures and superlattices with combined functional

properties. For example, rapid improvement of perovskite solar cells has made

them the rising star of the photovoltaics world and of huge interest to the academic

community. Therefore, since their operational methods are still relatively new,

there is great opportunity for further research into the basic physics and

applications around perovskites.

1.2 High electron mobility of BaSnO3

Transparent perovskite stannate BaSnO3 shows great potential as a high-

mobility electron transport material composed of abundant elements. Since the

report of an extraordinary high room temperature mobility of 320 cm2/Vs for La-

doped BaSnO3 single crystals, which is the highest value reported for perovskite

oxides, the material has rapidly attracted interest as high-mobility channel layers

in oxide thin film transistors [12, 13] and multi-functional perovskite-based

optoelectronic devices [14, 15]. To fully exploit the potential of (Ba,La)SnO3 for

device applications, current research concentrates on understanding and

improving the electron transport in epitaxial (Ba,La)SnO3 thin film The high

mobility in (Ba,La)SnO3 single crystals is attributed to the large dispersion of the

Sn 5s orbital-derived conduction band and the ideal 180 degree O-Sn-O bond

angle in the network of corner sharing (SnO6)2- octahedra in the cubic perovskite

structure. Quantitatively, the electron mobility is given by

𝜇 =𝑒𝜏

𝑚∗ (1.1)

3

where 𝑒 is the electron charge, m* is the electron effective mass and 𝜏 is the

momentum relaxation time denoting the average time of momentum loss by

scattering. A fundamental understanding of the electron transport in epitaxial

(Ba,La)SnO3 films thus requires a quantitative analysis of the effective mass m*

and scattering relaxation time, which both strongly dependent on carrier

concentration. Its room temperature mobility is more than an order of magnitude

higher than that of perovskite oxides with conduction bands derived from d

orbitals, of which SrTiO3 (STO) prototypical example as shown in Figure 1.1.

The small effective mass has been suggested as the primary cause for the high

mobility. However, here we will show that the mass is not the only reason, and

that BSO has a significantly lower scattering rate than other perovskite metal

oxide, for instance, STO.

The large difference in mobility between single crystals and epitaxial films was

attributed to the existence of the dislocations in the epitaxial films from the

mobility’s dependence on the carrier density. The effect of dislocations on the

electrical transport properties have hardly been studied for perovskite films.

Many perovskite films of interest have been either in a highly metallic regime or

in a highly insulating regime, where the effects of dislocations are not evident. In

a few cases where the carrier density is relatively low, the films have been

homoepitaxial, without a need to consider the effect of dislocation. [16] In this

letter, we have investigated the influence of dislocation for electrical properties of

BLSO films on SrTiO3 (STO) substrates in detail, as we change the dislocation

density.

4

Figure 1.1: Electron mobilities of (Ba,La)SnO3 single crystal and other

semiconductors plotted against the carrier density.

1.3 Thermal stability

The oxygen stability of BSO is confirmed by high temperature thermal

annealing in O2, Ar, and Air environments. The oxygen diffusion constant is

evaluated to be about 10-15 cm2 s-1 at 530 °C. This value is 2~8 orders of magnitude

lower than titanates, cuprates, and manganites. [17-21] Recently, owing to these

two superior properties of BSO, the published articles on the transport and

transport mechanisms of BSO are consistently increasing. [22-28] Based on these

two noble properties and with dopability in two different cation sites, we can

expand the research field as shown in Figure 1.2. The author strongly believes that

they are remarkable works in oxide electronics, if the demonstrations of a FET, p-

n junction, and 2-dimensional electron gas (2DEG) based on BSO is possible.

5

Figure 1.2: Resistance change of BLSO epitaxial films at different gases. (a) Time

dependence of the resistance at T = 530 C. (b) Temperature dependence of the

resistance with the environments of three different gases.

1.4 Transparent conducting oxide

(TCO) and transparent oxide

semiconductor (TOS)

Transparent oxide semiconductor (TOS) are a series of metal oxides, composed

of heavy metal cations (HMCs) with an outside shell electronic configuration of

(n-1)d10 ns0 (n≥4) and oxygen anions. TOSs (Transparent Oxide Semiconductors)

have wide bandgaps in which the ns-orbitals of the heavy metal cations primarily

constitute the bottom part of the conduction band and the oxygen 2p orbitals form

the top of the valence band Uniquely, the spatial spreading of the outside ns

6

orbitals with a spherical symmetry in the HMCs is much larger than that in light

metal cation such as aluminum, leading to a wider conduction band.

Transparent electronic materials with a large bandgap and little subgap

absorption are attracting a large interest. For example, transparent conducting

oxides (TCO) are used as essential materials for photovoltaic devices [29],

flexible electronics [30] and optoelectronic applications [31]. TCOs such as

indium tin oxide (ITO) [32], F-doped tin oxide (FTO) [33], and Al-doped zinc

oxide [34] exhibit large conductivities with their typical carrier densities larger

than 1020 cm-3 in the degenerately doped regime. On the other hand, there are

needs for materials with low carrier density and high mobility in the

semiconducting regime. Good examples are transparent oxide semiconductors

(TOS), most prominent of which is indium gallium zinc oxide (IGZO) being

developed as the thin film transistor (TFT) materials for display industry as shown

in Figure 1.3. However, a common problem faced by IGZO transistors is the

instability of the threshold voltage during device operation because of oxygen

unstable property of oxygen. Therefore, the main issue is to find appropriate

material for TCO application. Better understanding and control of its

conductivity and emission would enhance its performance in existing applications

and enable new fields.

Figure 1. 3: Flexible TTFTs. (a), Structure of TTFT fabricated on a plastic sheet.

(b), A photograph of the flexible TTFT sheet bent at R = 30 mm. The TTFT sheet

is fully transparent in the visible light region. (c), a photograph of the flexible

TTFT sheet.

7

Chapter 2

Epitaxial growth of

BaSnO3 thin films

2.1 High quality BaSnO3 sample

growth using pulsed laser deposition

As a materials growth technique, laser ablation was utilized for the first time in

the 1960's, after the first commercial ruby laser was invented. Nevertheless, as a

thin film growth method it did not attract much research interest until the late

1980 s, when it has been used for growing high temperature superconductor films.

Since then, the development of the pulsed laser deposition (PLD) technique has

been more rapid and the amount of research devoted to this topic has increased

dramatically. By using high-power pulsed UV-lasers and a vacuum chamber, a

variety of stoichiometric oxide films can be grown in a reactive background gas

without the need for further processing. The author also has used the pulsed laser

deposition system to grow high quality BaSnO3 (BSO) thin films as shown in

Figure 2.1 and 2.2. Pulsed laser deposition has proved to be a promising method

8

for producing complex metal oxides thin films. In comparison with other

methods, PLD has the capability of controlling many process parameters, such as

target-to-substrate distance, laser pulse width, laser frequency, energy, and

wavelength along with background reactive gas pressure and substrate bias and

temperature which can strongly influence on film properties.

The film growth process is quite simple in pulsed laser deposition system.

Intense laser pulses are focused in a vacuum chamber onto a target surface where

they are absorbed. Above a threshold power density depending upon the target

material, significant material removal from the target occurs in the form of an

ejected luminous “plume” containing separated ions. The plume then propagates

in a direction normal to the target with a possible interaction with an ambient gas.

Finally, the composition of individual ions can reach on the substrate.

In addition, PLD is incredibly precise. It can deposit for example a film of

YBCO that is one unit cell (1.2 nm) thick; experiments show that a single unit cell

of YBCO is a superconductor. In complex multicomponent material deposition

with conventional evaporation methods, the various cations come from different

sources. To produce the right mixture in the deposited film, the rate of arrival of

each species must be monitored and controlled. This becomes especially difficult

when large background gas pressures are used during the deposition. However,

PLD does not require such monitoring because the composition of the film

replicates the composition of the target. Moreover, in most ion beam based

techniques, the pressure of the background gas in the chamber puts severe

limitations on the operating parameters. For electron-beam evaporation, the

background gas pressure cannot exceed 10-4 mbar, and for sputtering the pressure

determines the rate of atom ejection from the target surface.

9

Figure 2.1: Pulsed laser deposition (PLD) system

Figure 2.2: A schematic of pulsed laser deposition system

10

In this experiment, all BSO thin films were grown by pulsed laser deposition

technique. A pulsed excimer laser with a repetition rate of 10 Hz was used to

deposit BSO films. The wavelength of the laser was 248 nm and the pulse

duration was nominally 25 ns. The base pressure in the deposition chamber was

6ⅹ10-6 Torr and the oxygen backfill pressure, during laser ablation, was

controlled 100 mTorr using a precision mass flow controller. The substrates were

resistively heated to growth temperature of individual sample and the films were

cooled down to room temperature with oxygen pressure of 500 Torr.

11

2.2 Structure properties of La-doped

BaSnO3 thin films

2.2.1 X-ray diffraction analysis (XRD)

For proof of the crystallinity of the La-doped BSO (BLSO) thin film, we

obtained the structural information of 4 % BLSO thin film by the X-ray diffraction

method. Figure 2.3(a) shows the x-ray diffraction (XRD) θ-2θ data of BLSO films

of from various growth temperatures. The range of 2θ is taken from 20° to 75°

in the measurement. It is noticeable that there is no other phases or randomly

oriented grains appeared in the scans, indicating that the films were epitaxially

grown pure cubic perovskite phase BaSnO3 on the STO (001) substrates. The

BLSO thin film which is grown even at 600 °C shows very good quality epitaxial

characteristics. The diffraction peaks of BLSO (003) which is difficult to observe

in the thin film, indicate the excellent crystallinity. The ω-scan rocking curves

on BLSO (002) peaks are also shown in Figure 2.3(b). The full widths at half

maximum (FWHM) are the indicator of crystalline quality because the more

aligned the grain is along c-axis, the sharper the rocking curve will be. Their full

width at half maximum (FWHM) were measured from 0.110 to 0.086 as the

growth temperature increases.

This full with half maximum value is remarkable value compare with

previously reported BLSO film. In the previous report, The FWHM value was

0.57 and it is interesting to note that a previous study found the mobility to be at

best 0.69 cm2/Vs although the reported BLSO films were epitaxially grown on the

same type of substrate. It implies the importance of crystallinity.

12

Figure 2.3: (a) theta-2theta diffraction pattern and (b) omega-scan of a 4% BLSO

thin film at various growth temperature of 600, 650, 700, and 750 °C.

For proof of the epitaxiality of the BLSO thin film at various doping level, we

obtained the structural information of individual BLSO thin film. Figure 2.4

shows the x-ray diffraction (XRD) θ-2θ data of BLSO films of from various La-

doping level from 0.4 to 4 atomic percent (at%). It is also, no other phases or

randomly oriented grains appeared in the scans, indicating that the films were

grown epitaxially on the STO (001) substrates.

13

Figure 2.4: Theta-2theta diffraction patterns of BLSO thin films at various La-

doping level.

The omega-rocking curves of the BLSO at various doping concentration are

presented in Figure 2.5. According to the data, it is hard to make a conclusion

whether the FWHM depend on the La contents. The FWHMs of various doping

depending samples do not seem to make any large differences between in the La-

doping level. It also shows small FWHM value, irrespective of doping

concentration. However, a previous study of BLSO ceramics, at higher doping

concentration, had already reported the presence of Ba2Sn2O7 phase in BLSO

system with La the content of which is larger than 5 %.

Although these are very low values for FWHM of the rocking curves, the slight

broadening may be related with defects such as dislocations and grain boundaries.

14

Figure 2.5: Omega-rocking curves of the BLSO thin films at various La-doping

level.

Figure. 2.6 shows phi-scans of the (110) reflection for the BaSnO3 (001) film

and the (110) reflection for the SrTiO3 (001) substrate. The extent of in-plane

alignment was examined by taking phi-scan of (110) reflection of BaSnO3 film

and SrTiO3 substrate. A peak is observed every 90° exhibiting the for fold

symmetry expected for BaSnO3, with no twinning, a good indication of no

significant strain experienced in the film.

The in-plane lattice constant and strain state of the films were examined by X-

ray reciprocal space mapping (X-RSM) around the (103) reflections as shown in

Figure. 2.7. It is seen that, due to a large lattice mismatch of 5.4 % between the

BLSO films and the STO substrates, the films are almost strain relaxed without

significant heteroepitaxial strain. It is calculated that the BLSO film has nearly

the same in-plane and out-of-plane lattice constants of 4.112 Å and 4.127 Å ,

respectively, implying that the in-plane direction BLSO films show slight

compressive strain, while they are slightly elongated along the out-of-plane

15

direction. The broader (103) reflection of the films grown at lower temperature

means higher density of defects exists, which is consistent with the XRD rocking

curve results. Therefore, from our x-ray diffraction study, we could conclude that

a small microstructural difference exist in the films grown at various temperatures.

Figure 2.6: Phi-scans of the (110) reflection for the 4% BLSO (001) film and the

(110) reflection for the SrTiO3 (001) substrate.

16

Figure 2.7: Reciprocal space map (RSM) around the (103) asymmetric Bragg

diffraction point for BLSO film on STO.

