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Development of Transition-Metal Doped Cu{sub 2}O and ZnO Dilute Magnetic Semiconductors

Permanent Link: http://ufdc.ufl.edu/UFE0021321/00001

Material Information

Title: Development of Transition-Metal Doped Cu{sub 2}O and ZnO Dilute Magnetic Semiconductors
Physical Description: 1 online resource (171 p.)
Language: english
Creator: Ivill, Mathew P
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2007

Subjects

Subjects / Keywords: cu2o, ferromagnetism, magnetoresistance, pld, spintronics, zno
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: The field of spintronics has recently attracted much attention because of its potential to provide new functionalities and enhanced performance in conventional electronic devices. Oxide materials provide a convenient platform to study the spin-based functionality in host semiconducting material. Recent theoretical treatments predict that wide band-gap semiconductors, including ZnO, can exhibit high temperature ferromagnetic ordering when doped with transition metals. This work focused on the possibility of using wide band-gap oxide semiconductors as potential spintronic materials. The structure, magnetic, and electronic transport properties of transition-metal doped ZnO and Cu2O were investigated. Mn and Co were used as transition metal dopants. Thin films of these materials were fabricated using pulsed laser deposition (PLD). The Mn solubility in Cu2O was found to be small and the precipitation of Mn-oxides was favored at high growth temperatures. Phase pure Mn-doped Cu2O samples were found to be non-magnetic. Samples were p-type with carrier concentrations on the order of 1014 & #8211;1016 cm & #8722;3. The effects of carrier concentration on the magnetic properties of Mn-doped ZnO were studied using Sn and P as electronic codopants. Sn acts as an n-type dopant providing extra electrons to the ZnO. P acts as a p-type dopant that supplies excess holes to compensate the native electron concentration in ZnO. The electron concentration was decreased using P, but the films remained n-type. An inverse correlation was found between the ferromagnetism and the electron concentration; the ferromagnetic coupling between Mn spins increased with decreasing electron concentration. The nature of ferromagnetism in Co-doped ZnO was also investigated. Ferromagnetism was found in films deposited at 400 degrees Celsius in vacuum, while films deposited in oxygen or at higher temperatures were non-magnetic. Films deposited under vacuum had rather high electron concentrations and were presumably doped with oxygen vacancies. The Co-doped films also exhibited peculiar magnetoresistance (MR) that had a strong dependence on the carrier concentration. At low temperatures, a progression from positive to negative MR was observed with increased electron concentration as the films crossed over the metal-to-insulator transition (MIT).
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 Mathew P Ivill.
Thesis: Thesis (Ph.D.)--University of Florida, 2007.
Local: Adviser: Norton, David P.

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2007
System ID: UFE0021321:00001

Permanent Link: http://ufdc.ufl.edu/UFE0021321/00001

Material Information

Title: Development of Transition-Metal Doped Cu{sub 2}O and ZnO Dilute Magnetic Semiconductors
Physical Description: 1 online resource (171 p.)
Language: english
Creator: Ivill, Mathew P
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2007

Subjects

Subjects / Keywords: cu2o, ferromagnetism, magnetoresistance, pld, spintronics, zno
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: The field of spintronics has recently attracted much attention because of its potential to provide new functionalities and enhanced performance in conventional electronic devices. Oxide materials provide a convenient platform to study the spin-based functionality in host semiconducting material. Recent theoretical treatments predict that wide band-gap semiconductors, including ZnO, can exhibit high temperature ferromagnetic ordering when doped with transition metals. This work focused on the possibility of using wide band-gap oxide semiconductors as potential spintronic materials. The structure, magnetic, and electronic transport properties of transition-metal doped ZnO and Cu2O were investigated. Mn and Co were used as transition metal dopants. Thin films of these materials were fabricated using pulsed laser deposition (PLD). The Mn solubility in Cu2O was found to be small and the precipitation of Mn-oxides was favored at high growth temperatures. Phase pure Mn-doped Cu2O samples were found to be non-magnetic. Samples were p-type with carrier concentrations on the order of 1014 & #8211;1016 cm & #8722;3. The effects of carrier concentration on the magnetic properties of Mn-doped ZnO were studied using Sn and P as electronic codopants. Sn acts as an n-type dopant providing extra electrons to the ZnO. P acts as a p-type dopant that supplies excess holes to compensate the native electron concentration in ZnO. The electron concentration was decreased using P, but the films remained n-type. An inverse correlation was found between the ferromagnetism and the electron concentration; the ferromagnetic coupling between Mn spins increased with decreasing electron concentration. The nature of ferromagnetism in Co-doped ZnO was also investigated. Ferromagnetism was found in films deposited at 400 degrees Celsius in vacuum, while films deposited in oxygen or at higher temperatures were non-magnetic. Films deposited under vacuum had rather high electron concentrations and were presumably doped with oxygen vacancies. The Co-doped films also exhibited peculiar magnetoresistance (MR) that had a strong dependence on the carrier concentration. At low temperatures, a progression from positive to negative MR was observed with increased electron concentration as the films crossed over the metal-to-insulator transition (MIT).
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 Mathew P Ivill.
Thesis: Thesis (Ph.D.)--University of Florida, 2007.
Local: Adviser: Norton, David P.

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2007
System ID: UFE0021321:00001


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DEVELOPMENT OF TRANSITION-METAL DOPED Cu20 AND ZnO DILUTE
MAGNETIC SEMICONDUCTORS



















By

MATHEW P. IVILL


A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL
OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT
OF THE REQUIREMENTS FOR THE DEGREE OF
DOCTOR OF PHILOSOPHY

UNIVERSITY OF FLORIDA

2007

































( 2007 Mathew P. Ivill




































To my family and friends, especially to my parents and Kathryn, for their love and


support.









ACKNOWLEDGMENTS

Graduate school has proven to be quite a wild ride, filled with moments of great

accomplishment, balanced with moments of insane frustration. Luckily, I've had the

privilege of collaborating with amazing people who have helped me while I worked on this

dissertation. It would be impossible to list everyone who has helped. Nevertheless, I will

do my best to name those who made my years (and all the many hours and late nights in

the lab) at the University of Florida an educational, memorable, and enjoyable experience.

First of all, I thank my advisor, Dr. David Norton, for taking me into his research

group. I am deeply grateful for his constant encouragement and support throughout my

graduate studies. He has helped me grow both academically and personally, and has been

an exceptional role model. Without his help this dissertation would not be possible. I also

thank the other members of my committee: Dr. Stephen Pearton, Dr. Cammy Abernathy,

Dr. Art Hebard, and Dr. Susan Sinnott for their help and guidance.

I am grateful to Dr. Brent Gila, a person who possesses a vast amount of practical

knowledge. I learned something new every time we chatted in lab.

I've spent many enjoyable hours working with my fellow group members. I am happy

to thank them all, both past and present, for their support and friendship in and out of

the lab, including Dr. Seh-Jin Park, Dr. Beong-Seong Jeong, Dr. Hyung-jin (Johnny) Bae,

Dr. Kyunghoon Kim, Dr. Jennifer Sigman, Dr. George Erie, Dr. Yaunjie Li, Dr. Seemant

Rawal, Mitesh Patel, Vijayram Varadarajan, Micheal Jones, Hyunsik Kim, Li-Chia Tien,

Patrick Sadik, Charlee Callender, Lii-Cherng (Daniel) Leu, Joe Cianfrone, Fernando Lugo,

Ryan Pate and Zivin Park. I am especially grateful to both Dr. Young-Woo Heo and Dr.

Yongwook Kwon, both of whom took me under their wing when I started graduate school,

and whose kindness and personality have been a great inspiration. Young-Woo taught me

many of the lab techniques I used throughout grad school, including how to deposit films

using PLD, and was always kind enough to lend a hand with my research, even when he

was extremely busy with his own. Yongwook also taught me many things; he introduced









me to LabView programming, electrical characterization, and also taught me some Kendo

along the way! I wish them all great success and happiness.

I thank the members of Dr. Abernathy's research group, including Dr. Rachel

Frazier, Dr. Jennifer Hite, and Dr. Gerald Thaler for discussions pertaining to magnetic

semiconductors. I also thank my collaborators and friends from the physics department,

Ritesh Das, Dr. Josh Kelly, and Dr. Ryan Rairigh for providing the SQUID measurements

presented in this dissertation and their insightful discussions interpreting the data.

Also Rajiv Misra who offered some very useful advice regarding electrical transport

measurements.

I thank Dr. John Budai and Dr. Matthew Chisholm for collaboration with high-resolution

XRD and TEM on selected samples. I also thank Dr. Valentin Craciun and Eric Lambers

from the Major Analytical Instrumentation Center (MAIC) for their help with material

characterization using XRD and XPS.

I thank Dr. Sarah Russell Gonzalez, one of the friendliest people I know, for

introducing me to the LaTeX typesetting program in which this dissertation was written.

I thank my parents and my brother for their love and unwavering support. They've

always been around when I needed them and they continue to support me in my decisions.

I would have never made it this far without them and I am very fortunate to have them in

my life.

Finally, I thank my fiancee, Kathryn Kennedy, for her endearing love and support

(and hours of proofreading). She was always there to cheer me up when experiments went

wrong, to inspire me when life became overly frustrating, and to celebrate with me when

things finally went right.









TABLE OF CONTENTS
page

ACKNOWLEDGMENTS ................... ............. 4

LIST O F TA BLES . . . . . . . . . 9

LIST O F FIG U RES . . . . . . . . . 10

A B ST R A C T . . . . . . . . . . 13

CHAPTER

1 INTRODUCTION .. .. .. .. .. .. .. .. .. .. .. .. 15

1.1 C harge and Spin . . . . . . . . 15
1.2 W hat is Spintronics? . . . . . . . 16
1.3 Dilute Magnetic Semiconductors ................... ..... 17

2 REVIEW OF DILUTE MAGNETIC SEMICONDUCTORS ........... 19

2.1 Magnetic Semiconducting Materials: A Short History ............ 19
2.2 DMS Theory: The Physical Origins of Ferromagnetism in DMS ...... 22
2.2.1 Dietl's Mean-field Theory ...... ....... . .. 22
2.2.2 First-principles Design: DFT Calculations . . 23
2.2.3 Ferromagnetism in Disordered Alloys .. . . 24
2.2.4 Ferromagnetism in a Spin-split Conduction Band . . .. 25
2.3 Experimental Progress in ZnO DMS .............. ..... 26
2.3.1 Mn-doped ZnO ................... . . 27
2.3.2 Co-doped ZnO .................... . . 28

3 THIN FILM DEPOSITION AND EXPERIMENTATION . . 34

3.1 PLD as a Tool for Thin Film Oxides ............. ..... 34
3.2 PLD System Used for This Work . ........... ... 35
3.2.1 The Growth Environment . ........... ... 35
3.2.2 The Laser Source .................. . ... 36
3.3 Typical Growth Procedure ................. . ... 37
3.3.1 Substrate Preparation ............... . . ... 37
3.3.2 Thin Film Growth .. .. . ... 38
3.4 Fabricating PLD Ablation Targets . .......... ... 39
3.5 Thin Film Characterization ................ . ... 40
3.5.1 X-ray Diffraction . . . 40
3.5.2 Magnetoresistance and Hall Effect Measurements . . .. 41
3.5.3 Electron Dispersive Spectroscopy (EDS) ....... . . 42
3.5.4 Optical Absorption ................. . .... 43
3.5.5 SQ U ID . . . . . ... . . ... 45









4 PROPERTIES OF Mn-DOPED Cu20 DMS .................. .. 49

4.1 Introduction .................... ............. 49
4.2 Cu2O: A Wide-bandgap P-type Semiconductor ....... . . 49
4.3 Experimental ... .. ............................. 51
4.4 Results and Discussion ..... .... 52

5 PROPERTIES OF ZnO CODOPED WITH Mn AND Sn . .... 70

5.1 Introduction . . . . . .. . . 70
5.2 Experim ental . . . . . ... . ... 71
5.3 Results and Discussion ..... .... 71

6 PROPERTIES OF ZnO CODOPED WITH Mn AND P . .... 82

6.1 Introduction . . . . . .. . . 82
6.2 Experim ental . . . . . ... . ... 82
6.3 Results and Discussion ..... .... 83

7 PROPERTIES OF COBALT-DOPED ZnO .. . 100

7.1 Introduction . . . . . ... . . 100
7.2 Experim ental . . . . . ... . . 100
7.3 Results and Discussion ..... .. . 101
7.3.1 Chemical Composition . .............. . 101
7.3.2 Structure and Phase Analysis . ........ . 101
7.3.2.1 Films with precipitation .......... . . 101
7.3.2.2 Films without precipitation ........ . . 102
7.3.3 Optical Properties ................... . . 103
7.3.3.1 Optical absorption . .......... 103
7.3.3.2 Photoluminescence . .......... 105
7.3.4 Magnetic Characterization . ........... . 106
7.3.5 Electrical Transport .................. . . 107
7.3.5.1 H all effect . . . . . . 107
7.3.5.2 M agnetoresistance . .......... 110
7.3.6 Effects of Annealing ................. . . 113

8 CONCLUSION . .. .. .. .. .. .. .. .. ... . 139

8.1 M n-doped Cu20 . . . . . . . . 139
8.2 M n-doped ZnO . . . . . ... . . 140
8.3 Co-doped ZnO . . . . . ... . . 142

APPENDIX

A HALL EFFECT SYSTEM AND EQUIPMENT ....... . . 144

A .1 Introduction . . . . . ... . . 144
A.2 Hall Effect Equipment .................... . . 144









A.3 Sample Geometry and Measurement Technique ...... . . 147
A.4 Hall Effect Method ................. . . .... 150
A.5 Limitations and Tips for Better Measurements ....... . . 151

REFERENCES ................ ................. 157

BIOGRAPHICAL SKETCH .................... . . ..... 171









LIST OF TABLES


Table


2-1 List of ZnO-based DMS experimental results . ......... 33

5-1 Resistivity as a function of Sn content in codoped ZnO:3%Mn films. ...... ..76

7-1 Possible cobalt-induced secondary phases. .................. .... 116

7-2 Transport data for a 30%Co-doped ZnO film with cobalt precipitation. . 132









LIST OF FIGURES


Figure


2-1 Schematic representation of magnetic exchange between two Mn ions mediated
by a delocalized hole. . ................ . .... 30

2-2 Predicted Curie temperatures based on Dietl's calculations . . 31

2-3 Illustration of bound magnetic polarons. ............. ...... 32

3-1 Schematic of pulsed laser deposition system. . ......... 47

3-2 Visualization of Bragg's law for x-ray diffraction. ........ . ..... 48

4-1 Cubic unit cell of Cu20. .................. ........... 57

4-2 Phase Stability curves for the Cu-Cu20-CuO system. ....... . . 58

4-3 X-ray diffraction data for epitaxial Cu20 on (001) MgO ............. ..58

4-4 X-ray diffraction data for Mn-doped Cu20 films grown on (001) MgAl204 in an
oxygen pressure of ImTorr. .................... . .... 59

4-5 X-ray diffraction data for Mn-doped Cu20 films grown on (001) MgAl204 in an
oxygen pressure of 0.lmTorr. ..... ..... 60

4-6 X-ray diffraction data for Mn-doped Cu20 films grown on (001) MgAl204 in
vacuum . . . . . . . . . 61

4-7 Phase assemblage for films grown under different conditions. . . ... 62

4-8 Magnetic behavior for an epitaxial Mn-doped Cu20 film grown at 300'C and
Im Torr of oxygen. .. .. .. .. .. .. .. .. .. . .. 63

4-9 Magnetic behavior for MgAl204 substrate. .................. .. 64

4-10 Magnetic behavior for an epitaxial Mn-doped Cu20 film .. . 65

4-11 Low temperature photoluminescence spectra for Mn-doped Cu20 films. . 66

4-12 Transport data for 1% Mn-doped Cu20 films. ................. 67

4-13 Temperature-dependent transport data for 1% Mn-doped Cu20 film. . 68

4-14 Field-varying transport measurements for a 1% Mn-doped Cu20 film. . 69

5-1 X-ray diffraction of ZnO films codoped with Mn and Sn ............. ..77

5-2 X-ray diffraction of an epitaxial ZnO film doped with 3%Mn and 0.1%Sn . 78

5-3 XPS spectra for ZnO:3%Mn film codoped with 0.01%Sn. ............ ..79









5-4 Plot showing the dependence of the coercive field on Sn concentration at different
SQUID measurement temperatures. .................. ..... 80

5-5 Magnetization measured at 300K for epitaxial ZnO:3%Mn films that are codoped
with 0.001% Sn, 0.01%Sn, 0.1% Sn, and no Sn. ......... . ..... 81

6-1 X-ray diffraction of ZnO films codoped with Mn and P both before and after
annealing . . . . . .. . . .. 90

6-2 High-resolution w-rocking curves on ZnO films codoped with Mn and P before
and after annealing. ................... ..... . .... 91

