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Room Temperature Deposited Amorphous Semiconducting Oxides and Perovskite Complex Oxide Heterostructures

Permanent Link: http://ufdc.ufl.edu/UFE0043649/00001

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Title: Room Temperature Deposited Amorphous Semiconducting Oxides and Perovskite Complex Oxide Heterostructures
Physical Description: 1 online resource (140 p.)
Language: english
Creator: Kim, Seonhoo
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2011

Subjects

Subjects / Keywords: amorphous-semiconducting-oxide -- complex-oxide-heterostructure -- oxynitride -- perovskite -- superlattice -- thin-film -- transparent-conducting-oxide
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: There is significant interest in the growth and properties of transparent conducting oxides (TCOs), specifically p-type TCOs. While n-type transparent semiconductors have been produced at low temperatures and on glass or plastic substrates, very few p-type TCOs have been reported - most of them are crystalline TCOs. This dissertation includes research of amorphous p-type transparent semiconducting oxide, zinc cobalt oxide (ZCO). Room temperature deposited ZCO thin films are investigated as p-type TOS which is confirmed by both positive Seebeck and Hall coefficients. Room temperature fabricated p-n heterojunction additionally confirms p-type conduction of ZCO films by using n-type indium gallium zinc oxide. In addition, nitrogen incorporated n-type oxide thin films such as indium aluminum zinc oxynitirde and zinc aluminum oxynitride are prepared at room temperature and investigated with structural, electrical, and optical properties. With the search for new p-type TCOs, other functional materials such as complex oxide thin films and heterostructures are studied to introduce interesting physical properties at the interface. Advanced growth techniques make it possible to prepare epitaxial heterostructures such as perovskite-related superlattices. This dissertation includes growth and characterization of LaVO3/SrTiO3 superlattices. In the artificial structures, two-dimensional electron gases (2DEGs) are observed at atomically controlled interfaces between two complex oxides.
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Statement of Responsibility: by Seonhoo Kim.
Thesis: Thesis (Ph.D.)--University of Florida, 2011.
Local: Adviser: Norton, David P.
Electronic Access: RESTRICTED TO UF STUDENTS, STAFF, FACULTY, AND ON-CAMPUS USE UNTIL 2013-12-31

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2011
System ID: UFE0043649:00001

Permanent Link: http://ufdc.ufl.edu/UFE0043649/00001

Material Information

Title: Room Temperature Deposited Amorphous Semiconducting Oxides and Perovskite Complex Oxide Heterostructures
Physical Description: 1 online resource (140 p.)
Language: english
Creator: Kim, Seonhoo
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2011

Subjects

Subjects / Keywords: amorphous-semiconducting-oxide -- complex-oxide-heterostructure -- oxynitride -- perovskite -- superlattice -- thin-film -- transparent-conducting-oxide
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: There is significant interest in the growth and properties of transparent conducting oxides (TCOs), specifically p-type TCOs. While n-type transparent semiconductors have been produced at low temperatures and on glass or plastic substrates, very few p-type TCOs have been reported - most of them are crystalline TCOs. This dissertation includes research of amorphous p-type transparent semiconducting oxide, zinc cobalt oxide (ZCO). Room temperature deposited ZCO thin films are investigated as p-type TOS which is confirmed by both positive Seebeck and Hall coefficients. Room temperature fabricated p-n heterojunction additionally confirms p-type conduction of ZCO films by using n-type indium gallium zinc oxide. In addition, nitrogen incorporated n-type oxide thin films such as indium aluminum zinc oxynitirde and zinc aluminum oxynitride are prepared at room temperature and investigated with structural, electrical, and optical properties. With the search for new p-type TCOs, other functional materials such as complex oxide thin films and heterostructures are studied to introduce interesting physical properties at the interface. Advanced growth techniques make it possible to prepare epitaxial heterostructures such as perovskite-related superlattices. This dissertation includes growth and characterization of LaVO3/SrTiO3 superlattices. In the artificial structures, two-dimensional electron gases (2DEGs) are observed at atomically controlled interfaces between two complex oxides.
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Statement of Responsibility: by Seonhoo Kim.
Thesis: Thesis (Ph.D.)--University of Florida, 2011.
Local: Adviser: Norton, David P.
Electronic Access: RESTRICTED TO UF STUDENTS, STAFF, FACULTY, AND ON-CAMPUS USE UNTIL 2013-12-31

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2011
System ID: UFE0043649:00001


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1 ROOM TEMPERATURE DEPOSITED AMORPHOUS SEMICONDUCTING OXIDES AND PEROVSKITE COMPLEX OXIDE HETEROSTRUCTURES By SEONHOO KIM A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT O F THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2011

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2 2011 SeonHoo Kim

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3 To my parents for their sincere support

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4 ACKNOWLEDGMENTS I thank my parents and sister for their constant encouragement and belief to a ccomplish this dissertation. I would like to thank my advisor, Dr. David Norton, for guidance and support to conduct scientific research I thank my supervisory committee members, Dr. Stephen Pearton, Dr. Ravi Singh, Dr. Franky So, and Dr. Fen Ren, who hav e guided me through the doctoral research. I also thank my colleagues in the department of Materials Science and Engineering and other departments, whom I have worked together.

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5 TABLE OF CONTENTS page ACKNOWLEDGMENTS ................................ ................................ ................................ .. 4 LIST OF TABLES ................................ ................................ ................................ ............ 8 LIST OF FIGURES ................................ ................................ ................................ .......... 9 ABSTRACT ................................ ................................ ................................ ................... 12 CHAPTER 1 INTRODUCTION ................................ ................................ ................................ .... 14 2 BACKGROUND AND MOTIVATION ................................ ................................ ...... 16 2.1 Amorphous Semiconducting Oxide s ................................ ................................ 16 2.2 Transparent Conducting Oxides ................................ ................................ ....... 17 2.3 Crystalline P type Oxides ................................ ................................ .................. 18 2.3.1 Binary P type Oxides ................................ ................................ ............... 18 2.3.1.1 Cu 2 O ................................ ................................ .............................. 18 2.3.1.2 NiO ................................ ................................ ................................ 19 2.3.2 Delafossite P type Oxides ................................ ................................ ....... 19 2.3.2.1 CuAlO 2 ................................ ................................ ........................... 20 2.3.2.2 CuGaO 2 ................................ ................................ ......................... 20 2.3.3 Spinel P type Oxides ................................ ................................ ............... 21 2.3.3.1 ZnCo 2 O 4 ................................ ................................ ......................... 21 2.3.3.2 ZnRh 2 O 4 ................................ ................................ ......................... 22 2.3.3.3 ZnIr 2 O 4 ................................ ................................ ........................... 22 2.4 Amorphous P type Oxide ................................ ................................ .................. 22 2.5 Ligand Field Splitting ................................ ................................ ......................... 23 2.6 Mott Insulator ................................ ................................ ................................ .... 24 2.6.1 Electron Correlation ................................ ................................ ................. 25 2.6.2 Lanthanum Vanadate, LaVO 3 ................................ ................................ .. 27 2.7 Band Insulator ................................ ................................ ................................ ... 28 2.7.1 d 0 Insulator ................................ ................................ .............................. 28 2.7.2 Other Closed shell In sulators ................................ ................................ .. 29 2.7.3 Strontium Titanate, SrTiO 3 ................................ ................................ ...... 29 2.8 Metal Insulator Transition ................................ ................................ .................. 30 2.8.1 Doping induced Metal Insulator Transition ................................ .............. 31 2.8.2 Bandwidth controlled Metal Insulator Transition ................................ ...... 31 2. 9 Examples Of Oxides With MI Transition Behavior ................................ ............ 32 2.9.1 Ternary Transition Metal Oxides ................................ ............................. 32 2.9.2 Binary Oxides ................................ ................................ .......................... 33 2.10 Complex Oxide Thin Films and Heterostructures ................................ ........... 34 2.10.1 Two dimensional Electron Gas ................................ .............................. 35

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6 2.10.2 Polar Interfaces ................................ ................................ ..................... 36 2.10.3 Perovskite Related Superlattices ................................ .......................... 37 3 EXPERIMENTAL DETAILS ................................ ................................ .................... 43 3.1 Thin Film Fabrication ................................ ................................ ........................ 43 3.1.1 Substrate Preparation ................................ ................................ .............. 43 3.1.2 Target Prepar ation ................................ ................................ ................... 44 3.1.2.1 ZnCo 2 O 4 target ................................ ................................ ............... 44 3.1.2.2 LaVO and SrTiO 3 targets ................................ ........................... 45 3.1.3 Pulsed Laser Deposition ................................ ................................ .......... 45 3.1.4 Sputter Deposition ................................ ................................ ................... 47 3.2 Structural and Compositional Characterization ................................ ................. 49 3.2.1 X Ray Diffraction ................................ ................................ ..................... 49 3.2.2 Auger Electron Spectroscopy ................................ ................................ .. 50 3.2.3 X ray Photoelectron Spectroscopy ................................ .......................... 51 3.2.4 Scanning Probe Microscopy ................................ ................................ .... 52 3.3 Electronic Characterization ................................ ................................ ............... 52 3.3.1 Seebeck ................................ ................................ ................................ .. 52 3.3.2 Hall Effect ................................ ................................ ................................ 53 3.3.3 Physical Property Measurement System ................................ ................. 54 3.4 Optical Characterization ................................ ................................ .................... 55 4 STRUCTURAL, ELECTRICAL, AND OPTICAL PROPERTIES OF N TYPE INDIUM ALUMINUM ZINC OXYNITRIDE THIN F ILMS ................................ .......... 59 4.1 Scientific B ackground ................................ ................................ ....................... 59 4.2 Experimental Methods ................................ ................................ ...................... 60 4.3 Results and Discussion ................................ ................................ ..................... 61 4.3.1 Chemical Composition and Surface Morphology ................................ ..... 61 4.3.2 Electrical and Optical Properties ................................ ............................. 62 4.4 Summary ................................ ................................ ................................ .......... 65 5 ELECTRICAL AND OPTICAL PROPERTIES OF N TYPE ZINC ALUMINUM OXYNITRIDE THIN FILMS ................................ ................................ ..................... 77 5.1 Scientific B ackground ................................ ................................ ....................... 77 5.2 Experimental Methods ................................ ................................ ...................... 78 5.3 Results and Discussion ................................ ................................ ..................... 78 5.4 Summary ................................ ................................ ................................ .......... 81 6 STRUCTURAL, ELECTRICAL, AND OPTICAL PROPERTIES OF P TYPE ZINC COBALT OXIDE THIN FILMS ................................ ................................ ....... 90 6.1 Scientific B ackground ................................ ................................ ....................... 90 6.2 Experimental Details ................................ ................................ ......................... 92 6.3 Results and Discussion ................................ ................................ ..................... 93 6.3.1 P type Conduction ................................ ................................ ................... 93

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7 6.3.2 Structural Properties ................................ ................................ ................ 94 6.3.3 Optic al Properties ................................ ................................ .................... 94 6.3.4 Transport Properties ................................ ................................ ................ 94 6.3.5 Oxide Heterojunction ................................ ................................ ............... 95 6.4 Summary ................................ ................................ ................................ .......... 96 7 METAL INSULATOR TRANSITION AT THE INTERFACE OF LaVO 3 /SrTiO 3 SUPERLATTICES GROWN ON TIO 2 TERMINATED SrTiO 3 ............................... 108 7.1 Scientific B ackground ................................ ................................ ..................... 108 7.2 Experimental Details ................................ ................................ ....................... 109 7.3 Results and Discussion ................................ ................................ ................... 110 7.3.1 Structural Characterization ................................ ................................ .... 110 7.3.2 Transport Property ................................ ................................ ................. 112 7.4 Summary ................................ ................................ ................................ ........ 113 8 CONCLUSIONS ................................ ................................ ................................ ... 121 APPENDIX: CRYSTAL SYSTEMS ................................ ................................ .............. 124 LIST OF REFERENCES ................................ ................................ ............................. 128 BIOGRAPHICAL SKETCH ................................ ................................ .......................... 140

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8 LIST OF TABLES Table page 5 1 Atomic concentration of nitro gen (N), oxygen (O), zinc (Zn), and aluminum (A l) ................................ ................................ ................................ ...................... 89 A 1 Atomistic parameters of CuAlO 2 ................................ ................................ ...... 124 A 2 Atomistic parameters of ZnCo 2 O 4 ................................ ................................ .... 125 A 3 Atomistic parameters of LaVO 3 ................................ ................................ ....... 126 A 4 Atomistic parameters of SrTiO 3 ................................ ................................ ....... 127

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9 LIST OF FIGURES Figure page 2 1 The delafossite structure of CuAlO 2 ................................ ................................ ... 39 2 2 The spinel structure of ZnCo 2 O 4 ................................ ................................ ......... 40 2 3 The orthorhombic perovskite structure of LaVO 3 ................................ ................ 41 2 4 The cubic perovskite structure of SrTiO 3 ................................ ............................ 42 3 1 Unit cell step surface of etched and annealed SrTiO 3 substrates obtained by AFM measurement ................................ ................................ ............................. 57 3 2 Schematic of a Van der Pauw configuration to determine the Hall vol tage (V H ) by Hall effect measurements. ................................ ................................ ...... 58 4 1 Indium, zinc, and aluminum atomic concentration as a function of Al 2 O 3 power of the InAlZnON film ................................ ................................ ................. 67 4 2 Differential Auger pattern of the InAlZnON film ................................ ................... 68 4 3 X ray photoelectron spectroscopic analysis of N(1s) of the InAlZnON film ......... 69 4 4 AFM image of InAlZnON film with the root mean square (RMS) value of 1.79 nm ................................ ................................ ................................ ...................... 70 4 5 Resistivity, carrier density, and mobility of InAlZnON films in the Al 2 O 3 power range of 150 W to 350 W ................................ ................................ .................... 71 4 6 Resistivity, carrier density, and mobility of InAlZnON films in the InZnO power range of 100 W to 200 W ................................ ................................ .................... 72 4 7 Resistivity, carrier density, and mobility of InAlZnON films as a function of nitrogen ratio in the Ar N 2 plasma ................................ ................................ ....... 73 4 8 Temperature dependence of the electrical conductivity for the InAlZnON film ... 74 4 9 Temperature dependence of the carrier density and the electron mobility for the InAlZnON film ................................ ................................ ............................... 75 4 10 Optical absorption spectrum of the InAlZnON. The optical bandgap energy (E g ) is determined to be ~3.55 eV. ................................ ................................ ...... 76 5 1 Resistivity, carrier density, and mobility as a function of the ZnO sputtering power of nitrogen incorporated Zn Al O films. ................................ .................... 82

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10 5 2 Optical absorption spectra of various ZnO sputtering powers in the range of 150 to 300 W. ................................ ................................ ................................ ..... 83 5 3 Resistivity, carrier density, and mobility as a function of the Al 2 O 3 sputtering power of nitrogen incorporated Zn Al O films. ................................ .................... 84 5 4 Optical absorption spectra of vario us Al 2 O 3 sputtering powers in the range of 150 to 300 W. ................................ ................................ ................................ ..... 85 5 5 Resistivity, carrier density, and mobility as a function of the N 2 volume ratio in the mixture of Ar N 2 ................................ ................................ ........................... 86 5 6 Optical absorption spectra of various N 2 volume ratios in the range of 0 to 5%. ................................ ................................ ................................ ..................... 87 6 1 Thermoelectric power measurement of the Zn Co O film depo sited in 50 mTorr ................................ ................................ ................................ .................. 97 6 2 Temperature dependence of the electrical cond uctivity of the Zn Co O film ...... 98 6 3 FE SEM surface image ( 85,000 magnification) of the Zn Co O film deposited in 50 mTorr. ................................ ................................ ........................ 99 6 4 FE SEM surface image (100,000 magnification) of the Zn Co O film deposited in 80 mTorr. ................................ ................................ ...................... 100 6 5 FE SEM surface image (100,000 magnification) of the Zn Co O film deposited in 100 mTorr. ................................ ................................ .................... 101 6 6 Optical absorption spectrum of the Zn Co O film depo sited i n 100 mTorr ........ 102 6 7 Electrical resistivity as a function of oxygen gas pressure in the range of 50 to 110 mTorr. ................................ ................................ ................................ .... 103 6 8 Carrier density as a function of oxygen gas pressure in the range of 50 to 110 mTorr. ................................ ................................ ................................ ........ 104 6 9 Hall mobility as a function of oxygen gas pressure. ................................ .......... 105 6 10 Current density applied voltage (J V) curve of amorphous oxide p n junction using a p type Zn Co O film and a n type InGaZnO film. ................................ 106 6 11 Schematic illustration of the p n junction structure. ................................ .......... 107 7 1 3 /SrTiO 3 superlattice grown at 600C in 510 5 Torr of oxygen on a (100) SrTiO 3 substrate. ................................ ..................... 114 7 2 HRTEM image (500,000 magnification) of [(LaVO 3 ) 10 /(SrTiO 3 ) 10 ] 5 superlattice ................................ ................................ ................................ ....... 115