17

2.2.1 Atomic force microscope (AFM).

Figure 2.8(a) shows the surface morphology of the 100 nm thick 4% BLSO film

grown at various temperatures by an atomic force microscope (AFM). The film

surfaces show a topographic structure made by grains although the overall

roughness of the films is very small ranging from 0.5 nm for 600 °C growth to 0.2

nm for 750 °C growth [Figure 2.8 (b)]. Apparent coarsening of the surface of the

BLSO film grown at lower temperature was observed. Additionally, at higher

growth temperature, larger grains were formed. Consistent with our XRD

analysis, we believe that the grain boundary and dislocation density is smaller for

the films grown at higher temperatures.

Figure 2.8: (a) Surface morphologies of BLSO films grown at various

temperatures. (b) Grain size and surface roughness as a function of growth

temperature.

18

2.2.3 Transmission electron microscopy (TEM)

The BaSnO3 films that we deposited on SrTiO3 with a lattice constant 3.905 Å

have 5.4 % lattice mismatch with the substrate. Lattice mismatched

heteroepitaxial layers, most of the mismatch may be accommodated by misfit

dislocation and threading dislocation during growth as shown in Figure 2.9.

Misfit dislocation and threading dislocation generation must be considered for

situations in which the film is grown well in excess of a critical thickness. For

lattice mismatched films, regardless of the growth mechanisms, such as Frank-

van der Merwe (layer-by-layer), Stranski-Krastanov (initial wetting followed by

islanding), or Volmer-Weber (incoherent islanding) increasing film thickness will

ultimately lead to misfit dislocation generation and concomitant threading

dislocations. The 100 nm thick 4% BLSO film was investigated by transmission

electron microscopy for crystal defects. The cross-sectional TEM is instrument

which can observe the crystal defect directly. In those images, the misfit

dislocations are shown as the black dots confined in the interface while, the

threading dislocations are indicated by the dark lines crossing from interface

between film and substrate to top surface of film. For estimating density of the

misfit dislocations, one can count 12 or 13 black dots in the lateral scale of 100

nm BLSO near the interface in Figure 2.10 (a) and (b). The misfit dislocation in

every 19 lattice spaces at the interface provides compensation for the large lattice

misfit between BSO film and STO substrate. 19 lattice spacing of BSO along the

[010] direction is 7.81 nm and this matches well with 20 spacing of STO along

same direction. This calculation is the consistent with the simple observation

from TEM image. TDs form when adjacent islands coalesce in the case there is

a lattice mismatch between the film and the substrate. When the density of island

nuclei were high enough and coalescence was to occur at an early stage in the

growth, a high density of TDs could arise with Burgers vectors parallel to the

interface. At lower growth temperature, the density of island nuclei is higher,

resulting in higher density of dislocations.

19

Figure 2.9: The origin of (a) misfit dislocations and (b) threading dislocations.

Figure 2.10: Cross-sectional TEM bright-field images of BLSO films grown at (a)

600 °C and (b) 750 °C.

For the misfit dislocations, the density of misfit dislocation can be easily

calculated in the film-substrate interfaces. In the case of BSO/STO interfaces, the

density is calculated to be 1.3ⅹ106 cm-1.

20

2.2.4 Etch-pit development (EPD)

From the TEM image, we can evaluate the density of misfit dislocation along

[010] lateral direction of BLSO film. However, for the density of threading

dislocations, it is not very clear to count the number of dark lines in the images.

As the second alternative plan, we were able to directly confirm the threading

dislocation density using our AFM. The AFM images of BLSO films are

obtained in Figure 2.11 (b) after etching with 0.4 % diluted nitric acid (NHO3).

The etching creates depressed holes near the dislocations core due to slightly faster

etching of the defective atomic structure. When counted from the AFM surface

image, the TD density was 8.9 × 1010 at growth temperature 750 °C, respectively.

Figure 2.11: Surface morphologies of BLSO films by AFM (a) before and (b) after

etching with 0.4% nitric acid for 10 seconds.

21

Chapter 3

Electrical transport

properties of La-doped

BaSnO3

3.1 Scattering mechanisms and carrier

mobility in BaSnO3

At room temperature, the electron mobility of the La-doped BaSnO3 (BLSO)

system is unprecedentedly high compared with that of other wide band gap

semiconductors. It has been recently published about high mobility in BLSO

materials: as high as 320 cm2/Vs at the carrier concentration 8.0 × 1019 cm−3 in

single crystals. It has the highest room temperature mobility among TCOs. Its

room temperature mobility is more than an order of magnitude higher than that of

perovskite oxides with conduction bands derived from d orbitals, of which SrTiO3

(STO).

In this chapter, the author will introduce the characteristics of BSO that impart

such a high mobility, and the fundamental limits on this mobility. The small

effective mass has been suggested as the primary cause for the high mobility.

However, It will be shown that the mass is not the only reason, and that BSO has

22

a significantly lower scattering rate than, for instance, STO.

3.1.1 Electrical mobility and effective mass

The electron mobility may be defined in terms of the conductivity effective mass

𝑚∗ and the scattering rate 𝜏-1 by

𝜇 =𝑒𝜏

𝑚∗ (3.1)

where e is an electron charge. Therefore, high electrical mobility can be realized

with a small electron effective mass and a small total electron scattering rate.

In solids, when the energy of a free carrier (E) is a function of the momentum (k),

the effective mass (m*) can be expressed as below:

m∗ = (ħ𝑑2𝐸(𝑘)

𝑑𝑘2 ) (3.2)

where h is Planck’s constant.

However, the non-parabolicity of the La:BaSnO3 conduction band is derived from

the dependence of electron effective mass on doping level.

Therefore, the energy dispersion E(k) of electrons near the conduction band

minimum can be described using a first-order parabolic approximation.

3.1.2 Total scattering rate

The free carrier scattering rate of each scattering source can be defined in terms

of the scattering cross-section by:

23

τ−1 = 𝑁𝜎𝑣 (3.3)

where N is the density of the scattering centers, σ is the total cross-section of the

scattering center, and ν is the velocity of free carriers. It is the mean thermal

velocity [νth = (3kBT/m*)1/2] for the non-degenerate semiconductor and constant

of Fermi velocity [ν = νF] for degenerate semiconductors. Therefore, once we

know the form of the total cross-section, we can estimate the scattering rate with

the density of the impurities. The total cross-section for isotropic elastic scattering

can be calculated from the differential scattering cross-section, σ(θ), by using the

below equation:

σ = 2π ∫ σ(θ′)(1 − cos(θ′) sinθ′𝑑θ′2π

0 (3.4)

The differential scattering cross-section can be calculated from the quantum

mechanical scattering problem i.e., a free carrier scatters from the initial state k

into the final state k’ due to the perturbing potential. In this case, the time

dependent Schrödinger equation can be expressed with the perturbed Hamiltonian

H.

HΨ𝑘(𝑟, 𝑡) = (H0 + H′)Ψ𝑘(𝑟, 𝑡) = −𝑖(ħ𝜕

𝜕𝑡Ψ𝑘(𝑟, 𝑡)) (3.5)

where H0 and H’ are the unperturbed Hamiltonian of a solid and the first-order

correction due to the perturbation, respectively. The solution of equation is given

by:

HΨ𝑘(𝑟, 𝑡) = ∑ 𝑎𝑘

𝑘

(𝑡)𝑒−𝑖2𝜋𝐸𝑘𝑡

ℎ 𝜑𝑘(𝑟) (3.6)

Based on the time-dependent perturbation theory, the transition probability per

unit time from k- to k’-state is given by:

24

𝑃𝑘𝑘′(𝑡) =|𝑎𝑘′(𝑡)|2

𝑡=

4𝜋2

ℎ|𝐻𝑘𝑘′|2𝛿(𝐸𝑘′ − 𝐸𝑘) (3.7)

Where

𝐻𝑘𝑘′ = ⟨𝑘′|𝐻′|𝑘⟩ =1

𝑁𝛺∫ 𝜑𝑘′

∗𝐻′𝜑𝑘′∗

𝑁𝛺

𝑑3𝑟 (3.8)

The matrix element Hkk’, has a finite value when the Fermi golden rule is

satisfied. Based on this, the differential scattering cross-section can be written as:

𝜎(𝜃′, 𝜙′) =

𝑁𝛺(2𝜋)3 𝑃𝑘𝑘′

(𝑑3𝑘′)𝑑𝜔

𝑣𝑘𝑁𝛺

=(𝑁𝛺)2𝑃𝑘𝑘′𝑑3𝑘′

(2𝜋)3𝑣𝑘𝑠𝑖𝑛𝜃′𝑑𝜃′𝑑𝜙′ (3.9)

where, vk and dω = sinθ'dθ'dϕ' are the initial velocity of the incident particle and

the solid angle between the incident wave vector k and the scattered wave vector

k’. Now, considering the case of isotropic elastic scattering and the equation (3.7)

and (3.8) into equation (3.9), the final form of the differential

scattering cross-section is given by:

𝜎(𝜃′) =(𝑁𝛺)2𝑘′2|𝐻𝑘𝑘′|

2

(ℎ𝑣𝑘′)2 (3.10)

Therefore, if we know the H’kk’, then we can calculate the analytical form of

the scattering rate. From now on, it is important to know the main origin for the

perturbing Hamiltonian corresponding to each scattering. There are many

scattering sources such as electron phonon scattering, ionized impurity scattering,

neutral impurity scattering, and dislocation scattering.

25

3.1.3 Acoustic phonon scattering

Scattering of electrons by longitudinal-mode acoustical phonons is described in

this section. The scattering of electrons by longitudinal-mode acoustical phonons

is the most important scattering source in intrinsic or lightly doped

semiconductors at room temperature. The scattering of electrons by longitudinal-

mode acoustical phonons can usually be treated as an elastic scattering because

the electron energy is much larger than the phonon energy and the change in

electron energy during such a scattering process is small compared to the average

energy of electrons. The ratio of phonon energy to mean electron energy is

usually much smaller than unity for T > 100 K. At very low temperatures, mean

electron energy may become comparable to the acoustical phonon energy, and the

assumption of elastic scattering may no longer be valid. Fortunately, at very low

temperatures other types of scattering such as ionized impurity and neutral

impurity scattering may become dominant. It is noted that acoustical phonons

may cause scattering in two different ways, either through deformation potential

scattering or piezoelectric scattering. An acoustical wave may induce a change

in the spacing of neighboring atoms in a semiconductor. This change in atomic

spacing could result in the fluctuation of energy band gap locally on an atomic

scale and is known as the deformation potential as shown in Figure 3.1. The

deformation potential is measured as the change of energy band gap per unit strain

due to the acoustical phonons. This type of scattering is usually the most

important scattering source for intrinsic or lightly doped silicon and germanium

at room temperatures.

Piezoelectric scattering is another type of acoustical phonon scattering. This

type of scattering is observed in III-V and II-VI compound semiconductors with

the zincblende and wurtzite crystal structures. The lack of inversion symmetry

in these semiconductors creates a strain-induced microscopic electric field

perturbation, which leads to piezoelectric scattering with emission or absorption

of an acoustical phonon. This type of scattering is important for pure III-V and

II-VI compound semiconductors at low temperatures. These two types of

26

acoustical phonon scattering are discussed next.

Figure 3.1: The change of conduction band edge and the deformation potential

due to thermal expansion or contraction of the lattice spacing.

The scattering rate and the electrical mobility governed by the deformation

potential scattering are given by:

τD−1 = 16𝜋3 (

𝑚∗𝐸𝑐𝑜𝑛𝑠𝑡.

ℎ2)

2

(𝑘𝐵𝑇

𝐶𝑙) 𝑣 (3.11)

𝜇𝐷 =1

16𝜋3 (ℎ2

𝑚∗𝐸𝑐𝑜𝑛𝑠𝑡.)

2

(𝐶𝑙

𝑘𝐵𝑇𝑣) (3.12)

In the case of non-degenerate semiconductors, ν = νth = (3kBT/m*)1/2. Therefore,

the scattering rate of the deformation potential scattering is proportional to T 3/2,

which indicates that the electrical mobility governed by the deformation potential

scattering is proportional to T -3/2. However, for the degenerate semiconductor, v

= vF is constant, indicating that a scattering rate is proportional to T and the related

electrical mobility is proportional to T -1.

27

3.1.4 Optical phonon scattering

Optical phonon scattering becomes the predominant scattering source at high

temperatures or at high electric fields. Both polar and nonpolar optical phonons

are responsible for this type of scattering. The scattering of electrons by nonpolar

optical phonons may be treated as one type of deformation potential scattering

process. Nonpolar optical phonon scattering becomes important for silicon and

germanium crystals above room temperatures when intervalley scattering

becomes the dominant process. Polar optical phonon scattering is the

predominant scattering mechanism for ionic or polar crystals such as II-VI and

III-V compound semiconductors. For these crystals the motion of negatively and

positively charged atoms in a unit cell will produce an oscillating dipole, and the

vibration mode is called the polar optical-mode phonon. Polar optical phonon

scattering is associated with the atomic polarization arising from displacement

caused by optical phonons. This is often the most important scattering

mechanism at room temperature for III-V compound semiconductors. Optical

phonon scattering is usually an inelastic process that cannot be treated by the

relaxation time approximation because the optical phonon energy is comparable

to that of mean electron energy at room temperature.