6-3 AFM scans on epitaxial ZnO:3%Mn, 2%P films before and after annealing. 92

6-4 Resistivity and carrier concentration behavior of P-doped ZnO films. . 93

6-5 Transport data for the as-deposited ZnO:3%Mn, 2%P film at 300K. ...... ..94

6-6 XPS spectra for ZnO:3%Mn film codoped with 2% P. ............. 95

6-7 Room temperature optical transmission for ZnO:3%Mn, 2%P films ...... ..96

6-8 Room temperature SQUID measurements for epitaxial ZnO:3%Mn,2%P films 97

6-9 SQUID measurement at 10K for epitaxial ZnO:3%Mn, 2%P films before and
after annealing. . .. .. .. .. .. .. .. .... . .... 98

6-10 Field-cooled and zero field-cooled magnetization measurements for a ZnO:3%Mn,
2%P film annealed at 600C in 02 ................. .. .... 99

7-1 EDS results for a select number of films grown under different conditions. . 116

7-2 XRD scans for a series of films grown in vacuum at 400C. . . .. 117

7-3 TEM micrographs of a sample doped with 30%Co ...... . ....... 118

7-4 Convergent beam TEM diffraction patterns of ZnO film doped with 30%Co 119

7-5 XRD scans for ZnO films doped with 30%Co. ................. 120

7-6 Thermodynamic predominance diagram for cobalt oxides. . . 121

7-7 UV-Vis transmission of Co-doped ZnO films ......... . 122

7-8 Optical band-gaps of Co-doped ZnO films. . ......... 123

7-9 PL results for Co-doped ZnO films .................. ..... .. 124

7-10 SQUID magnetization curves for Co-doped ZnO films deposited at 400'C in
vacuum . . . . . . . . . 125

7-11 SQUID magnetization curves for Co-doped ZnO films ...... . . 126










7-12 Anomalous Hall effect in 30%Co-doped ZnO. .. . 127

7-13 The magnetoresistance of an undoped ZnO film. ....... . . ..... 128

7-14 The magnetoresistance of 5%Co-doped ZnO films. ...... . . ..... 129

7-15 Magnetoresistance of 30%Co-doped ZnO films. ........ . ..... 130

7-16 Magnetoresistance of 30%Co-doped ZnO film at temperatures between 10K to
100K ... .. .. .. ... .. .. .. .. .. .. 131

7-17 Temperature dependent resistivity measurements for 5% and 30% Co-doped
ZnO film s .. .. .. .. .. .. ... .. .. . . . 132

7-18 Hall resistivity and magnetoresistance for a 30%Co-doped ZnO film with cobalt
precipitation ........................ .. . . .... 133

7-19 SEM micrographs at different magnifications of the surface of a Co-doped ZnO
film that has been annealed in H2/Ar at 500C for 60min. . . ... 134

7-20 0-20 XRD for a 30%Co-doped ZnO film before and after annealing in forming
gas at 500 C. ................... ....... . . .... 135

7-21 High-resolution XRD scans for 30%Co-doped ZnO film that has been annealed
in forming gas at 500C. .................... . . ..... 136

7-22 SQUID magnetometry for a 30%Co-doped ZnO film before and after annealing
in forming gas at 500C. .................... . . ..... 137

7-23 SQUID magnetometry for 30%Co-doped ZnO films deposited under differerent
conditions .. . . . . .. . . . . 138

A-i The resistivity and Hall measurement system. ................. .. 154

A-2 Circuit diagrams for 2-point and 4-point resistivity measurements. . ... 155

A-3 Circuit shunt capacitance and settling time ............. . .. 156









Abstract of Dissertation Presented to the Graduate School
of the University of Florida in Partial Fulfillment of the
Requirements for the Degree of Doctor of Philosophy

DEVELOPMENT OF TRANSITION-METAL DOPED Cu20 AND ZnO DILUTE
MAGNETIC SEMICONDUCTORS

By

Mathew P. Ivill

August 2007

Chair: Dr. David Norton
Major: Materials Science and Engineering

The field of spintronics has recently attracted much attention because of its

potential to provide new functionalities and enhanced performance in conventional

electronic devices. Oxide materials provide a convenient platform to study the spin-based

functionality in host semiconducting material. Recent theoretical treatments predict

that wide band-gap semiconductors, including ZnO, can exhibit high temperature

ferromagnetic ordering when doped with transition metals. This work focused on the

possibility of using wide band-gap oxide semiconductors as potential spintronic materials.

The structure, magnetic, and electronic transport properties of transition-metal doped

ZnO and Cu20 were investigated. Mn and Co were used as transition metal dopants. Thin

films of these materials were fabricated using pulsed laser deposition (PLD).

The Mn solubility in Cu20 was found to be small and the precipitation of Mn-oxides

was favored at high growth temperatures. Phase pure Mn-doped Cu20 samples were

found to be non-magnetic. Samples were p-type with carrier concentrations on the order of

10141016 cm3.

The effects of carrier concentration on the magnetic properties of Mn-doped ZnO were

studied using Sn and P as electronic codopants. Sn acts as an n-type dopant providing

extra electrons to the ZnO. P acts as a p-type dopant that supplies excess holes to

compensate the native electron concentration in ZnO. The electron concentration was

decreased using P, but the films remained n-type. An inverse correlation was found









between the ferromagnetism and the electron concentration; the ferromagnetic coupling

between Mn spins increased with decreasing electron concentration.

The nature of ferromagnetism in Co-doped ZnO was also investigated. Ferromagnetism

was found in films deposited at 400C in vacuum, while films deposited in oxygen or at

higher temperatures were non-magnetic. Films deposited under vacuum had rather high

electron concentrations and were presumably doped with oxygen vacancies. The Co-doped

films also exhibited peculiar magnetoresistance (MR) that had a strong dependence on

the carrier concentration. At low temperatures, a progression from positive to negative

MR was observed with increased electron concentration as the films crossed over the

metal-to-insulator transition (MIT).









CHAPTER 1
INTRODUCTION

1.1 Charge and Spin

The microelectronics industry has been at the forefront of information technology,

providing the devices necessary for fast and efficient information processing and storage.

The integrated circuit, packed with transistor arrays fabricated from semiconducting

silicon material, has remained the workhorse of the information processing industry over

the past 50 years. These devices utilize the charge properties of electrons (or holes) to

control the flow of current through the circuit. By the proper organization of transistors

on a chip, information can be computed by the digital logic of electronic charge. Increase

in the raw computational power and speed of these transistors over the past 50 years has

been propelled by one important trend-miniaturization. Steadily following the early

prediction by Gordon Moore (commonly called Moore's Law) miniaturization has led to

the number of transistors on a chip to double about every 18 months. Intel's cutting-edge

processors are currently fabricated with 65nm line-widths, and next generation 45nm

architectures are hitting the market soon [1].

On the other hand, the technology of information storage relies on another fundamental

property of electrons: the quantum mechanical electron spin. Magnetism in solids is a

direct consequence of the spin property of electrons. Electrons have two available spin

states, spin-up and spin-down. Permanent magnetic materials contain an imbalance in the

number of spin-up and spin-down electrons. Binary information may be encoded in the

form of non-volatile magnetic domains within the grains of ferromagnetic material. As the

size of these domains shrink, more information can be stored per unit area of material.

Increases in magnetic storage density have occurred at rates faster than any other industry

in history, with storage density increasing over 50 million times since the creation of the

first hard disk drive in 1957 [2]. The areal density of drives continues to increase at a rate

of about 100% per year and contemporary state-of-the-art disks have been developed with









densities over 100 Gbits/in2 [3]. Advances in perpendicular recording are projected to

push areal densities even higher by deterring the superparamagnetic limit to 1Tb/in2, and

research into novel recording schemes could push densities even higher [4, 5]. While the

charge and spin properties of the electron have separately spawned two of the most rapidly

improving technologies of our time, little has been done to combine the two to function

simultaneously in the same material, an idea that may lead to further enhancements in

information technology beyond the limits of miniaturization.

1.2 What is Spintronics?

The nascent field of spin electronics (or spintronics) stands to assimilate these two

fundamental properties of the electron-charge and spin-to form the basis for a new

class of device design [6-10]. Operating by the manipulation, transport, and detection of

charge carrier spins, spintronics is expected to improve upon traditional electronic and

photonic devices, allowing for enhancements in the form of reduced power consumption,

faster device operation, and new forms of information computation. Spintronics may

lead to devices such as spin-polarized LEDs, spin-FETs, and spin-based qubits for

quantum computers. Increased functionalities are also expected, such as integrated

magnetic/electronic operations on the same chip.

Currently, very few spintronic devices have appeared on the market but have already

made an astounding impact on technology. For example, metallic-multilayered structures

displaying large amounts of magnetoresistance (so-called Giant Magnetoresistance) have

replaced conventional hard-disk read heads, leading to huge increases in hard-disk storage

[6]. These devices consist of a sandwich structure where layers of ferromagnetic and

non-ferromagnetic material are alternately stacked. The resistance through the device

depends on the relative magnetic orientation of the ferromagnetic layers. When the layers

are magnetically aligned so that their directions of magnetization point opposite to one

another (180 degrees misalignment), the resistance through the device is large. Conversely,

when the magnetic layers are aligned parallel, the resistance is reduced. These sensors









have become the prominent form of solid-state magnetic sensing technology. As the size of

magnetized bits shrink and areal densities continue to increase, sensors must become even

more sensitive to the small changes in magnetization. The advent of GMR technology has

become a multi-billion dollar industry and revolutionized the read heads used in hard disk

technology.

1.3 Dilute Magnetic Semiconductors

Of particular interest is the creation and control of spin-polarized currents in

semiconducting material. Ferromagnetic semiconductors provide easier integration

of spintronics into existing semiconductor devices. For instance, highly efficient spin

injection is possible between semiconductor/semiconductor interfaces, whereas only spin

polarizations of a few percent are possible between a ferromagnetic metal/semiconductor

interface due to the conductivity mismatch [11]. Ferromagnetic semiconductors allow the

'tools' of conventional semiconductor technology to be utilized. These tools include p-n

junctions and heterostructures, which provide a convenient platform for a wide variety of

devices that allow for electronic gain and light emission. However, in order to fully realize

semiconductor-based spintronics, significant challenges related to the lifetime, control, and

detection of spin polarized carriers in semiconductors must be addressed. Materials that

can retain their ferromagnetism comfortably above room temperature are crucial to the

practical application of spintronic devices.

These reasons have created interest in developing a class of materials known as dilute

magnetic semiconductors (DMS). DMSs are semiconductors doped with a few percent of

magnetic atoms. The magnetic atoms occupy lattice sites and induce ferromagnetism

in the otherwise non-magnetic semiconductor host. DMSs typically have ordering

temperatures much lower than room temperature. There has been some success in

attaining room temperature ferromagnetism in various semiconductors, but the results

are non-reproducible among research groups and contradictory to some theoretical

predictions. Thus there is a great research opportunity for the study of ferromagnetism









in semiconductors for both real-world applications and for the added knowledge to

fundamental physics.









CHAPTER 2
REVIEW OF DILUTE MAGNETIC SEMICONDUCTORS

The idea of magnetism in semiconductors is not new. The question of whether

charge and spin can coexist in the same material to enhance material properties has been

addressed for many years. While magnetism in metallic and insulating materials was well

known, the possibility of magnetic ordering in semiconductors was not discovered until the

mid 1900s. From rather difficult beginnings, the field of magnetic semiconductors has seen

significant progress, especially in the advancement of magnetically doped semiconductors.

Significant challenges still remain related to the preparation and growth of materials,

understanding the physical origins of ferromagnetism in semiconductors, and raising the

magnetic ordering temperatures, just to name a few. This section will first provide a

glimpse into the history of experimental progress surrounding magnetic semiconductors.

Next, a review of the leading theories that delve into the physical origins of magnetic

coupling in DMS is given. Lastly, a brief review of the recent progress in magnetically

doped ZnO is presented.

2.1 Magnetic Semiconducting Materials: A Short History

The story of magnetic semiconductors originates from humble beginnings. In the

1960s and 1970s, semiconducting behavior in ferromagnetic material was uncovered with

the discovery of the europium chalcogenides (EuO) and chromium spinels (CdCr2S4,

CdCr2Se4). These materials are true ferromagnetic semiconductors in the sense that they

have magnetic atoms "built-into" the crystal sublattice. The elegant interplay between

band electrons and localized magnetic ions in these materials brought about extensive

research and scientific interest into the field. However, these materials have not progressed

beyond the field of academic research for several reasons [12, 13]. First of all, their crystal

structures are incompatible with conventional semiconductors, like Si and GaAs, making

their integration with contemporary electronics difficult. The synthesis of these materials

is also cumbersome and hard to reproduce, making industrial production of the crystals









expensive. And lastly, low ferromagnetic ordering temperatures (Tc < 100K) make them

less attractive for practical applications.

Moving in a slightly different direction, other work has focused on making non-magnetic

semiconductors magnetic by doping them with small amounts (typically a few percent)

of magnetic atoms. This class of materials has attracted renewed interest in the

development of magnetic semiconductors. Such compounds are known as diluted magnetic

semiconductors, or DMS, because of the dilute concentrations of magnetic impurities.

Notice, these materials are fundamentally different than the Eu chalcogenides and Cr

spinels since the magnetic atoms are artificially added into the lattice; the magnetic atoms

are not a part of the periodic crystal structure of the parent material.

Early studies of DMS materials began with Mn-doped II-VI alloys of the form

A"1 MnBvJ (where A" = Zn, Cd, Hg and BV1 = S, Se, Te). These materials were

heavily studied in the 1980s and comprehensively reviewed by Furdyna [14]. It is

worthwhile to review some aspects of these materials since ZnO also belongs to the

II-VI family of semiconductors. The ternary nature of these compounds makes them

amenable to tuning the lattice and band parameters by varying alloy composition,

making them an attractive candidate for the preparation of heterostructure devices. The

alloys crystallize into either the zinc-blende or wurtzite structure and are formed by sp3

tetrahedral bonding, incorporating the valence s-electrons from the group II metal and the

p-electrons from the group VI element. Elemental Mn has a half-filled 3d-shell and two

valence (4s2) electrons. Manganese atoms may substitute on the group II sites as Mn+2

by giving up these two valence electrons. High solubilities of Mn in the host materials

while maintaining the zinc-blende or wurtzite structures are possible, which is thought to

arise from the chemical similarity of Mn+2 to the group II element. The 3d-shell of Mn is

exactly half-filled and requires substantial energy to add an electron; this makes the 3d5

orbit act chemically similar to a 3d10 orbit. The magnetic properties of these alloys are

dictated by the exchange interactions between local atomic moments (provided by the









Mn) and the sp-band electrons, and have dramatic effects on the optical and electrical

properties of the material, such as giant Faraday rotation and bound magnetic polaron

formation. Driven mostly by superexchange mechanisms-an indirect exchange interaction

mediated through the anion-these systems exhibit high temperature paramagnetism,

low temperature spin-glass phase, and type III antiferromagnetic ordering [14]. Neutron

diffraction studies show that the antiferromagnetic ordering of these structures is limited

to short ranges, implying that the magnetic ordering is confined to the formation of

small cluster regions [14]. Recently, however, ferromagnetic ordering has been achieved

in low-dimensional quantum wells driven by hole-mediated exchange, but with low Curie

temperature (Tc < 2K) [15]. An additional obstacle to the practical applicability of II-VI

material is the capability of doping the material both n-type and p-type (bipolar doping).

Again, these materials are not practical since the materials that did show ferromagnetism

were restricted to low ordering temperatures (Tc just a few degrees above OK).

In the early 90s, a technological breakthrough in the advancement of DMS occurred

with the discovery of ferromagnetism up to -35K in Mn-doped InAs [16-18]. InAs

is an established III-V compound semiconductor material. Transition metal species

are known to have very low solubility in host III-V materials, but the problem was

overcome by non-equilibrium epitaxial growth using low temperature molecular beam

epitaxy (LTMBE). III-V materials find widespread application in the electronics and

optoelectronics industries as high-speed digital devices, visible and infra-red light-emitting

diodes and lasers, and magnetic sensors. The demonstration of ferromagnetism in InMnAs

offered the intriguing opportunity to study spin-based phenomena in these well established

semiconductor devices.

Eventually, the success of LTMBE growth of InMnAs led to the development

Mn-doped GaAs DMS. Segregation of Mn secondary phases, namely the MnAs phase, was

suppressed using low temperature growth (Tg=250'C). If, however, the temperature was

raised or the Mn flux was too high, phase segregation could occur. Mn acts as an acceptor









dopant when substituted on the group III sites leading to high hole concentrations which,

as explained later, is a necessity for ferromagnetism in the material. The coupling of the

charge and spin-based processes have been repeatedly proven in GaMnAs, including the

realization of spin-polarized light emission [19] and electrical and optical control over

the ferromagnetism [20]. Unfortunately, despite efforts to raise it, GaMnAs is limited by

its low Curie temperature of 170K (well below room temperature). Raising the Curie

temperature has been the biggest challenge for GaMnAs based DMS.