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11 7 3 HRTEM image (800,000 magnification) of [(LaVO 3 ) 10 /(SrTiO 3 ) 10 ] 5 superl attice ................................ ................................ ................................ ....... 116 7 4 Fast Fourier Transform (FFT) of the images of the (a) LaVO 3 layer, (b) SrTiO 3 layer, and (c) SrTiO 3 substrate. ................................ ................................ ........ 117 7 5 AFM Image of the [(LaVO3)10/(SrTiO3)10]5 superlattice grown on TiO 2 terminated SrTiO 3 substrate. ................................ ................................ ............ 118 7 6 AFM Image of the [(LaVO 3 ) 10 /(SrTiO 3 ) 10 ] 5 superlattice grown on non etched SrTiO 3 substrate. ................................ ................................ .............................. 119 7 7 Resistance of the [(LaVO 3 ) 5 /(SrTiO 3 ) 5 ] 5 superlattice as a function of temperature ................................ ................................ ................................ ...... 120

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12 Abstract of Dissertation Presented to the Graduate S chool of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy ROOM TEMPERATURE DEPOSITED AMORPHOUS SEMICONDUCTING OXIDES AND PEROVSKITE COMPLEX OXIDE HETEROSTRUCTURES By SeonHoo Kim D ecember 2011 Chair: David Norton Major: Materials Science and Engineering There is significant interest in the growth and properties of transparent conducting oxides (TCOs) specifically p type T C Os. While n type transparent semiconductors have been prod uced at low temperatures and on glass or plastic sub strates, very few p type TCO s have been reported m ost of them are crystalline TCO s. This dissertation includes research of amorphous p type transparent semiconducting oxide, zinc cobalt oxide (ZCO). Roo m temperature deposited ZCO thin films are investigated as p type TOS which is confirmed by both positive Seebeck and Hall coefficients. Room temperature fabricated p n heterojunction additionally confirms p type conduction of ZCO films by using n type ind ium gallium zinc oxide. In addition, nitrogen incorporated n type oxide thin films such as indium aluminum zinc oxynitirde and zinc aluminum oxynitride are prepared at room temperature and investigated with structural, electrical, and optical properties. W ith the search for new p type T CO s, other functional materials such as complex oxide thin films and heterostructures are studied to introduce interesting physical properties at the interface. Advanced growth techniques make it possible to prepare epitaxial heterostructures such as perovskite related superlattices. This dissertation

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13 includes growth and characterization of LaVO 3 /SrTiO 3 superlattices. In the artificial structures, two dimensional electron gases (2DEGs) are observed at atomically controlled int erfaces between two complex oxides.

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14 CHAPTER 1 INTRODUCTION group in the department of Materials Science and Engineering at the University of Florida. My motivation lies in int erest in the functional oxide materials for electronic applications since they show interesting physical characteristics such as superconducting, metallic, semiconducting, insulating, magnetic and optical. My doctoral research primarily focuses on the ele ctronic oxide thin film processing and properties. Specifically, it is attractive to find new transparent conductive oxides (TCOs) showing p type conductivity due to the limited number of p type transparent oxide semiconductors. In contrast, several n typ e TCOs can be produced rather cheap ly on low cost substrates such as glass and plastics at low temperature. However, the applications are limited to the n type transparent electrodes with controllable conductivity. Thus, fabrication of amorphous p type thi n films at or near room temperature is a fascinating goal for further transparent, flexible, and large area electronic device applications. Note that the p n junction is a key structure in most of semiconductor based devices. Oxide thin films ar e prepared either by sputter deposition or by pulsed laser deposition using a KrF excimer laser. The structural, electrical, and optical properties of oxide thin films and p n junctions are characterized. In Chapter 4, n type nitrogen incorporated indium aluminum zin c oxide thin films are synthesized and investigated at room temperature with compositional, structural, electrical, and optical measurements. In contrast to difficulty in controlling carrier concentration of indium zinc oxide and indium gallium zinc oxide thin films, the films show smooth change in electrical resistivity and carrier density. In Chapter 5, the effect of nitrogen incorporation

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15 on electrical and optical properties of n type zinc aluminum oxynitride thin films are investigated at room temperatu re. Nitrogen ratio during deposition has an influence on compositional Zn/Al ratio of oxynitride thin films, resulting in varying electrical and optical properties. Chapter 4 and Chapter 5 include study of n type oxide thin films deposited via sputtering a t room temperature. In Chapter 6, room temperature deposited zinc cobalt oxide thin films are investigated which shows p type conduction confirmed by positive Seebeck coefficients and Hall coefficients. Oxygen background gas pressure dependence of electric al properties is observed with field emission scanning electron microscope and Hall effect electronic measurements. For electronic device applications, a rectifying p n junction is fabricated by using p type zinc cobalt oxide layer and n type indium gall ium zinc oxide layer. Chapter 7 i n the dissertation is focusing on functional complex oxide thin films and heterostructures. Advanced deposition techniques enable the fabrication of ox ide layer systems with control of the individual layers and engineering of the interface at the nanometer scale. Especially, exciting new interface phenomena at perovskite related oxide interfaces such as heterointerfaces between LaAlO 3 and SrTiO 3 explor es research of atomically controlled superlattices including interfacial p olar discontinuity doping caused by elec tronic reconstruction. In Chapter 7 LaVO 3 /SrTiO 3 heterostructures are synthesized and studied for transport and structural properties with particular interest in temperature induced metal insulator transition and as ymmetric heterointerfaces.

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16 CHAPTER 2 BACKGROUND AND MOTIV ATION 2.1 Amorphous Semiconducting Oxide s Amorphous semiconductors have been attractive due to advantages in processing temperature and uniformity of device characteristics. The discovery of hydro genated amorphous silicon (a Si:H) by Spear and Lecomber [1] opened large area electronic technologies using thin film transistor (TFT) devices for technologies such as flat panel displays and flexible electronics. However, the properties of amorphous silicon limits its application. First, the carrier mob ility of amorphous silicon is low (less than 1 cm 2 /V s for electrons) since carrier transport in a Si:H is interrupted by multiple trapping in localized states generated by structural defects. This low mobility prevents the use of amorphous silicon materia ls in high speed electronic application. Second, amorphous silicon is not transparent in the range of the visible light, which limits its use in transparent device applications [2] These properties limit the applicability of amorphous silicon to high speed and transparent electronic devices. Amorphous oxide semiconductors o vercome many of the disadvantages of hydrogenated amorphous silicon for the following reasons. (i) Large magnitude of the carrier mobility observed in amorphous oxides can increase the channel mobility of TFTs, which can lead to high speed electronics with faster device oper ation. (ii) Amorphous oxide semiconductors are transparent in the range of visible light due to wide band gap ener gy, which allows development of transparent electronics. Several wide bandgap crystalline non oxide transparent semiconductors such as GaN [3 5] and SiC [6, 7] have been investigated. In addition, polycrystalline ZnO opened numerous applications in areas such as piezoelectric transducers, varistors, phosphors, and transparent conducting films [8] Amorphous

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17 semiconducting oxides have advanta ges in fabrication of field effect transistors since amorphous oxide thin films can be obtained at room temperature with glass or flexible plastic substrates applicable to polymer based devices. To date, several transparent semiconducting oxides have devel oped such as SnO 2 [9 11] In 2 O 3 [12, 13] Ga 2 O 3 [14, 15] and ZnO [8, 16 19] H. Hosono rep orted amorphous indium gallium zinc oxide which showed high conductivity and high electron mobility of 10~60 cm 2 /V s which is much larger than electron mobility in amorphous silicon, ~ 1 cm 2 /V s [20] As an active channel material of TFTs, amorphous oxides such as ZnO [ 21] InZnO [22] and InGaZnO [23] have been used to fabricate thin film transistors (TFTs) at room temperature. In contrast to imper fect flexible devices based on a Si:H [24 26] and organic semiconductors [27 29] Nomura et al. [23] reported high performance transparent flexible TFTs using flexible plastic substrates with n channel semiconducting oxide materials. 2.2 Transparent Conducting Oxides Transparent conducting oxides (TCOs) are an essential part of technologies that require both large area electrical contact and optical access in the visible range of the light spectrum [30] Research to develop TCO thin films, which are transpare nt like glass and conductive like metals, has been conducted using multicomponent oxides [31] The first observation that the optical transparency and electrical conductivity co exists was reported in 1951 in Cd oxide [32] To date, indium tin oxide (ITO) [33 35] has been widely used with its resistivity as low as 10 4 cm and a higher transmittance in the visible range. The TCO market is commercial ly dominated by aluminum doped ZnO (ZnO:Al) [36] and tin doped indium oxide [31] ITO offers one of the best combinations of high optical transparency and high electrical conductivity. Recent progress in thin film

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18 solar cells, flat panel display applications, and roll to roll coating operations for touch screens has been impressive but still requires improved TCO materials for transparent electrodes [37, 38] TCOs are c urrently being developed for the use as a transparent thin film transistor in transparent displays [23] The properties of conductivity and transparency are strongly interrelated, indicating a certain trade off between tw o phenomena The most common TCOs such as ITO and SnO 2 :F show the interdependence by the effect of deposition parameters on two properties of the thin films. R. Pommier et al. reported that sheet resistance (R sh ) and transmittance (T) of each deposited f ilm depend s significantly on the film thickness [39] 2.3 Crystalline P type Oxides As described above a large number of n type semiconducting oxide s has been discovered. However th e number of p type crystalline or amorphous semiconducting oxides is limited due to the charge localization in the valence band of oxides. Generally, oxides show strong localization of hole carriers at the valence band edge due to the large electronegativi ty of oxygen. For p type crystalline oxide semiconductors, modification of the energy band structure is required to reduce localization and to enhance hole carrier mobility. 2.3.1 Binary P type Oxides Among the binary conducting oxides, Cu 2 O and NiO indic ate a native p type conduction behavior. 2.3.1.1 Cu 2 O Cuprous oxide (Cu 2 O) is a well known p type oxide, which has a bandgap of 2.17 eV and shows a hole mobility exceeding 100 cm 2 /V s [40] Copper deficient Cu 2 O is p

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19 type due to the formation of copper vacancy V Cu acting as a hole carrier producer. Cu 2 O is intrinsically p type because the positively charged do nors (oxygen vacancy, V o ) acting as the hole killers ha ve no transition level in the gap and do not annihilate holes, and because the possible hole killer s, Cu i have both high formation energy and deep transition and is not capable to destroy holes create d by V Cu [41] The valence band maximum is formed by the interaction between filled d orbitals of Cu shells and filled 2 p orbitals of O shells. M. Izaki et al. [42] reported photovoltaic devices fabricated by using p type Cu 2 O. E. Fortunato et al. [43] reported bottom gate p type thin film transistors based on p type Cu 2 O. 2.3.1.2 NiO NiO is a p type oxide semiconductor with a bandgap from 3.6 to 4.0 eV [44, 45] Sato et al. obtained p type NiO with a resistivity of 1.410 1 cm and a hole concentration of 1.310 19 cm 3 and fabricated semitransparent thin fi lm p i n junction diodes using p type NiO and n type ZnO:Al layers [45] Nominally pure stoichiometric NiO is an ins ulator with a room temperature resistivity of the order of 10 13 cm [46] In undoped but nonstoichiometric NiO, conduction is usually p type, indicating that monovalent impurities or nickel vacancies are responsible, each of which leads to the formation of two Ni 3+ ions to obtain neutrality [44] In nonstoichiometric NiO, there must be an excess of oxygen which results in production of Ni vacancies. 2.3.2 D elafossite P type Oxi des In addition to binary conducting oxides, p type crystalline semiconducting oxides includes oxide with the crystal structure of delafossites such as CuAlO 2 [47, 48] and CuGaO 2 [49] and of spinels such as Zn M 2 O 4 ( M = Co, Rh, Ir) [50 52]

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20 2.3.2.1 CuAlO 2 In the delafossite structure of CuAlO 2 the copper ion (Cu + ) has fully filled d orbitals with ten electrons, which causes the closed shell valence state interacting with O 2 p orbitals. The closed shell is free from visible coloration arising from a d d transition in partial ly occupied transition metallic ions. As shown in F igure 2 1 t he CuAlO 2 is composed of alternate stacking of the dumbbell structure of linear O Cu O configuration and AlO 6 octahedral layers. Kawazoe et al. [47] reported the properties of thin films of p type CuAlO 2 with the conductivit y as large as ~1 S cm 1 at room temperature. Both the positive Hall coefficient, measured by van der Pauw electrode configuration, and the positive Seebeck coefficient suggests p type conduction. As no intentional doping of holes is introduced, the best ex planation of p type conduction might result from excess oxygen including cation vacancies and interstitial oxygen. Chemical analyses of CuAlO 2 delafossite samples show the Cu/Al atomic ratio in the film as 1.0. In the case of Cu 2 O, the Cu vacancy and the i nterstitial oxygen are the origin of the positive hole carriers [53] Yanagi et al. [54] reported that the indirect and direct allowed optical band gaps of CuAlO 2 are determined to be ~1.8 and ~3.5 eV respectively and that the upper valence band is primarily composed of admixed state of Cu 3 d and O 2 p orbitals. 2.3.2.2 CuGaO 2 CuGaO 2 [49] thin film s show the positive sign of Seebeck coefficient, in dicating the p type conductivity of the film. The dc conductivity of the CuGaO 2 film is estimated to be as large as 5.610 3 S cm 1 at room temperature with the optical band gap energy of ~3.4 eV. Photoemission spectroscopic measurement shows the valence b and edge starts around 0.5 eV, suggesting CuGaO 2 film to be a p type oxide and indicating

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21 contribution of Cu 3d components to the upper edge of valence band in CuGaO 2 [49, 55] 2.3.3 Spinel P type Oxides In the sp inel structure of Zn M 2 O 4 ( M = Co, Rh, Ir), the B site ions have a oxidation state of 3+ with a d 6 electronic closed shell [51] In the normal spinel structure, o xygen ions are in FCC close packing. A s shown in F igure 2 2 Zn ions are in tetra hedral A sites and M ions are in octahedral B sites. Transition metal sites are tetrahedrally and octahedrally coordinated by oxygen atoms. P type spinel oxides have hole carrier conduction paths caused by a band formation by hybridization of t 2g d orbital s and oxygen 2 p orbitals. For the valence band formation, keeping a d 6 low spin electronic configuration is critical. The optical bandgap generates from the ligand field splitting between empty e g 0 and fully occupied t 2g 6 in low spin configuration. 2.3.3. 1 ZnCo 2 O 4 H. Kim et al. [50, 56] claimed spinel ZnCo 2 O 4 thin films show bipolarity of the conduction behavior which is dependent on the oxygen partial pressure ratio. All six electrons of Co 3+ in the 3 d orbital occ upy t 2g states and form the low spin state of t 2g 6 in an octahedral crystal environment. The dependence of the conduction type on the oxygen partial ratio indicates that charge carriers are generate d from oxygen vacancies acting as donors and excess oxygen acting as acceptors. Non stoichiometry induces defects of the films. P type ZnCo 2 O 4 thin film shows the electrical conductivity of ~1.8 S cm 1 hole carrier concentration of 2.8110 20 cm 3 and mobility of 0.2 cm 2 /V s.