For a multivalley semiconductor such as silicon or germanium, intravalley

scattering (i.e., scattering within a single conduction band minimum) near room

temperature is usually accompanied by absorption or emission of a longitudinal-

mode acoustical phonon. In this case, calculated mobilities in materials.

However, at higher temperatures, intervalley scattering (i.e., scattering from one

conduction band minimum to another) may become the dominant scattering

process. Intervalley scattering is usually accompanied by absorption or emission

of a longitudinal-mode optical phonon. Since the energy of an optical phonon is

comparable to that of the average electron energy, scattering of electrons by

intervalley optical phonons is generally regarded as inelastic. In this case, the

change in electron energy during scattering is no longer small, and hence the

relaxation time approximation can be used only if certain assumptions are made

28

for this type of scattering.

BSO has a 5-atom unit cell that leads to a total of 15 phonon modes, three of

which are polar longitudinal optical (LO) modes. In polar crystals, LO phonons

tend to dominate scattering at RT compared to other phonons due to their strong

long-range coulomb interaction. In addition, we need to assess ionized impurity

scattering, since large concentrations of dopants are intentionally introduced in

order to achieve carrier densities as high as 1019 ~ 1021 cm-3. From far-IR and

Raman spectra, It is reported BSO has three polar LO (18, 51 and 88 meV) and

the corresponding three doubly-degenerate TO (17, 30 and 78 meV) mode

frequencies.

Table 1: Frequency of experimental IR and Raman peaks in comparison with the

calculated modes of cubic BSO at the zone center (C-point) and at the M-point

of the Brillouin zone.

29

Figure 3.2: (Color online) Calculated mobility versus temperature for LO-phonon

scattering (solid red circles) for n=1020 cm-3. The calculated mobility due to

scattering from the individual phonon modes is shown:LO1 (purple open squares),

LO2 (green open circles), and LO3 (orange triangles) with energies 18, 51 and 88

meV, respectively.

3.1.5 Ionized impurity scattering

The origin of ionized impurity scattering is the Coulomb interaction between

free carriers and ionized dopants. This scattering also can be regarded as an

elastic scattering. The mass of ionized dopant is much larger than that of free

carriers, resulting in a negligible change in electron energy. Therefore, the

relaxation approximation also can be valid for the ionized impurity scattering. If

the charge number of ionized dopants is Z, the Coulomb potential can be written

as:

VC(𝑟) =𝑞𝑍

4𝜋휀𝑟휀0𝑟 (3.13)

30

where q, ε0, and εr are the electrical charge, the vacuum permittivity, and the

relative dielectric constant, respectively. However, in doped semiconductors, we

should consider the Coulomb screening effect due to the free carriers.

In the case of non-degenerate semiconductors, we can assume that the thermally

excited free carriers obey the Maxwell-Boltzmann statistics at high temperatures.

Based on these conditions, the screened Coulomb potential (Yukawa potential)

can be written as:

VC(𝑟) =𝑞𝑍𝑒−𝑟/𝜆𝐷

4𝜋𝜀𝑟𝜀0𝑟 with 𝜆𝐷 = √

𝜀𝑟𝜀0𝑘𝐵𝑇

𝑞2𝑁𝐼 (3.14)

Here, λD is the Debye screening length, and NI is the density of ionized dopants.

Therefore, the perturbing Hamiltonian can be expressed by:

𝐻′ = 𝑞𝑉𝐶(𝑟) =𝑞2𝑍𝑒−𝑟/𝜆𝐷

4𝜋휀𝑟휀0𝑟 (3.15)

Based on the Bloch theorem under the periodic potential, the electron wave

function of k-state is given by:

𝜑𝑘(𝑟) =1

𝑁𝛺𝑢𝑘(𝑟)𝑒−𝑖𝑘𝑟 (3.16)

Therefore, the matrix element of the perturbing potential due to the ionized

dopants can be calculated based on equation (3.8), resulting in:

𝐻𝑘𝑘′ =1

𝑁𝛺∫ 𝑒−𝑖𝑘′𝑟 (

𝑞2𝑍𝑒−𝑟/𝜆𝐷

4𝜋휀𝑟휀0𝑟) 𝑒𝑖𝑘𝑟𝑑3𝑘 =

𝑞2𝑍𝜆𝐷2

2𝑁𝛺휀𝑟휀0(1 + 2|𝑘|sin (𝜃′2

))

(3.17)

31

Thus, based on equation (3.10), the differential scattering cross-section is given

by:

𝜎(𝜃′) = (2𝜋𝑞2𝑚∗𝑍𝜆𝐷

2

휀𝑟휀0ℎ2 )

2

(1 + 4|𝑘|2 sin2 (𝜃′

2) 𝜆𝐷

2 )

−2

(3.18)

The scattering rate of the ionized impurity scattering is obtained by substituting

(3.18) into (3.4), and the result is given by:

𝜏𝐼,𝑘−1 = 2𝜋𝑁𝐼𝑣 (

𝜋𝑞2𝑚∗𝑍

휀𝑟휀0ℎ2|𝑘|2)

2

[𝑙𝑛(1 + 𝛾2) −𝛾2

(1 + 𝛾2) ] (3.19)

𝛾2 =ℎ2휀𝑟휀0

2𝑞2𝑚∗(2𝜋2𝑛)1/3 (3.20)

If we consider 𝑘2 = 4𝜋2𝑚∗𝐸/ℎ2, the scattering rate can be written as:

𝜏𝐼,𝑘−1 = 2𝜋𝑁𝐼√

2𝐸

𝑚∗ (𝑞2𝑍

4𝜋휀𝑟휀0𝐸)

2

[𝑙𝑛(1 + 𝛾2) −𝛾2

(1 + 𝛾2) ] (3.21)

This equation which is known as the Brooks–Herring formula directly indicates

that the relaxation time τI,k is proportional to E3/2. In the case of nondegenerate

semiconductors, to obtain the electrical mobility governed by ionized impurity

scattering, we should calculate the average relaxation time τI over the energy.

As we assume that the free carriers obey Maxwell-Boltzmann statistics, the

average relaxation time can be calculated with the below equation:

32

𝜏𝐼 =64(𝜋𝑚∗)

12(휀𝑟휀0)2(2𝑘𝐵𝑇)

32

𝑍2𝑞4𝑁𝐼[𝑙𝑛(1 + 𝛾2) −

𝛾2

(1 + 𝛾2) ]

−1

(3.22)

Therefore, the electrical mobility limited by the ionized impurity scattering is

given by:

𝜇𝐼 =64𝜋1/2(휀𝑟휀0)2(2𝑘𝐵𝑇)3/2

𝑍2𝑚∗1/2𝑞3

1

𝑁𝑖𝑖[𝑙𝑛(1 + 𝛾2) −

𝛾2

(1 + 𝛾2) ]

−1

(3.23)

In degenerate semiconductors, where the Fermi level is well above the

conduction band minimum, the electrons with the Fermi energy EF are the most

relevant agents to determine the transport properties because the movement of the

electrons, the energy of which is lower than the Fermi energy, tends to be

cancelled out by that of other electrons which propagate with the same speed but

into the opposite direction. As already mentioned in the previous subsection, the

Fermi energy of the degenerate semiconductors can be written in terms of carrier

density n.

𝐸𝐹 =ℎ2

8𝜋2𝑚∗(3𝜋2𝑛)2/3 (3.24)

When replacing E in Equations 3.21 by the EF in Equation 3.24 yields the

degenerate form of Brook-Herring formula.

𝜇𝐼 =3(휀𝑟휀0)2ℎ3

𝑍2𝑚∗2𝑞3

𝑛

𝑁𝑖𝑖 [𝑙𝑛(1 + 𝛾2) −

𝛾2

(1 + 𝛾2) ]

−1

(3.25)

In this equation, it is clearly seen that the electrical mobility limited by the

ionized impurity scattering shows the temperature-independent character.

33

3.1.6 Neutral impurity scattering

In conventional semiconductors, mobility due to neutral impurity scattering has

been effectively described by an electron scattering due to hydrogen-like

atoms,[35] which are formed by capturing a free carrier in the conduction band by

an ionized dopant. In this case, the total differential cross-section of neutral

impurity scattering can be written as:

σN =20 aB

𝑘 with 𝑎𝐵 =

εrε0h2

𝜋𝑚∗𝑒2 (3.26)

where aB is the scaled Bohr radius of such a bound state. The electrical mobility

due to the neutral impurities μN is then likely to be expressed as:

𝜇𝑁 =2𝜋𝑒

10𝑎𝐵

1

𝑁𝑁

(3.27)

Here, NN is the number of neutral impurities per unit volume. In a conventional

semiconductor, the neutral impurity scattering is usually observable at low

temperatures when a non-degenerate regime is realized with low carrier density.

The origin of such neutral impurities is attributed to the hydrogen-like bound state

formed by captured free carriers around the ionized dopants.[35] This kind of

neutral impurity can be easily broken by thermal fluctuation. In addition to this,

electrical mobility due to neutral impurity scattering is larger than 104 cm2/Vs

when the NN is smaller than 1018 cm-3, indicating that the neutral impurity

scattering is much smaller than the other scattering sources.

34

3.1.7 Threading dislocation scattering

Dislocations in a semiconductor can act as scattering centers for both electrons

and holes. The scattering of electrons by a dislocation may be attributed to two

effects. First, a dislocation act as a charge trap center. Second, a dislocation

may be viewed as a line charge, and hence has an effect similar to that of a charged

impurity center. To deal with scattering of electrons by dislocations, one may

consider the dislocation line as a space charge cylinder of radius R and length L.

The probability that an electron is scattered into an angle dθ` by a dislocation

line can be expressed by

𝑃𝑑𝑖𝑠 =𝑑(𝑏/𝑅)

𝑑𝜃′=

1

2sin (

𝜃′

𝑑𝜃′) (3.28)

where b is the scattering impact parameter. The differential scattering cross-

section per unit length of dislocation line charge is thus given by

𝜎𝑑𝑖𝑠(𝜃′) = 𝑅 sin (𝜃′

2) (3.29)

The total scattering cross-section can be obtained by substituting (3.29) into (3.4)

and integrating over θ’ from 0 to π, which yields

𝜎𝑇 =8R

3 (3.30)

Therefore, the relaxation time due to scattering of electrons by dislocations is

given by

𝜏𝑑𝑖𝑠 =1

𝑁𝑑𝑖𝑠𝜎𝑇𝑣=

3

8 𝑁𝑑𝑖𝑠𝑅𝑣 (3.31)

35

The electron mobility due to scattering by dislocations can be obtained directly

from (3.31), yielding

𝜇𝑑𝑖𝑠 =𝑞 𝜏𝑑

𝑚∗= (

3𝑞

8 𝑁𝑑𝑖𝑠𝑅)

1

(3𝑚∗𝑘𝐵𝑇)1/2 (3.32)

where Nd is the density of dislocation lines.

36

3.2 Electrical transport of La-doped

BaSnO3 single crystal and thin film

3.2.1 BLSO single crystal

For the BLSO single crystals, we already mentioned that there are dominant

two scattering source at room temperature; optical phonon scattering and ionized

impurity scattering. As explained in 3.1.2, BSO has a single nondegenerate CB

and therefore only intraband scattering occurs. Therefore, we focus on LO

phonons and neglect other phonons. Stanislavchuk et al. determined the

experimental three polar LO (154, 421 and 723 cm-1); or 18, 51 and 88 meV) and

the corresponding three doubly-degenerate TO (135, 245 and 628 cm-1; or 17, 30

and 78 meV) mode frequencies, and for the high-frequency dielectric constant

(ε∞=4.3).

Figure 3.2 shows the contributions to the mobility due to the individual LO

phonon modes. As expected, at low temperatures (0-100 K), only the lowest

frequency mode (LO1) is occupied and contributes to limiting the mobility.

Starting at 100 K the LO2 mode (51 meV) gets populated. The highest energy (88

meV) LO3 mode starts contributing to scattering at temperatures above 250 K.

LO-phonon contributions to mobility are often fitted to an expression that is

inversely proportional to the BE distribution as derived by Low and Pines for a

single LO mode. For materials with multiple LO modes, such as the perovskite

oxides, the fits are performed by adding the reciprocal mobilities due to each mode

with some assumption or knowledge about which modes dominate in the

temperature range of interest. At room temperature, the electron mobility limited

by the phonons in this analysis of temperature dependent mobility curves and been

calculated to be about 600 cm2/Vs as shown in Figure 3.3.

37

Figure 3.3: (Color online) (a) Calculated drift mobility versus electron density

(cm-3) at 300 K (RT) for LO-phonon scattering, LO (orange dotted line), and

ionized impurity scattering, impurity (blue dashed line), as well as the total drift

mobility, total (black solid line). (b) Comparison of the screened (orange dotted

line) and unscreened (green open circles on dotted line) values for LO versus

electron density. Obviously, LO phonon does not strong depend on the carrier

density since the phononic scattering is not scattering by charged impurities as

shown in Figure 3.3 (b).