2.2 DMS Theory: The Physical Origins of Ferromagnetism in DMS

Understanding the physical mechanism behind magnetic ordering in DMS materials

is an essential ingredient to their further development. Indeed, if both a conceptual and

quantitative foundation to the origin of ferromagnetism in these materials is developed,

they may provide the direction necessary to a successful recipe for the fabrication of higher

Tc materials. At the present time, however, there is an incomplete understanding of the

origin of ferromagnetism in DMS material and the subject remains an issue of active

debate. This section will discuss the contemporary theories on the subject and the path

they provide for subsequent research.

2.2.1 Dietl's Mean-field Theory

The motivation for studying ZnO for spintronics began with the work of Dietl et al.

[21]. Dietl and coworkers employed a mean-field model of ferromagnetism, as originally

described by Zener, to the case of III-V and II-IV compound semiconductors. Zener's

model proposed that ferromagnetism is driven by the exchange interaction between

carriers and localized moments. Zener's early model was abandoned because it was found

inappropriate to describe the magnetism of transition metals. However, Dietl found that it

could be used to accurately predict the ferromagnetic Curie temperatures of GaMnAs and

ZnMnTe.

The model assumes that the ferromagnetic exchange interactions occur between

localized spins doped into the semiconductor matrix and are mediated by charge carriers.









These spins are assumed to be randomly distributed throughout the host semiconductor

lattice. Specifically, the doped Mn ions reside on group II or III sites and provide the

localized spins. Conceptually, the affect may be envisaged as a feedback between the

magnetic spins and the carriers by a spin-spin process; the localized magnetic spin induces

carrier polarization which then induces magnetic polarization and so on (Figure 2-1)

[22, 23]. The model suggests that high values of Tc are obtainable in p-type material,

while the Tc of n-type material should be constrained to lower temperatures. This can

be attributed to both the large p-d exchange integral (No/) and density of states of the

valence band, while the conduction band's s-d exchange (Noa) and density of states are

significantly smaller [24]. Note that in the case of III-V semiconductors, Mn also acts

as an acceptor dopant whereas it substitutes isovalently in the II-VI semiconductors.

Dietl extended his calculation to predict the Curie temperatures of other semiconductor

systems and oxides. The predictions were based on a Mn concentration of 5% and a hole

concentration of 3.5x102cm 3. The results are summarized in figure 2-2 as a function

of bandgap. Of particular importance, are the predicted Curie temperatures in excess of

room temperature for GaN and ZnO.

Diet's theory has proven useful in understanding the experimental results for

GaMnAs. However, it does not appear to be consistent with the experimental results

of transition metal doped wide band-gap semiconductors, such as the predictions for

GaN and ZnO. This stems from several reasons, including the difficulty in experimentally

preparing p-type ZnO material and the many observations of ferromagnetism in n-type

ZnO DMS. Nevertheless, Dietl's original theory has led to multiple experimental and

computational studies of transition metal doping in ZnO.

2.2.2 First-principles Design: DFT Calculations

Sato and Katayama-Yoshida have employed first-principles design to investigate

ferromagnetism in both semiconductor and oxide spintronics [25-27]. The magnetic

stability of transition metal doped ZnO was calculated using density functional theory









(DFT) within the framework of the local density approximation (LDA). The random

distribution of transitional metal ions over the lattice (creating disorder in the alloy) was

inherently included in the calculations by the coherent potential approximation (CPA).

Magnetic stability was calculated by comparing the total energy difference between the

ferromagnetic and spin-glass state, the lower of the two representing the ground state of

the system. In the case of Mn, their results are consistent with Dietl's theory that the

ferromagnetic state is stabilized with the addition of hole doping, and without holes the

spin-glass state is favored. However, V, Cr, Fe, Co, and Ni impurities were predicted to

be ferromagnetic without the need of additional charge carriers. Electron doping further

stabilized the ferromagnetic state in these alloys. Their work also points to a contribution

of d-states at the Fermi level, hinting at some delocalization of d-states. It was suggested

that this could lead to ferromagnetic ordering through a double-exchange interaction,

in which ferromagnetic alignment is stabilized by the hopping of 3d electrons between

neighboring TM sites. This mechanism is driven by partially unoccupied up-spin (or

down-spin) states in the 3d-band and is therefore not possible in the case of Mn, which

exhibits a half-filled 3d band. In the case of p-type doping, however, the transference of

weakly-bound 3d electrons between Mn ions may be mediated by the presence of holes.

The valence band p-states hybridize with the 3d-states of Mn and itinerant holes can

retain their d-like character. This stabilizes the ferromagnetic phase for Mn doping.

2.2.3 Ferromagnetism in Disordered Alloys

An additional theoretical approach considers whether ferromagnetic ordering between

the localized spins can originate from localized carriers. Again, the model is developed

in the mean-field treatment but accounts for positional disorder in the alloy. Numerical

studies within the mean-field treatment show that the nature of ferromagnetism is strongly

affected by this disorder and that the Tc can be pushed to higher temperatures with

increasing randomness in the position of Mn ions [28].









Ferromagnetism in this localized carrier regime can be explained through the

formation of bound magnetic polarons (BMPs) [29, 30]. A BMP is a quasi-particle

comprised of the localized carrier and the magnetic atoms encompassed within its radius

(Figure 2-3). The localized carrier is bound to its associated defect (such as a donor

atom if the carrier is an electron) in a hydrogenic orbital of radius, rh = c(m/m*)ao,

where c is the high frequency dielectric constant, m* is the effective mass, and ao is the

Bohr radius (53pm) [31]. This radius can be large (~8Ain ZnO) extending over several

lattice constants, and can encompass a number of magnetic dopants depending on their

concentration. The exchange interaction between the bound carrier and the magnetic

moments tends to align the moments parallel to one another inside the BMP. At high

temperatures, the BMPs may be isolated from one another. However, as the temperature

is lowered, the BMP radius grows and the individual BMPs begin to overlap. Overlapping

BMPs become correlated and their spins align, producing long range ferromagnetic

interactions [29]. At a critical temperature, the overlapping BMPs are percolated

throughout the sample and the transition to ferromagnetism occurs. The BMP model

is equally applicable to n-type or p-type material [30].

The BMP model allows for ferromagnetism in an insulating or semi-insulating regime.

This is especially attractive in the case of ZnO where high p-type doping, as required by

Dietl's model, is inherently difficult.

2.2.4 Ferromagnetism in a Spin-split Conduction Band

Coey et al. have proposed another model for ferromagnetism in DMS materials

based on a spin-split donor impurity band [31]. The model is consistent with the observed

magnetization for n-type transition-metal doped ZnO. In this model, donor defects (which

could arise from either oxygen vacancies or zinc interstitials in the case of ZnO) overlap at

large concentrations to form an impurity band. The impurity band can interact with local

magnetic moments through the formation of bound magnetic polarons (BMP). Within

each BMP, the bound carrier interacts with the magnetic dopants inside its radius and









can align the spins of the magnetic dopants parallel to one another. Ferromagnetism is

achieved when the BMPs start to overlap to form a continuous chain throughout the

material, thus percolating ferromagnetism in the DMS. However, Coey showed that in

this model, to achieve a high Tc, a fraction of the polaronic charge must delocalize (or

hybridize) onto each magnetic dopant. In a band scheme, this occurs when the impurity

band overlaps with unoccupied d-levels of the magnetic dopant. It was shown that for Sc,

Ti, and V, the spin-up states of the 3d TM metal are aligned with the impurity levels,

resulting is significant alignment. Similarly for Fe, Co, and Ni doping, the spin-down

states perform the same function. Interestingly, Mn and Cr doping would not lead to

strong magnetization due to small hybridization.

Within the framework of Coey's model, Kittilstved et al. have performed detailed

spectroscopic experiments on cobalt-doped ZnO [32]. Their results show that the singly

ionized Co+ state lies close to the conduction band, similar in energy to a shallow donor

state. Since the energies are similar, charge transfer can occur between the cobalt atoms

and the donor impurities, thus leading to the hybridization necessary for ferromagnetism.

Kittilstved et al. has also shown that this leads to an inherent polarity difference for

ferromagnetism in cobalt and manganese-doped ZnO. Whereas ferromagnetism in

cobalt-doped ZnO is closely tied to the presence of shallow donors, manganese-doped

ZnO is closely tied to the presence of shallow acceptors. The difference lies in the location

of the singly ionized Mn+3 state, which sits close to the valence band in ZnO. This idea is

described further at the end of the next section.

2.3 Experimental Progress in ZnO DMS

On the experimental front, there has been a wide distribution in the magnetic

properties reported for transition metal doped ZnO. Experiments have now covered a

broad range of parameters, including various transition-metal dopants (every element

in the first row of the transition metal series has now been surveyed), compositional

variations, preparation techniques and growth conditions, and post-growth processing.









The observed results are often conflicting and non-reproducible between research groups.

The discrepancy in the observed properties (and their interpretation) likely stems from

different growth techniques and conditions, and insufficient characterization. Most of

the difficulties arise in determining if the material is a true DMS (TM atoms randomly

substituting Zn lattice sites) or if ferromagnetism originates from TM clustering or

dopant-induced secondary phases. In any case, the results indicate that the underlying

mechanisms of ferromagnetism in ZnO DMS are quite sensitive to growth conditions and

must be clearly delineated by careful analysis.

Describing all the experimental trials in ZnO DMS over the past several years would

be tedious and overwhelming. There are already several reviews covering the subject

[33-36]. A compilation of some results is listed in Table 2-1 [36]. Instead, to provide

a flavor of the experimental progress, a brief summary of the important achievements

surrounding transition-metal doped ZnO is provided.

2.3.1 Mn-doped ZnO

By far, the two most studied magnetic dopants in ZnO have been Mn and Co.

Fukumura and coworkers were the first to study Mn-doped ZnO DMS using PLD [37].

A large solubility of 35% Mn was achieved while retaining the wurtzite structure of ZnO

(reminiscent of the earlier studies on II-Mn-VI compounds discussed earlier). This is

over the thermodynamic solid-solubility limit of Mn in ZnO and is a testament to the

non-equilibrium conditions obtainable by thin film growth. They later showed the heavily

doped alloy to exhibit spin-glass behavior with a spin-freezing temperature of -13K

due to strong antiferromagnetic exchange coupling between neighboring Mn atoms [38].

The high solubility of Mn achieved in ZnO motivated other experimental efforts into

the synthesis of ZnMnO. While some groups reported ferromagnetism, others observed

antiferromagnetic, spin-glass, or paramagnetic behavior (for example, refer to reference

[36]).









Sharma and coworkers were the first to report ferromagnetism above room temperature

in dilute Mn-doped ZnO bulk and thin film samples [39]. Bulk pellets with a nominal

concentration of 2 at% Mn (EDS showed the actual concentration to be much lower:

-0.3 at%) sintered below 700'C were found the have a Curie temperature of 420K.

Additionally, thin films deposited by PLD with 2.2 at% Mn were shown to exhibit

ferromagnetism at room temperature. However, using similar preparation techniques to

Sharma et al., Kundaliya and coworkers [40] convincingly demonstrated that the observed

high-temperature ferromagnetism resulted from a metastable phase (oxygen-vacancy-stabilized

Mn2- ZnO3_5), and not from the proposed carrier-mediated ferromagnetism between Mn

atoms. There is also discrepancy in the reported overall distribution of Mn atoms. For

example, a homogenous distribution of Mn was observed by Cheng and Chien [41], while

Jin et al. [42] found clustering of Mn atoms. Clearly, thorough characterization is needed

to fully appreciate and understand the origin of ferromagnetism in these materials.

2.3.2 Co-doped ZnO

One of the early works on cobalt-doped ZnO DMS was by Ueda et al. [43]. They

found the material to be ferromagnetic above 280K with 5-25%Co and 1%Al (added

as an n-type dopant) without the addition of secondary phases. Differences in the

magnetization were attributed to differences in the conductivity; films with higher

carrier concentrations (_1020cm 3) showed ferromagnetic features with higher Ms and Tc.

Since then, additional experimental studies have investigated the properties and origin

of ferromagnetism in cobalt-doped ZnO. Again, the results are conflicting with reports

of ferromagnetism in phase pure films [44, 45], ferromagnetism from clusters [46], and no

observed ferromagnetism [47].

The first report of reversible (controlled) switching of ferromagnetism in any

DMS was demonstrated by Schwartz and Gamelin in cobalt-doped ZnO [48]. The

reversibility was mediated by the incorporation and removal of Zn interstitials. The

Zn interstitial (Zni) is a known n-type dopant that produces a shallow donor level below









the conduction band. Diffusing Zni into the lattice lowers the conductivity and activates

room temperature ferromagnetism. Removing Zni, by heating in air, returned the films

to an insulating state and subsequently quenched the ferromagnetism. The process was

reversible over many cycles. This reversibility is evidence that free carriers activate

ferromagnetism in cobalt-doped ZnO. The process was observed in both MOCVD grown

films and ZnO:Co nanoparticle films prepared by spin coating. Strong hybridization of Zni

donor states with Co+2 states near the conduction band (which, as explained earlier, is

theoretically believed to cause ferromagnetism) was used to explain the magnetic ordering.

Conduction electrons, derived from the Zni donors, delocalize over several Co+2 ions and

ferromagnetically align their spins through a double exchange interaction.

Importantly, from the same group, Kittilstved was able to demonstrate a chemical

polarity difference between the ferromagnetism in ZnCoO and ZnMnO [49]. Specifically,

p-type ZnMnO led to ferromagnetism, while ferromagnetism in ZnCoO was activated by

n-type doping. Doping of the ZnMnO was done on a local level by N-capping ZnMnO

nanoparticles with amines. ZnCoO nanoparticle films were made n-type by capping with

oxygen. Reversing the capping layers, ZnCoO:N and ZnMnO:O, led to the disappearance

of ferromagnetism in both sets of films. Optical absorption, MCD, and photoconductivity

measurements were employed to understand this inherent polarity difference [32]. For

n-type ZnCoO, the authors showed that a resonance in the charge transfer (Co+' Co+2

+ ecB, AE 0.27eV) and donor state energies can lead to a large hybridization necessary

for ferromagnetism. For ZnMnO, a similar resonance was observed but derived from the

Mn+3 state close to the valence band (Mn+3 Mn+2 + h+B, AE 0.22eV) with acceptor

state energies.

















Mn+2 h+ Mn+2





J s(i) s(j)
site i site j

Figure 2-1. Schematic representation of magnetic exchange between two Mn ions mediated
by a delocalized hole. Adapted from [22].















-- Group IV
- III As
III P


III Sb
U Oxides


Ge
InAs

GaSb


I GaAs
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0.0 0.5


1.5 2.0 2.5 3.0 3.5 4.0


Semiconductor Band Gap (eV)

Figure 2-2. Predicted Curie temperatures based on Dietl's calculations [21] for 3%Mn and
a hole concentration of 3.5x1020cm3 (After [50]).


AlAs
U-
















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Figure 2-3. Illustration of bound magnetic polarons. An electron bound within its
hydrogenic orbital couples to magnetic impurities causing them to align
parallel to one another inside the orbit radius (Adapted from [31]).


o M





















































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CHAPTER 3
THIN FILM DEPOSITION AND EXPERIMENTATION

Pulsed Laser Deposition (PLD) was used for all the film deposition presented in this

dissertation. This section will cover the experimental apparatus and methods for these

film depositions. A brief discussion of the basic PLD process will be given first. This is

followed by a description of the specific PLD system used for deposition, including the

growth chamber and UV excimer laser. Then the exact experimental steps used to deposit

films in this work will be described, including the preparation of oxide targets (the source

materials) and the typical growth procedure for obtaining epitaxial thin films.

3.1 PLD as a Tool for Thin Film Oxides

PLD has evolved into a successful and widely-used research tool for the fabrication

of thin-film complex oxides [75, 76]. PLD is a physical deposition method. High-intensity

laser pulses (upwards of 100s MW cm 2) are used to vaporize atomic species from a

target of desired chemical composition. The target material is typically a solid source,

but liquid sources, such as organic liquids, are also feasible [77]. The laser energy is

absorbed by the target and a rapidly expanding plume, containing electrons and ground-

and excited-state neutral atoms and ions, is ejected from the target surface [76]. This

plume is highly directional. It is emitted perpendicular to the target surface and has an

angular distribution given by f(0)=(cos0)", where f(0) is the distribution of ablated flux

and n'5-25 [76]. The ablated material is collected onto a heated substrate located several

centimeters from the target.

PLD growth offers several advantages in the realm of thin film deposition [76]:

1. Congruent (stoichiometric) transfer of the target material to the deposited film.
The various chemical components of complex oxides, or other multicomponent
materials, are evaporated simultaneously. This allows the control of film composition
by simply preparing targets of the desired composition. Typically, the compositional
stoichiometry of the target is accurately reproduced in the film.