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22 2.3.3.2 ZnRh 2 O 4 H. Mizoguchi et al. [51] reported the ZnRh 2 O 4 normal spin el is p type semiconductor with a band gap of ~2 eV. Photoemission spectroscopy spectra and optical absorption spectra show that the band gap generates from the splitting of d orbitals due to the field strength of the oxygen ligands. The positive Seebeck c oefficient of 0.14 mV K 1 at room temperature confirms the major carrier is a positive hole. The electrical conductivity of the film is as large as 0.7 S cm 1 In an octahedral ligand field effect, the Rh 4d levels are split into empty e g 0 levels and fully occupied t 2g 6 levels with low spin con f iguration, acting as the conduction band bottom and the valence band top, respectively. 2.3.3.3 ZnIr 2 O 4 M. Dekkers et al. [52] reported the spinel ZnIr 2 O 4 thin films show p type conduction with the positive Seebeck coefficient of 53.9 V/K. The room temper ature conductivity of polycrystalline ZnIr 2 O 4 is determined to be 3.39 S cm 1 and the optical band gap is estimated to be 2.97 eV. 2g and e g states increases from 3d metal Co to 4d metal Rh to 5d metal iridium. Jrgensen int is a relative number depending on the ligand and g is a function of the metal [57] The value of f is about 0.9 and varies up to 0.37 from the position of oxygen within th e spectrochemical series [58] The value of g increases as the met al change from 3d to 4d to 5d in the unit of eV. 2.4 Amorphous P type O xide Narushima et al. [59] present ed the first p type amorphous oxide semiconductor, a ZnORh 2 O 3 which has the electrical conductivity of 2 S cm 1 at room temperature and

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23 the bandgap energy of 2.1 eV. The positive Seebeck coefficient and p n heterojunction confirms it shows p type conduction. The zinc rhodium oxide thin film is claimed to be amorphous because no sharp peak is observed in X ray diffraction measurement. Fine ordered structures are observed from the am orphous film by using transmission electron microscopy. T. Kamiya et al. [60] investigated the structural properties of p type amorphous a ZnORh 2 O 3 thin film. HR TEM images of the film exhibits a lattice like structure at the nanometer scale when the chemical composition is rodium rich. A rand om continuous network structure of the film is formed with edge sharing and corner sharing RhO 6 octahedra. It is claimed that hole transport paths made of edge sharing and corner sharing RhO 6 network is stable even in the amorphous structure because the or bitals (d xy +d yz +d zx ) forming the valence band maximum are not distorted largely in a disordered network structure. 2.5 Ligand Field Splitting Ligand (Crystal) field theory [61] explains the effect of the chemical environment on d n configurations of transition metal oxides 3 d transition metals ion with octahedral coordination has five d orbitals (d z2 d x2 y2 d xy d yz d zx ). These five orbitals are split into two distinct energy states described as e g and t 2g according to the sym metry behavior [62] : two orbitals have lobes of maximum probability pointing directly at the near neighbor oxygen atoms, whereas the other three have nodal planes in these directions [63] T he octahedral surrounding generates a ligand field splitting t 2g energy state and the higher e g energy state The electron configuration of transition metal ions in octahedral sites follows Pauli differe nt orbitals with parallel spins minimizes the electrostatic repulsion between

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24 electrons [64] In ligand fields, the low spin or high spin configuration can be obtained according to the exchange energy and the relative m filling of orbitals with parallel spins is favorable to exchange energies, whereas complete filling of orbitals from the bottom up lowers the energy when the splitting is large [63] What is important is that Co 3+ among 3d transition metal ions generally has the low spin configuration [65] The low spin configuration from Co 3+ occupying octahedral sites lea ds to separate fully occupied t 2g 6 levels forming the valence band and empty e g 0 levels forming the conduction band. By the orbital splitting in low spin configuration, positive hole carriers are expected to travel through t 2g 3d orbitals of Co 3+ interacti ng with ox ygen 2p orbitals. For instance, LaCoO 3 of the rhombohedrally distorted perovskite structure is a p type semiconducting oxide with a small band gap of 0.2 eV at 300 K where Co 3+ ions are located in octahedral sites [66] 2.6 Mott Insulator The Mott insulating state occurs when the kinetic energy gain is not large enough and the electron cannot move to another atom caused by the strong Coul omb repulsion energy [67] The electron electron correlation, strong Coulom b repulsion between electrons, is the origin of the insulating behavior [68] Near the Mott insulator regime fluctuations of spin, charge, and orbital correl ations are strong and metallic phases occur in what is called the anomalous metallic phase. Mass enhancement in V 2 O 3 is attributed to a typical anomalous fluctuation [69] As band theory predicts a metallic state when all bands are no t fully occupied or empty, Mott insulator is different from the usual band insulator in the internal degrees of freedom, spin, and orbital. For example, LaVO 3 is a Mott insulator with the electron configuration of 3 d 2 and LaTiO 3 is a Mott insulator with th e electron configuration of 3 d 1 of a Ti 3+ based oxide,

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25 suggesting the electrons in 3 d orbitals are itinerant or weakly localized [70, 71] In contrast, SrVO 3 with the electron configuration of 3 d 1 is a good metal c aused by the d electron transfer interaction modulated with strong overlapping between V 3 d and O 2 p states [72] 2.6.1 Electron C orrelation Here, the effect of electron electron interaction in correlated oxides is introduced in terms of Mott insulating, (anti)ferromag n etism, and superconductivity. The Mott insulating state is of basis for different physical systems such as the superexchange interaction and the superconducting states. When any itinerant electron moves to another site of an unoccupied band, it needs a certain amount of energy. No movement is generate d if the Coulomb energy is larger th an the energy gained by delocalization. Thus, it is not conductive even when the band is not fully filled. The Coulomb repulsion between electrons leads to correlated behavior and it can be expressed as (2 1) where U is the Coulomb energy between t wo electrons on a single site and W is the bandwidth of the system, the energy gained by delocalizing electrons. Note that U and W are affected by a change in interatomic distances, by introduction of an alloying element, and by variation of temperature or pressure. In other words, the Mott insulating state occurs when Coulomb repulsion is larger than the bandwidth. Mott criterion is introduced in terms of the Bohr radius and the electron separation length [73, 74] A Mott transition occurs in a single vale nce system from strong electron correlation (U>W) to weak correlation (WW gives rise to

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26 insulating behavior and U
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27 S uperconductivity may occur when electrons are removed by hole doping in the Mott insulator. BCS theo ry [79, 80] demonstrates that there can be an attractive interaction between electrons caused by the phonons electron lattice interaction The electrons form bosonic Cooper pairs because of the interaction. This superconducting state is called the BCS state. For high T c superconductivity electron interactions matters and the superconductivity, for example, can be tuned by changing the carrier density [81, 82] 2.6.2 Lantha num Vanadate, LaVO 3 As shown in F igure 2 3 LaVO 3 has a GdFeO 3 type orthorhombic structure with lattice parameters a=5.555548 b=7.84868 and c=5.55349 at room temperature [83] and is an antiferromagnetic insulator with a band gap of 1.1 eV [84] The structure can be considered as pseudocubic perovskit e structure with a lattice parameter of a c =3.92 [85] LaVO 3 is an interesting Mott insulator with the electron configuration of d 2 dominated by strong Coulomb repulsion, which shows a metal insulator transition with antiferromagnetic order for (La,Sr)VO 3 [86 88] H. Seim et al. synthesized non stoichiometry lanthanum vanadium oxide (LaVO 3 ) samples and investigated the effect of non stoichiome try in La 1 x VO 3 on unit cell dimension thermal and magnetic properties [89, 90] In pure LaVO 3 carrier transport is generate d via the excitation of holes from acceptor states which is located at 0.12 eV above t he valence band. In Sr doped LaVO 3 [91] V 4+ states with Sr 2+ ions form a band of filled hole traps with the transition from semiconducting to metallic conduction when the Fermi level acquires a mobility edge produc ed by Anderson localization [92] in the valence band. The transition is characterized with hopping conduction at low temperature. High

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28 temperature conductivity measurements of lanthanum strontium vanadate (La x Sr 1 x VO 3 ) present clear evidence of a mobility edge associated with the Anderson transition and the thermopower measurements exhibit little evidence of polaron formation [93] A.V. Mahajan et al. reported further exploration of structural, electronic, and magnetic properties of La x Sr 1 x VO 3 [94] In contrast to La x Sr 1 x VO 3 the transport mec hanism is driven from small polaron formation above the MI transition in doped LaMnO 3 [95] A Mott Hubbard gap in LaVO 3 between the upper and lower Hubbard bands (V 3d bands) and a charge transfer (CT) gap between the upper Hubbard band and the occupied O 2p band are determined to be 1.1 eV and 3.6 eV, respectively [84] 2.7 B and Insulator In contrast to Mott insulators that should show conductance under conventional band theory, band insulators are insulators in which the electrons in the valence band are separated by a large band gap from the conduction band. The large enough band gap prevents intrinsic semiconduction. Semiconductors have a small enough band gap between the valence band and the conduction band that excitations can overcome the gap. The introduction of doping material increases conductivity in semiconductors. 2 .7.1 d 0 I nsulator d 0 insulators are stoichiometric oxides with the d 0 electron configuration such as TiO 2 and KTaO 3 The band gap of d 0 insulators originates from the filled oxygen 2p valence band of bonding orbitals and the empty metal d conduction band o f antibonding orbitals. No optical absorption occurs with photon energies less than the band gap and diamagnetic property is observed due to no unpaired electrons. d 0 insulators often show semiconducting properties due to nonstoichiometry.

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29 2.7.2 Other Clos ed shell I nsulators Closed shell insulators are observed w hen the metal d band is full entirely or partially. d 10 insulators such as Cu 2 O (3d 10 ) have fully filled d orbitals. d 6 insulators with low spin configuration such as LaRhO 3 (4d 6 ) and d 8 square plan ar ions such as PdO are insulators in ground states. 2.7.3 Strontium Titanate SrTiO 3 Strontium titanate, SrTiO 3 crystallizes in the simple cubic perovsktie structure (O h 1 ) [96] with lattice parameter of a=3.905 at room temperature and has a fundamental absorption edge at 3.22 eV [97, 98] As shown in F igure 2 4 SrTiO 3 (STO) is composed of alternating stacks of SrO and TiO 2 atomic layers where Sr atoms are surrounded by 12 oxygen atoms and Ti atoms are surrounded by 6 oxygen atoms. Frederi kse et al. first reported the transport properties of SrTiO 3 that exhibits a band type conduction with an electron effective mass larger than the free electron mas s and low temperature mobilities larger than 1000 cm 2 /V sec [99] As oxygen vacancies generally act as electron donors, oxygen vacancies in strontium titanate (SrTiO 3 ) thin films are particularly important due to a tendency to maintain high carrier mobility even at high carrier density showing a metallic conducting behavior [100] The profile of the oxygen vacancy concentration in the SrTiO 3 superlattice films grown by pulsed laser deposition exhibits an abruptness on the atomic scale [101] Oxygen deficient STO has a metallic conduction and superconducting transition at T=0.3 K [102, 103] The oxygen vacancy related defect in STO substrates is monitored to investigate the eff ects of heat treatment in controlled O 2 and vacuum environment [104] In addition to reduced SrTiO 3 the electron Hall mobility and electron concentration of niobium doped SrTiO 3 are measured to study the scattering mechanisms [100] The mobility results are

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30 analyzed in terms of scattering by polar optical lattice modes at high tempera ture and scattering by ionized defects at low temperature, respectively. S trontium titanate (SrTiO 3 ) is one of the most popular substrates for oxide thin films for studies of novel properties at interface which are not observed in either of the bulk phase An atomically well defined and clean surface on the substrate enables one to accomplish layer by layer growth of various oxide thin films, which is prepared by selective etching of SrO atomic layer in NH 4 F HF solution and by annealing ab ove 600 C in ultr a high vacuum [105] A wet treatment in buffered NH 4 F HF solution (BHF or BOE) is used for atomically flat terraces with one unit cell high step [106] As atomic smoothness and surface termination play an important role in fabricating high temperature superconducting oxides and two dimensional epitaxy of heterostructures such as superlattices, there are a number of publications of surface tr eatment of STO substrates including annealing and wet etching methods [106 110] In addition to preparation of TiO 2 terminated STO substrates, generation of SrO terminated perovskite surface was reported, which imp roves the quality of cuprate films [111, 112] Beside the conventional wet etch procedure, an alternative wet etching in HCL HNO 3 (3:1) acidic solution is used to obtain high quality atomically flat TiO 2 terminated surfaces of STO with a controllability of the defect states that cause the presence of the Ti 3+ states at the surface [113] 2.8 Metal Insulator Transition Since Mott expected a first order metal insulator tra nsition (MIT) without structural phase change induced by a strongly correlated electronic Coulomb energy [73, 114] the first o rder metal insulator transition for a Mott insulator [115] is observed by pressure and temperature with a structural phase transition from the monoclinic

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31 structure of an insulator to the tetragonal structure of a metal [116 118] The conductivity of the electron gases showing MI transition is modulated by applying a voltage between the LaAlO 3 /SrTiO 3 interface and the SrTiO 3 substrate through a quantum phase transition from an insulating to a metallic state [119] Cen et al. [120] reported an interfacial MI transition at room temperature with the creation and erasure of nanoscale conducting regions at the LaAlO 3 /SrTiO 3 interface by using a conducting atomic force microscope probe. An abrupt MI transition without a structural phase transition is observed when a DC electric field is applied to an epitaxial VO 2 film based two terminal device [121] Here, metal insulator transitions of the perovskite related oxides in Mott Hubbard systems [122] are introduced. 2.8.1 Doping i nduced Metal Insulator T ransition Doping in transition metal perovskite oxides (AMO 3 ) is caused by substitution of divalent cations (A') for trivalent cations (A) which changes the valence state s of transition metal (M) ions from 3+ to 4+. The effect of the substitution (A 1 x A' x MO 3 ) is doping by controlling the 3d band filling. Doping induced (bandfilling controlled) MI transitions result from the disappearance of the carrier density. For example the optical properties of V O and Ti O perovskites have been reported as a function of bandfilling in La 1 x Sr x TiO 3 [123] and La 1 x Sr x VO 3 [87] 2.8.2 Bandwidth c ontrolled M etal I nsulator T ransition Bandwidth (W) control in transition metal perovskites (A 1 x A' x MO 3 ) is achieved by variations in the rare earth (A) site. The influe nce of the decrease in the M O M bond angle is that the splitting (U W) between upper and lower Hubbard bands increases. In other words, the increase in the splitting is caused by increasing the Coulomb energy (U) and by reducing the M O M overlap that gov erns the d conduction bandwidth (W).