The second source of scattering is the ionized La impurities, which, owing to

the charged Coulomb potential, are screened by the carriers and, thereby, produce

38

the carrier density dependence of electron mobility. In degenerate semiconductors,

the relation between the electron mobility and carrier density is well described by

the degenerate form of Brooks-Herring formula which is already mentioned in

3.1.3. In Equation 3.23, Z and Nii is the charged state (1 in this case) and the

density of the ionized impurities, respectively. The relation between n and Nii in

Equation determines the shape of the mobility curve. For the BLSO single

crystals, it might be safe to assume the full activation of the La impurities (n = Nii)

owing to no evident charge traps to be reckoned with. Finally, the two mobilities,

phonon and ionized impurity are to be combined by Matthiessen's rule.

1

𝜇𝑆𝐶,𝑇𝑜𝑡𝑎𝑙=

1

𝜇𝑝ℎ+

1

𝜇𝑖𝑖 (3.33)

The combination produces the solid curve in Figure 3.4, which agrees well with

the experimental data.

39

Figure 3.4: Electron mobility of the BLSO single crystals. The black dashed

curves are electron mobilities limited by each single scattering source: phonons

and ionized impurities. The red curve is the combination of the two black dashed

ones by the Matthiessen's rule.

40

3.2.2 BLSO thin film

In single crystals, it found that La-doped BaSnO3 has a high electron mobility

of over 300 cm2/Vs at room temperature (Figure 3.4). However, the highest

mobility of 70 cm2/Vs in our films is still much lower than the values in the single

crystals as shown in Figure 3.5. That indicate that there is additional scattering

source that limit the electron mobility of BLSO film compare with single crystal.

Figure 3.5: Electron mobility of the BLSO single crystals and epitaxial films on

STO substrate.

From the TEM analysis, we already revealed high density of threading

dislocations on the film.

41

Figure 3.6: Formation energy of intrinsic and extrinsic defects in BSO and BLSO

systems. (a) Intrinsic defects in Sn-rich environment. (b) Intrinsic defects in

O-rich environment. (c) Extrinsic defect in Sn-rich environment. (d) Extrinsic

defect in O-rich environment.

According to Figure 3.6(b), in the degenerated BLSO system in which the

Fermi level is located at well above the conduction band, not only double-charged

Ba vacancies can be easily generated but also quadruple-charged Sn vacancies are

generated with the almost equal likelihood. This has an important implication for

the charged state of the cores of the threading dislocations in the BLSO epitaxial

films. For the sake of simplicity, if only edge dislocations are to be considered,

the dislocation cores can be regarded as edge lines of inserted partial planes. The

charged state of the cores depends on the atomistic defects along the edge lines.

Considering the result of the formation energy calculation, it is both VBa2- and

VSn4- that are probably created along the cores the threading dislocations in the

BLSO epitaxial films. If this scenario is true, the cores of the threading

dislocations in the BLSO epitaxial films can be modelled by a line defect with a

charge density 6q/c where c is the c-axis parameter of the BLSO epitaxial films.

The density of the charge traps provided by the threading dislocation too is

calculated to be 6qNth/c. Applying the results obtained Ntd = 6 ⅹ1010 cm2 and c =

42

4.127 A, 6Ntd/c = 8.7 ⅹ1018 cm3.

Compare with single crystal case, it may be assumed that all the La atoms

are readily ionized to donate an electron to the reservoir of free carriers.

However, in our films, the doping level is significantly reduced as

compared to the La nominal doping concentration.

Figure 3.7: Comparison of carrier concentration obtained by Hall measurements

and La impurity concentration. The dashed line indicates 100% doping efficiency.

43

This is because defect that is able to trap the electron charges in the BLSO

epitaxial films is the La-antisite defect that describes the substitution of La for Sn

instead of Ba in the BSO system. In the BSO system, two cations, Ba and Sn, have

+2 and +4 valence states, respectively. Hence two natural ways of doping the

material. One is the substitution of Ba2+ by dopants with +3 valence state; the other

is the substitution of Sn4+ by dopants with +5 valence state. The problem of the

former is that the dopant with +3 valence state can also substitute the Sn4+. The

role of a La-antisite defect in the BLSO system is trapping one mobile carrier

because La impurities in Sn-site are likely to act as deep acceptors rather than

shallow ones. To estimate the density of the La-antisite defects in the BLSO

epitaxial films, a comparison is to be made between the nominal concentration of

La impurities and the measured carrier density.

44

Figure 3.8: Electron mobility of the BLSO thin film. The black dashed curves are

electron mobilities limited by each single scattering source: phonons, ionized, and

threading dislocation impurities. The black curve is the combination of the three

black dashed ones by the Matthiessen's rule.

In Figure 3.8, we show an analysis similar to that used for bulk samples in

Figure 3.4 to compare our calculated Hall mobility with the experimental values.

When comparing with experimental measurements, it is important to recognize

that scattering mechanisms at low carrier concentration is definitely dominant by

threading dislocation scattering limit, however, at high carrier concentration, we

should take additional mechanisms into account for instant due to the neutral

45

impurity scattering. If one electron is trapped near La3+, the bound state of La3+

+ e- would be effectively similar to Ba2+ so that it might act as a neutral impurity

in the system. However, from the experimental result, we could not distinguish

between La antisite and neutral impurity defect. Therefore, we just tried to

estimate the neutral impurity density in BLSO thin film compare with La impurity

concentration. We predict the neutral impurity concentration of 0.5 % of La

impurity concentration and then plot that result in Figure 3.9. Consequently it is

quite well fit with experimental results.

Figure 3.9: Electron mobility of the BLSO thin film. The black dashed curves are

electron mobilities limited by each single scattering source: phonons, ionized,

threading dislocation, and neutral impurity impurities. The black curve is the

combination of the three black dashed ones by the Matthiessen's rule.

46

In addition to that the author investigated the temperature-dependent electrical

transport properties of 4% La-doped BaSnO3. Figure 3.10 shows the temperature

dependence of resistivity (ρ-T) measured from the BLSO films at growth

temperature from 600 °C to 750 °C. The four curves are almost parallel with

each other, suggesting that, as the growth temperature decreases, additional

impurity scattering increases. Furthermore, all samples show metallic behavior,

indicating that they are all in the degenerately doped regime. Using van der

Pauw method, we investigated the carrier density and Hall mobility of BLSO films,

as in Figure. 3.5. The mobility of BLSO films decrease from 70 to 34 cm2/Vs as

the growth temperature is decreased, suggesting that there are higher densities of

defects which interrupt the movement of carrier at lower growth temperature.

Figure 3.10: (a) The resistivity-temperature curves and (b) Carrier concentration

and Hall mobility measured from BLSO/STO(001) films grown at various

temperatures.

The simultaneous reduction of carrier density and Hall mobility when grown at

lower temperature is unusual. This phenomenon can happen when an acceptor-

like trap level is formed between the conduction and valance band due to defects;

charge carrier is controlled by the different density of defects and the traps act as

strong scattering centers. The reduction in the electron mobility in the epitaxial

BLSO/STO films can be easily explained by considering the TDs as the scattering

centers for electrons [36-38], as we confirmed by AFM and TEM images. It was

47

suggested that dangling bonds in the dislocation core act as traps corresponding

to intermediate energy levels [39]. Acceptor-like trap [40, 41] is negatively

charged so it act as Coulomb scattering center and reduce the electron mobility.

Similar reductions of mobility in films compared to that of single crystals have

been observed in GaN films [42, 43]. The overall mobility vs. carrier density is

illustrated in Figure 3.11, comparing GaN and BLSO film data. Two pairs of {A,

C} and {B, D} of GaN films have same dislocation density but different carrier

density. [40] When compared among the same dislocation density films, GaN

films have μ ∝ n 1

2 tendency, as expected for the scattering by dislocations. In

our previous publication [44], we have reported a similar mobility dependence on

the carrier density of epitaxial BLSO films all grown at 750 °C, containing

presumably same dislocation densities. On the other hand, when compared

among the same carrier density films, the transverse mobility of GaN films is

affected by threading dislocation density in the form of μ ∝ 𝑁𝑑𝑖𝑠−1 in the range of

non-degenerate doped region where 𝑁𝑑𝑖𝑠 is dislocation density of film. Almost

same, but slightly smaller dependence of the mobility on the dislocation density

was observed in our BLSO films. The slightly smaller effect of dislocation

density on the mobility may come from the fact that our BLSO films have higher

carrier density in the degenerately doped regime. Our epitaxial BLSO films

grown at 750 °C with dislocation density of 4.9 × 1010 cm−2 showed very little

conductivity when the carrier density is lowered below 3.0 × 1019 cm−3, while

GaN films with similar amount of dislocation density show large mobility even

with the carrier density in the range of 1017 ~ 1018 cm−3 . We believe this

suggests that the dislocations of BLSO films trap significantly larger electron

carrier densities than those of GaN films. In single crystals, GaN [45] shows μ ~

100 and BLSO [44] has μ ~ 300 at the carrier concentration level around 8.0 ×

1019 cm−3. In thin films, however, GaN has higher mobility than BLSO at the

similar carrier concentration and dislocation density even though BLSO has

higher dielectric constant. [46] It seems that larger carrier density is required for

BLSO films than GaN films to shield the effect of dislocation scattering,

48

consistent with the fact that the dislocation traps more carriers in BLSO films than

in GaN films.

The effects of dislocations on the electrical transport property in BLSO films

are much larger than that in GaN: dislocations trap more carrier density and more

carrier density is required to shield the scattering by the dislocation. In some

way that is not understood yet, the dislocations in perovskite structure have

significant impact on the mobility. Unfortunately, we were able to find no

previous study on the effect of dislocations on transport in perovskite structure.

More detailed atomistic structure of dislocations in perovskites needs to be studied

using, for example, high-resolution TEM. On the other hand, we will need to

find a lattice matched substrate for BLSO in order to fully take advantage of the

high mobility and chemical stability for further progress of science and technology

based on perovskite BaSnO3 system.

49

Figure 3.11: Mobility of BLSO films are plotted as a function of carrier density n

at room temperature. GaN data are shown for comparison. {A, B} films and

{C, D} films have same carrier density but different dislocation density.

50

Chapter 4

High mobility SnO2 thin

film with various buffer

layers SnO2 is a multifunctional semiconductor with exotic electrical and optical

properties, excellent physical and chemical stability. In past years, the SnO2 has

been widely applied in chemical sensor, catalysis, transparent electrodes, lithium

battery, due to its direct band gap of 3.6 eV (for ZnO: 3.3 eV, GaN: 3.2 eV), high

exciton binding energy of 130 meV, and high carrier mobility of about 300 cm2/Vs

at room temperature. SnO2 is a promising host material for constructing next

generation ultraviolet light emitting diodes (UV LEDs) and photodetectors.

Furthermore, it is a natural nonpolar semiconductor with tetragonal rutile structure,

and there is no undesirable built-in electrostatic field that reduces the electron-

hole wave functions overlap and diminishes the internal quantum efficiencies of

LEDs, which have obsessed the GaN based polar semiconductor LEDs.

Many of the emerging applications for functional wide band gap

semiconductors require highly crystalline epitaxial films. Different fabrication

methods have been reported to grow the (single crystal-like) SnO2 epitaxial films.

We intensively focused on high quality epitaxial growth using pulsed laser

51

deposition method and researched the effect of buffer layer.

4.1 Motivation of SnO2 research for

transparent oxide semiconductor

TOSs are a series of metal oxides, composed of heavy metal cations (HMCs)

with an outside shell electronic configuration of (n-1)d10 ns0 (n≥4) and oxygen

anions. TOSs (Transparent Oxide Semiconductors) have wide bandgaps in which

the ns-orbitals of the heavy metal cations primarily constitute the bottom part of

the conduction band and the oxygen 2p orbitals form the top of the valence band.

Uniquely, the spatial spreading of the outside ns orbitals with a spherical

symmetry in the HMCs is much larger than that in light metal cation such as

aluminum, leading to a wider conduction band.

Transparent electronic materials with a large bandgap and little subgap

absorption are attracting a large interest. For example, transparent conducting

oxides (TCO) are used as essential materials for photovoltaic devices, flexible

electronics, and optoelectronic applications. TCOs such as indium tin oxide

(ITO), F-doped tin oxide (FTO), and Al-doped zinc oxide exhibit large

conductivities with their typical carrier densities larger than 1020 cm-3 in the

degenerately doped regime. On the other hand, there are needs for materials with

low carrier density and high mobility in the semiconducting regime as an active

material for switching. Good examples are transparent oxide semiconductors

(TOS), most prominent of which is indium gallium zinc oxide (IGZO) being

developed as the thin film transistor (TFT) materials for display industry.

FTO thin films are widely used in the photovoltaic industry as a TCO for its low

cost compared to ITO and its stability in air. [47] The SnO2-x is an n-type

semiconductor with a wide bandgap (~3.6 eV) and it has been reported to possess

52

excellent physical and chemical stability [48, 49]. In single crystals large exciton

binding energy of 130 meV (60 meV in ZnO and 25 meV in GaN) as well as high

carrier mobility of about 250 cm2 V-1 s-1 with its carrier density of 7×1015 /cm3

were reported [50]. In its epitaxial thin film form, Ta-doped SnO2 on TiO2

substrates [51] and Sb-doped SnO2 on r-plane sapphire substrates [52] were

investigated. In these studies the undoped SnO2 had a considerable carrier density

of about 1×1018 cm-3, rendering investigation of the semiconducting regime

difficult.