2. Almost anything can be ablated, including high melting temperature oxides.









3. The laser beam is capable of imparting considerable amounts of kinetic energy to the
ablated chemical species. The additional kinetic energy can lead to larger sticking
coefficients and enhanced adatom mobility on the growth surface, permitting epitaxy
at reduced growth temperatures.

4. Deposition in a background gas is permissible. The highly energetic plume can
readily undergo gas-phase reactions with ambient gases, such as 02, H20, and NO2,
providing an addition degree of freedom in the growth parameters of thin film oxides.

3.2 PLD System Used for This Work

3.2.1 The Growth Environment

A commercial PLD system built by Neocera was used for the thin film depositions

described in this dissertation. A diagram of the system is provided in Figure 3-1. The

system is composed of one main vacuum chamber which is used for growth. This chamber

is evacuated by a pair of pumps that are situated underneath the chamber and can be

sealed-off by the gate valve. The laser used for target ablation is separate from the system.

The other components of the system are contained within the growth chamber.

The Neocera system uses a Pfeiffer turbo pump to create and maintain the low

pressures needed for growth. The turbo pump consists of a series of rapidly spinning

blades to transfer gas out of the chamber. Essentially, gas molecules collide with the

blades and are kinetically driven toward the pump exhaust. This is accomplished through

multiple stages until the exhaust gas is compressed to the fore-line pressure of the backing

pump. The turbo pump is water cooled by a recirculating chiller to suppress overheating.

The system's turbo pump is backed by an oil-less four-chamber diaphragm pump. This

backing pump serves two purposes. Before the turbo pump can be switched on, the

backing pump rough pumps the growth chamber down from atmosphere to an inlet

pressure manageable by the turbo pump. The growth chamber is evacuated to a pressure

of at least 240 Torr before the turbo pump is switched on. The backing pump serves

its second purpose by maintaining the fore-line pressure for the turbo pump to operate

efficiently and removing the exhaust gas into the atmosphere. Since there is no load-lock,









the chamber has to be evacuated from atmosphere for every growth run. The ultimate

base pressure of the system is 2 x 10 6 Torr.

The pressure inside the growth chamber is monitored by a pair of vacuum gauges.

Pressures above ImTorr are monitored with a Granville convectron gauge. Pressures below

ImTorr are monitored using a cold-cathode gauge.

The ablation targets are mounted onto a rotating carousel. A total of six targets can

be mounted at one time and each target can be accessed during growth by rotating the

carousel to the position of the desired target. This allows for the growth of film stacks

that can be used for device fabrication, compositionally graded layers, superlattices,

etc. The ablation target is also rotated around its center axis during deposition. Target

rotation helps maintain uniform target ablation and film deposition.

The substrates are mounted onto a solid block, resistive heater. The heater is

mounted to a vacuum flange that can be removed to mount the samples and provide

access to the growth chamber to exchange targets.

Laser access into the growth chamber is granted through an optical window made

from SUPRASIL fused silica, which has a large optical transparency to UV radiation in

order to minimize attenuation of the laser energy.

3.2.2 The Laser Source

A Lambda-Physik Compex 205 KrF pulsed laser is used as the ablation source.

This laser produces a coherent beam with a 248nm wavelength. The laser energy and

pulse repetition can be varied to suit a particular experiment. Most of the growth in this

dissertation was done at energies around 300mJ with repetition rates ranging from 1 to

10 Hz. The laser output is directed into the PLD chamber by a series of four mirrors.

One of the mirrors is mounted to a microprocessor driven rotational stage. This stage

can be programmed to repeatedly scan the laser across the surface of the ablation target.

There is an approximate 10% loss of beam energy at each mirror. The beam is focused

using two plane focusing lenses to focus both the horizontal and vertical width of the









beam. Typically, the beam was focused to a spot size of 4-6 mm2. The Compex 205 is a

chemical excimer laser that emits ultra violet (UV) light at a wavelength of 248nm. The

term 'excimer' is derived from EXCIted diMER. The Compex uses a KrF excimer, derived

from ultra-high purity Kr and F2 gases, as the source of light emission. The excimer is

formed by supplying electrical energy to the gas mix. Since Kr is a noble gas, the bonded

KrF complex is a highly unstable excited state and quickly decomposes to the unbonded

ground state. The stored chemical energy is released in the form of a photon with the

characteristic decomposition energy (the lasing energy):

2KrF(g) = 2Kr(g) + F2(g) + energy(248nm)

3.3 Typical Growth Procedure

A typical growth run using the PLD system is described next. Describing the

deposition process in detail can alleviate ambiguity at each step and provide a means of

tracking down differences in film preparation between users. The process is described in

two steps, substrate preparation and film growth.

3.3.1 Substrate Preparation

The sapphire substrates used for ZnO film growth typically came in the form of 2"

diameter wafers. Square substrates 6.5mm x 6.5mm were cleaved from the larger wafers

using a diamond scribe. The MgA1204 and MgO substrates used for the Cu20 films came

pre-cut in 1cm x 1cm squares and were cut into four 5mm x 5mm squares for film growth.

Before each deposition the substrates were chemically degreased in sequential baths

of TCE, acetone, and methanol. The degreasing method was the same for all substrates.

First, about 15mL of each solvent was poured into separate 50mL beakers. The bottoms

of the beakers were then placed into an ultrasonic bath of tap water to aid the solvent in

removing any contamination or residue inside the beaker. The beakers were ultrasonicated

for at least 5 minutes. Next, each beaker was dumped out and refilled with about 20mL of

fresh solvent. The substrates were first placed into the beaker of TCE and ultrasonically

rinsed for at least 5 minutes. To keep the substrates off the bottom of the beaker, the









substrates were held with plastic tweezers (which were pinched closed using a binder

clip). The substrates were then subsequently cleaned in the acetone beaker and then the

methanol beaker. In order to keep the substrates clean of dried solvent residue, they were

immediately blown dry with N2 after being removed from the methanol.

After cleaning the substrates they were mounted to the heater block. The heater

block was cleaned before attaching the substrates. The block was scraped clean of dry

silver paint (left from the previous deposition) using a stainless steel razor blade, etched

in 1:1 Nitric acid/Di H20, and wiped with acetone. The substrates were attached to the

block with silver paint (Ted Pella 'Leitsilber 200' silver paint) and allowed to dry for

20min inside a fume hood. A glass petri dish was placed upside-down over the substrates

to protect them from dust during the drying time.

3.3.2 Thin Film Growth

After drying, the samples and heater assembly are mounted inside the chamber by

reattaching the heater flange to the system. The gate valve is opened and the backing

pump is switched on to rough pump the growth chamber from atmosphere. When the

pressure drops below 250 Torr, the turbo pump is then turned on. The system is allowed

to pump for a few hours until a pressure below 1 x 10 5 Torr is reached. The temperature

controller is ramped to the desired substrate temperature at a rate of 10'C/min. Oxygen

is flowed in to fill the chamber with 10mTorr of 02. This is thought to help drive off any

hydrocarbon contamination from the substrate during heating. When the set temperature

is reached, the sample shutter is closed to block the substrates from the target, the oxygen

flow is stopped, and the target is pre-ablated at 300mJ for at least 1000 laser pulses.

Pre-ablation cleans the target surface before actual film growth. Since the chamber

pressure rises from vaporized species off the chamber wall during heating, the chamber

is left at temperature in 10mTorr 02 until the base pressure drops below 1 x 10 5 Torr.

To start the film deposition, the shutter is flipped open, oxygen is introduced into the

chamber at the desired growth pressure, the laser energy and reprate are set and the laser









is switched on. During the growth, the laser beam is scanned and the target is rotated

to provide uniform ablation across the target surface. Once the growth is completed, the

temperature is ramped back to room temperature. The oxygen overpressure is usually

maintained at the growth pressure during cool down. The samples are then removed from

the system and stored in a desiccator until needed.

3.4 Fabricating PLD Ablation Targets

Ceramic targets of known composition were used as the source material for the PLD

film growth. Films were deposited by ablating these targets inside the growth chamber.

The ablation targets were fabricated using the solid-state reaction of mixed oxide powders.

The high purity source powders (>99.99% pure) were purchased from Alfa-Aesar. Targets

were prepared based on atomic percent concentrations of the dopant species (in contrast

to weight percent). For example, a Zno97Mno030 target contains 3at% Mn doped into the

ZnO. Standard dimensional analysis using atomic weights from the periodic table were

used to convert at% to weight percent so the appropriate amounts of powder could be

weighed. The powders were weighed in a plastic weighing dish using an electronic scale.

After weighing, the powders were transferred to an alumina mortar. Methanol was added

to the powders to aid with mixing and the powders were ground and thoroughly mixed

using an alumina pestle. The powders were then allowed to dry in air. A uniaxial press

was used to compact the powders into a 1" diameter disc using a stainless steel die. The

applied pressure was typically 2-4 metric tons. Each disc was placed on a thin piece of

alumina and placed inside a high temperature box furnace. The pellets were then covered

with an alumina crucible. The box furnace was raised to 1000'C at a rate of 10'C/min.

The targets were sintered at 1000'C in air for 12 hours and allowed to cool to room

temperature.









3.5 Thin Film Characterization

3.5.1 X-ray Diffraction

X-ray diffraction (XRD) is a versatile tool for studying the crystalline nature of

materials. Analysis of grain size, film quality, and phase identification are possible

through XRD techniques. I used powder XRD to investigate the structural information

of the deposited films. Specifically, crystalline phases of material were determined by

theta-2theta diffraction scans.

The films were measured with a Philips APD 3720 diffractometer. The system uses a

Cu x-ray source that emits primarily Cu Ka, photons with a wavelength of 1.5405 A. A

Ni filter absorbs most of the other characteristic wavelengths from the source, although

some Cu KI2 and KO photons escape toward the sample. These wavelengths are important

because they generate extra diffraction peaks in the collected data.

In XRD, the incident x-rays interact with the material through constructive and

destructive interference from the periodic lattice planes of the crystal. The condition for

constructive interference is governed by Bragg's Law:

nA = 2d sin0

where A is the x-ray wavelength, d is the distance between atomic planes, and 0 is

the angle between these planes and the x-ray source (Figure 3-2). Therefore only specific

d-spacings satisfy the Bragg condition at a certain angle.

The Philips system is configured in Bragg-Brentano geometry. The x-ray source is

pointed at the sample and held in a fixed position. The sample rotates around the angle

0, and the detector rotates around 20. In this 0-20 geometry, only planes parallel to the

sample surface can satisfy the Bragg condition. This provides crystalline orientation

information of the deposited film.

The x-ray intensity is plotted as a function of 20 to reveal the diffraction pattern

for a sample. Phase identification is possible by comparing the diffraction peaks to









known material standards in the JCPDS catalog. With this method, recognition of any

dopant-induced secondary phases within the semiconductor host was determined.

Additional x-ray characterization for select samples was carried out on a Philips

X'pert high-resolution diffractometer. w-rocking curves were used to assess the lattice

parameter and crystalline quality of the deposited films. The epitaxial quality between

films was quantified by the w-rocking curve FWHM (full-width at half maximum). b-scans

were used to measure the in-plane orientation and film texture to confirm epitaxial

registry with the substrate.

3.5.2 Magnetoresistance and Hall Effect Measurements

Transport measurements provide valuable insight into the electronic structure

of materials. Electrical properties are sensitive to material imperfections, including

electrical dopants, atomic impurities, lattice strain, and structural defects. The transport

properties of films grown in this work were probed by magnetoresistance and Hall effect

measurements.

Samples were measured using either Van der Pauw or Hall bridge geometries. Van

der Pauw samples were approximately 6.5mm x 6.5 mm square. Hall bridge specimens

were patterned during film growth by depositing through a stainless steel shadow mask.

Contacts were soldered to the samples using indium metal. The films were mounted

inside a Quantum Design Physical Property Measurement System (PPMS) to control

the ambient temperature and magnetic field near the sample, and the electrical data was

collected using a Keithley high-impedance Hall effect system. Details of the configuration

are provided in the appendix.

The Hall effect provides valuable information about the charge carriers in a material.

The Hall effect is caused by the interaction of moving charge carriers with an applied

magnetic. This interaction is dictated by the Lorentz force and leads to an accumulation

of charge at the sample edge. The accumulated charge distribution induces a potential

drop (the Hall voltage) between the sample edges until the Lorentz force is balanced. The









carrier type, mobility, and density can be determined from measurements of the induced

Hall voltage. Hall effect measurements become especially valuable in the study of magnetic

materials. In a ferromagnetic material, the Hall resistivity is described by the empirical

expression:

pX = RM +RoB

Ro is termed the ordinary Hall coefficient and is caused by the Lorentz force. R, is

the anomalous Hall coefficient that results from a finite sample magnetization. While the

ordinary term is caused by classical physics, the anomalous term derives from quantum

mechanical effects. The anomalous term results from spin-dependent scattering between

charge carriers and moment carrying centers. The physical origin of this interaction is the

spin-orbit coupling of the charge carrier as it passes by a magnetic impurity. The strength

of this term is given by the strength of the spin-orbit coupling and the relative density of

spin-up and spin-down electrons.

Magnetoresistance can provide key insights into the transport properties of magnetic

semiconductors, including the potential landscape of impurities, lattice disorder,

and electron-electron interactions [78]. Magnetoresistance was typically measured

simultaneously with the Hall data.

3.5.3 Electron Dispersive Spectroscopy (EDS)

EDS was used to determine chemical compositions. The EDS analysis was performed

inside a JEOL 6335F field-emission scanning microscope fitted with a liquid nitrogen

cooled EDS detector.

EDS measures the emitted x-ray spectrum of a material when bombarded by a

beam of energetic electrons. The high energy electrons are focused onto a sample where

they transfer their kinetic energy into the lattice through a series of collisional events. A

small fraction of these events are capable of ionizing an atom by ejecting an inner shell

electron. The atom then reverts back to its ground state by filling the vacancy with an

electron from a higher energy orbital, and either a photon or Auger electron is emitted.









The emitted photons have a characteristic energy depending on the atom that emits

it. Therefore, elemental information can be determined. The relative amount of each

element can be calculated by comparing the peak heights and applying the corresponding

ZAF corrections, where Z is the atomic number, A is absorption, and F is the x-ray

fluorescence.

3.5.4 Optical Absorption

Information related to the electronic band structure of semiconductors can be inferred

from their interaction with light. The absorption of an optical photon necessarily involves

the transition of an electron from an occupied energy state to a vacant state of higher

energy. Energy and momentum must be conserved in the process.

Consider a semiconductor with a filled valence band and empty states at the bottom

of the conduction band. If the photon energy is below the lowest allowed electronic

transition, in this case the band-gap energy, the photon is not absorbed and passes

through the material; the material is transparent to the particular photon energy. If,

however, the photon energy is greater than the gap energy, the photon is absorbed by

an electronic transition from the valence to the conduction band. Therefore, the onset of

optical absorption is hv > Eg. The semiconductor acts as a low-pass optical filter.

The absorption coefficient, a, for direct interband transitions is given by the relation

[79]:

(ah) = A,(hv Eg)2

where, hv is the photon energy, Eg is the band-gap, and Ao is a parameter associated

with the transition probability and refractive index, a is calculated from the absorbance

data using a = 3A, where A is the absorbance and t is the film thickness. If a plot

of (ahlv)2 vs. hv reveals a straight line, the sample has a direct gap absorption. The

band-gap is determined by extrapolating the linear portion to zero.

Transitions where the wave vector is not conserved, Ak#/O, are also possible. These

are indirect transitions across the gap. In order to conserve the total momentum, these









transitions require exchanging momentum with the lattice through the absorption or

creation of a phonon. Compared to direct transitions, the required electron-phonon

interaction makes these transitions less likely to occur. The absorption coefficient for

indirect transitions is [79]:

(ahv) = Aj(hv Eg)2

Besides band-to-band absorptions, other transitions are also possible, including

impurity-to-band (discussed below), impurity-to-impurity, excitonic, and intraband

transitions [79].

Perturbations in the band structure of the material can lead to the formation of

band tails [79]. The perturbations are caused by imperfections in the crystal lattice,

such as impurities, structural defects, and lattice disorder. The band tailing extends the

conduction and valence band states into the energy gap. Electronic transitions between

the tails cause an exponentially increasing absorption coefficient known as Urbach's

rule: d(ln a)/d (hv) = In the absorption spectrum, these Urbach tails appear as a

broadening of the lowest energy transitions.

Ionized impurities will interact with band carriers through the Coulomb interaction.

A positively charged donor, for example, will attract conduction band electrons and

repel valence band holes causing a local distortion in the band structure. This effect can

be more or less strong depending on how many impurities are clustered together at a

particular spot in the material. Additionally, impurities can alter the density of states.

When doped in large concentrations, the discrete impurity states can evolve into impurity

bands. Optical transitions between these bands are permissible.

For the particular case of transition metal oxides, intra-ion transitions between

d-states, or d-d transitions, provide an additional avenue for optical excitation. The

electronic configuration of the metal ion is perturbed by the local chemical environment,

the so-called crystal field splitting. The orbital degeneracy of the d-states is lifted the

states are split into different energy levels. The splitting is strongly influenced by









coordination geometry. For example, octahedral sites result in a completely different

arrangement of energy states than, say, a tetrahedral symmetry would produce.