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32 The M O M bond angle decreases from 180 while the tolerance factor it will be described later decreases from 1, which relates the size of the rare earth oxygen bond length (d R O ) to the size of transition metal oxy gen bond length (d M O ). Photoemission and optical studies have investigated the effects of bandwidth decrease which is increasing electron correlations (U/W) on the spectral weight transfer [124, 125] As a consequ ence, bandwidth controlled MI transitions are governed by the divergence of the carrier effective mass. For example, RNiO 3 (R is a rare earth) series exhibits a bandwidth controlled MI transition [126] Tolerance factor: R O M O ) It becomes 1 in the ideal perovskite structure (AMO 3 ). Deviation of the tolerance factor is a measure of the mismatch of the equilibrium (A O) and (M O) bond length. Isovalent O M interactions [127] 2.9 Examples Of Oxides W ith MI Transition Behavior 2.9.1 Ternary Tr ansition Metal O xides In La 1 x Sr x VO 3 the hole doping by substitution of the Sr atom for th e La atom site leads to an MI transition around x=0.2 [86, 91, 128] Changes in electronic properties on the doping induced MI transition in hole doped Mott insulators of La 1 x Sr x VO 3 result from the collapse of the Mott gap and resultant spectral weight transfer [87] Sr(Ti 1 x V x )O 3 composition spread films on oxygen deficient STO buffer layer/STO substrates using pulsed laser deposition in vacuum shows sharp MI transitions in the composition range of x=0.5 to x=1.0. The sharp transition is cons idered to be related to a certain phase transition by forming a dead layer with an energy gap between Sr(Ti 1 x V x )O 3 and SrTiO 3 [129]

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33 In addition, the superlattice of SrTiO 3 and SrVO 3 exhibits a MI transition at around 150 K with an anomalous nonlinear conduction [130] while both SrTiO 3 [100] and SrVO 3 [131] shows metallic conduction. The bulk perovskite nickelate family (RNiO 3 R is a rare earth) shows bandwidth controlled MI transition caused by the increase or decrease in the Ni O Ni bond angle with larger or smaller rare earth cation, resp ectively [68] NdNiO 3 exhibits a first order metal insulator transition with decreasing temperature, accompanied by an orthorhombic to monoclinic transition of the structural phase [132] In the epitaxial ultrathin films of NdNiO 3 the MI transition is quenched under the compressive strain, while it remains under the tensile strain, suggesting a sizable charge transfer gap remains in the ground state [133] The strain induced structural changes has an influence on the MI transition of t he epitaxial SmNiO 3 films [134] NdNiO 3 and PrNiO 3 thin films grown by pulsed laser deposition we re investigated to demonstrate the effect of substrate on electrical transport properties [135] The Sm x Nd 1 x NiO 3 thin film solid solutions prepared by PLD using NdNiO 3 and SmNiO 3 targets make the position of temperature of MI t ransition (T MI ) tunable [136] The (Nd 0.8 Y 0.2 )NiO 3 solid solution thin films displays MI transition near room temperature, indicating poss ible tunability of T MI The essential role of strain in perovskite manganite epitaxial thin films of (Nd 1 x Pr x ) 0.5 Sr 0.5 MnO 3 is demonstrated in the form of anisotropic and substrate dependent crossover of the MI transition through the tetragonal symmetry br eaking [137] 2.9. 2 Binary Oxides Vanadium oxide : Vanadium dioxide (VO 2 ) shows a metal insulator transition (MIT) near room temperature (T MIT = 340 K in bulk crystals) [116] The sputtered thin films without post annealing show a semiconductor to metal tra nsition at 63 C

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34 accompanied by a change in the electrical resistivity of nearly 10 5 [138] The effect of a strong electric field on the metal insulator phase transition in vanadium dioxide shows that this transition can be induced by an electric field [139] An abrupt MI transition is also observed in VO 2 thin films induced by a switching voltage pulse [140] The tetragonal lattice exhibits the metallic phase above T MIT [141] and the monoclinic lattice exhibits the semiconducting phase under T MIT [142] Either the electron correlation or the phase transformation leads to the metal insulator transition [143 145] The metal insulator transition characteristic s are strongly related to the electronic structure changes in vanadium oxide thin films, which emphasize that the energy band structure of the valence levels is responsible for the MI transition [146] While the metal insulator transition of VO 2 exhibits the first order phase transition, accompanied by an abrupt resistivity change, the metallic and insulating phases co exists and the different phases change dynamically in the process of the phase transition [147] 2.10 Complex Oxide Thin Films and Heterostructures In complex oxides, the complexity in crystal structure along with the localized nature of many electronic oxide properties yields opportunities to manipulate functionality via local modification and control of structure and chemistry. Perovskite related complex transition metal oxides have been widely investigated for various physical pheno mena such as ferroelectricity, metal insulator transition, antiferromagnetism, double exchange ferromagnetism, high temperature superconductivity, colossal magnetoresistance, and charge and orbital ordering. The wide variety of physical properties is most ly due to the electron correlation characteristic. One of the reasons is the directionality of the outer electron orbitals such as 3d or 4d of transition metals [148] This directionality has an effect on the degree of

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35 electron electron interaction and the change of the crystal structure. Another is that many tran sition metal ions are multivalent in the compound. Electrons in the outer shells lead to the electron correlation without influence on the core states of the transition metal ions. For example, LaTiO 3 is Mott insulator [149] La 1 x Sr x MnO 3 shows the ferromagnetic property [150] and YBa 2 Cu 3 O 7 which is derivative of perovskite crystal structure, is superconducting [150] In LaMnO 3 the colossal magnetoresistance (CMR) is related with a transition fro m a high temperature paramagnetic insulator to a low temperature ferromagnetic metal by hole doping from substitution of Ca or Sr atoms for La atoms [151] 2.10.1 Two dimensional E lectron G as Dingle et al. [152] investigated modulation dop ed GaAs AlGaAs heterojunctions where AlGaAs layers are doped. Two dimensional electron gas (2DEG) in GaAs results from electron accumulation from dopant impurities in AlGaAs due to spatial se paration between conduction electrons and impurity ions. Single heterojunctions and superlattices of GaAs AlGaAs exhibits the high electron mobility behavior [153 155] of two dimensional electron gas [156, 157] Quantization of the Hall resistance of the two dimensional electron gas at the interface of GaAs/Al x Ga 1 x As heterojunction prepared by using the molecular beam epitaxy techniques on semi insulating GaAs substrates is observed at 4.2 K [158] The q uantum Hall effect in a high mobility two dimensional electron gas in polar ZnO/Mg x Zn 1 x O heterosturctures grown by laser molecular beam epitaxy is observed and the electron density is controlled by tuning the magnesium content in the potential barrier and the growth polarity [159] In ZnMnO/ZnO heterostructures, two dimensional

PAGE 36

36 electron gas is formed at the interface and the temperature dependence of the magnetoresistance behaviors exhibits in the heterostructure [160] The quasi two dimensional electron gas is observed at the interface between LaAlO 3 and SrTiO 3 that are insulating perovskites [161, 162] The high mobility of the electrons at the interface is as large as 10 4 cm 2 /V s at 4.2 K, indicating higher densities of 2DEG at heterostructure interface than those in III V semiconductors [161 163] In addition to LaAlO 3 / SrTiO 3 interface, metallic electron gases are observed at the interface between LaTiO 3 and SrTiO 3 [164] The electron gas formed at the interface between LaAlO 3 and SrTiO 3 behaves a two dimensional superconductor with the transition temperature of ~200 mK [165] 2.10.2 Polar I nterfaces Polar interface exhibits a charge discontinuity between the LaO AlO 2 stacking and the vacuum because the vacuum has zero effective charge an d LaAlO 3 has orientation of alternating layers of LaO + and AlO 2 T he charge discontinuity leads to a surface reconstruction in bulk materials such as LaAlO 3 [166] and SrTiO 3 [167] Due to the polar nature of the LaAlO 3 a potential diverges as the thickness of the LaAlO 3 layer increases which may c ause an electronic reconstruction above a certain critical thickness [119] T he polar discontinuity at the heterostructure interface of electron correlated oxid es including non polar SrTiO 3 and polar LaTiO 3 [164] leads to an electro nic construction. In addition, the polar heterojunction interface of semiconductors such as Ge GaAs heterojunctions [168] may exhibit an ionic reconstruction. As described in Chapter 2.10.1 the high mobility two dimensional electron gas in polar ZnO/Mg x Zn 1 x O heterosturct ures is obtained by controlling magnesium concentration in the potential barrier and the growth polarity [159] H. Tampo et al. [169]

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37 observed 2DEG in Zn polar ZnMgO/ZnO heterostructures with strong confinement of electrons at the interface. T he electron mobility is as large as 250 cm 2 /V s at room temperature by increasing Mg composition of heterostructures. G.A. Baraff et al. [170] reported self consistent calculation of electronic structure at the ideal interface between intrinsic GaAs terminated on a (100) Ga plane, and intrinsic Ge. At the interface, there are in complete bonds at the termination of the group IV Ge layer and the commencement of III V alternations of GaAs. The electronic structure at an abrupt GaAs Ge interface can be considered as interface of polarity or valence discontinuities Kunc et al. [171] calculated the total energy, charge density, and electronic state of polar Ge GaAs interface. 2.10.3 Perovskite Related Superlattices Perovskite transition metal oxides are attractive for fabrica tion of the artificial superlattice due to the chemical stability and simple structure. It is significant for the superlattice investigation to understand the interface properties of perovskite oxides Here, complex oxide heterostructures of LaAlO 3 (LAO) a nd SrTiO 3 (STO) will be focused on. LaAlO 3 /SrTiO 3 interfaces : As described before, the polar discontinuity exists at the interface between LAO and STO since polar LAO has 1 alternating layers composed of positively charged planes of LaO + and negatively c harged planes of AlO 2 while non polar STO has neutral layers composed of SrO 0 and TiO 2 0 [161] The discontinuity distributes hal f a hole per unit cell area to the AlO 2 SrO interface and half a electron per unit cell area to the LaO TiO 2 interface. The hole doped AlO 2 SrO interface turns out to be insulating while the LaO TiO 2 interface is conducting with high mobility. No net free carriers of the insulating interface may results from hole carrier

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38 compensation by oxygen vacancies and from the difficulty of hole doping of closed shell transition metal ions [172, 173] Electrostatic potential b uilds up with a dipole shift and diverges in the limit of infinite LAO thickness in an electronically unreconstructed case. To avoid this polar catastrophic situation, charge redistribution at the interface is needed by injection of q/2 into the TiO 2 laye r [173] This electronic reconstruction [174] to compensate the interfacial polar discontinuity may explain the high conductivity of the interface between LAO and STO. The charge transfer between perovskite LAO and STO heterointerfaces makes conduction la yers with high electron mobility that show Schubnikov de Haas oscillations [161] Thiel et al. [119] observed a threshold thickness of LAO; the interface is not conductive until four monolayers of LAO thickness. Huijben et al. [175] observed a critical separation distance of LAO layers below which conductivity and carrier density at the interface decreases. One possible reason of carrier creation at the interface is oxygen vacancies. The transport properties at the LaO + TiO 2 0 interface varies as a function of the oxygen partial pressure during growth [161] due to the link between the vacancy concentration and the conductivity. Rijnders et al. [176] reported the effect of oxygen vacancies on the LAO STO interface by showing the sheet resistance characteristic as a function of temperature. The effect of oxygen vacancy may not be ruled out but it is not dominant origin of the carriers because both heterointerfaces in LAO STO systems do not show n type conductivity. The interface s between insulating oxides with metallic and sup erconducting behavior s have been demonstrate d with the role of oxygen vacancies [177, 178]

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39 Figure 2 1 The delafossite structure of CuAlO 2 Blue spheres represent copper (Cu) ions, light blue spheres represent aluminum (Al) ions, with red spheres representing oxygen (O) ions.

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40 Figure 2 2 The spinel structure of ZnCo 2 O 4 Grey spheres represent zinc (Zn) ions in tetrahedral sites, blue spheres represent cobalt (Co) ions in octahedral sites, with red spheres representing oxygen (O) ions.

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41 Figure 2 3 The orthorhombic perovskite s tructure of LaV O 3 Gree n spheres represent lanthanum (La) ions, yellow spheres represent vanadium (V) ions with red spheres representing oxygen (O) ions

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42 Figure 2 4. The cubic perovskite structure of SrTiO 3 Green spheres represent strontium (Sr) ions, light blue sp heres represent titanium (Ti) ions, with red spheres representing oxygen (O) ions.

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43 CHAPTER 3 EXPERIMENTAL DETAILS 3.1 Thin Film Fabrication 3.1.1 Substrate Preparation For the deposition of oxide and oxynitride thin films, the substrates include c axis or i Al 2 O 3 and 4907 C orning glass. For the growth of complex oxide thin films and heterosturctures, the substrate is a axis oriented SrTiO 3 The crystal structure of ABO 3 perovskites is composed of alternating stacks of AO and BO 2 atomic layers. A terraced surface with step heights of around half a unit cell can be obtained by cutting or cleaving the crystal. The crystal structure has two different AO and BO 2 termination on the surface. As received SrTiO 3 substrates have mixed SrO and TiO 2 terminated domains on the surface. As a STO platform with a high crystalline quality and TiO 2 single surface termination is critical for the fabrication of a novel oxide heterointerface showing a two dimensional electron gas (2DEG) behavior, it is necessary to control the surface termination to obtain single terminated surfaces. To prepare TiO 2 terminated surfaces, the STO surfaces are treated as follows. STO substrates are rinsed in deionized water for 5 minutes, and then are etched using buffe red oxide etch (BOE), also known as buffered HF or BHF, for 15 seconds. Etched STO substrates are rinsed in deionized wager for 5 minutes, and then are dried by blowing compressed nitrogen gas. As a last step, the substrates are annealed for 3 hours at 900 C in an oxygen ambient using a tube furnace. Figure 3 1 shows terrace steps with height of unit cell on the etched and annealed STO substrates. Before mounting substrates on the substrate holder, all types of substrates are cleaned using trichloroethylene acetone, and methanol for 10 minutes each in the

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44 ultrasonic bath. Compressed nitrogen gas is used for blowing and drying cleaned substrates. The silver paste (Ted Pella product #16035) is used for adhesion of substrates on the substrate heater and for he at conduction from the heater to the substrates. Before loading substrate into the vacuum chamber, the silver paste is dried enough using a glassware to block particles. 3.1.2 Target Preparation The targets are fabricated by mixing dry powders of precursor materials such as ZnO, Co 3 O 4 La 2 O 3 V 2 O 5 SrTiO 3 to meet the stoichiometric atomic ratios. Mixed powders are ground together using a clean mortar and pestle method without using any solvent, and then pressed at 2000 5000 psi into a cylindrical mold with 1 inch diameter. Lastly, the pressed target is sintered in an air ambient in a closed box furnace for 10 ~ 15 hours. 3.1.2.1 ZnCo 2 O 4 target There are two different powders of CoO(II) and Co 3 O 4 (Cobalt(II,III) oxide) to make a polycrystalline ZnCo 2 O 4 targe t. Both are commercially available from Alfa Aesar. As polycrystalline ZnCo 2 O 4 is a normal spinel structure, Co 3 O 4 (spinel structure) powder is preferred to CoO (rock salt structure) powder. While CoO only includes Co 2+ ions, Co 3 O 4 contains a mixed oxidati on state of 2+ and 3+. As explained, Co 3+ ions in octahedral sites produce the low spin configuration. Furthermore, the melting point of CoO is relatively high at 1830C compared to Co 3 O 4 which has decomposition temperature of 900C. To prepare the target Co 3 O 4 and ZnO powders are mixed to satisfy stoichiometric Zn/Co ratio of the mole fraction of the chemical formula, ZnCo 2 O 4 After the powders are mixed using mortar and pestle, the target is pressed in an 1 inch diameter die. Next, the pressed target is sintered in the box furnace at ~1000C in air

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45 ambient since the decomposition temperature of Co 3 O 4 is around 900C. The following is an expected equation of a chemical reactio n during solid state synthesis ; 3ZnO +2 Co 3 O 4 2 O 4 (3 1) Oxygen ambi ent can be used to prevent oxygen deficiency during synthesis. 3.1.2.2 LaVO and SrTiO 3 target s For LaVO 3 thin film growth a polycrystalline LaVO target is prepared using La 2 O 3 and V 2 O 5 from Alfa Aesar. Powders with precise weight are mixed to sati sfy an equation of a chemical reaction; La 2 O 3 + V 2 O 5 2LaVO 4 (3 2) In the box furnace in an air ambient, the mixed and pr essed powder is sintered at 700 C for 10 hours, is cooled down to room tempera ture, and then sintered at 1000 C for 12 hours addi tionally. LaVO 3 thin films are prepared by varying oxygen pressure with the target. For SrTiO 3 t arget fabrication, the SrTiO 3 powder commercially available from Alfa Aes ar is used and sintered at 1000 C for 12 hours using a box furnace in an air ambient. 3.1.3 Pulsed Laser Deposition Pulsed laser deposition (PLD) is a simple and versatile tool for thin film and multilayer research using a pulsed laser beam which strikes a target of the desired composition in a vacuum chamber [179] A pulsed laser rapidly removes material from a surface of a solid target and forms an energetic plasma plume, which moves onto a substrate. In brief, the mechanisms of PLD are laser ablation, plasma formation, plume propagation, and nucleation and growth. Du ring laser ablation, the photon energy converts into electronic excitations and then into thermal, chemical, and mechanical

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46 energy [180] The photon energy is also absorbed by a background gas such as oxyge n molecules and by optical elements such as mirrors and lens in the beam path. The wavelength of the laser such as KrF (248 nm) and ArF (193 nm) is an important factor since it determines the penetration path. When the energy is not mostly absorbed in a ve ry shallow region near the surface of the solid target, particulates at the deposited film surface can increase significantly. While plume propagates with different velocities of neutral atoms, ions, and electrons, some degree of thermalization is required in order for smooth film growth and for re sputtering prevent ion which generates by the most energetic ions in the plume [181] The stoichiometric removal from the solid target is of paramount important in PLD technique. During growth (laser ablation), a pre ablation step is applied in order to obtain a steady state. However, the stoichiometric removal does not guarantee translation into the growth of stoichiometrically removed materials due to different sticking coefficients and re s puttering. In addition, non stoi chiometric targets are used for compensation of atoms due to the difference in volatility. In case of oxides, oxygen content is needed to be controlled properly even if oxygen background pressure is applicable for oxide film growth. The effect of substrate heating on plume propagation i n oxygen atmosphere shows the relationship between the substrate temperature and the plume dynamics of complex oxides [182] With background oxygen gas, PLD is generally used for the oxide deposition to oxygenate the films. The deposition depends on following parameters: (i) the laser source and wavelength, (ii) the structural and chemical composition of the target material, (iii) the chamber pressure and chemical composition of the background gas, and (i v) the substrate temperature and distance between the target and the substrate.