In order to evaluate the potential of doped SnO2 for TOS in addition to TCO,

study of the transport properties of doped tin oxide thin films in the low doping

regime, typically less than 1018 cm-3, seems important. Although many

polycrystalline doped SnO2 thin films have been studied [53, 54] in the carrier

density range of 1×1016 ~ 4×1017 cm-3, their mobility were low, which makes

important the study of epitaxial thin films in the low doping regime. The study

of epitaxial thin films can lead to understanding of the limiting mechanism for

their mobility values. In this report we studied the carrier density and mobility in

the epitaxial SnO2-x films on r-plane sapphire in the low doping regime, the

controllability of the doping level by oxygen vacancies, and the stability of the

oxygen vacancy doping. From these studies we also suggest the limiting

mechanisms for the mobility value in the low doping regime.

4.2 Structural properties of SnO2 on r-

plane Al2O3 substrate

In our research for high mobility at low carrier concentration, we decided to use

on r-plane sapphire substrate because of its interfacial configuration. In Rutile

SnO2 structure, oxygen atoms form edge-shared octahedron chains along the (001)

53

direction. If we cut SnO2 (101) plane, then the interface configuration can be

seen like Figure 4.1. You can see high similarity in oxygen octahedral

configurations between the r-plane sapphire surface and the SnO2 (101) surface.

Figure 4.1: Interface structure of (110) SnO2 and R-plane Al2O3

4.2.1 Epitaxial SnO2 film growth condition

SnO2-x thin films were grown by pulsed laser deposition (PLD) technique. A

pulsed excimer laser with a repetition rate of 10 Hz was used to deposit SnO2-x

films. The wavelength of the laser was 248 nm and the pulse duration was

nominally 25 ns. The substrates used were (1012) Al2O3 (r-plane). 30 nm thick

SnO2-x films were deposited on top of a fully oxidized insulating buffer layer of

varying thicknesses from 100 to 1000 nm to study the effect of the buffer layer

thickness. In addition, we deposited the SnO2-x thin films on top of 500 nm thick

SnO2 buffer layers while controlling the oxygen pressure from 30 to 45 mTorr.

54

The buffer layers were deposited at the oxygen pressure of 100 mTorr and they

were insulating. The growth temperature was 700 and 800 °C.

Hall-effect measurements were performed using van der Pauw method to

measure the carrier density. Lattice parameters and crystallographic quality were

determined by x-ray diffraction (XRD) analysis using PANalytical X'Pert Pro

diffractometer with a Cu Kα. Dislocation analysis was carried out using a JEM-

3000F transmission electron microscope (TEM) in a dark field mode. To check

the stability of oxygen vacancy doping, we have measured the conductance

changes of 100 nm SnO2-x layer during thermal cycles in argon, oxygen and air

ambient. The temperature was ramped up from 25 to 500 °C in an hour,

maintained at 500 °C for 5 hours, and then ramped down to 25 °C in 5 hours.

4.2.2 X-ray diffraction analysis (XRD)

Figure 4.2: (a) XRD theta-2theta pattern of SnO2 films on r-Al2O3 substrate. (b)

XRD phi scan of 1um-thick SnO2 film on r-Al2O3 substrate.

55

From the x-ray diffraction analysis, only (101) and (202) reflection peaks are

observed in Figure 4.2(a). That means, deposited films are pure tin dioxide with

Rutile structure. Figure 4.2(b) shows φ-scans of the (110) reflection for the SnO2-

x (101) film and the (0006) reflection for the Al2O3 (1012) substrate. Two SnO2

film peaks separated by 100 ° and one substrate peak in the middle have been

observed. On the surface of the r-plane sapphire, when SnO2 [010] // Al2O3

[1210], the SnO2 film peaks of (110) and (110) reflection are expected in φ-scans

at ± 50 ° relative to the Al2O3 (0006). Therefore, the epitaxial relations of SnO2

grown on r-plane sapphire are SnO2 (101) [010] // Al2O3 (1012) [1210]. This

orientation relation is consistent with the previously published experiments [55,

56]. We can obtain the d101 lattice spacing of 2.66 Å and 2.64 Å , corresponding

to the SnO2 films with thickness of 0.1 and 1 um, respectively. Since the bulk

value of d101 for SnO2 is 2.64 Å , the epitaxial strains of the SnO2 thin films seem

to get relaxed as the thickness increases. The full width at half maximum (FWHM)

of the rocking curve of the SnO2 (101) diffraction peak was 0.87 ° and 0.64 ° for

the SnO2 films with thickness of 0.1 and 1 um, respectively. The slight decrease

of FWHM value may be related to decreasing density of defects such as

dislocations and grain boundary as the films become thicker.

4.2.3 Thickness dependence of electrical transport properties

We deposited tin dioxide at various thickness. These films, below 10 nm

thickness, are insulating. In contrast, regarding other films which are above 10

nm thickness, the mobility increases very rapidly as shown in Figure 4.3. That

means, these films contain dead layer at the bottom of films which has low

mobility. This experiment presents that high density of dislocations are

concentrated at the interfacing part of the films because of large lattice mismatch.

56

Figure 4.3: Hall mobility μ, carrier concentrations n, resistivity ρ of SnO2 films on

r-plane sapphire substrate as a function of film thickness at oxygen pressure of 10

mTorr.

4.2.4 Transmittance electron microscopy (TEM)

Figure 4.4 shows dark field images of the interface between SnO2 film and the

sapphire substrate under a transmission electron microscope (TEM). The Figure

4.4 (a) and (c) are images taken in the [010] direction, while the Figure 4.4 (b) and

(d) are images taken in the [101] direction. The parallel fringes in the images

taken in the [010] direction are crystallographic shear planes (CSPs), which are

usually called anti-phase boundaries (APBs), whereas no dislocations or planar

defects are observed when the film is viewed in the [101] direction. From the

high-resolution TEM images of Figure 4.4 (c) and (d) we can observe how the

each atomic planes are aligned at the interface. The offset of atomic plane

57

positions at the interface is along the SnO2 [101] direction as shown in Figure 4.4

(c). Introduction of CSP in every four lattice spaces at the interface provides

compensation for the large lattice misfit between film and substrate [57]. Four

SnO2 lattice spacing plus one anti-phase boundary spacing (half a regular spacing)

along the [101] direction is 2.57 nm and this matches well with 5 spacing of Al2O3

(1014). These CSPs tend to disappear as the films grow thicker, as shown in

Figure 4.4 (a). The decrease of FWHM of the XRD rocking curve in thicker films

agrees well with this finding. When viewed parallel to SnO2 [101] direction, no

such APBs were observed at the interface, implying that the SnO2/Al2O3 interface

is fully coherent. The 0.474 nm lattice spacing of SnO2 (010) matches well with

0.476 nm lattice spacing of Al2O3 (1210).

58

Figure 4.4: Dark field image of SnO2/r-Al2O3 interface viewed in the SnO2[010]

direction. Dark field image of SnO2/r-Al2O3 interface viewed in the in the SnO2[-

101] direction. High-resolution TEM of SnO2/r-Al2O3 interface viewed in the

SnO2[010] direction and in the in the SnO2[-101] direction.

59

4.3 Ruthenium-doped SnO2 buffer layer

From the X-ray diffraction analysis and Hall measurement, we can estimate that

interfacing layer shows worse crystallinity and mobility than upper layer of film.

Hence, to get a high mobility SnO2 layer, it is found that a buffer layer is essential.

Prior to deposition of the conducting SnO2 layers, an insulating buffer layer was

needed, which eliminated the influence of high density of dislocations on the

interface between film and substrate. We have searched the various buffer layers

to reduce the interfacial defects and it is found that Ruthenium doped SnO2 (Ru-

doped SnO2) buffer layer demonstrates good insulating and crystalline properties.

From the X-ray diffraction analysis, phase pure SnO2 film is grown on Ru-doped

SnO2 on r-plane sapphire substrate (see Figure 4.5). Through the omega scan,

SnO2 and Ru-doped SnO2 shows good crystallinity in the inset of Figure 4.5. As

a result of using Ru-doped SnO2, we can get a mobility of almost about 80 cm2 V-

1 s-1 at carrier density of 3.7×1018 cm-3. In this experiment, we also found that

the thickness of films does not have an effect on the mobility very much even if

the film's thickness is thin. As Figure 4.6 reveals, the mobility is very sensitive

to the thickness dependence of films which do not have a buffer layer. This

proves that Ru-doped SnO2 can act as an excellent buffer layer. By using a

ruthenium-doped tin dioxide buffer layer, we deposited SnO2 at different oxygen

pressures in order to control oxygen vacancy. Figure 4.6 shows that the carrier

density is quite sensitive to oxygen pressure for growth. We analyzed the

mobility of films depending on carrier densities. These data of thin films show

that the mobility is directly proportional to the square root of the carrier density.

The similar reductions in mobility have been observed through binary compound

of GaN [45] and ternary compound of BSO films [58]. This dependence usually

has been proved by grain boundaries and dislocations. There will be small angle

grain boundaries and threading dislocations, which act as double-Schottky

barriers for the electron transport. Higher density doping induces thinner barrier

width and smears out the barriers resulting in enhanced mobility. That means,

60

our SnO2-x layer still has high density of dislocation even it has good crystallinity

and mobility. We will need to find a lattice matched substrate or a more

appropriate buffer layer for SnO2 in order to research dislocation free

characteristics of SnO2 material and fully take advantage of the high mobility for

further progress of material science and display technology.

Figure 4.5: XRD theta-2theta pattern of SnO2 and Ru-doped SnO2 films on R-

Al2O3 substrate

61

Figure 4.6: (a) Hall mobility μ, carrier concentrations n, resistivity ρ of 30 nm

SnO2 films with and without 1 um Ru-doped SnO2 buffer layer on r-plane

sapphire substrate. (b) Hall mobility μ of 30 nm SnO2 films on 1 um Ru-doped

SnO2 buffer layer on r-plane sapphire substrate plotted as a function of carrier

concentrations n.

62

4.4 Rutile TiO2 buffer layer

As a candidate of possible buffer layers, Titanium dioxide was used for this

experiment because of the following reasons; First, it is rutile structure, same as

tin dioxide. Second, since Titanium dioxide has a different lattice mismatch with

sapphire substrate from tin dioxide. Third, Insulating titanium dioxide film can

be grown easily. From the X-ray diffraction analysis, Titanium dioxide is also

grown (1,0,1) direction on r-plane sapphire substrate as shown in Figure 4.7. We

deposited titanium dioxide at10, 50, and 100 mTorr oxygen deposition pressure

to investigate the crystallinity. Through the omega scan, Titanium dioxide shows

best crystallinity at 10mT among them. Thus, In the Titanium dioxide buffer

layer case, the 10mT Titanium dioxide buffer layer seems to be a good buffer layer.

However, crystallinity of Tin dioxide does not seem to be influenced by

crystallinity of Titanium dioxide buffer layer from Table 2. So, the mobility of

tin dioxide with TiO2 buffer layer exceed the mobility of just grown on r- plane

sapphire substrate.

Table 2: Electrical properties of SnO2-x on 10 mTorr, 50 mTorr, and 100 mTorr

oxygen growth pressure of TiO2 on r-plane sapphire substrate.

63

Figure 4.7: (a) XRD theta-2theta pattern of SnO2 and TiO2 films on r-Al2O3

substrate. (b) XRD omega scan of TiO2 and SnO2 film on r-Al2O3 substrate,

respectively.

64

4.5 Fully oxidized SnO2 buffer layer at

700 °C

We decided to use fully oxidized tin dioxide as buffer layer. Through an

experiment, we found that Tin dioxide can be insulating over 50 mTorr oxygen

deposition pressure as shown in Figure 4.8. So, we thought that tin dioxide which

is grown at 100 mTorr oxygen pressure could be good insulating buffer layer to

get a high mobility. However, the result was not satisfactory. The mobility was

34.6 cm2/Vs and carrier density was 10 to the power of 19 in per cubic centimeters.

This mobility is almost half value comparing with just grown on sapphire substrate

and the carrier density is increased an order of magnitude.

Figure 4.8: XRD θ-2θ pattern of 1 um reflection of (101) SnO2-x in 10, 30, 50, and

100 mTorr oxygen pressure. (b) XRD φ scan of 500 nm SnO2 film on R-plane

Al2O3 substrate in an oxygen pressure of 100 mTorr.

65

This unsatisfying result is caused by condition of buffer layer. As theta-2theta

scan shows, phase-pure tin dioxide films were grown on r-plane sapphire substrate.

However, through the omega scan, the poor crystallinity is proved in the fully

oxidized tin dioxide film. This poor crystallinity may come from small grain size

and high density of dislocations And this buffer layer affect the upper conducting

tin dioxide layer, directly. So we studied to look for other appropriate buffer

layers.