The splitting are characteristic of the metal ion and the lattice. Therefore, transition

metals doped into oxide lattices will produce characteristic absorption lines that can

be used to identify the coordination and valence state of the dopant. This becomes

particularly useful at high doping concentrations which may be near the solid solubility

limit.

In this work, a Perkin-Elmer Lambda 800 UV/Vis double-beam spectrometer was

used for optical absorption measurements. Samples were mounted to rigid opaque panels

that contain a small hole, 3mm in diameter, for light to pass through. This ensures that

only light passing through the sample is detected. It also keeps the transmitted beam

size constant between samples. The samples were held normal to the incident light.

Absorbance spectra were measured using unpolarized light with wavelengths between

200nm and 900nm. The data was used to determine the band gap as a function of doping.

The optical signature of magnetic dopants was also used to confirm their solubility and

valence state in the host lattice.

3.5.5 SQUID

Magnetic measurements were performed by my collaborators in Dr. Art Hebard's

research group in the University of Florida's physics department. The measurements were

done using a Quantum Design MPMS SQUID (Superconducting Quantum Interference

Device) magnetometer. Currently, SQUIDs provide the most sensitive resolutions for

magnetic field measurements.

Magnetization versus field loops were taken at various temperatures. Ferromagnetic

materials will produce a finite hysteresis in these loops. Therefore, loop hysteresis was

used to verify ferromagnetism in my samples. The samples were mounted inside a plastic

drinking straw (which has no ferromagnetic components) and placed perpendicular to the

applied field.









Additionally, Field Cooled/Zero Field Cooled (FC/ZFC) measurements were used

to track the magnetic response as a function of temperature. FC measurements were

performed by measuring the magnetization as the sample is cooled down to 10K in a small

applied field. ZFC measurements were performed by cooling the sample to 10K in zero

field and then applying a small field as the sample was heated back to room temperature

while measuring the magnetization.

The diamagnetic response of the sample and substrate were subtracted from the M

vs. H data. This was performed for each measurement. The magnet was swept to 5T

to reveal the diamagnetic and paramagnetic response of the sample and substrate. The

slope of the high field response is the background magnetic susceptibility, X = The

susceptibility multiplied by the applied field is then subtracted from each data point.


































oGas Inlet


Turbo Pump


Figure 3-1. Schematic of pulsed laser deposition system.




































2Film
Film


Incoming
X-rays


Diffracted
X-rays

D d


\' 0'
N-


Figure 3-2. Visualization of Bragg's law for x-ray diffraction. Adapted from [80].









CHAPTER 4
PROPERTIES OF Mn-DOPED Cu20 DMS

4.1 Introduction

Semiconducting oxides offer the potential for exploring and understanding spin-based

functionality in a semiconducting host material. Dietl's theoretical prediction suggests

that carrier-mediated ferromagnetism should be favored for heavily-doped p-type ZnO.

However, this poses a serious challenge experimentally since ZnO is naturally an n-type

semiconductor due to intrinsic defects and difficult to dope even moderately p-type. This

poses the opportunity of using alternative wide-gap semiconductors that are naturally

p-type to fabricate high Curie temperature DMS materials.

In this section, the properties of Mn-doped Cu20 are explored. Cu20 is a p-type,

direct wide bandgap semiconductor that may hold interest in exploring spin behavior in

an oxide DMS. Activities focused on understanding Mn incorporation during thin-film

synthesis, as well as electrical transport and magnetic characterization. Ferromagnetism

is observed in select Mn-doped Cu20 films, but appears to be associated with Mn304

secondary phases. In phase-pure Mn-doped Cu20 films, no evidence for ferromagnetism is

observed above that attributed to the substrate.

4.2 Cu20: A Wide-bandgap P-type Semiconductor

Cu20 (Cuprite) is one of the earliest know semiconductor materials. The development

of copper-Cu20 rectifiers dates back to the early 1920's, decades before silicon devices

would dominate the semiconductor market. The rectifiers were easily fabricated by

oxidizing pure copper at high-temperature inside a furnace, and were advantageous

over earlier "point contact" (cat-whisker) rectifiers since the interface remained free of

contamination and could be made uniformly over large areas of copper [81]. Cu20 also

has a large theoretical solar cell efficiency (18%) and was heavily studied for photovoltaic

properties since the 1970's; however practical application of Cu20 solar cells is limited due

to difficulties with improving the semiconductor's electrical properties [82, 83]









Cu20 is one of the few binary p-type semiconducting oxides. It has a direct band-gap

of -2.0 eV. The structure is cubic with a = 4.27A(Figure 4-1). Oxygen atoms sit on

a bcc lattice and are encaged by a tetrahedron of Cu atoms. Each Cu atom is two-fold

coordinated resulting in the rare occurrence of O-Cu-O linear bonding. Each Cu-O bond

length is about 1.84A. The states at the top of the Cu20 valence band are predominantly

of Cu 3d character with O 2p states lying lower in the band [84]. The dominant method

for p-type conduction is typically attributed to the formation of Cu vacancies in the

lattice: Cu'+ Cuo + h+

Using both DFT and DFT+U calculations, Nolan and Elliott have calculated the

formation energy of the Cu vacancy to be on the order of 0.4-1.7eV [84]. They find the

vacancy creates an acceptor state around 0.2eV above the Fermi level, and that the

holes are delocalized around the vacancy. Experimentally, the acceptor states have been

determined to be 0.26eV [85], 0.5eV [86], 0.4eV [87] above the valence band. The states at

the top of the Cu20 valence band are predominantly of Cu 3d character with O 2p states

lying lower in the band [84]. Oxides typically have small hole mobilities since the O 2p

valence bands are localized; however, the fully-occupied 3d10 states in Cu20 are mobile

when converted to holes [84]. Reasonable hole mobilities on the order of 100 cm2v is 1

have been found experimentally [82, 88].

The number of reports exploring the magnetic properties of transition-metal doped

Cu20 are rather limited. There is also variation in the observed ferromagnetism between

groups, similar to the contemporary state of research into other DMS materials, like

TM-doped ZnO. Wei and coworkers found bulk pellets of Cu20 doped with a nominal

concentration of 1.7 at%Mn to be ferromagnetic up to 300K [89]. They report a magnetic

saturation of 0.4pB/Mn at 10K which dropped to ~0.05pB/Mn at room temperature.

These results were calculated using the 1.7 at% nominal concentration, but their EDS

results indicated a Mn concentration of -0.3-0.5 at% which would result in a higher

saturation (up to 2.5pB/Mn at 10K). The same group also found room temperature









ferromagnetism with a saturation moment of 0.6/B/Mn in Cu20:Mn films deposited

by near-room temperature electrodeposition [90]. In both cases, the Cu20 films doped

with Mn showed lower resistivity than undoped films by a factor of 2. Mn-doped C20

films were depostied by Pan et al. using rf magnetron sputtering [91]. The films were

paramagnetic above 5K and showed some weak ferromagnetism near 5K with a moment

of 5.3pB/Mn. Kale et al. used PLD to study 5 at% cobalt-doped Cu20 films that were

codoped with 0.5 at% Al, V, and Zn [92]. Ferromagetism was only observed in films

codoped with Al, but persisted up to room temperature with a moment of 0.44PB/Co.

Doping Zn or V did not show ferromagnetism but caused the resistivity to decrease and

increase, respectively. Since there was no correlation between the ferromagnetism and

resistivity, they suggested the ferromagnetism may be realted to orbital defects introduced

by the Al (which has s and p valence orbitals). In contrast, Antony et al. did observe

ferromagnetism at 400K in 5%Co-doped Cu20 nanoclusters [93]. DFT calculations,

employed by Sieberer and coworkers, show that the ferromagnetic properties of Cu20

doped with Co or Mn is dependant upon the presence of defects (copper or oxygen

vacancies) [94]. The found defect-free Mn-doped Cu20 to be anitferromagnetic, while

long-range ferromagnetism could occur when defects are present. For the case of Cu20

doped with Co, the exchange constants are mostly ferromagnetic in defect free material

(only nearest-neighbor sites were found to have antiferromagnetic coupling when the

Hubbard U is increased). Cu vacancies were found to increase the Tc, while oxygen

deficiency introduced strong oscillations in the magnetic exchange.

4.3 Experimental

Pulsed-laser deposition was used for film growth. Manganese-doped CuO targets were

fabricated using high-purity CuO (99.995%), with MnO2 (99.999%) serving as the doping

agent. The targets were pressed and sintered at 860'C for 12hr, followed by 950'C for

2hr in air. Targets were fabricated with a composition of Cul Mno10. A KrF excimer

laser was used as the ablation source. A laser repetition rate of 5 Hz was used, with a









target to substrate distance of 4 cm and a laser pulse energy density of 1-3 J/cm2. The

growth chamber exhibits a base pressure of '10 6 Torr. Film thickness was approximately

300 nm. Single crystal (0 0 1) MgA1204 and MgO were used as the substrate material in

this study. Four-circle X-ray diffraction was used to determine crystallinity and dopant

solubility. SQUID magnetometry was used to characterize the magnetic properties of the

deposited films. In particular, it was used to determine the presence of ferromagnetism

and measure the Curie temperature. Photoluminescence was used to characterize the

optical properties of the material.

4.4 Results and Discussion

The epitaxial growth of Cu20 thin films is dependent upon achieving oxidation

conditions in which the Cu ion assumes a 1+ valence. Figure 4-2 shows the thermodynamic

stability curves for Cu-Cu20-CuO as a function of oxygen partial pressure and temperature

[95]. Using this, we find that epitaxial Cu20 films can be grown from high-purity CuO or

Cu targets using pulsed-laser deposition in an oxygen ambient. Figure 4-3 shows the X-ray

diffraction data for a (0 0 1) Cu20 film grown on (0 0 1) MgO with P(02)=4x10-4 Torr,

T = 750'C, using a Cu target. Similar results were obtained on perovskite substrates,

yielding p-type films with a carrier density of 1015 cm3 and a room temperature mobility

of 26 cm2 v 1 s1. The primary focus of this work was to investigate the synthesis and

properties of epitaxial Cu20 doped with Mn. The selection of Mn as the transition metal

dopant is based on the best available evidence in determining which magnetic impurities

are likely to yield ferromagnetism. Manganese doping has resulted in interesting magnetic

phenomenon in II-VI and III-V semiconductors, including spin glass or antiferromagnetic

behavior for a number of systems, and is predicted to yield a high Curie temperature

in ZnO as previously discussed. In the Cu20 structure, the Cul+ cation is twofold

coordinated with an ionic radius of 0.46 A. Mn2+ does not normally exhibit a twofold

coordination in bulk materials. However, the radius for fourfold coordinated Mn2+ is 0.66

A, which is close to that for the fourfold coordinated Cul+ (0.60 A). As a 2+ cation on









a 1+ site, Mn would be expected to compensate for the p-type conductivity observed

in Cu20. The resistivity of Mn-doped Cu20 films was significantly higher that that for

undoped films.

The phase stability and solid solubility of Mn in Cu20 was investigated as a function

of deposition temperature and oxygen pressure. X-ray diffraction was used to determine

conditions that limit segregation of secondary phases. Film growth was carried out over a

temperature range of 300-700'C and an oxygen pressure of 10 3, 10 Torr, or in vacuum.

The base pressure of 10 5 Torr consists mostly of water vapor. Figures 4-4 to 4-6 show

the X-ray diffraction data for Mn-doped films grown under these conditions. Several

items should be noted. First, the Cu20 phase is dominant over most of this range. This

is surprising as the thermodynamic stability behavior suggests that CuO should be the

stable phase for T < 600 C, P(02)=10-3 Torr, and T < 550 C, P(02)=10-4 Torr. Two

possibilities exist in explaining this discrepancy. First, the doping of Cu20 with Mn may

shift the phase stability line. Second, epitaxy at low temperature may be sufficient to

stabilize the Cu20 phase. A segregation of Mn oxides in the Mn-doped films was also

examined for the various film-growth conditions. The X-ray diffraction results indicate

the presence of antiferromagnetic Mn203 as an impurity Cu20 phase for T > 500 0C.

For films grown in vacuum, a weak peak that could be associated with either Mn203 or

ferromagnetic Mn304 is observed. Phase-pure Cu20 films were obtained at T <400 C,

indicating the metastable incorporation of Mn in the Cu20 matrix. Figure 4-7 shows the

phase assemblage as a function of growth conditions.

The magnetic properties of Mn-doped samples were measured using a Quantum

Design SQUID magnetometer. Measurements were made on films grown at low temperatures,

in which no Mn304 impurity phase peaks were evident in the X-ray diffraction patterns,

as well as films grown at elevated temperatures. In order to characterize the magnetic

properties of the Mn-doped samples, field-cooled and zero field-cooled magnetization

measurements were performed from 4.2 to 300 K. By taking the difference (AM) between









these two quantities, the para- and diamagnetic contributions to the magnetization can

be subtracted, leaving only a measure of ferromagnetic behavior. Figure 4-8 shows the

AM difference as a function of temperature for a Mn-doped sample grown at 300 C in 1

mTorr of 02. At low temperature, a small, but finite field-cooled minus zero field-cooled

magnetic signal persists up to -250 K as seen in figure 4-8. However, the magnetization

signal is small, and can be attributed to a background magnetic signal from the MgA1204

substrate. Figure 4-9 shows the temperature dependent AM and M vs. H behavior for

the substrate material without a Cu20 film. For this and other phase-pure samples, no

ferromagnetism could be detected above that attributed to the substrate.

We also investigated the magnetic properties of Mn-doped Cu20 grown at 700 'C.

These epitaxial Mn-doped Cu20 films are clearly ferromagnetic with a Tc of -48 K as

seen in figure 4-10. A key requirement in understanding ferromagnetism in transition

metal doped semiconductors is to delineate whether the magnetism originates from

substitutional dopants on cation sites, or from the formation of a secondary phase that

is ferromagnetic. The importance of this issue cannot be understated. The concept of

spintronics based on ferromagnetic semiconductors assumes that the spin polarization

exists in the distribution of semiconductor carriers. Localized magnetic precipitates might

be of interest in nanomagnetics, but is of little utility for semiconductor-based spintronics.

The question of precipitates vs. carrier-mediated ferromagnetism is complex, and is a

central topic of discussion for other semiconducting oxides that exhibit ferromagnetism,

in particular the Co-doped TiO2 system [96, 97]. Several issues must be addressed in

order to gain insight into the possible role of secondary phase precipitates in the magnetic

properties of transition metal doped semiconductors, specifically for Cu20 films. First, one

should identify all candidate magnetic phases possible from the assemblage of elements.

The coincidence of Tc with a known candidate secondary ferromagnetic phase indicates a

likely source of at least part of the magnetic signature. For the present material, metallic

Mn is antiferromagnetic, with a NIel temperature of 100 K. In addition, nearly all of









the possible Mn-based binary and ternary oxide candidates are antiferromagnetic. The

exception to this is Mn304, which is ferromagnetic with a Curie temperature of 46 K

[98, 99]. X-ray diffraction measurements on the sample considered in figure 4-10 show

evidence for the Mn304 phase. Obviously, the simplest explanation for ferromagnetic

behavior in this material is Mn304 precipitates.

In addition to magnetization, the optical properties were also examined. The

photoluminescence properties of the films were measured at room temperature using

a He-Cd laser (325 nm) and taken over a wavelength range of 350-800 nm. The power

density was 1 W/cm2. A 0.3 nm scanning grating monochromator with a Peltier-cooled

GaAs photomultiplier was utilized. The plot in figure 4-11 shows photoluminescence

spectra for Mn-doped Cu20 films deposited at 400 and 700 'C. The peak at 610 nm

corresponds to the 1 s exciton [100-102]. This peak is rather weak and broad, but is most

evident in the film grown at 700 'C. Note also the peak at -735 nm, which has previously

been associated with extrinsic defects in the Cu20 material [103]. Most of the additional

broad peaks could be attributed to the background photoluminescence from the substrate.

The emergence of intrinsic photoluminescence in the film grown at 700'C is consistent

with the segregation of Mn from the Cu20 matrix. It is also possible that luminescence

from either Mn2+ or Mn4+ also contributes to the spectra observed.

With the advent of new equipment available in the lab for measurements, resistivity

and Hall data were collected on films that were aged -3-4 years. There was no visual

decomposition of the films, but a detailed study of the microstructure was not conducted.

Measurements were performed in Van der Pauw configuration using a varying magnetic

field with a maximum strength of IT. Soldered indium contacts were placed on the corners

of the sample. Figure 4-12 shows the collection of measurements on various samples. The

films appear to become slightly more resistant with higher growth temperature, but the

trend is small and may be insignificant. The Hall mobility increases with higher growth

temperature, but there does not seem to be a significant relationship with growth pressure.