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47 In the study of room temperature deposited zinc cob alt oxide thin films, a KrF exci mer laser with 248 nm wavelength is used as ablation source. Oxygen gas is used as a backgrou nd gas. The working pressure is varied from 50 mTorr to 200 mTorr. Oxygen with ozone gas can be utilized if oxygen gas itself is not enough to fully oxygenate the deposited film even in high pressure. It is because oxygen deficiency which may cause oxygen vacancies in the films can be a major blockage in achieving p type conducting oxide. In addition, laser energy and repetition rate are critical variables to acquire p type amorphous oxide films. Laser energy is handled to vary from 130 mJ/pulse to 200 mJ/p ulse and laser repetition rates of 1~10 Hz is tested. Lastly, sapphire, spinel wafers, glass and plastic films are considered as deposition parameters even when the films are deposited at room temperature. The films in the range of 100 to 500 nm on substra tes are deposited at room temperature for characterization. To acquire p type amorphous ZCO films and to control electrical properties, those parameters should be controlled overall. In the research of complex oxide thin films and heterostructures, a KrF l aser is used as an ablation source and oxygen gas is used as a background gas. The working pressure varies from 110 5 to 110 3 Torr, and the substrate temperature va ries in the range of 500 to 800 C during growth of LVO and STO layers. The targets are mo unted on a rotating holder and placed in a vacuum chamber evacuated to a base pressure of low 10 Torr. Laser energy varies from 80 mJ/pulse to 150 mJ/pulse with the repetition rate of 1~2 Hz. 3.1.4 Sputter Deposition Sputter deposition is a vacuum deposition technique which is widely used in the semiconductor industry. Sputter deposition has an ad vantage in depositing materials which are hard to be evaporated due to high melting points. During sputtering, material

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48 ejects from the target (the source material) due to collision of the target by energetic particles which have larger energy than the sur face binding energy, and the collision cascade generates. Then, an ejected atom from the target material deposits onto the substrate. It is a physical vapor deposition (PVD) method which is explained by momentum exchange between the primary particles such as ions in the plasma and atoms in the target material because of bombardment. Substrates are located in the vacuum chamber that is pumped down to the base pressure. An inert gas such as argon (Ar) is often used as the sputtering gas to reach the working p ressure. The atomic weight of the sputtering gas is preferred to be similar with the atomic weight of the target materials for efficient momentum exchange. Sputtering begins when a negative charge applied to the target, which cause a plasma or glow dischar ge. Positive charge gas ions in the plasma travel to the negative biased target material and generate the collision. Atoms or particles from the target materials are deposited as a thin film on the substrates. Sputtered films show good uniformity, density, purity, and adhesion. In the research, magnetron sputtering is used for metal deposition or oxide (or oxynitride) thin film deposition, which has almost no restriction in the target materials. It has a DC mode and a RF mode. DC sputtering is used for co nducting materials like metals, not for non conducting materials since the positive charge builds up on the material and stops sputtering. RF sputtering (or pulsed DC) is used for both conducting and non conducting materials including semiconductors and is olators. The magnetic field is used to trap secondary electrons close to the target, which increase electron path length and increase probability of electrons striking Argon. It enhances the ionization efficiency causing a higher sputter rate, indicating t hat the plasma can

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49 maintain at a lower pressure. The sputtered atoms are neutral so that the magnetic trap does not affect. 3.2 Structural and Compositional Characterization 3.2.1 X Ray Diffraction Complex oxide thin films and heterostructure s are optimall y prepared in order to obtain good crystal quality, preferred orientation, and atomically smooth surface of thin films by varying the laser energy, the substrate temperature, and the oxygen pressure. After the growth, the structure and growth orientation o f thin films and heterostructures are characterized using X ray diffraction (Philips APD 3720) in UF Major Analytical Instrumentation Center (MAIC) X ray diffraction is coherent elastic scattering of x rays by atoms in a crystal and has been in use in two main reasons to identify the crystalline solid with its unique characteristic and to determine the structure from X ray crystallography. X ray wave interference, commonly known as X ray diffraction, is evidence for the periodic atomic structure of cryst als. n = ger) is introduced to describe the condition for constructive interference from successive crystal planes and to explain why the faces of crystals reflect X ray beams at certain angles of incidenc ). The d is the distance between the ray beam. In Philips APD 3720 system, a copper (Cu) anode, which has characteristic used as the X ray source, showing different spectral line wavelengths such as 1 2 1 and so on. These wavelengths are known and emitted at known ratio of intensity. The system operates in a Theta from 20 to 80 and works under high power 40 kV 20 mA. The incident X rays may reflect in many directions but will be

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50 measured at one detecting location where angle of incidence equals to angle of reflection by moving the detector. An out of can plots X ray intensity in the MAIC facility is also available to measure rocking curves. As the phi scans are characteristic of t he orientation of the films, the profile shows the in plane orientation between the film and the substrate which determines if the films is epitaxially grown on substrates The omega rocking curve examines the degree of crystallinity of the epitaxial film by measuring the value of the full width at half maximum (FWHM) as the diffraction rocking curves taken in the Bragg geometry include the variation of the interplanar distance. The superlattice periodicity can be examined by using the equation (3 3) 1 2 are the locations of two adjacent superlattice peaks [183] 3.2.2 Auger Electron Spectroscopy Auger electron spectroscopy (AES) is an analytical tool for examining the chemical composition of solid surface. The major advantages include the high sensitivity for chemical analysis in the range of several angstrom near the surface, a rapid data acquisition speed, and a detecting ability all elements except hydrogen and heliu m, and a capability of high spatial resolution. The Auger process is initiated by creation of a core hole by exposing samples to a beam of high energy electrons (typically in the range of 2 to 10 keV), which have sufficient energy to ionize levels such as K, L 1 and L 2,3 levels. After atomic ionization by removal of a core electron and electron emission,

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51 X ray fluorescence or Auger emission generates while the ionized atom relax back to a lower energy state. During the excitation process, Auger transition, at least two energy states and three electrons must take part in an Auger process, demonstrating that H and He atoms do not generate Auger electrons. Despite the advantages, AES has a main limitation in examination of non conducting samples. Charging effe ct occurs when the number of secondary electrons moving out of the sample surface is different in the number of incident electrons. It leads to the shift of the measured Auger peaks since neutralization method is not applicable. After preparation of oxide (or oxynitride) thin films, the Perkin Elmer PHI 660 system in the MAIC facility has been used for the compositional analysis 3.2.3 X ray Photoelectron Spectroscopy X ray Photoelectron spectroscopy (XPS) is a quantitative surface analysis technique obtaini ng chemical information of the surface of solid materials instrument. It examines the elemental composition, chemical state, and electronic state of the elements at the surface. In contrast to Auger Electron spectroscopy, XPS analyzes the surface region of both insulators and conductors in the depth range of a few micrometers to a few millimeters. When exposing X ray to the sample, the core level electron of the atom comes out of the sample surface and the energy of the emitted electron is collected by an e lectron energy analyzer as a function of the binding energy (BE), which is characteristic of the element. One of the major advantages of the XPS is the ability to differentiate between different oxidation states and chemical environments. From the plot obt ained by using the Perkin Elmer 5100 XPS system in the MAIC facility, the peak shape and precise position determine the chemical state of the element.

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52 3.2.4 Scanning Probe Microscopy For characterization of oxide thin films and heterostructures, Atomic For ce Microscopy (AFM), Field emission scanning electron microscopy (FE SEM), and High resolution transmission electron microscopy (HR TEM) are used to examine the surface morphology and the surface roughness. The AFM Dimension 3100 in the MAIC facility is us ed to obtain the image of the thin film surface and the root mean square (RMS) value of surface roughness. In general 5 5 m 2 areas at 512 512 points are scanned. JEOL 6335F FEG SEM is used to investigate the image of the thin film oxide surface with magni fication above 85,000. Focused ion beam (FIB) dual bean strata DB235 is used to prepare TEM samples since the TEM requires very thin samples typically around 100 nanometers. FIB is for cross sectional investigation of complex oxide heterostructure to exam ine interfaces between two different perovskite oxide thin film layers. After FIB process, JEOL JEM 2010F scanning transmission electron microscope is used to obtain high resolution TEM images over 500,000 magnification. 3.3 Electronic Characterization 3. 3.1 Seebeck The Seebeck coefficient, which is also called the thermoelectric power, is a material property defined as the Seebeck voltage per unit temperature. The Seebeck coefficient is understood on the basis that electrons are both carriers of electrici ty and heat. If the temperature difference exists over the sample, a net diffusion of electrons from the hot region to the cold region generates, and then the voltage or electric potential is induced creating an opposing electric field. In equilibrium, the electric field results in the electric potential that is called the Seebeck voltage. The phenomenon is explained well but the underlying solid state physics is complicated to understand that

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53 the Seebeck coefficient has different signs for different metals Note that the sign and magnitude are related to an asymmetry of the electron distribution around the Fermi level. To investigate the carrier type of conduction of oxide thin films, the Seebeck coefficient is determined from a temperature gradient that is converted into a net diffusion of electrons, and it indicates a p type conduction when positive values are obtained if p p > n n When electrons are scattered by lattice vibrations, impurities, and defects, the mean free path of electrons and Seebeck coef ficient (3 4) are finally determined. 3.3.2 Hall E ffect In order to characterize the type and the number of carriers, the Hall effect [184] is widely used for transport properties of semiconducting materials. The Hall effect occurs when the movement of carriers in the material generates in a magnetic field, and the charge carriers is under the influence of the Lorentz force in a direction pe rpendicular to both the electric field and the magnetic field. The induced electric field, Hall field, that resulted from accumulation of the charge carriers causes the Hall voltage (V H ) between the top and the bottom of the thin film. Since the semiconduc ting material has both p type carriers (holes) and n type carriers (electrons) with different concentrations and mobilities, the Hall coefficient (R H ) is given by (3 5)

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54 The positive Hall coefficients confirms a p type carr ier conduction when p p 2 > n n 2 which is more difficult to be realized than to get the positive Seebeck coefficient since the electron mobility is much larger than the electron mobility [18 5] For the verification of p type conduction and for the determination of the electrical resistivity, the carrier concentration, and the Hall mobility, Van der Pauw method [186 ] is employed at room temperature using a Lakeshore 7507 Hall Effect Electronic measurement system in the UF Nanoscale Research Facility (NRF) as shown in F igure 3 2 By using the method, the Hall voltage (V H ) (3 6) where V 13 = V 13, P V 13, N V 24 = V 24, P V 24, N, V 31 = V 31, P V 31, N V 42 = V 42, P V 42, N is obtained from the diagonal difference of the voltages between positive and negative magnetic fields. The polarity of the Hall voltage in various magnetic fields indicates the conducting carrier type of the film a semiconducting thin film material is p type if V H is positive, and it is n type if V H is negative. 3.3.3 Physical Property Measurement System The Quantum Design PPMS represents a unique concept in labo ratory equipment: an open architecture, variable temperature field system, designed to perform a variety of automated measurements ( http://www.qdusa.com/products/ppms.html ) Use the PPMS with our spec ially designed measurement options, or easily adapt it to your own experiments. Sample environment controls include temperature range of 1.9 400 K in the presence of applied magnetic field up to 7 tesla. Its advanced expandable design combines many featu res in one instrument to make PPMS the most versatile system of its kind. The PPMS is capable of measuring the electric properties such as dc and ac

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55 resistivities as well as current voltage characteristics (I V curve), Hall measurements, and magnetic momen ts for any thin film, bulk material and nanowires. For general purpose resistance measurements and I V curve characterization, a two point electrical measurement is normally used. However, when the probe resistance or the probe contact resistance is relati vely high, or the resistance of the measured sample is relatively low, a four point probe measurement leads to more accurate results. When using 4 wire or Kelvin measurement, these parasitic resistances can be neglected for the two voltage probes because t he voltage is measured with a high impedance voltmeter, which draws very little current. Since negligible current flows in these probes, only the voltage drop across the sample is measured, which increases an accuracy compared to two point probe method. 3. 4 Optical Characterization The electronic optical band gaps (E g ) of semiconducting oxides are obtained from optical transmission spectra using Lambda 800 UV Vis spectrophotometer (Perkin Elmer instrument). The spectrometer is used to measure transmittance or absorbance calculated with parameters such as film thickness (l) and measured absorbance (A) or measured transmittance (T) by using Beer Lambert law (3 7) where I 0 is the intensity of the incident light at a given wavelength, and I is the 2 or 1/2 as a function of photon energy (eV) to demonstrate if the b and gap is direct or indirect, respectively. The value of band gap (E g ) is determined by extrapolating the

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56 2 1/2 = 0. This simple linear extrapolation procedure for E g is valid for an empty or nearly empty conduction band (or valence band). In case of degenerate semiconducting materials with large carrier density, the Fermi energy level (E f ) can move into the conduction band (or valence band) and shift the absorption edge towards higher energies. The Moss Burstein effect [187, 188] explains the shift in absorption edge and the overesti mation of E g

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57 Figure 3 1. Unit cell step surface of etched and annealed SrTiO 3 substrates obtained by AFM measurement. [ Adapted from Cianfrone, J. 2011. Functional complex oxide thin films and related superlattices grown via pulsed laser deposition Ph.D dissertation (Page 55, Figure 3 2). U niversity Florida, Gainesville Florida.]

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58 Figure 3 2. Schematic of a Van der Pauw configuration to determine the Hall voltage (V H ) by Hall effect measurements.

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59 CHAPTER 4 STRUCTURAL, ELECTRIC AL, AND OPTICAL PROP ERTIE S OF N TYPE INDIUM ALUMINUM ZINC OXYNIT RIDE THIN FILMS 4.1 Scientific B ackground In recent years, there has been significant interest in the use of semiconducting oxide thin films for a variety of electronic, magnetic, and optoelectronic applications [23, 30, 189 195] Technologies of interest have ranged from thin film transistors [196, 197] to light emitting diodes [198, 199] To date, indium zinc oxide [200] and indium gallium zinc oxide [201] have been studied as active channel layers for transparent thin film transistors Within the broad family of electronic oxides, oxynitride thin films have also been of interest. In the area of diffusion barriers in metallization, tantalum oxynitride films have been examined [202] Indium oxynitride [203] zinc oxynitride [204] indium tin oxynitride [205] and titanium oxynitride [206] films have intere sting properties relevant to optoelectronic devices. Cadmium germanium oxynitride [207] films have been used for NH 3 and H 2 S gas sens ors, while indium oxynitride [208] films are useful in NO 2 gas sensors. Oxynitride films are typically prepared by physical deposition in a background of oxygen and nitrogen. RF magnetron sputtering in the Ar N 2 mixed plasma can be used [202 210] Lithium phosphorous oxynitride [211] thin films have been fabricated by pulsed laser deposition in an ambient N 2 gas; LaSrTiON [212] thin films have been fabricated by solution based spin coating. In many cases, tuning of film properties can be achieved by controlling nitrogen gas ratio in the Ar N 2 plasma [200, 201, 204, 205, 208] RF magnetron sputtering technique has been predominantly used to prepare oxide or oxynitride thin films by mixing gases suc h as Ar, N 2 O 2 etc. By varying the composition of gases in plasma, film properties can be optimized to fabricate electronic

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60 and optoelectronic devices. However, since some oxide films were sensitive to oxygen gas [200, 201] electrical properties significantly varied with small change of oxygen partial pressure in plasma. We expect, therefore, nitrogen gas in plasma can be an option to control thin film properties effectively. This work was basically conducted at room temperature in order for further application for low temperature and large area device applications. In Chapter 4 the synthesis and electrical properties of InAlZnON thin films deposited by rf magnetron sputtering is reported. The films were synthes ized by co sputtering from aluminum oxide and indium zinc oxide (IZO) targets. Nitrogen is introduced into the films by using an argon/nitrogen gas mixture during rf magnetron sputter deposition. All films were deposited at room temperature. Chemical compo sition and nitrogen content in the films were measured by Auger electron spectroscopy (AES) and X ray photoelectron spectroscopy (XPS). Surface morphology was studied with atomic force microscope (AFM) and optical absorption coefficients were obtained by using UV Vis spectrophotometer Varying t he sputter power of the IZO and Al 2 O 3 targets resulted in continuous control of electrical properties. Tuning of the film carrier density could be achieved by varying the nitrogen concentration in the Ar N 2 plasma. Electrical conductivity, carrier density, and mobility were examined at temperatures ranging from 100 K to 300 K. 4.2 Experimental Methods InAlZnO and InAlZnON thin films were deposited on glass substrates by rf magnetron sputtering. The 3 inch diameter IZO and Al 2 O 3 targets were used for co depo sition at room temperature. The substrates were cleaned in an ultrasonic bath for 5 minutes each in trichloroethylene, acetone, and methanol, and were blown dry by