4.6 Fully oxidized SnO2 buffer layer

at 800 °C

Lastly, for high charge mobility, the buffer layers which is grown at 100 mTorr

oxygen deposition pressure were insulating with its resistivity larger than 5 Mega

Ω cm. As a result of using fully oxidized tin dioxide buffer layer, the electron

mobility values of SnO2-x films increased from 23 to 106 when the buffer layer

thickness increased from 100 nm to 700 nm, which we attribute to reduction of

the large density of defects. However, at thicker thickness of 1 μm, the electron

mobility was found to slightly decrease, which may be related with difficulty of

good epitaxy and the surface roughness in our very thick films. This proves that

the fully oxidized tin dioxide buffer layer can act as a good buffer layer.

66

Figure 4.9: (a) Hall mobility μ, carrier concentrations n of 30 nm SnO2-x films on

r-plane Al2O3 substrate on SnO2 buffer layer of various thicknesses.

Above TEM results clearly show that the high density planar defects are

concentrated at the interface, mainly due to the large lattice mismatch between the

film and the substrate. The lattice mismatch between the SnO2 (101) film and the

(1012) sapphire substrate is about 11 % along the SnO2 [101] direction and 0.4 %

along the SnO2 [010]. Since we expect such CSPs will have a large influence on

the carrier density and the mobility of the SnO2-x films, it becomes necessary to

deposit an insulating SnO2 buffer layer to investigate the electronic properties of

the epitaxial SnO2-x films. The SnO2 buffer layers were grown in the oxygen

pressure of 100 mTorr. The buffer layers were insulating with its resistance larger

than 1011 2-x layers were grown

in the oxygen pressure of 30 mTorr and their conductivities and Hall carrier

densities were measured using the van der Pauw method. In Figure 4.9 the Hall

carrier density and mobility of each 30 nm thick SnO2-x films on the fully oxidized

SnO2 buffer layers with thicknesses ranging from 100 to 1000 nm are plotted. The

mobility values of SnO2-x films increased from 23 cm2 V-1 s-1 to 106 cm2 V-1 s-1 and

67

the carrier density increased from 9×1017 cm-3 to 3×1018 cm-3 when the buffer layer

thickness increased from 100 nm to 700 nm, which we attribute to reduction of

large density of CSPs. However, at thicker thickness of 1 um, the mobility was

found to slightly decrease, which may be related with difficulty of good epitaxy

and the surface roughness in our very thick films.

Figure 4.10: Hall mobility μ of 30 nm SnO2-x films on 500 nm fully oxidized

insulating SnO2 buffer layer on r-plane Al2O3 substrate plotted as a function of

carrier concentrations n.

In SnO2-x films carrier doping is controlled by the amount of oxygen vacancies.

The carrier density was varied from 2.1×1018 cm-3 to 1.8×1017 cm-3 by changing

the oxygen deposition pressure from 30 to 45 mTorr, as shown in Figure 4.10. In

order to minimize the effect of CSPs, we studied the doping dependence of the

mobility of SnO2-x thin films grown on top of 500 nm thick SnO2 buffer layers.

It is quite clear that the carrier density can be well controlled by the oxygen

68

pressure during growth. Although we backfilled the chamber with 600 Torr of

oxygen before cooling the sample, the high pressure oxygen does not seem to alter

the doping amount. This seems related with the oxygen stability in SnO2-x which

we will discuss further in the next paragraph. The mobility of the films decreases

as the carrier densities decrease. The mobility is very closely proportional to the

square root of the carrier density, as seen in Figure 4.10. This dependence is

usually explained by scattering from charged defects such as dislocations. As the

carrier density increases, the charged defects become screened better by more

carriers and the effective scattering cross section decreases, yielding higher

mobility. Similar reduction in mobility has been observed in a binary compound

of GaN and recently in a ternary compound of BaSnO3 films, although the main

scattering centers in those cases were threading dislocations. In our epitaxial

SnO2-x films on r-plane sapphire, the CSPs probably act as the charged scattering

centers. However, in spite of the CSPs, our SnO2-x films show the mobility of 20

cm2/Vs in the low carrier density of 1×1017 cm-3.

69

Figure 4.11: (a) XRD θ-2θ pattern of 500 nm SnO2-x on r-plane Al2O3 substrate

and (b) blow-up of (101) reflection of SnO2-x in 10, 30, 50, and 100 mTorr oxygen

pressure. (c) XRD φ scan of 500 nm SnO2 film on r-plane Al2O3 substrate in an

oxygen pressure of 100 mTorr.

The left figures 4.11 show the XRD patterns of 500 nm SnO2-x films deposited

on r-plane sapphire substrates at 800 ◦C in oxygen pressures of 10, 30, 50, and

100 mTorr. From the theta-2theta x-ray diffraction analysis, the (101) diffraction

70

peaks shift to a higher angle with the increasing oxygen pressure. The lattice

spacing for the film grown in 100 mTorr oxygen pressure, 2.641 A, is very close

to that of the bulk SnO2. This finding suggests that the oxygen deficiency in the

crystal lattice plays a dominant role in enlarging the lattice constant. Similar

results were also observed in ZnO which is a binary compound metal oxide and

also in SrRuO3 and SrTiO3 which are ternary compound metal oxide thin films.

Figure 4.12: (a) XRD omega scan of 500 nm (101) SnO2-x on r-plane Al2O3

substrate and (b) full with half maximum value of 500 nm (101) SnO2-x in 10, 30,

50, and 100 mTorr oxygen pressure, respectively.

We have also found the full width half maximum (FWHM) of the rocking curve

of the SnO2 (101) peak slightly increases from 0.5◦ to 0.57◦ when the oxygen

pressure changed from 100 mTorr to 10 mTorr during the deposition in Figure

4.12.

71

4.6 Oxygen stability of SnO2-x

When oxides are doped by oxygen vacancies, one has to wonder how stable the

doping level is in the ambient environment. To address this issue, we investigated

the thermal oxygen stability of SnO2-x film by measuring the conductance changes

of 100 nm SnO2-x during three thermal cycles under argon, oxygen and air

atmosphere in series, as shown in Figure 4.13. The conductance change at 500 °C

under Ar gas showed behavior of oxygen out-diffusion, while oxygen in-diffusion

behavior is seen in O2 atmosphere. The time scale of the diffusion process is

several hours, as long as that of BaSnO3, which was recently reported [59]. This

is not surprising since the oxygen stability of the SnO6 octahedra is not going to

be much different whether in the rutile SnO2 structure or in the perovskite BaSnO3

structure. However, it is noteworthy that the oxygen vacancies in SnO2-x do not

cause faster diffusion than in fully oxidized BaSnO3. It seems that the SnO2-x

octahedron with its oxygen vacancy, once formed, are rather stable. This is

probably why the backfilling high pressure oxygen does not change the doping

level in SnO2-x. The conductance change of SnO2-x in air at 500 °C is only about

0.5 %, demonstrating its exceptional thermal stability.

72

Figure 4.13: Conductance changes of the 100 nm SnO2-x during a thermal

annealing cycles under argon, oxygen and air atmosphere.

73

Chapter 5

Transparent thin film

transistor based on

SnO2 with amorphous

gate oxides Transparent thin-film transistors (TTFTs) based on wide bandgap oxide

semiconductors, including zinc oxide (ZnO), indium gallium zinc oxide (IGZO),

indium zinc oxide, and indium zinc tin oxide have been intensively studied in the

field of thin-film electronics during the past decades [1]–[7], and are rapidly

approaching commercialization as a replacement for amorphous Si in high-

performance electronics applications, such as active-matrix organic light-emitting

diodes, active-matrix liquid-crystal displays, and flexible displays [7]. In

particular, IGZO is one of the most promising channel materials because of its

high carrier mobility, good TTFT performance, low processing temperature, and

high transparency. However, there have been significant attempts to find a new

oxide semiconductor as an alternative to Zn-based semiconducting oxides, owing

to their oxygen unstable properties. And, Still, for TTFTs and next-generation

74

microelectronic device applications with low power consumption and higher

performance, it is necessary to develop new semiconducting oxide materials

exhibiting better transistor properties; that is, higher mobility, a smaller

subthreshold swing (ss), a higher ratio of on-to-off drain current (Ion/Ioff), and

longer stability.

5.1 Transparent thin film transistors

(TTFT)

The metal-insulator-semiconductor field effect transistor consists of a layer of

metal, a layer of insulating material and a layer of semiconducting material as

shown in Figure 5.1. The current density in the active layer can be written as:

𝐽 = 𝜎𝐸 (5.1)

where 𝜎and 𝐸 are the electrical conductivity and electric field, respectively.

From the known equation 𝜎 = nqμ, we can write the Eq. (5.1) as:

𝐽 = 𝑛q𝜇𝐸 (5.2)

The IDS can be written as:

𝐼 = 𝑛q𝜇𝐴𝐸 (5.3)

Where A (penetrated area by electron) = Wt. W and t denote the channel width

and the channel thickness, respectively.

A substitution of 𝑛 = 𝑁

𝐿𝐴 gives

𝐼𝐷𝑆 = 𝑊𝐶𝑜𝑥(𝑉𝐺𝑆 − 𝑉𝑡ℎ)𝜇𝐸 (5.4)

Considering the potential φ(x) in an arbitrary point, if we apply both VGS and VDS

to the device, the IDS can be expressed as:

75

𝐼𝐷𝑆 = 𝑊𝐶𝑜𝑥(𝑉𝐺𝑆 − 𝑉𝑡ℎ − φ(x))𝜇𝑑φ

𝑑𝑥 (5.5)

By integrating from 0 to L, under the boundary condition φ(x = 0) = 0 and

φ(x = L) = VDS, we can calculate the IDS as:

𝐼𝐷𝑆 = 𝑊

𝐿𝜇𝐶𝑜𝑥(𝑉𝐺𝑆 − 𝑉𝑡ℎ −

𝑉𝐷𝑆

2)𝑉𝐷𝑆 (5.6)

Figure 5.1. Schematics of conventional FETs

At small VDS (VDS≪VGS-Vth), IDS is proportional to VDS. We call this region the

“linear region”. As VDS increases, IDS deviates from its linear behavior. And then

finally, at high VDS, the current is saturated. We can evaluate the saturated current

condition by calculating the derivative of IDS to VDS. We can obtain the VDS needed

to saturate IDS as:

VDS = VGS - Vth (5.7)

This VDS is denoted as 𝑉𝐷𝑆𝑠𝑎𝑡.

Substitution of Eq. (5.7) into Eq. (5.6) gives

𝐼𝐷𝑆𝑠𝑎𝑡=

𝑊

2𝐿𝜇𝐶𝑜𝑥(𝑉𝐺𝑆 − 𝑉𝑡ℎ)2 (5.8)

From Eq. (5.6) and Eq. (5.8), the field effect mobility can be obtained by

calculating the transconductance (𝑔𝑚 ≡∂𝐼𝐷𝑆

∂𝑉𝐺𝑆).

76

In the linear region,

∂𝐼𝐷𝑆

∂𝑉𝐺𝑆=

𝑊

𝐿𝜇𝐶𝑜𝑥𝑉𝐷𝑆 (5.9)

In the saturation region,

(∂√𝐼𝐷𝑆

∂𝑉𝐺𝑆)2 =

𝑊

2𝐿𝜇𝐶𝑜𝑥 (5.10)

In real device measurements, the field effect mobility values evaluated from

equations (5.9) and (5.10) increases dramatically when above threshold voltage,

since the accumulated carriers screened the electric field created by defects such

as bulk traps, interface traps, vacancies, and more. So in the appropriate high level

of electric field, the field effect mobility approaches the bulk mobility of the active

layer. If a high electric field is applied, electrons are gradually accumulated near

the surface. Therefore, the field effect mobility doesn’t increase and is restricted

by the surface roughness and interface quality. For this reason the value of the

field effect mobility is generally smaller than the bulk mobility of the active layer

and depends on the surface quality. So an atomically flat active layer surface and

a high quality interface between the gate dielectric and the active layer are needed

for high performance FET devices. Furthermore, from equation 5.7 and 5.8, we

can evaluate the threshold voltage from the experiment results by using the

extrapolation of IDS-VGS and 𝐼𝐷𝑆1/2 - VGS in both linear and saturated regions,

respectively. Generally, as shown in Figure 5.2 (a), the IDS is not zero below Vth.

Below Vth, IDS decreases exponentially with decreasing gate voltage. This non-

zero current is known as the subthreshold current, originating from the diffusion

of carriers due to thermal energy. We can divide the transfer characteristic curve

into two regions known as diffusion and drift regions, as shown in Figure 5. 2 (b).

In the carrier diffusion region, the IDS is proportional to exp(qVs/kT). In the carrier

drift region, the IDS is proportional to 𝑉𝐺𝑆 − 𝑉𝑡ℎ and (𝑉𝐺𝑆 − 𝑉𝑡ℎ)2 for the linear

and saturation region, respectively. In the subthreshold region, we define an

77

important parameter, the subthreshold swing (S), which is the necessary VGS to

increase the IDS by one order.

Figure 5.2: (a) Plot of 𝐼𝐷𝑆1/2- VGS , (b) IDS - VGS output, and (c) IDS-VGS transfer

characteristic of HfO2/SnO2 field effect transistor.