This may be an indication of better crystalline quality with higher growth temperatures,

which leads to enhanced mobility of the carriers. While the mobility increases, the carrier

concentration drops off with higher temperature. It is also important to remember that

different secondary phases are evolving in the microstructure as the growth conditions are

varied which presumably have an impact on the observed electrical properties.

Temperature dependent Hall effect data was studied on a newly deposited film grown

at 400C in ImTorr 02 on a MgA1204 substrate. Hall data was taken in a Van der Pauw

configuration with four 80nm thick platinum contacts sputtered on the corners of the

sample. Pt was chosen for it's high barrier height (~5.6eV) to make an ohmic contact to

the p-type film. The sample was measured in a temperature range from 200K to 400K.

The room temperature resistivity was 885 ohm-cm, but quickly became too resistive (>106

ohm-cm) to measure accurately below 200K. The acceptor state activation energy can be

calculated assuming an Arrhenius-type activation of the form: log p = log A + 2 oa

where p is the resistivity in ohm-cm, A is a constant, Ea is the activation energy, k is

Boltzmann's constant, and T is temperature in degrees Kelvin. The slope was calculated

by a linear least-squares fit through the data (inset of figure 4-13). Ea was determined

to be 0.25eV. This value is similar to the value (0.2eV) calculated for acceptor states

arising from Cu vacancies using first-principles DFT and DFT+U theories [84]. It is

also consistent with the 0.22-0.25eV values found for Cu20 films doped with Ni cations

[104]. The temperature dependent resistivity, mobility, and carrier concentration is given

in figure 4-13 with the log p vs. 1/T activation fit. The value of the resistivity and hall

resistivity at 300K as a function of field is given in figure 4-14. The positive slope of the

hall resistivity clearly indicates p-type behavior and is linear throughout the field sweep.

No indication of anomalous Hall effect is observed. The resistivity has a slight negative

MR (-0.4%) at 7 Tesla. The film became too resistive to measure at lower temperature.











































Figure 4-1. Cubic unit cell of Cu20. Oxygen atoms occupy bcc lattice positions that are
surrounded by a tetrahedron of Cu atoms. Cu atom are therefore linearly
coordinated to two oxygen atoms. This forms a rare occurrence of lineal
O-Cu-O bonding.


~LII



















-50-

*/ Cu20
/

-100 -

/ Cu


-150
200 400 600 800 1(
Temperature (oC)


Figure 4-2. Phase Stability curves for the Cu-Cu20-CuO system.


(a) Cu,O


10


20 30 40


50 60


'10L_5 IP 5!I !!! IL~! !!! I IIl !l!1!,
-105 -75 -45 -15 15 45 75 105


2-theta (deg) plu (deg)

Figure 4-3. X-ray diffraction data for epitaxial Cu20 on (001) MgO, showing both (a)
out-of-plane and (b) in-plane orientation.


















P(02) = 1 mTorr


700C
6000C
5000C
4000C
3000C


20 30 40 50 60 70 80


20 (deg)


Figure 4-4. X-ray diffraction data
an oxygen pressure of


for Mn-doped
ImTorr.


Cu20 films grown on (001) MgAl204 in


Cu20
(111)


Cu20
(200)


2000
1800
1600
1400
1200
1000
800
600















Cu20
(200)


Mn203 (222)
(222) ,,, nI


MgAI204
(400) P(02) = 0.1 mTorr

Cu20
(220)


2000
1800
1600
1400
1200
1000
800
600


...0 ...... 300 0C


3I
30


4I
40


5I
50


6I
60


7I
70


8I
80


2e(deg)
Figure 4-5. X-ray diffraction data for Mn-doped Cu20 films grown on (001) MgAl204 in
an oxygen pressure of 0.lmTorr.


7000C

6000C
5000C
4000C


to 1 APo -.- I ..- 9 ,I


16% 1L A 1

























Mn203
C (222)
(110) M
(110) \


2000-

1800-

1600-

1400-

1200-

1000-

800-

600-

400-

200-

0-


I I


MgAI2O4
(400)


50


I I


vacuum

Cu2O
(220)








700 C

S600 C

500 C


I I


I I


20 (deg)

Figure 4-6. X-ray diffraction data for Mn-doped Cu20 films grown on (001) MgAl204 in
vacuum.


2I I3
20 30





















L A


..TLI


300 400 500


1 ? !


A CuO2
* Mixed Phase
Cu20 + CuO
E Mixed Phase


Cu20
* Mixed
Cu20


+ Mn203
Phase
+ Mnx-10x


600 700


Temperature (oC)


Figure 4-7. Phase assemblage for films grown under different conditions.


k


o--L













2.2x10 -
2.0x10-6 T CuoMn
1.8x106
1.6x10-6 5000e
1.4x106
1.2x10 -
E 1.Ox10-
S8.0x1077
6.Ox10 .
< 4.Oxl10-
2.0x10 a .0
0.0
-2.Ox10" -
-4.0xl 07
-4. x10" .I .. I I r I I I
0 50 100 150 200 250 300

Temperature(K)
Figure 4-8 Magnetic behavior for an epitaxial Mn-doped CusO film grown at 300C and
ImTorr of oxygen





























i0-s MOAI204 + Aa paint


106-e 10K -ms=
1 B" = --'='


0.0 -




10 -
11 fl


-1000 -500 0
H(Gauss)


500 1000


1.0x1C05

5.0x10"







-1. 0x10"


N1011021 10/12/02

MaAJ204 + Aa Daint


100K ".





--



-1000 -500 0 500 1000
H(Gauss)


MgAI204 with Ag paint M100902D


0 50 100 150 200 250 300
Temperature(K)


1.6x10lO
1-4x10&
1.2x10-'
1.0x10-6
8.0x10-7
" 6.0x107
E 4.0x107
S2.Ox107
0.0
-2.0xl 07
-4.0x1 0-7
-6.0x10 '
-8.0x 10'


Temperature(K)


Figure 4-9. Magnetic behavior for MgA1204 substrate.


I Ox


5.0x





S.Ox


-1Ox'


MaA1204 + Aa paint


1o000e


-2.1x10-

-2.2x10

-2.3x1 05

-2.4x10-5

-2.5x10

-2.6x1 0'

-2.7x10-5

-2.8x10
















4.0x10-5
S 1x10 10K
5E 0 ---"
3.0x10-5 o -

-1x10 .-

E 2.0x105 -2x105

-5
<-5 -1000 0 1000
1.0x10 H(Gauss)


0.0 -
5000e

0 50 100 150 200 250 300
Temperature(K)

Figure 4-10. Magnetic behavior for an epitaxial Mn-doped CusO film grown at 700'C in
vacuum.

















0.02.
T = 7000C





. 0.01-
C:
a 0.009
. 0.008 T =400C
0.007
0.006

0.005
500 550 600 650 700 750 800 850
wavelength (nm)

Figure 4-11. Low temperature photoluminescence spectra for Mn-doped Cu20 films.
















0 Vaccum
E 0 1 mTorr
1 0 mTorr


D


0 0


300 400 500 600 700


400 500


400 500


Growth Temperature (oC)

Figure 4-12. Transport data for 1% Mn-doped Cu20 films. The resistiviy (a), mobility (b),
and hole concentration (c) are plotted as a function of growth temperature
for different oxygen pressures.


























S2slope =1280 74

Ea = 0 254 eV
2 21 so as o 4 s so s
10f/T (K 1)

0


0
CD
10 0C


1014 3


1013


200 250 300 350


(b)


O--


200 250 300 350 400


Temperature (K)



Figure 4-13. Temperature-dependent transport data for 1% Mn-doped Cu20 film. (a) The
resistiviy vs. temperature. The inset shows resistivity vs. 1/T with a linear
fit to calculate the activation energy of acceptor states above the valence
band. (b) The mobility [hollow circles] and the hole concentration [filled
triangles].






































(a)








000

,0 0
/'



300K ; "







-80k -60k -40k -20k 0 20k 40k 60k 80k
Applied field (Oe)


(b) 300K
0
o


0
0



O0
0


-80k -60k -40k -20k 0 20k 40k 60k 80k
Applied field (Oe)


Figure 4-14. Field-varying transport measurements for a 1% Mn-doped Cu20 film.
(a) Hall resistivity. Dotted line is a linear fit used to calculate the Hall
coefficient. The inset shows the derivative of the data to emphasize there
is no curvature from the anomalous hall effect. (b) Magnetoresistance.









CHAPTER 5
PROPERTIES OF ZnO CODOPED WITH Mn AND Sn

5.1 Introduction

As discussed in Chapter 2, several recent theories regarding the origin of ferromagnetism

in ZnO DMS emphasize the importance of holes in mediating the exchange interaction

between doped Mn atoms. Dietl's mean-field calculations predict that room temperature

ferromagnetism is possible in Mn-doped ZnO that is heavily doped with holes, while

carrier-mediated ferromagnetism in n-type material should be limited to lower temperatures.

The recent work by Kittilstved and coworkers demonstrate that the Mn+3 charge transfer

energy lies close to the valence band, similar in energy to ZnO acceptor states. It is

thought this can lead to large hybridization necessary to support ferromagnetic ordering.

The advancement of spintronics as a technology depends upon the development and

understanding of semiconductors that can support spin-polarized carrier operation at or

above room temperature.

In this chapter, the synthesis and magnetic properties of Mn-doped ZnO epitaxial

films codoped with Sn are examined. Codoping allows independent control over the

magnetic and electronic properties of the material by doping for each separately. In

II-VI materials, Mn+2 is isovalent and does not introduce carriers. By codoping II-VI

semiconductors, Mn provides the localized spins while an additional dopant can be used to

control the carrier concentration. This provides a convenient platform to study the effects

of carrier concentration on the observed magnetic properties in ZnO DMS. As a group

IV cation, Sn can exist in either the 4+ or 2+ valence, although the 4+ valence is most

common. As such, it can serve either as a doubly ionized donor or isoelectronic impurity.

For the ZnO films deposited in this work, Sn behaves as a donor. The magnetization

dependence on the carrier density is investigated.

Superconducting quantum interference device magnetometry measurements indicate

that the films are ferromagnetic with an inverse correlation between magnetization and









electron density as controlled by Sn doping. Magnetism in low free-electron density

material is consistent with the bound magnetic polaron model, in which bound acceptors

mediate the ferromagnetic ordering. Increasing the electron density decreases the acceptor

concentration, thus quenching the ferromagnetic exchange. This result is important in

understanding ferromagnetism in transition-metal-doped semiconductors for spintronic

devices.

5.2 Experimental

Epitaxial Mn, Sn-doped ZnO films were grown by conventional pulsed-laser

deposition. Laser ablation targets were prepared from high-purity powders of ZnO

(99.999%), with MnO2 (99.999%) and Sn02 (99.95%) serving as the doping agents. The

pressed targets were sintered at 1000'C for 12 h in air. The targets were fabricated with

a nominal composition of 3 at.% Mn and 0, 0.1, 0.01, and 0.001 at. % Sn. A Lambda

Physik KrF excimer laser was used as the ablation source. The laser energy density

was 1-3 J /cm2 with a laser repetition rate of 1 Hz and target-to-substrate distance of

6 cm. The growth chamber exhibits a base pressure of 10 5 Torr. Films were deposited

onto single-crystal, c-plane oriented sapphire substrates. Film growth was conducted

over a temperature range of 400-600'C. An oxygen pressure of 20 mTorr was used for

all film growth experiments. Film thicknesses were approximately 300-400 nm. X-ray

diffraction was used to determine the crystallinity and secondary phase formation.

Superconducting quantum interference device (SQUID) magnetometry was used to

characterize the ferromagnetic behavior of the doped films, focusing on the films grown at

4000C.

5.3 Results and Discussion

The phase stability and solid solubility of Mn in the ZnO films were investigated as

a function of growth temperature for films with varying Sn concentrations. Figure 5-1

shows the x-ray diffraction scans for films deposited under the given growth conditions.

In all cases, the dominant film peaks correspond to c-axis perpendicular ZnO. Note that,









for some of the films, a peak located around 64.55' appears. At first glance, this peak

was assigned to the (400) reflection of Mn304. The (400) diffraction peak of Mn304 has

a 20 value of 64.65' which closely matches the observed peak. However, the formation of

Mn304 at 400'C and not at higher temperatures in 0.1%Sn doped samples is peculiar and

should be questioned. Note that previous reports from Fukumura indicated that epitaxial

ZnO films with a Mn concentration as high as 35% could be achieved while maintaining

the wurtzite structure using pulsed-laser deposition [37]. The observed peaks are small

(a few hundred counts above background) and rather sharp. A similar peak is observed

in undoped ZnO films, which is included in Figure 5-1. Therefore, it is believed this peak

is associated with the ZnO and represents the KO artifact from the ZnO (004) peak. The

XRD used in these experiments has a Ni filter for attenuating Cu KO radiation, but the

KO line is clearly penetrating because the KO peak from the ZnO (002) peak is present

at 31'. The KO peaks can be checked by taking the d-spacing of the ZnO (002) planes

and calculating the respective 20 value for KO radiation (A=1.3922 A) using the Bragg

equation nl= 2d sinO. The 20 value for the ZnO(004) KO peak is determined to be around

64.6', which is commensurate with the observed peak. Note that even if the peak does

represent the formation of Mn304 in the film, the phase would not contribute to a high

temperature ferromagnetic signal since the Curie temperature is only 50K as mentioned

previously. Also notice that the precipitation of Sn-containing phases is not observed in

the diffraction scan, nor is it expected even if present as the nominal concentration of Sn

in the films is <0.1%.

The epitaxial nature of the ZnO films was determined using four-circle high-resolution

x-ray diffraction. Figure 5-2(a) shows an w rocking curve about the ZnO (002) peak for

the film grown on c-plane sapphire substrate at a growth temperature of 500'C and Sn

concentration of 0.1%. A '0 divergence slit was placed over the x-ray source and a 1mm

receiving slit was placed in front of the detector. The ZnO (002) rocking curve displays

a full width at half maximum (FWHM) of 0.47'. The in-plane alignment is evident in









the phi scan and pole figure of the ZnO (101) plane shown in figure 5-2(b). The 60'

peak intervals are consistent with the hexagonal symmetry of the epitaxial ZnO wurtzite

structure.

The Mn valence state in the ZnO lattice was investigated using x-ray photoemission

spectroscopy (XPS). Figure 5-3 shows the core level XPS spectra for a ZnO film doped

with 3%Mn and 0.01%Sn. The data was charge corrected by shifting the O is peak to

530.1 eV. The film was sputtered with Ar+ for 4 min. to remove surface contamination.

The Mn 2P3/2 binding energy is 640.7 eV. This is consistent with values assigned to Mn+2

in ZnO [105, 106]. There is a ~6eV energy difference between the Mn 2p3/2 and its higher

binding energy satellite peak. The binding energy and satellite peak are consistent to the

reported for single crystal MnO by Langell et al. [107]. They found the satellite structure

to be particularly sensitive to the oxide stoichiometry. In the case of Mn203 (the Mn+3

valence) the binding energy was higher at 641.1 eV and the satellite structure tended

to decrease for the higher oxide phases [107]. The binding energies for metallic Mn and

Mn+4 sit close to the Mn+2 value, but the energy for Mn has been seen at 637.7 eV and

that for Mn+4 at 642.4 eV in ZnO [105]. The binding energy and satellite structure for

the measured film suggests that most of the Mn doped into the ZnO is in the +2 valence

state.

The room-temperature resistivity for the Mn-doped ZnO films with varying

concentrations of Sn was determined using a four-point van der Pauw geometry. Defect

chemistry calculations for Mn-doped ZnO indicate that Mn+2 forms a donor level -2.0

eV below the conduction-band edge [108]. This deep donor state with Mn substitution on

the Zn site in ZnO has no direct effect on the electron concentration at room temperature.

However, defect chemistry calculations also indicate a reduction in Zn interstitials with

Mn doping. Zn interstitials are generally accepted as the primary shallow donor defects in

nominally undoped ZnO. This will yield an increase in resistivity for Mn-doped films as

compared to undoped material [108-110]. The Mn-doped ZnO films with no Sn exhibit a









resistivity on the order of 102 -cm with a carrier density of mid-1016/cm3. This carrier

density is substantially lower than that seen for undoped epitaxial films and is consistent

with the reduction of shallow donors. Limited results on the doping behavior of Sn in

ZnO indicate that it introduces a donor state [111-115], although in some II-VI compound

semiconductors, Sn is an amphoteric dopant, substituting on either the II or VI site

[116, 117]. For ZnO, the expectation is that Sn will substitute on the Zn site due to a

close match in ionic radii between Zn+2 (0.074 nm) and Sn+4 (0.069 nm). For the epitaxial

films considered in this work, Sn behaves as a donor. The resistivity of ZnO:Mn films with

various Sn content is shown in Table 5-1. The resistivity of the films drops rapidly with Sn

doping with a minimum of 0.185 2 cm for a Sn concentration of 0.1%. The most common

valence state of Sn is +4, yielding a doubly ionized donor if doped substitutionally on the

Zn site. Hall measurements indicate that the films are increasingly n-type with Sn doping

up to 0.1 at. %. It should also be noted that other work has shown that the addition of Sn

to ZnO ceramics also yields an enhancement in crystallinity [111, 112].