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61 nitrogen gas. The sputter deposition chamber was pumped down to less than 5x10 6 Torr. The s putter targets were precleaned in an Ar or Ar N 2 mixed plasma for 10 minutes prior to deposition. A sputtering gas pressure of 10 mTorr was used for the deposition of all films. The nitrogen gas partial pressure was varied from 0 to 7 mTorr. Glass substrat es were mounted on a substrate holder. Substrate temperature remained below 30 C during the deposition. The power applied to the IZO target was varied from 100 W to 200 W; power to the Al 2 O 3 target was varied from 150 W to 300 W. The electrical properties of the films, specifically the Hall coefficient, electrical resistivity, carrier density, and mobility, were determined by Hall effect measurements using a Lakeshore 7507 system employing a van der Pauw method at room temperature. For Auger analysis, film s were pre sputtered with argon for 3 minutes. For XPS analysis, aluminum was used as the X ray source; the film was sputtered with argon for 1 minute prior to acquiring the data. Surface roughness was analyzed by AFM Dimension 3100 in tapping mode. The el ectrical conductivity, carrier density, and mobility of the films were measured from 100 K to 300 K by a four probe technique. Indium was used to form an ohmic contact. Optical transmission spectra were obtained by a Lambda 800 UV Vis spectrophotometer (Pe rkin Elmer instrument). The thicknesses of the films were determined by using an alpha step Tencor profilometer. Thickness of all oxynitride films ranged between 300 nm and 400 nm. 4.3 Results and Discussion 4.3.1 Chemical Composition and Surface M orpholog y The dependence of chemical composition in the films as a function of the Al 2 O 3 target sputter power was examined by using AES. Figure 4 1 shows the results for the case when the IZO sputter power was fixed at 150 W, and the nitrogen partial pressure

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62 was 0.5 mTorr during the deposition. Aluminum concentration in the InAlZnON films increased as Al 2 O 3 sputter power increased. As will be discussed later, this also resulted in a lowering of the carrier density of the films. Both indium and zinc concentration d ecreased as Al 2 O 3 sputter power increased from 150 W to 350 W. The presence of nitrogen in the oxynitride films was investigated by AES. Figure 4 2 shows the AES spectrum for a film deposited in 5% nitrogen/95% argon with an IZO sputter power of 150 W and an Al 2 O 3 sputter power of 350 W. The differentiated Auger electron energy peak at 381 eV corresponds to nitrogen [208] Unfortunately, a quantitative determination of nitrogen concentration was not possible from AES due to the low intensity of the nitrogen peak. The existence of nitrogen in the oxynitride film was also confirmed by XPS. Figure 4 3 shows an XPS spectrum for a film prepared with an IZO sputter target p ower at 200 W, an Al 2 O 3 sputter power of 200 W, and a nitrogen partial pressure of 4 mTorr. The spectrum shows the N(1s) peak at 398 eV [ 213] Figure 4 4 shows the surface morphology of the oxynitride film as observed by AFM. The film was deposited with a N 2 partial pressure of 0.5 mTorr, an IZO sputter power of 200 W, and an Al 2 O 3 sputter power of 200 W. The scan area shown in the image is 5 5 m 2 The root mean square roughness for the film was determined to be 1.79 nm. 4.3.2 Electrical and O ptical P roperties Hall measurements were used to determine carrier type, carrier concentration, and mobility. In this study, all films were n type as determined by room temperature Hall effect measurement. Figure 4 5 shows the electrical resistivity and carrier density for oxynitride films deposited at various Al 2 O 3 sputter target power with the IZO sputter power fixed at 150 W and a nitrogen pressu re of 0.5 mTorr. The carrier density

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63 decreased as the Al 2 O 3 target power increased from 150 W to 350 W. The decrease in the carrier density is consistent with the increase in resistivity in that range. As shown in Figure 4 1, this decrease in carrier conce ntration correlates with an increase in aluminum concentration. By controlling aluminum concentration with Al 2 O 3 target power, the carrier density could be systematically varied from 10 15 to low 10 19 cm 3 Figure 4 5 shows the corresponding mobility of th e films. The mobility var ies from 1.5 to 4 cm 2 /V s, showing a moderate increase with increasing carrier concentration. This increase in mobility with increasing carrier concentration behavior is qualitatively similar to that seen in InZnO films [214] The transport properties were also examined a s a function of the IZO sputter power. Figure 4 6 shows the resistivity and carrier density as a function of the IZO target power. The Al 2 O 3 sputter power was fixed at 250 W; the nitrogen partial pressure was 0.5 mTorr. As the IZO target power increased fr om 100 W to 200 W, the carrier density increased from high 10 15 cm 3 to low 10 19 cm 3 reflected in a decrease in the electrical resistivity. The relative aluminum concentration was reduced by increasing IZO sputter target power, resulting in a n increase d carrier density in the indium aluminum zinc oxynitride films. The mobility was measured to be 1.5 to 3 cm 2 /V s as shown in Figure 4 6. A.J. Leenheer et al. [200] reported the contr ol of the electrical resistivity and carrier density in conductive InZnO films by varying the oxygen gas pressure in the Ar plasma. As the amount of oxygen in the plasma increased, the carrier density decreased due to reduced oxygen vacancies of films. The present study shows that, by co sputtering an aluminum oxide target, the carrier density can be systematically controlled through the Al concentration.

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64 The effects of nitrogen on electronic properties of deposited films were also examined. Figure 4 7 show s the resistivity, carrier density, and mobility of films deposited at constant IZO (200 W) and Al 2 O 3 (200 W) sputter power with varying N 2 partial pressure. By increasing the partial pressure of nitrogen gas ratio in the Ar N 2 plasma, the carrier density and mobility were reduced and the electrical resistivity increased. The carrier density decreased from 10 19 cm 3 to 10 17 cm 3 when the N 2 partial pressure was varied from 0 to 4 mTorr in the sputter gas mixture. The carrier density increased to 10 18 cm 3 w hen the nitrogen gas partial pressure was increased to above 6 mTorr. In previous work on indium oxynitride [208] and indium tin oxynitride [205] films, both carrier density and mobility decreased with increasing nitrogen incorporation. In that work, it was speculated that nitrogen occupation of oxygen vacancies caused the reduction in carrier density [205] Nitrogen incorporation also decreased the mobility of the InAlZnON films. Above 20% N 2 the mobility of oxynitride films was an order of magnitude lower than that of the InAlZnO film with no nitrogen. The temperature dependent transport properties of the oxynitrid e films were also examined. Figure 4 8 shows the film conductivity over a temperature range of 100 K to 300 K for a film deposited with a N 2 partial pressure of 0.5 mTorr, an IZO sputter power of 150 W, and an Al 2 O 3 sputter power of 250 W. The conductivity of the film was 0.102 S cm 1 at room temperature. The plot indicates semiconducting behavior and Arrhenius type behavior with a thermal activation energy of 62 meV, which is similar to other oxide films [215] This is consistent with the Fermi level lying close to the conduction band for electron conduction.

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65 Both the carrier conce ntration and mobility were measured as a function of temperature. Figure 4 9 shows the temperature dependence of the carrier density. The carrier density was 4.1410 17 cm 3 at room temperature. The carrier density decreases as temperature is decreased. Fig ure 4 9 also includes the temperature dependence of the electron mobility. At room temperature, the mobility was 1.56 cm 2 /V s. The electron mobility decreased as temperature decreased, suggesting that ionized impurity scattering limits the mobility. The ch ange in the mobility as a function of temperature was larger than that of the carrier density. Figure 4 10 shows the optical absorption spectrum obtained from the optical transmission measurement in the range of 250 nm to 800 nm. The optical bandgap (E g ) of the oxynitride film deposited with a N 2 partial pressure of 0.5 mTorr, an IZO sputter power of 200 W, and an Al 2 O 3 sputter power of 250 W was estimated to be ~3.55 eV by 2 =0, indicating the film is transparent. 4 .4 Summary Indium aluminum zinc oxynitride (InAlZnON) films were fabricated by rf magnetron sputtering using both the IZO target and Al 2 O 3 target in Ar/N 2 plasma. All the films were determined to be n type electric conduction by Hall measurement at room te mperature. Atomic concentration of indium, zinc, and aluminum was examined by AES analysis. By varying Al 2 O 3 power, the amount of aluminum was controlled in InAlZnON film. Nitrogen incorporation into oxynitride films was investigated by AES and XPS analyse s. Surface roughness of oxynitride film was examined by AFM. Carrier densities of as deposited films were reduced by increasing Al 2 O 3 power with the fixed IZO power. In addition, carrier densities were augmented by increasing the IZO power with the fixed A l 2 O 3 power. By increasing the amount of nitrogen gas in the Ar N 2

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66 plasma, the carrier density and mobility were reduced. The carrier density and mobility of InAlZnON films were lower than those of InAlZnO film. Temperature dependent properties such as the conductivity, carrier density, and mobility were investigated in this work. The conductivity of the oxynitride film exhibited Arrhenius behavior with an activation energy of 62 meV. Optical bandgap energy of InAlZnON film was determined to be ~ 3.55 eV, wh ich means transparent. These transparent InAlZnON films would extend electronic and optoelectronic applications for low temperature process due to easily controllable properties by sputter deposition parameters at room temperature

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67 Figure 4 1. Indium, zi nc, and al uminum atomic concentration as a function of Al 2 O 3 power of the InAlZnON film

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68 Figure 4 2. Differential Auger pattern of the InAlZnON film

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69 Figure 4 3. X ray photoelectron spectroscopic analysis of N(1s) of the InAlZnON film

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70 Fig ure 4 4. AFM image of InAlZnON film with the root mean square (RMS) value of 1.79 nm

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71 Figure 4 5. Resistivity, carrier density, and mobility of InAlZnON films in the Al 2 O 3 power range of 150 W to 350 W

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72 Figure 4 6. Resistivity, carrier density, an d mobility of InAlZnON films in the InZnO power range of 100 W to 200 W

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73 Figure 4 7. Resistivity carrier density, and mobility of InAlZnON films as a function of nitrogen ratio in the Ar N 2 plasma

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74 Figure 4 8. Temperature dependence of the ele ctrical conductivity for the InAlZnON film

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75 Figure 4 9. Temperature dependence of the carrier density and the electron mobility for the InAlZnON film

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76 Figure 4 10. Optical abso rption spectrum of the InAlZnON. The optical bandgap energy (E g ) is determined to be ~3.55 eV.

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77 CHAPTER 5 ELECTRICAL AND OPTIC AL PROPERTIES OF N TYPE ZINC ALUMINUM OXYNITRIDE THIN FILMS 5.1 Scientific B ackground Transparent conducting oxides (TCOs) are attractive materials due to electrical conductivity and visual transp arency, and they are widely applied in opto electronic devices such as flat panel di splays or thin film solar cells [190, 216 220] With controllable electrical conductivity and optical transparency in the visible region, some oxide thin films have been used as active materials for rectifying junctions and field effect transistors [21, 23, 201, 221 225] The zinc oxide (ZnO) film has been studied as a wide band gap semicondu ctor, which shows n type conduction dominated by oxygen vacancies and Zn interstitial atoms [224, 226 229] Moreover, Xiong et al. [19] reported intrinsic p type ZnO films by using reactive sputtering by adjusting the oxygen partial pressure in the plasma. To achieve a high conductivity as wel l as a high transmittance in the visible range, the doped ZnO films with the group III elements such as aluminum (Al), indium (In), gallium (Ga), and boron (B) have been investigated [230 232] It has a potential t o replace InSnO 2 (ITO) because doped ZnO is cheaper, more temperature stable, and nontoxic with similar electrical and optical properties [233] Among them, Al doped ZnO (ZnO:Al) films have become attractive because of surface texturability for efficient light trap ping in silicon thin film solar cells [234, 235] To date, studies of ZnO:Al films prepared by using various sputtering targets of different Al doping concentration have been reported to optimize and to improve st ructural, electrical, and optical properties [36, 236, 237] The large conductivity values of ZnO:Al films have been achieved with an Al concentration of 2 3 at % [238] In addition, Perkins et al. [195] reported zinc aluminum oxide thin films prepared by co

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78 sputtering from two targets. Furthermore, nitrogen source such as N 2 and N 2 O has been used to study electrical and optical properties of ZnO [239] and ZnO:Al [230] films by annealing them after deposition. However, the effect of nitrogen incorporation of zinc aluminum oxide (Zn Al O) thin films has not been studied systematically. In Chapter 5 we have examined the influence of nitrogen introduction during sputtering on electrical and optical properties of Zn Al O thin films deposited on sapphire substrates at room temperature. In addition, the effect of N 2 volume ratio in the mixed working gas of argon and nitrogen (Ar N 2 ) was a lso investigated. 5.2 Experimental Methods Nitrogen incorporated Zn Al O thin films were prepared on single crystal sapphire (0001) substrates by radio frequency (RF) magnetron sputtering with a base pressure at around 5 10 6 Torr. The 3 inch diameter Zn O and Al 2 O 3 targets were used for co sputtering deposition at room temperature in the mixture of Ar N 2 The working pressure was kept at 10 mTorr and the sputtering power of two targets varied from 150 W to 300 W. The thickness of films was in the range of 400 500 nm and the temperature was remained below 30 C during the deposition. The electrical properties were obtained by Hall effect measurements using a van der Pauw method at room temperature. Optical properties were measured by a Lambda 800 UV Vis s pectrophotometer. For Auger analysis, films were pre sputtered with argon for 3 minutes. The thickness of films was determined by using an alpha step Tencor profilometer. 5.3 Results and Discussion Elect ric al and optical p roperties : The electrical properti es of nitrogen incorporated Zn Al O thin films deposited in 5 % N 2 /95 % Ar at room temperature are shown in Fig ure 5 1 The sputtering power of ZnO varies from 150 W to 300 W while the

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79 sputtering power of Al 2 O 3 fixed at 250 W. The resistivity increases from 3.4310 3 to 4.0110 5 ohm cm due to decrease in carrier density from 2.6010 15 to 8.1510 12 cm 3 as the sputtering power of ZnO increases. This decrease in electron carrier concentration correlates with decrease in aluminum concentration, which is consisten t with the contribution of substitutional Al 3+ ions on Zn 2+ sites and Al interstitial atoms The electron mobility varies between 2.5810 1 and 2.01 cm 2 /V s with no clear trend as the ZnO sputtering power varies. Figure 5 2 shows the optical absorption spe ctra of Zn Al O films in various ZnO target powers. The optical band gap (E g ) determined by Tauc plots reduces as the ZnO target power increases, which indicates that the effect of Al 2 O 3 decreases. This result may be caused by the large band gap of Al 2 O 3 ( ~6.2 eV) [240] in comparison with that of ZnO (~3.3 eV) [241] Figure 5 3 shows the Al 2 O 3 sputtering power dependence of the electrical resistivity and carrier dens ity of the films deposited in 5 % N 2 /95 % Ar at room temperature. The sputtering power of Al 2 O 3 varies in the range of 150 to 300 W with the fixed sputtering power of ZnO at 150 W. The decrease in resistivity from 510 5 to 5.44 ohm cm with increasing the Al 2 O 3 sputtering power results from increasing carrier density 1.9510 14 to 3.3410 18 cm 3 due to higher Al doping concentration in the films, which is cons istent with explanation fo r Figure 5 1. The resistivity of ZnO films is large in the order of 10 5 ohm cm when it is not doped intentionally and it decreases by doping of aluminum to ZnO grains in the crystalline film [242] The optical band gap shown in Figure 5 4 increases as Al 2 O 3 target power increases, suggesting the enhanced influence of Al 2 O 3 of Zn Al O films.