S = 1/(𝑑𝑙𝑛(𝐼𝐷𝑆)

𝑑𝑉𝐺𝑆) (5.11)

This value can be the standard for evaluating device performance in low-voltage,

low-power applications. The S can be written as:

S = ln10/(𝑑𝑙𝑛(𝐼𝐷𝑆)

𝑑𝑉𝑆

𝑑𝑉𝑆

𝑑𝑉𝐺𝑆) (5.12)

Then we can obtain Eq. (5.13) by derivation from Eq. (5.12)

S =𝑘𝑇

𝑞 ln10(1 +

𝐶𝑑𝑒𝑝

𝐶𝑜𝑥) (5.13)

where k, T, and q are Boltzman’s constant, temperature, and electron charge,

respectively. So theoretically, we can estimate the minimum value of S as 60

mV/dec at room temperature.

78

Modulation capacity of dielectrics

When we choose the gate dielectric to demonstrate FET devices, we should

consider some material parameters such as the breakdown field, dielectric

constant, and modulation capacity (see Section 5.2). In this section, the author will

discuss the carrier modulation capacity using the dielectric material. In fact, the

induced carrier density in the active layer from the gate dielectric is defined by a

function of VGS, VT, and VDS. Fortunately, we can estimate the modulation capacity

by the simple relation;

Q ≈ CV (5.14)

where Q = Nq, C = eA/d, and V = Ed. This is possible since we generally consider

high VGS conditions to calculate the maximum modulation carrier density. We can

simplify equation (5.1) as

𝑛2𝐷 = 휀𝐸

𝑞 (5.15)

Where q, ε, and E are the electron charge, permittivity of the gate dielectric,

and applied electric field, respectively.

79

5.2 Investigation of HfO2 gate

dielectrics

To confirm the reliability of HfO2, the author needed to check their dielectric

properties. The author first fabricated capacitor devices as follows. An ITO line

as a contact layer was grown by the PLD method using a Si mask at growth

temperature 150 °C, then HfO2 was each deposited as a dielectric layer by the

atomic layer deposition (ALD) method at 250 °C. Finally, an ITO line as a contact

layer was grown again above the HfO2 layers.

Figure 5.3: Optical microscope image and measurement schematic of ITO/oxide

dielectric/ITO capacitor.

The cross-sectional diagram and optical microscope image of ITO/oxide

dielectric/ITO capacitor is shown in Figure 5.3. After the fabrication of

ITO/HfO2/ITO capacitors, the author investigated the dielectric constant.

80

Dielectric constant of HfO2

The capacitance (C) – Frequency (F) characteristics are presented Figure 5.4.

From this measurement, the dielectric constant of HfO2 have been calculated by

the fundamental relation,

𝐾 = 𝐶𝐴

𝑑 (5.16)

where C, d, A are the measured capacitance, thickness, and area of junction,

respectively. The calculated dielectric constants are 15.7 for HfO2. These values

are very consistent with known properties in previous articles. [60, 61] From these

material parameters, the author has demonstrated an FET device based on SnO2

with amorphous HfO2.

Figure 5.4: C-F characteristic of an ITO/HfO2/ITO capacitor are shown. Filled

circle and unfilled square markers denote the capacitance and dissipation factor,

respectively.

81

5.3 Devices utilizing HfO2 gate

dielectric

5.3.1 Device fabrication process

The TFT devices using an HfO2/SnO2 structure on r-plane sapphire substrate

are fabricated as follows as shown in Figure 5.5. The FET device using HfO2/SnO2

is fabricated in two different thickness such as 15 nm and 150 nm in the active

layer to compare the effect of interfacial defects on the output and transfer

characteristics. First the SnO2 active layer is deposited on r-plane sapphire

substrate by PLD with Silicon stencil mask at 800 °C. Then, As the source and

drain contacts, ITO is deposited at temperature of 150 °C by PLD using a SUS

mask. Then an HfO2 gate dielectric layer is deposited by ALD on the channel

layer. After the gate dielectric layer deposition, the author monitors the

conductance of the SnO2 active layer and observed that the conductance generally

increases by more than three order of magnitude. The author is able to obtain the

original conductance value of SnO2 by annealing at 400 °C in an oxygen

atmosphere. The author believes that the conductance increase in the SnO2 active

layer may be due to hydrogen absorption on the surface during the ALD process.

Finally, ITO is deposited on the device as a gate electrode with a small overlap to

the source and drain contact layer to minimize the leakage current and maximize

the possible applied bias. The optical microscope image of an HfO2/SnO2 FET

device are presented in Figure 5.6.

82

Figure 5.5: The fabrication process of HfO2/BLSO devices.

Figure 5.6: (a) optical microscope image of an Al2O3/BLSO FET device.

The channel length and width of the HfO2/15 nm SnO2 FET are 100 and 100

μm, respectively. And the channel length and width of the HfO2/150 nm SnO2

FET are 100 and 200 μm, respectively. And the thickness of the HfO2 layer is 62.5

nm of the both FETs.

83

5.3.2 Output characteristics of HfO2/SnO2 FET

The output characteristics of HfO2/SnO2 FETs are shown in Figure 5.7 (a) and

(b), where the drain-source current (IDS) is plotted against the drain-source voltage

(VDS) for various fixed gate-source voltages (VGS). As seen above, IDS is

proportional to VDS in the region where VDS is smaller than VGS-Vth in the linear

region. IDS deviates from linear behavior as VDS increases. The current is saturated

in the region of VGS-Vth < VDS.

Figure 5.7: The output characteristics, IDS-VDS curves, of (a) 15 nm and (b) 150

nm SnO2 TFTs for various gate bias VGS at room temperature. Gate voltages of

TFTs are varied from 8 V to -2 V in 1 V steps.

Of course, these behaviors are parallel to the ideal output characteristics (IDS vs.

VDS) of a MOSFET which are described by the linear region, non-linear region,

and current saturation region. The output characteristics, IDS-VDS curves, of 15 nm

and 150 nm SnO2 TFTs for various gate bias VGS at room temperature. Gate

84

voltages of TFTs are varied from 8 V to -2 V in 1 V steps. The devices clearly

show the behavior of a n-type accumulation mode FET; under positive and

negative VGS, IDS increases and decreases, because positively increasing and

negatively increasing VGS accumulates and depletes the carriers in the active layer

as commented above, respectively.

5.3.3 Transfer characteristic of HfO2/SnO2 FET

The transfer characteristic, such as Ids-Vgs, 𝜇𝐹𝐸, Ion/Ioff ratio, and subthreshold

swig (S), are shown in Figure 5.8. (a) and (b) represents the transfer characteristics

of the 15 nm and 150 nm SnO2 active layer FETs, respectively. As expected

results by consideration of results shown in Figure 5.8., sufficient current

enhancement due to carrier accumulation in an active layer could be achieved by

applied positive gate bias in both devices. However, author couldn’t see the full

depletion of active layer by applied negative gate bias in 15 nm and 150 nm SnO2

FET. The author believes that this behavior comes from both high concentration

of carrier on SnO2 active layer and electric field cancel by interface trap charges.

And the author calculates the field effect mobility of devices by using the general

relation as commented in chapter 5.1. The calculated maximum mobility of 15 nm

and 150 nm SnO2 active layer are about 16 and 73 cm2/Vs, respectively, for VDS

= 1 V.

85

Figure 5.8: The transfer characteristics, IDS-VGS and 𝜇𝐹𝐸-VGS, of (a) 15 nm and (b)

150 nm SnO2 TFTs are represented as the solid line and the symbols, respectively.

The Ion/Ioff ratio at VGS of 15 nm and 150 nm SnO2 shows 105 and 106, respectively.

One of most important parameter in FET Ion/Ioff ratio, which is defined as the

ratio of the maximum to minimum IDS, is 105 and 106 for 15 nm and 150 nm SnO2

FETs, respectively. The field effect mobility of 150 nm SnO2 device is better than

15 nm SnO2 device. Also the subthreshold swing (S), is found to be 1.4 Vdec-1

and 0.48 Vdec-1 for 15 nm and 150 nm SnO2 FETs, respectively. By comparison

of 15 nm and 150 nm SnO2 FETs we can conclude that the use of lower threading

dislocation active layer state could decrease S and the operating voltage since

lower density of interface traps. As seen, these devices show obviously higher

performance than the previous reported pure SnO2 device results. It is provided

evidence of superior material properties of SnO2 such as high mobility and

atomically flat surface with oxygen stability and shows the potential of application

of transparent high mobility TTFT device. So the development of alternative gate

dielectric with free charge traps at the interface is needed.

86

7. Summary In summary, we studied the effect of dislocations on the mobility of BLSO films

by changing the dislocation density. We have changed the dislocation density

from 4.9 × 1010 cm−2 to 8.9 × 1010 cm−2 by lowering the growth temperature

from 750 °C to 600 °C. And also, we investigated the mobility dependence of

BLSO thin film in order to reveal the mobility limit mechanism of our BSO

material.

We show an analysis similar to that used for bulk samples to compare our

calculated Hall mobility with the experimental values. When comparing with

experimental measurements, it is important to recognize that scattering

mechanisms at low carrier concentration is definitely dominant by threading

dislocation scattering limit, however, at high carrier concentration. From that

result, we assumed that the neutral impurity can be reside in our film. Therefore,

we should take additional mechanisms into account for instant due to the neutral

impurity scattering. We predict the neutral impurity concentration of 0.5 % of La

impurity concentration. Consequently it is quite well fit with experimental results.

Similar reductions of mobility in films compared to that of single crystals have

been observed in GaN films. When compared with GaN films in the mobility vs.

carrier density plot, the dislocations of BLSO films seem to trap more carrier

density and require higher carrier density for shielding. This work suggest, for

the first time, that the dislocations in perovskite materials, especially in the low

doped regime, have a large impact on the electrical transport property.

Other material but also based on Tin, SnO2, we studied the effect of buffer layer

on the mobility of SnO2. We investigate the effect of TiO2, Ru-doped SnO2, and

87

fully oxidized SnO2. From our research, TiO2 does not significant effect on SnO2

active layer and SnO2 at grown 700 C. On the other hand, Ru-doped SnO2 and

fully oxidized SnO2 act a good buffer layer. From researching buffer layers, we

found that Ru-doped SnO2 buffer layer shows good insulating properties. By

using a Ru-doped SnO2 buffer layer, we can get high mobility of almost of 80 cm2

V-1 s-1 at a carrier density of 3.7×1018 cm-3. We have controlled the carrier density

from 3.7 ×1018 cm-3 to 3.4 ×1016 cm-3 by changing the oxygen deposition pressure

from 10 to 20 mTorr. By using the fully oxidized insulating SnO2 buffer layer,

we obtained high electron mobility of 106 cm2V-1 s-1 at a carrier density of

3.0×1018 cm-3. We have controlled the carrier density from 2.1×1018 cm-3 to

9.1×1016 cm-3 by changing the oxygen pressure during deposition from 30 to 47.5

mTorr. The electron mobility of epitaxial SnO2-x thin films seems limited by the

strong scattering from the defects, very similar to the cases of GaN and BaSnO3.

Moreover, we found that SnO2-x has excellent oxygen stability in spite of oxygen

vacancies. The high electron mobility and stability of SnO2-x in the low doping

regime demonstrates its potential as a transparent oxide semiconductor.

In applicative works, the author has demonstrated the FETs based on SnO2 with

various active layer thickness. It is well-known that SnO2 transport is limited by

threading dislocations, generated by large lattice lattice mismatch between SnO2

and r-Al2O3. The author tried to find an appropriate buffer layer to demonstrate

FET devices. Demonstrated the FETs based on SnO2 with amorphous HfO2

mentioned in chapter 5. The pure epitaxial SnO2 and fully transparent thin film

transistor are the first work in the knowledge of the author. In demonstration of

amorphous gate HfO2/SnO2 FETs, the author shows the dielectric properties such

as dielectric constant and breakdown field. All devices have shown the sufficient

performance due to the stable characteristic of SnO2. Especially, the superior

performance is achieved in 150 nm active layer SnO2 FET. This performance

could be comparable with famous FET devices based on ZnO and IGZO.

Additionally, we expect that it could be fabricate poly or amorphous SnO2 active

88

layer and it may show better behavior because of large dispersion of the Sn 5s

orbital-derived conduction band.

89

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96

국문 초록

페로브스카이트 산화물은 단순하면서도 응용적인 면에서 A와 B site 에

모두 도핑이 가능하다는 유연한 특성이 학문적인 연구 분야와 산업적인

연구 분야에 방대한 가능성이 있다는 장점을 가지고 있는 물질이다. 최근

우리 연구 그룹은 La-doped BSO (BLSO) 단결정이 8.0 × 1019 cm−3의 전하

밀도에서 320 cm2/Vs 정도의 전하 이동도를 갖는다는 사실을 논문에

게재하였다. 반면에 La-doped BSO 박막은 6.0 × 1019 cm−3의 전하밀도에서

27 cm2/Vs 정도로 상대적으로 낮은 전하 이동도를 갖는 다는 사실도 확인

하였다. 이런 단결정과 박막의 큰 전하 이동도의 차이가 dislocation에 의한

영향이라는 것을 전하 밀도와 전하 이동도의 관계의 연구를 통하여

알아내었다. 이러한 방법을 통하여 박막의 dislocation 을 연구한 경우는

매우 드물다.