The magnetic properties of the films were measured using a Quantum Design SQUID

magnetometer. The diamagnetic responses of the substrate and host semiconductor were

subtracted from the magnetization plots. The primary focus of the measurements was

to determine how the magnetic properties of the films changed as a function of electron

density as controlled by Sn concentration. Samples that showed minimal amounts of

possible Mn304 precipitation via x-ray diffraction were used for the SQUID measurements.

All the M vs. H loops show hysteretic behavior with a finite coercivity and loop closure.

From the hysteresis curves, an increase in loop width is observed with increasing Sn

concentration. Figure 5-4 shows the coercive field as a function of Sn concentration,

suggesting domain pinning as the Sn doping is increased. It is unclear why the addition

of Sn enhances the hysteretic magnetization response in the Mn-doped films. It might

indicate enhanced pinning of domains due to the Sn dopants.









Most interesting is the saturation magnetization behavior as Sn content is increased.

As noted earlier, increasing Sn concentration increases electron density and conductivity.

Figure 5-5 shows the room-temperature magnetization versus field behavior for the ZnO

samples containing 3% Mn and Sn contents of 0, 0.1, 0.01, and 0.001%. Magnetization

is given as the magnetic moment per Mn dopant ion. Initially, there is an increase in

magnetization with minimal Sn doping. However, with increasing Sn doping, there is

an inverse correlation between the Sn content and saturation magnetization. As the

electron density increases with Sn doping, the magnetization decreases. The maximum

magnetization corresponds to a magnetic moment per Mn ion of -0.5 PB/Mn. This is

consistent with the bound magnetic polaron model in which only a fraction of the Mn ions

are expected to order ferromagnetically due to competing superexchange antiferromagnetic

interactions between neighboring Mn ions [118]. The inverse correlation of saturation

magnetization with electron density is interesting and provides some insight into the

mechanism for ferromagnetism in Mn-doped ZnO. Overlap of the Mn d-states with

the valence band suggests that holes are necessary in order to induce ferromagnetic

order. For semi-insulating films to exhibit ferromagnetism, the bound magnetic polaron

model provides a mechanism whereby holes that are localized at or near the Mn ions are

responsible for mediating ferromagnetism. The addition of electrons to the system will

move the Fermi energy level up in the band gap, resulting in a decrease in hole density

and a reduction in magnetization. This is consistent with Kittilstved and coworker's

observation where ferromagnetism was induced when the holes from the acceptor states

hybridize with the charge transfer state Mn. Ferromagnetism was observed when the ZnO

was locally doped p-type, but no ferromagnetism was observed when doped n-type. This

appears consistent with early work on trivalent-doped (Zn,Mn)O where no ferromagnetism

was observed for heavily n-type films. It may also explain the discrepancy from other

studies of Mn-doped ZnO films in which the intrinsic defect-mediated donor states are

high in density. It is important to note the need to maintain a Mn concentration low









enough to avoid MnMn antiferromagnetic interactions, which are likely to dominate high

Mn-doped ZnO films.

In conclusion, the magnetic and transport properties of Mn-doped ZnO thin films

codoped with Sn were examined. Results indicate that the films are ferromagnetic with an

inverse correlation between magnetization and electron density as controlled by Sn doping.

The results are most consistent with the bound magnetic polaron model in which bound

acceptors mediate the ferromagnetic ordering. Increasing the electron density decreases the

acceptor concentration, thus quenching the ferromagnetic exchange. This result is relevant

to understanding ferromagnetism in transition-metal doped semiconductors.

Table 5-1: Resistivity as a function of Sn content in codoped ZnO:3%Mn films.
Sn concentration
0.0% 0.001% 0.01% 0.1%
Resistivity 195 320 17 0.185
(Q-cm)










































30 40 50
2e (degrees)


30 40 50 60 70 80
2e (degrees)


30 40 50 60 70 80 46 48 50 52 54 56 58 60 62 4 66 68 70 72 74
2e(degrees) 2e(degrees)


Figure 5-1. X-ray diffraction of ZnO films codoped with Mn and Sn grown at oxygen
partial pressure of 20mTorr and growth temperature of 400, 500, and 600'C.
The diffraction pattern for an undoped ZnO film grown at 400'C in vacuum is
also shown. The peak at 20=64.55' is clearly evident in the undoped film and
attributed to the KO artifact from the ZnO (004) reflection.














ZnO (002)



\ FWHM=047


co (degrees)


ZnO (101)


V


0 50 100 150 200 250 300 350
'D (degrees)


Figure 5-2 X-ray diffraction of an epitaxial ZnO film doped with 3%Mn and 0 1%Sn that
was deposited at 500C and p(02)=20mTorr (a) an w-rocking curve of the
ZnO (002) peak with a FWHM of 0 470 (b) in-plane I-scan and pole figure of
the ZnO (101) planes


104






6 0x105




4 0x10'



C
2 20x105
C



























1030 1025 1020
Binding Energy (eV)


Binding Energy (eV)


Binding Energy (eV)


Figure 5-3. XPS spectra for ZnO:3%Mn film codoped
spectrum, (b) Mn 2p spectrum, and (c) O


with 0.01%Sn. (a) Zn 2p3/2
Is spectrum. The film was


sputtered in Ar for 4 min. and charge corrected to the O is peak.

























100-
---10K
100K
1lOOK
90- --- 300K


80-
0
70-
LL

S60-
o /

50


40-


30
0 0001 001 01
Sn Concentration (%)



Figure 5-4. Plot showing the dependence of the coercive field on Sn concentration at
different SQUID measurement temperatures.






















ZnO:3%Mn


300K


74eL


-C No Sn
-- 0.1% Sn
0.01% Sn
0.001% Sn


Applied Field (Oe)


Magnetization measured at 300K for epitaxial ZnO:3%Mn films
codoped with 0.001% Sn, 0.01%Sn, 0.1% Sn, and no Sn. There


be an inverse correlation of the Sn content with the saturation magnetization.


Figure 5-5.


that are
appears to









CHAPTER 6
PROPERTIES OF ZnO CODOPED WITH Mn AND P

6.1 Introduction

In the previous chapter, the magnetization dependence on carrier concentration was

investigated using Sn as a donor dopant. The magnetization had an inverse correlation

with electron density, suggesting that higher magnetization was favored as the Fermi level

moved down in the bandgap. If the magnetization decreases with electron concentration,

it might be expected that doping with holes will increase the magnetization. This chapter

investigates the magnetic properties of Mn-doped ZnO codoped with P, a known p-type

dopant in ZnO.

Superconducting Quantum Interference Device (SQUID) magnetometry measurements

indicate that the films are ferromagnetic with an inverse correlation between magnetization

and electron density as controlled by P doping. In particular, under conditions where

the acceptor dopants are activated, leading to a decrease in free electron density,

magnetization is enhanced. The result is consistent with hole-mediated ferromagnetism

in Mn-doped ZnO, in which bound acceptors mediate the ferromagnetic ordering.

Increasing the electron density decreases the acceptor concentration, thus quenching

the ferromagnetic exchange. This result is important in understanding ferromagnetism in

transition metal doped semiconductors for spintronic devices.

6.2 Experimental

Epitaxial Mn, P doped ZnO films were grown by conventional pulsed-laser deposition.

Laser ablation targets were prepared from high purity powders of ZnO (99.999%), with

MnO2(99.999%) and P205 (99.95%) serving as the doping agents. The pressed targets

were sintered at 1000C for 12 hr in air. The targets were fabricated with a nominal

composition of 3 at.% Mn and 2 at.% P. A Lambda Physik KrF excimer laser was used

as the ablation source. The laser energy density was 1-3 J/cm2 with a laser repetition

rate of 1Hz and target-to-substrate distance of 6 cm. The growth chamber exhibits a base









pressure of 10 5 Torr. Films were deposited onto single crystal, c-plane oriented sapphire

substrates. The growth temperature was 400'C. An oxygen pressure of 20 mTorr was

used for all film growth experiments. Film thicknesses were approximately 300 to 400 nm.

X-ray diffraction was used to determine the crystallinity and secondary phase formation.

SQUID magnetometry was used to characterize the ferromagnetic behavior of the doped

films. The impurity concentrations were measured by EDS and determined to be within

10% of the nominal concentrations.

6.3 Results and Discussion

The phase stability and solid solubility of Mn in the ZnO films were investigated

before and after annealing for films with P codoping. Figure 6-1 shows the x-ray

diffraction scans for films deposited under the given growth conditions. In all cases, the

dominant film peaks correspond to c-axis perpendicular ZnO. Note that, for these films,

segregation of the Mn304 phase is not observed in the diffraction data. As mentioned

earlier, the Mn304 phase is ferromagnetic with a Curie temperature less than 50 K.

Previous reports from Fukumura et al. indicated that epitaxial ZnO films with a Mn

concentration as high as 35% could be achieved while maintaining the wurtzite structure

using pulsed laser deposition [37]. Upon annealing, a shift in the d-spacing for the ZnO

is observed. This may indicate a segregation of either P or Mn in the films with thermal

processing.

Figure 6-2 shows high-resolution w-rocking curves taken around the ZnO (002) on a

separate but similarly grown film. The similar shift in d-spacing appears after annealing.

The scans were taken using a 2 divergence slit over the x-ray source and a 1mm brass

slit over the detector optics. The FWHM before and after annealing barely changes,

suggesting the crystallinity does not change sufficiently with the anneal. Furthermore,

in-plane alignment is measured with b-scans that show the six-fold symmetry of the ZnO

wurtzite structure and confirm epitaxial registry with the substrate. The surface roughness

of the film slightly decreases as measured by tapping-mode Atomic force microscopy









(AFM) (Figure 6-3). The grain structure has also notably coalesced and coarsened into

larger grains.

Defect chemistry calculations for Mn-doped ZnO indicate that Mn+2 forms a donor

level 2.0 eV below the conduction band edge [108]. This deep donor state with Mn

substitution on the Zn site in ZnO has no direct effect on the electron concentration at

room temperature. However, defect chemistry calculations also indicate a reduction in

Zn interstitials with Mn doping. Zn interstitials are generally accepted as the primary

shallow donor defects in nominally undoped ZnO. This will yield an increase in resistivity

for Mn-doped films as compared to undoped material [108-110]. The Mn-doped ZnO

films with no P exhibit a resistivity on the order of 102 Q-cm with a carrier density of

mid-1016/cm3. This carrier density is substantially lower than that seen for undoped

epitaxial films, and is consistent with the reduction of shallow donors. Limited results on

the doping behavior of P in ZnO indicate that it introduces a donor state for as-deposited

films, an acceptor state when annealed.

The behavior of phosphorus in ZnO epitaxial films both as-deposited and upon

annealing has been reported in detail elsewhere [119]. For the as-deposited films, the

inclusion of phosphorus yields a significant increase in electron density, resulting in ZnO

that is highly conductive and n-type. The shallow donor behavior in the as-deposited films

is inconsistent with P substitution on the O site, and presumably originates from either

substitution on the Zn site or the formation of a phosphorus-bearing complex. Previous

work has shown that the defect-related carrier density in nominally undoped ZnO can be

reduced via high temperature annealing in oxygen or air. In the case of undoped material,

the reduction in donor density is presumed due to either a reduction in oxygen vacancies,

Zn interstitials, or perhaps out-diffusion of hydrogen that is incorporated in the ZnO

lattice during synthesis. In order to reduce electron density annealing in oxygen can be

performed. Figure 6-4 shows the resistivity of films annealed at various temperatures.

Note that the resistivity of the as-deposited phosphorus doped films is significantly lower









than that for the nominally undoped film. For as-deposited films, a shallow donor state

dominates transport. As the films are annealed at increasing temperatures, the resistivity

of the phosphorus-doped films increases rapidly. This is particularly evident for films

subjected to annealing temperatures of 600'C or higher. When annealed at 700'C, the

phosphorus doped ZnO films become semi-insulating with a resistivity approaching 104

Q-cm. The conversion of transport behavior from highly conducting to semi-insulting with

annealing should be attributed to at least two factors. First, the defect associated with

the shallow donor state in as-deposited films appears to be relatively unstable. This would

explain the increase in resistivity, but would alone predict a saturation of resistivity at

the value given by the undoped material. The dependence of post-annealed resistivity

on phosphorus content suggests that a deep level associated with phosphorus dopant is

present. This is, in fact, consistent with the expected results that P substitution on the

oxygen site yields a deep acceptor. Figure 6-4 also shows the carrier concentration for

P-doped films. For all of the data shown in Fig. 6-4, the Hall sign was negative, indicating

n-type material. Both carrier density and Hall mobility data for some annealed samples

are absent in the plots. From the measurements yielding unambiguous Hall voltage, both

the carrier density and mobility in phosphorus-doped films are observed to decrease with

annealing. This is consistent with a reduction in the shallow donor state density and

activation of a deep (acceptor) level in the gap.

Figure 6-5 shows the longitudinal (p..) and transverse Hall (py) resistivity measured

for a ZnO film doped with 3%Mn and 2%P at room temperature. The film was measured

inside a Quantum Design Physical Property Measurement System (PPMS) modified with

an external high-impedance Hall effect set-up from Keithley electronics. The film was

measured in the Van der Pauw configuration with four indium contacts soldered to the

corners of a square sample. The as deposited film is n-type with a carrier concentration of

2.8x1015 cm3 and mobility of 0.56 cm2v is- The p., in zero-field is 3,975 2-cm. There

is a small negative magneto-resistance (MR) that reaches a value of about -0.35% at 7









Tesla. This small negative MR is also seen in undoped ZnO films at room temperature.

The resistivity is a strong function of temperature and quickly becomes too resistive, >105

2-cm, to measure at lower temperatures. After annealing the sample for 60min in a tube

furnace at 600' in latm 02, the sample resistance was measured to be >1092, which was

too high to measure accurately with practical settling times.

Phosphorus is believed to create acceptor states after being thermally activated by

annealing [119]. This requires the substitution of P 3 onto an 02 lattice site to form the

acceptor state. However, phosphorus exists in a variety of valence states, including +3,

+5, and -3. To investigate the P incorporation into the film, XPS was used to investigate

the charge state of the P ion. Figure 6-6 shows the XPS spectra for a ZnO:3%Mn, 2%P

film before and after annealing. There is a broad structure between 128 and 132 eV, and

then another peak centered around 133.6 eV. The P+5 valence state in P205 has a binding

energy of 135.6 eV [120]. The the P 3 ion of Zn3P2 has a lower binding energy of 128.7 eV

[121]. Note that pure phosphorus has a binding energy of 130.5eV [121]. Therefore, the

higher binding energy peak of Figure 6-6 is likely related to the +5 valence. The broad

structure at lower energy could be the existence of P 3 bonded to Zn cations. Therefore,

the data suggests a coexistence of phosphorus charge states. There does not appear to be

a significant shift in the binding energies after annealing.

Before annealing, the binding energy of the Mn 2p3/2 peak is 641 eV and there is a

definite satellite structure at higher energy. The spectrum is consistent with Mn in the

+2 valence state. After annealing, there is a slight peak shift to higher energy and the

satellite structure significantly decreases. Langell et al. have reported similar behavior for

single crystal MnO after annealing in oxygen, which was attributed to the formation of

higher oxidized Mn phases [107]. It's probable that there is a mixture of Mn valence after

annealing.

Optical absorption was measured using a Perkin-Elmer Lambda 800 UV/Vis

dual-beam spectrometer. The absorbance of each sample was measured using unpolarized









light with wavelengths ranging between 200nm to 900nm. Figure 6-7 displays the optical

transmission for the ZnO:3%Mn, 2%P film before and after annealing, along with an

undoped ZnO reference sample. The undoped sample shows a sharp exciton absorption.

This absorption is smeared out in the doped films and a broad absorption is seen around

3 eV in the doped films. There are two absorption peaks assigned to Mn2+ in ZnO at 2.95

and 3.26 eV from the 6A1(S)'4T2(G) and 6A1(S)4 A1,4E(G) transitions, respectively

[122]. These are spin-forbidden d-d transitions and should be weak. Therefore, the in-gap

absorption around 3 eV is most likely attributed the Mn2+ 6A1(S)A4T2(G) transition.

Direct interband transitions follow the Tauc relation ahv = Ao(hv-Eg) where a is the

absorption coefficient, hv the photon energy, and Ao is a parameter associated with the

transition probablility and refractive index. (ahv)2 vs. hv is plotted in the inset of figure

6-7. Straight lines are fit through the linear regions of the plots to extract the band gap

of each sample. The band gap widens with Mn-doping. Both the in-gap absorption and

blue-shift have been reported previously in Znl-_Mn.O films [37]. After annealing, the

band gap decreases slightly. This could be caused by Mn segregating out of the lattice at

high temperature, which would be consistent with the peak shift in the XRD data and the

disappearance of the satellite peak in XPS.