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80 Figure 5 5 shows the effect of nitrogen volume ratio in the mixture of Ar N 2 on the electrical properties of the films deposited at room temperature with the ZnO sputter power of 250 W and the Al 2 O 3 sputter power of 250 W. The resistivity increases from 2.81 to 2.9910 5 ohm cm as nitrogen volume ratio increases, which results from the carrier density decrease from 9.610 18 to 8.1710 13 cm 3 Since nitrogen incorporated Zn Al O films show n type conduction and N 2 is a n type dopant [ 16] this result may not look reasonable. Probably, N source from N 2 may act as a p type dopant during sputtering deposition. As shown in Table 5 1, Auger analysis confirms the incorporation of nitrogen into the Zn Al O film. The amount of N inclusion in the films increases with the increase of ni trogen volume ratio except at 5 %. Note that substitutional N at an O site is an acceptor and substitutional N 2 at an O site is a double shallow donor [16] As shown in Table 5 1, however, Zn/Al atomic concentration ratio obtained from Auger analysis presents an increase in deposition rate with increasing nitrogen volume ratio Only nitrogen partial ratio varies an d all other deposition parameters are same, which suggests that nitrogen has an effect on the ZnO deposition rate during co sputtering of ZnO and Al 2 O 3 targets based on the increase of the net growth rate during co deposition. The reason why the total rate of co sputtering changes is not clear but an increase of Zn/Al ratio leads to decrease the carrier density and to increase the resistivity by reducing Al doping concentration. Figure 5 6 shows the variation of the optical band gap as a function of nitroge n volume ratio. It shows that the band gap gets narrower with increasing N 2 ratio which results in the decrease of Al concentration as shown in Table 5 1. The widening band gap with increasing Al concentration can be explained based on the Burstein effect [188]

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81 Figure 5 7 shows a relationship between the electrical resistivity and the Zn/Al atomic ratio to understand the effect of nitrogen partial pressure during co deposition since the Al content is dependent on deposition conditions such as the nitrogen volume ratio and the sputtering power. All data points are obtained from the films shown in Figure 5 1 ~ 5 6 The resistivity continuously increases with increasing the Zn/Al r atio from 2.81 to 4.0110 5 ohm cm which is caused by decreasing carrier density from 9.610 18 to 8.1510 12 cm 3 Nitrogen volume ratio in the range of 0 to 4% has an effect on Al atomic concentration of nitrogen incorporated Zn Al O films during depositio n, and the change of Zn/Al ratio shows a trend including the ratio of other films tuned by sputtering target powers in the N 2 Ar mixture. The Hall mobility of the films introduced in the paper varies in the range of 0.11 to 2.01 cm 2 /V s with no trend. Figu re 5 7 do not show a clear linear trend but include fluctuation of data points, which suggests that the formation of aluminum oxide or Al suboxide during deposition may cause the number of active Al dopants in the films. 5.4 Summary W e have studied the eff ect of nitrogen incorporation on the electrical resistivity and the optical band gap energy of Zn Al O thin films prepared by co sputtering of ZnO and Al 2 O 3 targets. The increase of Al atomic concentration of the films leads to increase the band gap energy as well as to decrease the resistivity by increasing the carrier density. Al atomic concentration is controlled by varying the sputtering power of two targets and by varying nitrogen volume ratio in the mixture of Ar N 2 Increasing N 2 ratio in the mixture causes decreasing Al concentration by changing the deposition rate of the ZnO target. The electrical resistivity as a function of Zn/Al atomic ratio is presented to see the effect of sputtering powers and nitrogen ratio on Zn/Al ratio during co depositio n.

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82 Figure 5 1. Resistivity, carrier density, and mobility as a function of the ZnO sputtering power of nitrogen incorporated Zn Al O films

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83 Figure 5 2. Optical absorption spectra of various ZnO sputtering powers in the range of 150 to 300 W.

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84 Figure 5 3. Resistivity, carrier density, and mobility as a function of the Al 2 O 3 sputtering power of nitrogen incorporated Zn Al O films

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85 Figure 5 4. Optical absorption spectra of various Al 2 O 3 sputtering powers in the range of 150 to 300 W.

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86 Figure 5 5. Resistivity, carrier density, and mobility as a function of the N 2 volume ratio in the mixture of Ar N 2

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87 Figure 5 6. Optical absorption spectra of various N 2 volume ratios in the range of 0 to 5%.

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88 Figure 5 7. Resistivity and carr ier density as a function of the Zn/Al atomic ratio

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89 Table 5 1. Atomic concentration of nitrogen (N), oxygen (O), zinc (Zn), and aluminum (Al) obtained from Auger analysis including the Zn/Al atomic ratio. Net deposition rates of co sputtered films are me asured to investigate the nitrogen effect during deposition at room temperature.

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90 CHAPTER 6 STRUCTURAL, ELECTRIC AL, AND OPTICAL PROP ERTIES OF P TYPE ZINC COBALT OXIDE THIN FI LMS 6.1 Scientific B ackground Amorphous and polycrystalline semiconductors have l ong been exploited for large area electronic applications. Amorphous silicon has been prominent in electronic device technologies since the creation of hydrogenated amorphous silicon [1] in 1975. However, amorphous silicon has limitations. A relatively low bandgap makes it opaque in the visible spectrum. Amorphous silicon also has a relatively low carrier mobility. In contrast, many of the amorphous and polycrystalline oxide semiconductors are optically transparent due to large bandgaps. Many exhibit a relatively large carrier mobility, particularly in the n type materials [195, 243, 244] For example, indium gallium zinc oxide (In Ga Zn O) is an n type amorphous semiconducting oxide currently being explored for thin film transistor technology due to its high electr on mobility of 10~60 cm 2 /V s [20] N type semiconducting oxides such as ZnO [21] InZnO [22] and InGaZnO [23] have been used as active materials for rectifying junctions [192, 221, 222, 245] and n channel field effect transistors (FET) [201, 224, 225, 246] fabricated at room temperature. In contrast to the relativel y large number of n type amorphous or polycrystalline oxide semiconductors, there are fewer p type semiconducting oxides, particularly those that can be fabricated at or near room temperature. The formation of p type ZnO is challenging [247, 248] Narushima et al. [59] reported the first p type amorphous oxide semiconductor, a ZnORh 2 O 3 which has the bandgap energy of 2.1 eV and the electrical conductivity of 2 S cm 1 at room temperature. The limited number of p type amorphous or polycrystallin e semiconducting oxides is a consequence of the charge

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91 localization in the oxide valence band that is principally comprised of oxygen 2 p orbitals. Typical oxides exhibit strong localization of holes at the valence band edge due to the large electronegativi ty of oxygen. For p type crystalline oxide semiconductors, modification of the energy band structure is needed to reduce localization and enhance hole mobility. Localization of holes at the valence band edge can be mitigated by using cations with the close d shell electronic configuration. To date, several p type crystalline semiconducting oxides have been reported, and include delafossites such as CuAlO 2 [47] and spinels such as Zn M 2 O 4 ( M = Co, Rh, Ir) [50 52] In the delafossite structure of Cu AlO 2 the copper ion (Cu + ) has a d 10 electronic configuration which is the closed shell valence state overlapping with O 2 p states. In the spinel structure of Zn M 2 O 4 ( M = Co, Rh, Ir), the B site ions have a d 6 c losed shell [51] Th e closed shell hybridization with ligand O 2 p orbitals is key to yielding p type conduction. P type spinel oxides are composed of transition metal ions in an octahedral crystal environment, keeping a d 6 electronic configuration that forms a hole carrier co nduction path [63] Furthermore, the bandgap between empty e g 0 and fully occupied t 2g 6 originates from the ligand field splitting generated by the octahe dral environment of transition metal ions. In low spin electron configuration, t 2g d orbitals interacting with oxygen 2 p orbitals yield a band that favors p type conduction. In the case of amorphous ZnORh 2 O 3 films, T. Kamiya et al. [60] reported that the network structure made of RhO 6 octahedra i s stable even in an amorphous phase. Hole transport paths result from the edge sharing and corner sharing RhO 6 network.

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92 In Chapter 6 we report on p type conductivity in zinc cobalt oxide (Zn Co O) films deposited at room temperature using pulsed laser dep osition (PLD). The X ray diffraction (XRD) scans of Zn Co O films deposited at room temperature shows no crystalline peaks. Film deposition is carried out at relatively high oxygen pressure. Both Seebeck coefficient and Hall coefficient measurements indica te that Zn Co O films can be p type oxide semiconductors. The properties of junction between Zn Co O and n type semiconducting oxides are also reported. 6.2 Experimental Details The Zn Co O films were deposited at room temperature on single crystal sapphir e (0001) substrates by pulsed laser deposition using a KrF 254 nm excimer laser and polycrystalline ZnCo 2 O 4 target. The laser energy was 185 mJ/pulse and the laser repetition rate was 2 Hz. The ablation target was fabricated from a mixture of commercial Co 3 O 4 and ZnO powders (Alfa Aesar). Powders were pressed into 1 inch diameter target and sintered at 1000 o C in the box furnace. The films were deposited at room temperature with oxygen as a background gas. Film thickness measured by a profilometer was 100 nm 200 nm. The Seebeck coefficient was obtained from the thermoelectric power measurement near room temperature. The Hall coefficient, electrical conductivity, carrier concentration, and Hall mobility were determined by Hall effect electronic measurement s (Lakeshore 7507) using a van der Pauw method at room temperature. Temperature dependent conductivity was obtained from a typical four point probe measurement. Optical transmission spectra were obtained by a Lambda 800 UV Vis spectrophotometer (Perkin Elm er instrument). Surface morphology was examined by using FE SEM with the acceleration voltage between 10 and 15 kV. For p n junction formation between Zn Co O and InGaZnO, the InGaZnO film was

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93 deposited by the RF sputtering system in Ar using an InGaZnO 4 t arget. The RF power was 150 W and the working pressure was 5 mTorr. The substrate temperature was room temperature. A shadow mask was used for Au metal deposition on the Zn Co O layer and for the InGaZnO/Ti/Au deposition in series. The dimension of the dep osition mask feature temperature in air. 6.3 Results and Discussion 6.3.1 P type Conduction Figure 6 1 shows the thermoelectric power measurement near 300 K for a Zn Co O film deposited at an oxygen pressure of 50 mT orr. The Seebeck coefficient is type conduction for the Zn Co O film. P type conduction is also confirmed by the positive Hall coefficient of 4.8210 2 cm 3 /C which is obta ined by using the four point van der Pauw method at room temperature. The electrical conductivity of the film deposited in 50 mTorr is determined to be 21.8 S cm 1 at room temperature. Figure 6 2 shows the temperature dependence of the electrical conductiv ity in the range from 160 K to 350 K for the film deposited at 50 mTorr. The thermal activation energy is determined to be 73 meV from the Arrhenius type plot indicating semiconductive behavior. This is consistent with the Fermi level lying close to the va lence band for hole conduction. The values of the electrical conductivity obtained for the Zn Co O film is larger than that for p type polycrystalline ZnCo 2 O 4 films [52] and that for p type amorphous zinc rhodium oxide [59] The optical properties of the films were examined using optical transmittance measurements.

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94 6.3.2 Structural P roperties The surface morphology was examined using Field Emission Scanning Electron Microscope (FE SEM). Figure 6 3, 6 4, and 6 5 shows FE SEM micrograph images of Zn Co O films deposited at various oxygen pressures. As shown in Figure 6 3 th e Zn Co O film deposited in 50 mTorr exhibits a smooth and continuous surface morphology. As shown in Figure 6 4 the film deposited in 80mTorr displays a granular morphology made of small si ze grains. As shown in Figure 6 5 the film deposited in 100 mTor r is composed of larger size grains. Note, however, that X ray diffraction of these films yields no apparent diffraction peaks. The microstructure observed for Zn Co O films deposited in oxygen pressures above 110 mTorr suggests possible segregation of the constituents. 6.3.3 Optical P roperties Figure 6 6 shows the optical absorption spectrum obtained from the optical transmission measurement in the range of 250 nm to 900 nm. The optical bandgap (E g ) of the Zn Co O film deposited in 100 mTorr is estimated to be ~2.3 eV, which is similar with the measured optical bandgap of the polycrystalline ZnCo 2 O 4 film [52] The band gap, E g of the Zn Co O films is independent of oxygen pressure used during deposition for the oxygen pressure range of 50 mTorr to 110 mTorr. 6.3.4 Transport P roperties The tr ansport properties of the films were examined as a function of deposition pressure. Figure 6 7 (mTorr) during growth. The lowest resistivity (4.5910 2 cm) is obtained for the film deposited cm) is obtained for the film deposited 110 mTorr. As the pressure of oxygen background gas increases, the

PAGE 95

95 electrical resistivity increases. The increase in apparent phase segregation with increasing oxygen may le ad to weakened connectivity between conductive regions. This discontinuity of Zn Co O film may play a role in the increase of electrical resistivity. Some Zn Co O films deposited in oxygen pressure greater than 110 mTorr have high resistivity on the scale of 10 4 cm. Figures 6 8 and 6 9 show the dependence of the carrier density and the mobility on the oxygen pressure used during deposition. The carrier density and mobility are obtained from Hall measurement at room temperature. In addition to the positive Seebe ck coefficient of the 50 mTorr film, positive Hall coefficients for films deposited in the range of 50 mTorr to 110 mTorr are consistent with the Zn Co O films being p type. The apparent hole carrier concentration decreases continually from 1.4910 20 cm 3 (50 mTorr) to 2.6010 16 cm 3 (110 mTorr) as oxygen pressure increases, which is consiste nt for films deposited at 50 mTorr may be an anomaly due to applying a single carrier channel model to a highly compensated material. As shown in Figure 6 9 the mob ility varies from 0.12 cm 2 /V s to 1.6 cm 2 /V s. 6.3.5 Oxide Heterojunction In order to further explore the properties of Zn Co O films, junctions were formed between Zn Co O films (p type layer) and amorphous InGaZnO film s [ 23] (n type layer). For the junction, the p type Zn Co O layer is deposited at room temperature by using PLD with an oxygen pressure of 100 mTorr. The n type InGaZnO layer is deposited by using radio frequency (RF) sputtering at room temperature. The elec trical resistivity of cm. The carrier density and the electron mobility in the InGaZnO film are determined by Hall measurements to be

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96 1.7810 17 cm 3 and 9.14 cm 2 /V s, respectively. Figure 6 10 shows the cur rent voltage characteristics of the junction. Figure 6 11 shows the junction structure, indicating that Au metal on the top of p type Zn Co O layer is used as a p side electrode and Au/Ti metals on the top of n type InGaZnO layer are used as a n side elect rode. Ohmic contact is observed between Au metal and p type Zn Co O film. For the device fabrication, a shadow mask is used for Au deposition and Au/Ti/InGaZnO deposition using RF spu ttering system. As shown in Figure 6 10, the junction is rectifying, with a threshold voltage of approximately 2.5 V. This is somewhat larger than the optical bandgap energy (~2.3 eV) of the Zn Co O film. The on off current ratio at 7 V is around 10 2 The I V curve in the positive voltage range for voltage exceeding turn on is linear, indicating that the p n heterojunction diode is non ideal due to highly resistive films of p type layers. This rectifying junction additionally confirms that the Zn Co O film is p type oxide semiconductor. 6.4 Summary W e have examined the properti es of semiconducting Zn Co O films deposited at room temperature. The Zn Co O film deposited at room temperature shows p type conduction as confirmed by the positive Seebeck coefficient and the positive Hall coefficient. The electrical conductivity as larg e as 21 S cm 1 is obtained from the film deposited in 50 mTorr. The increase in the electrical resistivity is observed as oxygen pressure increase. The carrier density decreases with increasing oxygen pressure. Semiconducting oxide p n junctions fabricated at room temperature show rectifying characteristics by using p type Zn Co O. The threshold voltage is somewhat larger than the bandgap energy. P type Zn Co O with controllable electrical properties may prove useful in thin film electronics and optoelectro nics.

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97 Figure 6 1. Thermoelectric power measurement of the Zn C o O film deposited in 50 mTorr with t

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98 Figure 6 2 Temperature dependence of the electrical conductivity of the Zn C o O film deposited in 50 mTorr with the thermal activation energy of 73 meV.

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99 Figure 6 3. FE SEM surface image (85,000 magnification) of the Zn Co O film deposited in 50 mTorr.

PAGE 100

100 Figure 6 4. FE SEM surface image (100,000 magnification ) of the Zn Co O film deposited in 80 mTorr.

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101 Figure 6 5 FE SEM surface image ( 100,000 magnification) of the Zn Co O film depo sited in 100 mTorr.

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102 Figure 6 6. Optical absorption spectrum of the Zn C o O film deposited in 100 mTorr with t he optical bandgap energy of ~2.3 eV.