이 논문을 통해서 우리는 dislocation이 STO 기판 위에서 BLSO 박막이

어떠한 영향을 받는지에 대하여 자세하게 연구하였다. 우리는 이 연구에서

threading dislocation 의 연구에서 전하이동도가 최대 70 cm2/Vs 정도에서

dislocation의 밀도가 증가함에 따라서 감소한다는 사실을 확인하였다. 이런

dislocation에 의한 특성은 기존에 잘 알려진 GaN 와 비슷한 성격을 가지고

있다는 사실도 이 연구를 통하여 확인되었다.

또한 우리는 R-plane 사파이어 기판 위에 성장된 SnO2-x 의 전기적

특성에 관한 연구를 수행하였다. 이 연구는 Pulsed Laser Deposition 기술을

이용하여 진행을 이행하였다. 이 연구를 통하여 박막의 두께에 따라 박막의

전하 이동도가 매우 급격히 바뀐다는 사실을 확인하였는데, 이는 박막과

계면 사이의 두 가지 dislocation에 의한 영향이라는 사실을

전자현미경검사법 (Transmittance electron microscopy) 을 통하여 확인하였다.

따라서 우리는 다양한 완충층 (buffer layer) 에 관한 연구를 진행하였다. 이

97

연구를 통하여 SnO2-x가 완충층의 두께에 매우 민감하게 반응한다는 사실을

확인하였다. 완충층의 두께가 100 nm 에서 700 nm로 증가할 때, SnO2-x 의

전하 이동도가 23에서 106 cm2/Vs으로 증가한다는 것을 확인 하였고, 이것은

박막과 기판 계면에 존재하는 dislocation에 의한 것이라는 것을 알아내었다.

추가적으로, 우리는 500 nm의 완전 산화된 SnO2 완화층 위에 다양한 산소

분압에서의 SnO2-x가 전자 밀도를 조절 할 수 있다는 사실을 확인 하였고,

이를 통하여 전기적 특성 연구하였다. 그리고 우리는 100 nm SnO2-x박막을

이용하여 열적 안정성 실험을 하였다. 이 실험을 통하여 우리의 SnO2-

x박막의 뛰어난 열적 안정성을 가진 다는 사실을 알 수 있었다. 이런 산소의

열적 안정성과 높은 전하 이동도는 투명 산화물 반도체 연구에 중요한

가능성을 가지는 물질이라고 여겨진다. 또한 저자는 이 물질을 이용하여

전계 효과 트랜지스터(Field effect transistor)를 구현하였고, 전하 이동도,

Ion/Ioff 비율, 그리고 subthreshold swing이 각각 72.1 cm2/Vs, 6.0×106, 그리고

0.48 V/dec 이라는 사실을 확인 할 수 있었다.

키워드 : 투명 전도성 산화물, 투명 산화물 반도체, 페로브스카이트 산화물,

BaSnO3, 넓은 밴드갭, 산소 안정성, 높은 전자 이동도, SnO2, 전계 효과

트랜지스터, 산화물 계면

Student number : 2012-30102

98

List of publications

1. H. Mun, H. Yang, Jisung Park, C. Ju, and K. Char, “High electron mobility in

epitaxial SnO2−x in semiconducting regime”, APL Mater. 3, 076107 (2015)

2. U. Kim, C. Park, T. W. Ha, R. Y. Kim, H. Mun, H. M. Kim, H. J. Kim, T. H.

Kim, N. W. Kim, J. Yu, K. H. Kim, J. H. Kim, and K. Char, “Dopant-site-

dependent scattering by dislocations in epitaxial films of perovskite

semiconductor BaSnO3” APL Mater. 2, 056107 (2014)

3. B. Kim, S. Kwon, H. Mun, S. An, and W. Jhe, “Energy dissipation of

nanoconfined hydration layer: Long-range hydration on the hydrophilic solid

surface “, Scientific reports 4, 6499 (2014)

4. H. Mun, U. Kim, H. M. Kim, C. Park, T. H. Kim, H. J. Kim, K. H. Kim, and

K. Char, “Large effects of dislocations on high mobility of epitaxial perovskite

Ba0.96La0.04SnO3 films”, Appl. Phys. Lett. 102, 252105 (2013)

5. H. J. Kim, U. Kim, T. H. Kim, J. Kim, H. M. Kim, B.-G. Jeon, W.-J. Lee, H.

Mun, K. T. Hong, J. Yu, K. Char, K. H. Kim, “Physical properties of transparent

perovskite oxides (Ba,La)SnO3 with high electrical mobility at room temperature",

Phys. Rev. B 86, 165205 (2012).

6. H. J. Kim, U. Kim, H. M. Kim, T. H. Kim, H. Mun, B.G. Jeon, K. T. Hong,

W.-J. Lee, C. Ju, K. H. Kim, K. Char, “High mobility in a stable transparent

perovskite oxide", Appl. Phys. Express 5, 061102 (2012).

99

Presentations

1. H. Mun, H. Yang, J. Park, C. Ju, and K. Char, “High electron mobility in

epitaxial SnO2-x in semiconducting regime”, Workshop on Oxide Electronics 20,

Nanjing, China, October 2016. (Poster)

2. H. Mun, H. Yang, J. Park, C. Ju, and K. Char, “High electron mobility in

epitaxial SnO2-x in semiconducting regime”, Materials Research Society Meeting,

Phoenix, AZ, April 2016. (Poster). (Oral)

3. H. Mun, H. Yang and K. Char, “Effect of a Ru-doped SnO2-x buffer layer on

thin-film transistors based on SnO2-x channel layer”, American Physical society

Meeting, San Antonio, TX, March 2015. (Oral)

4. H. Mun, J. Park, C. Ju, H. M. Kim, U. Kim, and K. Char, “Electronic transport

properties of epitaxial SnO2 on r-plane sapphire substrate by PLD”, American

Physical society Meeting, Denver, CO, March 2014. (Oral)

100

감사의 글

이 글을 쓰는 시점까지를 뒤돌아 보니, 참 감사를 드려야 할 분들이 많다는

생각이 듭니다. 저에게 중요한 영향을 주신 분들이 많지만 학위 과정 중 가장

큰 영향을 주신 분이 저의 지도교수님인 차국린 교수님이 아닌가 생각 됩니다.

처음 차국린 교수님 연구실로 대학원을 진학을 하게 되었을 때, 저는 연구의

목적도 그리고 목표도 불분명한 상황에 있었습니다. 그런 이유로 석사 과정

동안 많은 어려움과 낙담을 경험했던 것 같습니다. 그 시기에 포기 하지 않도록

이끌어 주신 배려가 저로 하여금 무사히 박사 과정을 마치게 할 수 있었던

정신적인 부분이 아닌가 생각을 하고 있습니다. 또한, 학문적으로 박사과정

동안 BaSnO3 연구를 처음 시작 하였을 때, 산화물 반도체에 관한 지식이 부족한

저에게 BSO3 물질의 가능성과 중요성이 무엇인지에 대하여 고민하게 해

주셨습니다. 이 연구를 기반으로 주석 산화물인 SnO2 산화물 반도체에 대하여

대한 전반적이면서도 깊은 이해를 하도록 지도해주신 부분은 앞으로의 저의

연구에 밑거름이 되리라 생각하고 더욱 열심히 노력 하도록 하겠습니다. 또한,

바쁘신 시간을 내어 박사 학위 논문 심사를 해 주시며, 많은 조언을 해 주신

노태원 교수님, 이탁희 교수님, 박제근 교수님, 그리고 박준석 박사님께도 깊은

감사들 드립니다.

오랜 연구실 생활 동안, 저에게 격려와 도움을 준 선배, 친구, 그리고

후배들에게도 감사의 마음을 전합니다. 처음 연구실에 들어와서 함께 시간을

보내며 저에게 교수님 다음으로 많은 부분을 가르쳐 주신 제욱이형, 연구실에서

많은 추억을 쌓게 해 준 선배이자 오랜 친구인 훈민이와 미안한 마음이 많이

드는 찬종이에게 고맙다는 말을 하고 싶습니다. 그리고 박사 과정 동안 초창기

BSO3 연구를 많이 하면서 연구적인 도움을 많이 받고 친해진 우성이와 항상

옆자리에서 많은 이야기를 나누어준 혁우 그리고 연구하고 졸업을 준비하면서

도움을 받은 철권이와 지성이에게 항상 고마움을 느끼고 있습니다. 그리고

선배로서 부족한 부분이 많지만 항상 신경을 써주고 도와 준 후배들이 있어서

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무사히 연구를 마칠 수 있었다고 생각합니다. 지금 후배들의 가장 큰 선배이자

항상 웃어주고 저와 이야기를 많이 나누어 준 주연이, 항상 묵묵히 일을 하고

다른 후배들에게 좋은 선배로서 본보기인 배울 점이 많다고 생각이 들었던 영모,

짧은 시간 동안 밖에 함께 연구를 하지는 못했지만 SnO2를 함께 연구하면서

정이 많이 들었던 현석이와 상원이, 연약해 보이지만 속으로는 강하고 생각이

깊은 유정이, 항상 친동생처럼 잘 챙겨주고 신경 써주며 장난기가 많긴 하지만

미워할 수 없는 하훈이, 새로운 연구실에 들어와서 적응하느라 힘든 와중에도

항상 꼼꼼하게 일하며 배려를 많이 해 준 예주, 그리고 가장 최근에 연구실에

들어와서 열심히 하는 모습이 보기 좋은 성윤이에게 고마움을 전합니다.

그리고, 단결정 BSO3 연구를 통하여 박막 연구에 큰 도움을 주신 김기훈

교수님과 많은 조언으로 도움을 주신 김기훈 교수님 연구실의 형준이형,

태훈이형, 웅제씨, 그리고 FET 제작을 위하여 ALD 증착 연구에 도움을 준

홍승훈 교수님 연구실의 동국씨에게도 고마움을 전합니다.

그리고 고마운 친구들로는, 다른 연구실 친구이지만, 대학원 생활 동안에

항상 함께 해서 정말 즐거웠고 앞으로 오랜 우정의 이어 갈 수 있는 친구가 되어

준 완희, 소영, 그리고 봉수형에게 말로 다 할 수 없는 고마움을 전합니다. 요즘

다들 바빠서 잘 보지는 못하지만 가끔 만나서 이야기 하는 것 만으로도 큰

즐거움이 되었던 정민, 순기, 철민, 희준, 기영이형, 성민이형, 은수, 경민이, 동렬,

호성, 형주, 철웅, 길환이형, 성훈이형, 경호형, 그리고 대학원 친구이면서

운동도 함께하고 항상 도움이 되어준 건욱이형, 민준, 규환, 정훈, 재현, 정구,

현학, 태영이에게도 고맙다는 말을 하고 싶습니다. 그리고 저와 가장 오래

우정을 나누어 온 고향 친구 현이, 대우, 환중이, 상엽이, 미연이, 미현이, 경숙이,

명숙이, 지민이, 현정이, 순옥이, 선화, 선영이에게도 모두 표현 할 수 없지만

항상 고마워 하는 마음을 전해봅니다. 그리고 저를 학창 시절부터 아껴주시고

지금까지도 항상 격려해 주시는 고등학교 은사님이신 정하동 선생님께도 감사

드립니다.

모든 분들께 제가 글로 적는 것 보다 항상 더 큰 마음으로 고마워하고, 더 많은

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분들께 감사하게 생각하고 있습니다.

오랜 박사 과정 동안 곁에서 묵묵히 힘이 되어준 사랑하는 가족들에게는

어떻게 고마움을 표현 해야 할 지 모르겠습니다. 힘든 박사과정을 하면서 알

게 된 것 중에 하나가 가족의 사랑이었던 것 같습니다. 항상 그 자리를 변함없이

지켜주시고, 무슨 일을 하든 항상 저를 믿어 주시며 힘이 되는 말씀을 많이 해

주시는 아버지, 어머니께 사랑한다는 말을 드리고 싶습니다. 그리고 항상

따뜻하게 저를 대해주는 큰누나와 생각이 깊어서 항상 다른 가족을 걱정해 주는

작은누나에게도 한번도 말해보지 못한 사랑한다고 말하고 싶습니다. 그리고

저를 어릴 적부터 친동생처럼 아껴주신 큰 매형과 항상 제 일에 관심을 가져

주시고 의논해 주시는 작은 매형께도 감사한다는 말을 전합니다.

석사와 박사 과정 동안 저를 곁에서 가장 많이 아껴주는 사랑하는 부인

진희에게도 사랑한다는 이야기를 남기고 싶습니다. 언제나 웃는 얼굴로

남편의 투정을 받아 주고 배려해 주는 좋은 부인이 함께 있어서 힘든 박사 학위

기간도 즐겁게 마무리 할 수 있었던 것 같습니다.

5년 동안의 박사 과정은 제 인생에서 매우 의미 있는 시간 이었습니다. 많은

좋은 분들을 알게 되었고, 많은 것들을 배울 수 있는 뜻 깊은 시간이었습니다.

이제 박사 과정을 마치고 사회에 나가 새로운 미래를 위해서 더욱더

노력하겠습니다. 저를 아껴 주시는 모든 분들께 이 기회를 빌어 다시 한 번 깊은

감사를 드립니다.

2016년 2월

문 효 식

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