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103 Figure 6 7 Electrical resistivity as a function of oxygen gas pressure in the range of 50 to 11 0 mTorr.

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104 Figure 6 8. Carrier density as a function of oxygen gas pressure in the range of 50 to 110 mTorr.

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105 Figure 6 9. Hall mobility as a function of oxygen gas pressure.

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106 Figure 6 10. Current density applied voltage (J V) curve of amor phous oxide p n junction using a p type Zn Co O film and a n type InGaZnO film.

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1 07 Figure 6 11. Schematic illustration of the p n junction structure.

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108 CHAPTER 7 METAL INSULATOR TRAN SITION AT THE INTERF ACE OF LAVO 3 /SRTIO 3 SUPERLATTICES GROWN ON TIO 2 TERMI NATED SRTIO 3 7.1 Scientific B ackground Perovskite related complex oxide thin films and heterostructures exhibit novel properties such as high temperature superconductivity, colossal magnetoresistance, ferroelectricity, and metal insulator transition. Advan ced growth techniques and well defined TiO 2 terminated SrTiO 3 substrates make it possible to fabricate perovskite based oxide heterostructures with the structural control at the nanoscale via molecular beam epitaxy (MBE) and pulsed laser deposition (PLD). Atomically controlled superlattices such as LaTiO 3 /SrTiO 3 [164] and LaAlO 3 / SrTiO 3 [161, 162, 173] include the n type LaO/TiO 2 interface with the formation of two dimensional electron gases (2DEGs) at the interface between complex insulating oxides. Formation of the LaAlO 3 /SrTiO 3 heterostr uctures on silicon (Si) wafers integrates oxide electronics with Si technology [249] Ohtomo and Hwang reported a conducting two dimensional electron gas layer at the heterointerface between two perovskite insulators, SrTiO 3 and LaAlO 3 [161, 162] The (LaO) + /(TiO 2 ) 0 interface between LaAlO 3 and SrTiO 3 shows the metallic behavior with the carrier density of ~10 17 cm 2 and the mobility of ~10 4 cm 2 V 1 s 1 due to polar discontinuity at the interface [161, 162, 2 50] Above a threshold thickness of LaAlO 3 layers, the electronic reconstruction occurs to compensate for the interfacial polar discontinuity and prevent s the catastrophic situation arising from the divergence of the electric potential in the limit of inf inite LaAlO 3 thickness [119, 250] A sheet carrier density is higher than what is expected, half an electron per unit cell (~3.2 10 14 cm 2 ), due to the creation of oxygen vacancies in the SrTiO 3 substrate during th e growth [161,

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109 163, 251, 252] M. Basletic et al. reported the carrier density profile in the LaAlO 3 /SrTiO 3 heterostructure by using a conducting tip atomic force microscope [253] To date, most recent experimental studies have focused on LaAlO 3 /SrTiO 3 superlattices. Further investigation of the interfacial polar discontinuity has been report ed between Mott insulator of LaVO 3 and band insulator SrTiO 3 Y. Hotta et al. [254] investigated polar discontinuity doping by electronic reconstruction at the heterointerfaces inducing n type and p type polar discontinuity between LaVO 3 and SrTiO 3 L. F. Kourkoutis et al. [255] observed the growth asymmetry at the polar LaVO 3 /SrTiO 3 heterointerfaces. In Chapter 7 we focus on the interfacial study of the LaVO 3 /SrTiO 3 heterostructures with the structural and electronic properties by x ray diffracti on, high resolution transmission electron microscopy with fast Fourier transform atomic force microscopy, and physical property measurement system Temperature dependence of the electrical resistance shows the metallic behavior with the upturn of the resi stance at low temperature. 7.2 Experimental Details The LaVO 3 (n unit cell)/ SrTiO 3 (n unit cell) superlattices were prepared on TiO 2 terminated or as received SrTiO 3 (100) single crystal substrates by pulsed laser deposition using a KrF excimer laser with ce ramic targets of LaVO 4 and SrTiO 3 The laser energy wa s 120 mJ /pulse and the laser repetition rate was 2 Hz. The LaVO 4 target wa s fabricated by mixing commercial La 2 O 3 and V 2 O 5 powders (Alfa Aesar) and the SrTiO 3 target wa s prepared from a SrTiO 3 powder (A lfa Aesar). In an air ambient, the LaVO 4 target and the SrTiO 3 target were sintered at 700 C for 10 hours and at 1000C

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110 for 12 hours, and at 1000 C for 12 hours, respectively. The growth temperatu re of the superlattices was 600 C and the oxygen pressure wa s 510 5 Torr for all processes. The long range periodicity and high crystallinity of the superlattice films were examined by x ray diffraction using a scan method. The cross sectional and topographic images were examined by using high resolution transmission electron microscopy (HRTEM, JEOL JEM 2010F) and atomic force microscopy (AFM, Dimension 3100). The transport properties of the LaVO 3 /SrTiO 3 sup erlattices such as temperature dependent resistance and current voltage characteristic were examined by using the q uantum design physical property measurement system (PPMS). 7.3 Results and Discussion 7.3.1 Structural Characterization Figure 7 ray diffraction profile of the heterostructure. LaVO 3 /SrTiO 3 superlattice peaks exhibit good regularity of the periodic structure. The zero order peak of the film is located at 22.85 degree next to the peak of (100) SrTiO 3 substrate with additional sa the superlattice shows no other peaks except the Bragg reflections of the substrate, constituent film, and satellites peaks [256] From equally spaced satellite reflections, the periodicity of the superlattice is evaluat ed to be 72. 5 using a diffraction evaluation of Schuller analysis [183, 257] which is slightly larger than the attempted period of 70.5 resulting from 9 monolayers of both LVO and STO layers The discrepancy may be caused by a deviation of the lattice constants of oxides from the bulk values when grown as thin film layers [258] Figure 7 2 shows a bright field HRTEM image of [(LaVO 3 ) 10 /(SrTiO 3 ) 10 ] 5 superlattice with the magnification of 500,000. For HRTEM analysis, a cross sectional

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111 sample is obtained by using focused ion beam (FIB) technique from the LaVO 3 /SrTiO 3 heterostructures along the [100] direction of SrTiO 3 The growth direction is from top to bottom. Between two different polar interfaces, the TiO 2 /LaO interface looks more diffuse than the VO 2 /Sr O interface. The VO 2 /SrO interface is more atomically abrupt on the LaVO 3 layer. This asymmetric interface may result from the presence of Sr surface segregation which results in cation interdiffusion at the interface [255] As Sr atoms are energetically preferred to be on the surface, a fraction of Sr atoms moves to th e surface during LaVO 3 growth on the SrTiO 3 layer. As a consequence, the SrTiO 3 layer growth on the LaVO 3 layer forms the VO 2 /SrO interface which shows more atomically sharp interface compared to the TiO 2 /LaO interface formed by the LaVO 3 growth on the SrT iO 3 layer. Note that polar interfaces lead to both electronic and structural reconstruction [164, 173] In addition, F igure 7 3 shows another HRTEM image of the superlattice with the magnification of 800,000. Figur e 7 4 (a), (b), (c) shows Fast Fourier Transform (FFT) of the images of the LaVO 3 layer, SrTiO 3 layer, and SrTiO 3 substrate. The periodic superlattice structure in HRTEM image gives rise to sharp spots in the resulting diffraction pattern. Diffraction patte rns from FFT show that the superlattice film is epitaxially grown along [100] zone axis on the (100) STO substrate. The surface morphology and roughness of the LaVO 3 /SrTiO 3 superlattice films are e xamined by AFM as shown in F igure 7 5 and 7 6 F igure 7 5 s hows that the surface of the oxide heterostructure is obtained from the [(LaVO 3 ) 10 /(SrTiO 3 ) 10 ] 5 superlattice which is grown on TiO 2 terminated STO substrate showing single terrace steps with the height of ~0.4 nm. The average root mean square (RMS) roughn ess is determined to

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112 be 0.11 nm over a 3 square micrometer area. F igure 7 6 image is observed from the [(LaVO 3 ) 10 /(SrTiO 3 ) 10 ] 5 superlattice grown on non etched STO substrate that includes both SrO and TiO 2 termination. The RMS value is evaluated to be 0.13 nm over a 3 square micrometer which is close to the value of F igure 7 5 indicating quite smooth surface. 7.3.2 Transport Property The transport properties of the LaVO 3 /SrTiO 3 heterostructures are examined as a function of temperature by van der Pauw resi stance measurement method using Quantum design PPMS Figure 7 7 shows the temperature dependent resistance of the [(LaVO 3 ) 5 /(SrTiO 3 ) 5 ] 5 superlattice suggesting that the interface between two insulating oxides exhibits metallic behavior. In the [100] growt h direction, LaVO 3 is composed of the alternating LaO and VO 2 layers which are charged +1 and 1 while SrTiO 3 consists of the alternating SrO and TiO 2 layers which are neutral in charge. The heterostructure forms polar interfaces such as the TiO 2 /LaO inte rface and the VO 2 /SrO interface The +1 valent LaO layer gives electrons to adjacent TiO 2 layers which increases the electron carrier density near the interface. Since the VO 2 /SrO interface is expected to be the p type interface and shows insulating behavi or [161, 254] the TiO 2 /LaO interface has a dominant effect on the metallic behavior arising from the electronic reconstruction caused by polar discontinuity. T he conducting ( La,Sr)TiO 3 and (La,Sr)VO 3 formation cau sed by interdiffusion does not have a significant effect on the metallic behavior at the interface [254] Note that we may not rule out the effect of growth induced oxygen vacancies at the surface (interface) of the SrTiO 3 substrate [161, 163, 252] because it is not easy to get fully oxidized interface by oxygen post annealing due to the LaVO 4 film formation [259] The system remains metallic in the range from 300 to 13.5 K with the

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113 As temperature reduces lower to 4 K, it sh ows insulating behavior which indicates temperature induced MI transition This insulating behavior may arise from the formation of LaTi 1 x V x O 3 near the TiO 2 /LaO interface. Note that the resistivity of LaTi 1 x V x O 3 phase shows an upturn at low temperature the upturn may arise from electron localization due to impurity scattering or grain boundary effects [260] The upturn may also be associated with magnetic 1 ) [261] The effect of the contact resistance on the insulating behavior at low temperature can not be disregarded, which is observed from current voltage characteristic. 7.4 Summary In conclusion we have epitaxially grown p eriodical LaVO 3 /SrTiO 3 superlattices and observed asymmetric interfaces in heterostructures. VO 2 /SrO interfaces show more atomically sharp than TiO 2 /LaO interfaces due to surface segregation effect. As deposited superlattices remain single terrace steps as TiO 2 terminated STO substrates do. Temperature dependent electrical resistance profiles show two dimensional electron gases at heterointerfaces with insulating behavior below 13.5K, indicating temperature induced metal insulator transition.

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114 Figure 7 1. XRD LaVO 3 /SrTiO 3 superlattice grown at 600C in 510 5 Torr of oxygen on a (100) SrTiO 3 substrate.

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115 Figure 7 2. HRTEM image ( 5 00,000 magnification) of [(LaVO 3 ) 10 /(SrTiO 3 ) 10 ] 5 superlattice LaVO 3 SrTiO 3 STO substrate

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116 Figure 7 3 HRTEM image ( 8 00,000 magnification) of [(LaVO 3 ) 10 /(SrTiO 3 ) 10 ] 5 superlattice LaVO 3 SrTiO 3 STO substrate

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117 Figure 7 4 Fast Fourier Transform (FFT) of the images of the (a) LaVO 3 layer, (b) SrTiO 3 layer, and (c) SrTiO 3 substrate (c) SrTiO 3 substrate (a) LaVO 3 layer (b) SrTiO 3 layer

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118 F igure 7 5 AFM Image of the [(LaVO3)10/(SrTiO3)10]5 superlattice grown on TiO 2 terminated S r TiO 3 substrate 500 nm (a) on TiO 2 ter minated STO

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119 F igure 7 6 AFM Image of the [(LaVO 3 ) 10 /(SrTiO 3 ) 10 ] 5 superlattice grown on non etched S r T iO 3 substrate 500 nm (b) on non etched STO

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120 Figure 7 7 Resistance of the [(LaVO 3 ) 5 /(SrTiO 3 ) 5 ] 5 superlattice as a function of temperature

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121 CHAPTER 8 CONCLUSIONS This dissertation has covered functional oxide materials such as transparent conductive oxides (TCOs) and functional com plex oxides with structural, electrical, and optical properties. InAlZnON has been investigated for its structural, electrical, and optical properties. Oxynitride thin films are prepared by co sputtering from aluminum oxide and indium zinc oxide targets at room temperature. Nitrogen is introduced into the films by using nitrogen gas in the plasma and nitrogen incorporation is measured by AES and XPS. Surface morphology is determined to be R a = 1.79 nm which is smooth for device application by AFM. Varying th e sputter target power leads to varying chemical composition of In, Zn, and Al with the ability to control electrical resistivity and carrier density of thin films. Varying nitrogen concentration in the plasma enables tuning carrier density. Temperature de pendence of the dc conductivity shows Arrhenius behavior with the activation energy of 62 meV. InAlZnON films are determined to be transparent with the optical band gap energy of 3.55 eV. ZnAlON has been investigated with interest in nitrogen incorporation effect on electrical and optical properties. ZnAlON thin films are prepared by co sputtering from aluminum oxide and zinc oxide targets at room temperature. Varying either sputter target power allows controlling resistivity and carrier density as well as tuning optical band gap energy. Increase in nitrogen volume ratio in the plasma also plays a role in decreasing aluminum atomic concentration, indicating the ability to control electrical and optical properties of nitrogen incorporated zinc aluminum oxide films. It is observed that

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122 increase of Zn/Al ratio in the films results in increase of resistivity and decrease of carrier density. P type conduction of zinc cobalt oxide films has been synthesized via pulsed laser deposition at room temperature with the b ackground gas of oxygen. Both positive Seebeck coefficients and positive Hall coefficients convince p type conductivity in zinc cobalt oxide thin films. While FE SEM images show granular morphology, X ray diffraction scan shows no crystalline peak, indicat ing as deposited films at room temperature may be amorphous. The electrical conductivity of the film deposited in 50 mTorr is determined to be around 21 S cm 1 at room temperature by using the thermoelectric power measurement and Hall effect measurement. A s oxygen background gas pressure increases, the electrical resistivity increases with decrease of carrier density. The microstructures of the films suggest that the connectivity between conductive regions gets weakened as oxygen pressure increases. Room te mperature fabricated oxide p n junction shows rectifying characteristic with a threshold voltage of 2.5 V, which additionally confirms zinc cobalt oxide thin films show p type conduction. Superlattices of LaVO 3 /SrTiO 3 have been investigated with interest i n metal insulator transition behavior at the heterointerface between LVO and STO. These superlattices show not only periodic structures, as measured by XRD, but also atomically controlled interfaces, as observed by high resolution TEM. It has been observed by AFM that superlattices grown on TiO 2 terminated STO substrate maintain single terrace steps with the RMS roughness value of 0.11 nm. Temperature dependent resistance shows metallic behavior with an upturn at low temperature. Possible reasons of tempera ture induced MI transition include LaTi 1 x V x O 3 formation near the TiO 2 /LaO

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123 interface and electron localization due to impurity scattering or grain boundary effects. The upturn may also be associated with magnetic ordering and the effect of contact resistan ce. Future research may focus on verifying grounds of insulat ing behavior at low temperature by magnetic property measurements as a function of temperature and infrared optical reflectivity (or conductivity) measurements.

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124 APPENDIX CRYSTAL SYSTEMS Table A 1 Atomistic parameters o f CuAl O 2

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125 Table A 2. Atomistic parameters of ZnCo 2 O 4

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126 Table A 3. Atomistic parameters of LaVO 3

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127 Table A 4. Atomistic parameters of SrTiO 3

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140 BIO GRAPHICAL SKETCH SeonHoo Kim was b orn and lived in Seoul, Republic of Korea. He received his Bachelor of Engineering degree cum laude in Materials Science and Engineering from Seoul National University in 2002. After graduation, he worked at two companies t o broade n background of semiconductor materials and ceramic processing through industrial experience. To continue his studies, he applied to the University of Florida and enrolled in the fall of 2006 in the Materials Science and Engineering. He conducted research with Dr. David Norton in the department of Materials Science and Engineering at the University of Florida since summer of 2007 and completed his doctoral dissertation in December of 2011.