The magnetic properties of the films were measured using a Quantum Design SQUID

magnetometer. The diamagnetic responses of the substrate and host semiconductor

were subtracted from the magnetization plots. The primary focus of the measurements

was to determine how the magnetic properties of the films changed as a function of

electron density as controlled by P doping. Samples that showed minimal amounts of

Mn304 precipitation via x-ray diffraction were used for the SQUID measurements. Figure

6-8 shows the room temperature magnetization as a function of applied magnetic field

for epitaxial ZnO:3%Mn films both without and with P co-doping. For the Mn-doped

film with no P, saturation in the magnetization is observed, but with little evidence for

hysteresis in the M vs. H curves. The as-deposited ZnO film doped with both Mn and P









showed a reduction in magnetization/Mn ion. This is consistent with the proposed models

for Mn-doped ZnO, where ferromagnetic ordering is not favored by electron doping. Most

interesting is the saturation magnetization behavior as the P doped samples are annealed.

As noted earlier, increasing P concentration in as deposited films initially increases

electron density and conductivity. Figure 6-8 shows the room temperature magnetization

versus field behavior for the ZnO samples containing 3%Mn and 2 %P annealed in oxygen.

Magnetization is given as the magnetic moment/Mn dopant ion. Initially, there is a

decrease in magnetization with P doping. However, with annealing, there is an inverse

correlation between the electron carrier concentration and saturation magnetization.

Similar results are seen at 10 K as shown in Figure 6-9. Initially, as the electron density

increases with P doping, the magnetization decreases. The inverse correlation of saturation

magnetization with electron density is interesting and provides some insight into the

mechanism for ferromagnetism in Mn-doped ZnO. Overlap of the Mn d-states with

the valence band suggests that holes are necessary in order to induce ferromagnetic

order. For semi-insulating films to exhibit ferromagnetism, the bound magnetic polaron

model provides a mechanism whereby holes that are localized at or near the Mn ions are

responsible for mediating ferromagnetism. However, the most important observation is

that the activation of acceptor states for hole formation is necessary in order to achieve

ferromagnetism. The holes may be delocalized, but with low mobility, thus yielding low

conductivity. In this case, the carrier mediated mechanism may suffice without the need

of invoking bound polarons as inherent to the ferromagnetic ordering. In either case, the

addition of electrons to the system will move the Fermi energy level up in the band gap,

resulting in a decrease in hole density and a reduction in magnetization. This appears

consistent with early work on trivalent doped (Zn,Mn)O where no ferromagnetism was

observed for heavily n-type films. It may also explain the discrepancy from other studies

of Mn-doped ZnO films in which the intrinsic defect-mediated donor states are high in

density. It should be noted that the amount of magnetization in the material remains









relatively low at all temperatures as seen in the temperature dependent magnetization

data in figure 6-10. It is important to note the need to maintain a Mn concentration low

enough to avoid Mn-Mn antiferromagnetic interactions, which are likely to dominate high

Mn doped ZnO films.

The results of this study are consistent with previous studies on the carrier type

dependence in Co- and Mn-doped ZnO nanocrystalline films [32, 49]. In that case,

ferromagnetism was observed in Mn-doped ZnO nanocrystals only when nitrogen, a

group V acceptor dopant, was introduced during the synthesis process. In contrast,

no ferromagnetism was observed in Co-doped ZnO nanocrystals when processed with

nitrogen. Based on this and other properties, the authors conclude that ferromagnetism in

ZnO is closely tied to the charge transfer electronic structure of the transition metal

dopant. For Mn, ferromagnetism is induced when the holes from the acceptor ion

delocalize onto Mn2+. Again, our results are consistent with this conclusion.










104
SZnO: 3%Mn, 2%P
o
3 0
10
o0 0
*-- o
13 A Annealed

4
M 10

103
C
O ON
C N 2

-10
SAs Grw As Grown o06


25 30 35 40 45 50 55 60 65 70 75
29 (deg)
lx104

ZnO (002)

d=0.2591 nm
03 Annealed
.p 5x103

(D
4--

SAs Grown d=0.2605nm
0

33.0 33.5 34.0 34.5 35.0 35.5 36.0
20 (deg)

Figure 6-1. X-ray diffraction of ZnO films codoped with Mn and P both before and after
annealing. The target was ZnO doped with 3% Mn and 2%P

















As Grown


29 = 34.3990
FWHM=1.90640


ZnO (112)


S(deg)


19 20


ZnO (002)


29=34.5190
FWHM=1.94810


14 15 16 17 18


ZnO (101)


o 10 10 200 20 300
(deg)


19 20


0 (deg)


Figure 6-2. High-resolution w-rocking curves on ZnO films codoped with Mn and P before
and after annealing. The scans are taken around the ZnO(002) reflection. The
insets show <-scans around the indicated peaks, showing the film's in-plane
alignment with the substrate.


25k -



20k-


ZnO (002)


30 Annealed













RMS roughness: 2.38 nm


(a) As Grown









200mu




=(b) Annealed -









200nm


Figure 6-3.


AFM scans on epitaxial ZnO:3%Mn, 2%P films before and after annealing.
Scans were taken in tapping-mode. The surface roughness decreased and
the grains have coalesced into larger grains with annealing. (AFM imaging
software provided by [123]).


RMS roughness: 1.91 nm























P-doped ZnO

U--


10



1 o



0.1
,)

0.01



0.001


Annealing Temperature(C)


Resistivity and carrier concentration behavior of P-doped ZnO films both
as-deposited and annealed in oxygen. The data shown is for 1 at.%P doped


Figure 6-4.



















25-

20 (a)

15-

10-

05-
E
S00o- 300K S

-0 5 -

-10-
Hall Coeff -2226 3 (cm3 C )
-15 Carriers 28x105 (cm3)

-20- Mobility 056 (cm2V1 s1)

-2 5-

-80k -60k -40k -20k 0 20k 40k 60k 80k
Applied Field (Oe)



3985
Initial Sweep ) 02
o Retum Sweep 0 -(b
3980-
01

300K o
3975- 00
0 0 000 o00O

E0 0
7 00 0O
O l* -02

3965- 0 O0 o

S*0

3960-
-04


3955 -0 5
-80k -60k -40k -20k 0 20k 40k 60k 80k
Applied Field (Oe)



Figure 6-5. Transport data for the as-deposited ZnO:3%Mn, 2%P film at 300K. (a)
Transverse Hall (py) and (b) longitudinal (p..) resistivity.














































Binding Energy (eV)


130 129


5 5k

137 136 135 134 133 132 131
Binding Energy (eV)


Binding Energy (eV)


(d) Mn 2p /42p_












60 6B55 650 645 640
Binding Energy (eV)


Figure 6-6. XPS spectra for ZnO:3%Mn film codoped with 2% P. (a) O is spectrum, (b)
Zn 2ps spectrum, (c) P 2p spectrum, (d) Mn 2p spectrum. The data was
charge corrected to the 0 is peak at 530.1 eV.


(c) P 2p












































15 20 25 30 35 40 45 50
Photon Energy (eV)



Room temperature optical transmission for ZnO:3%Mn, 2%P films both
as-deposited and after annealing at 600'C in oxygen. Data for an undoped
ZnO film is also included. Inset: Tauc plots with linear fits to determine the
optical band-gap for each film.


Figure 6-7.






































-01



-0 2


-1200


-o- 3% Mn, No P

-i- 3% Mn, 2% P As-grown
--- 3% Mn, 2% P Annealed





300K


Low e- density


High e- density


-400 0 400


Applied Field (Oe)


Room temperature SQUID measurements for epitaxial ZnO:3%Mn,2%P films
before and after anneal. Also shown is a ZnO:3%Mn film with no P.


Figure 6-8.
























-0-As Grown: 3% Mn, 2% P
-- 2 Annealed: 3% Mn, 2% P


-020--
-1200


Applied Field (Oe)


Figure 6-9. SQUID measurement at 10K for epitaxial ZnO:3%Mn, 2%P films before and
after annealing.




















-8.0x10o7



-1.2x10-6



-1.6x10-6



-2.Ox10-6


0 50 100 150 200 250 300


Temperature(K)


Figure 6-10. Field-cooled and zero field-cooled magnetization measurements for a
ZnO:3%Mn, 2%P film annealed at 600C in 02.









CHAPTER 7
PROPERTIES OF COBALT-DOPED ZnO

7.1 Introduction

The two previous chapters investigated the properties of Mn-doped ZnO as a

potential material for spintronic applications. In this chapter, cobalt is investigated as

a magnetic dopant in ZnO. The magnetic and magneto-transport behavior of doped films

were examined. Cobalt concentration is varied over a wide range, from 0 to 30 at.%Co. A

combination of x-ray diffraction, optical absorption, and transmission electron microscopy

were used to examine the solubility of cobalt in the ZnO lattice and phase segregation of

cobalt metal.

Films deposited at 400'C in vacuum were found to be ferromagnetic, while films

deposited in oxygen or at higher temperatures were found to be nonmagnetic. Segregation

of cobalt metal occurs in films doped with 15 at.% or greater Co concentrations when

deposited in vacuum, and the precipitates are found to be oriented within the lattice. The

segregation can be suppressed by depositing in higher base pressures, but the process is

not fully reproducible. Peculiar MR is observed in the films and the MR changes as the

carrier concentration crosses the metal-to-insulator transition.

7.2 Experimental

Cobalt-doped ZnO films were deposited via pulsed laser deposition (PLD) onto

c-plane oriented sapphire substrates. The ablation targets were prepared through the

solid-state reaction of mixed oxide powders. Appropriate amounts of ZnO (Alfa Aesar,

Puratronic, 99.9995%) and Co304 (Alfa Aesar, Puratronic, 99.9985%) powders were

ground and mixed in methanol, dried in air, pressed into pellets, and sintered at 1000C

for 12 hours in air. The targets were mixed to give proportions of Znl -Co.O with

x=0.00, 0.02, 0.05, 0.15, and 0.30. A KrF excimer laser (248nm wavelength) was used for

target ablation using a repetition rate of 1Hz and a laser energy density of 1-3 J/cm2. A

temperature range of 400'C-600'C and oxygen pressures up to 2x10 5 Torr were used in









the experiments. The vacuum (base pressure) of the chamber was ~7.0x10 6 Torr. Film

thicknesses were 200-300nm as measured by mechanical profilometry.

7.3 Results and Discussion

7.3.1 Chemical Composition

Energy Dispersive Spectroscopy (EDS) was used to measure the percentage of cobalt

in the films. Figure 7-1 shows the cobalt concentration in the film as a function of target

composition for films deposited under different conditions. The cobalt concentration is

generally larger in the films than the prepared targets. This likely occurs because Zn

has a higher vapor pressure than Co, thus less Zn is incorporated into the film during

deposition. The cobalt concentration is also higher with increased substrate temperature.

7.3.2 Structure and Phase Analysis

7.3.2.1 Films with precipitation

Crystal structure and phase analysis were characterized using X-ray diffraction (XRD)

in Bragg-Brentano geometry. Figure 7-2 shows the 8-20 x-ray diffraction patterns for

a series of films grown at 400'C in vacuum (base pressure 7x10 6 Torr) with varying

cobalt concentration. The primary peaks correspond to the wurtzite ZnO (002) indicating

good texture with the c-plane of the sapphire substrate. As the cobalt concentration is

increased above 10%, the appearance of a new peak begins to develop around a 20 value of

44.4 degrees. The peak is small (only a few hundred counts above background) and does

not correspond to any ZnO or substrate peaks. Both the small intensity and 20 position

make identification of the peak difficult using XRD as there are several cobalt containing

phases with similar 20 values of around 44.4 degrees, including the spinel family of cobalt

oxides and cobalt metal. However, exact determination of the peak is critical since the

presence of ferromagnetic cobalt metal could contribute to the magnetic signature of the

films. The cubic and spinel cobalt oxides are antiferromagnetic, though some papers report

that small nanocluster powders of cobalt oxides are ferromagnetic due to uncompensated









surface spins [124, 125]. Table 7-1 is a list of the possible phases with their respective 28

diffraction values and magnetic character.

High resolution XRD and TEM were used to characterize the secondary phase

observed in the powder XRD scans. Cross-sectional TEM was used to more precisely

delineate the nature and location of the extra phase as shown in figure 7-3. Most of

the precipitation occurs near the film/substrate interface in the form of small (-5nm)

particulates. However, for most of the film, the cobalt dopant appears to reside in the

ZnO lattice without precipitation. Convergent beam TEM diffraction patterns of the film

and nano-precipitates are shown in figure 7-4. The particles appear to be oriented with

the lattice and have d-spacings of doo2= 0.20nm and d 210=0.13nm. This is consistent

with metallic cobalt and the particles are tentatively assigned as such. The peak intensity

was too low to collect information using high-resolution XRD on the as grown film.

However, after annealing the samples in hydrogen (4%H2/Ar balance) at 500'C, the

amount of cobalt increases at the expense of ZnO causing partial degradation of the film,

and high-resolution XRD could be employed which is discussed in a later section. The

change in microstructure at low annealing temperature suggests that the cobalt that is

substitutional in the ZnO lattice is not stable at moderately high temperatures.

7.3.2.2 Films without precipitation

By modifying the growth conditions, the secondary phase can be suppressed. Figure

7-5(a) shows XRD scans for samples grown at 400'C at different pressure conditions.

The cobalt phase forms at low base pressures. By depositing in higher base pressures

(>10 5 Torr) or adding a small amount of oxygen, the Co phase disappears from the XRD

scans. It should be noted that the process is not fully repeatable. The cobalt phase is still

present by XRD for some depositions but seems less probable at the higher base pressures.

As a rough estimate, the phase is suppressed in 80-90% of the samples grown in higher

base pressures. It is assumed the pressure at vacuum is composed mostly of water vapor.









Figure 7-5 also shows XRD scans for films deposited at 500'C and 600'C in vacuum.

At 500'C the CoO phase begins to form along with the cobalt metal. At 600'C the cobalt

metal phase disappears and the CoO phase becomes more prominent. Thermodynamically,

CoO is the most stable phase at these temperatures and pressure. This can be seen

from the thermodynamic predominance diagram for the cobalt oxides given in figure 7-6.

Interestingly, one would expect the formation of Co304 at lower growth temperatures

and not the metallic cobalt phase determined from TEM. This suggests the cobalt metal

phase is stabilized by the ZnO lattice and not by thermodynamic considerations. This is

consistent with the well oriented particles observed in TEM.

7.3.3 Optical Properties

7.3.3.1 Optical absorption

Evidence for Co substitution in the ZnO lattice can be inferred from optical

absorption measurements. Figure 7-7 shows transmission data for Co-doped ZnO films

that do not show precipitation by XRD. An undoped ZnO film is included as reference.

Each film was normalized by dividing by its maximum observed transmission (T/Tmax)

to compare the intensity of absorption peaks between films. Three adsorption peaks are

apparent in the doped films. These peaks are characteristic d-d transition levels attributed

to Co+2 occupying tetrahedral lattice positions, and indicate that cobalt is substituting

as Co+2 on Zn lattice sites in the films [45, 126, 127]. The intensity of these absorptions

in the doped films also increases with increased Co concentration suggesting most of the

Co is soluble in the lattice. Specifically, the peaks located at energies of 1.9 eV (651nm),

2.0 eV (608nm), and 2.2 eV (564nm) correspond to the 4A2 2E(G), 4A2A4T1(P), and

4A2 2A1(G), respectively [128].

The band-gap of the alloys was calculated by plotting (ahv)2 vs. hv and extrapolating

the linear portion of the plot to (ahv)2 = 0. The plots are given in figure 7-8. At low

cobalt doping-the undoped film and the film doped with 2%Co-the absorption edge

is well defined and can be fit reliably. The exciton peak is clearly visible in the undoped









film indicating good quality ZnO. However, as the cobalt doping is increased, a low energy

absorption onset appears and the doping smears out the linear region giving it a more

rounded shape. Linear fits to both the high and low energy regions are given in figure 7-8

as a function of cobalt concentration. The high energy slopes give a band-gap energy that

linearly increases (blue-shifts) with nominal cobalt concentration. The low energy slopes

give an energy that roughly remains constant around 2.8-2.9 eV.

Some reports in the literature observe a red-shift in the band gap energy as the

cobalt concentration is increased [129-132]. The red-shift is typically attributed to the

sp-d exchange between the ZnO band electrons and localized d-electrons associated with

the doped Co+2 cations. The interaction leads to corrections in the energy bands; the

conduction band is lowered and the valence band is raised causing the band gap to shrink

[133].

On the other hand, other papers have reported a blue-shift in the band gap of ZnO

with cobalt doping. Peng et al. reported a blue-shift in the band-gap of the material

and a red-shift of the band tails, which is similar to our observations [134]. Yoo et al.

also observed a blue-shift in Al and Co codoped ZnO films which were attributed to the

Burnstein-Moss effect from an increase in the carrier concentration [135].

The blue-shift in the gap energy is probably not caused by the Burnstein-Moss effect

in these films. In a heavily doped n-type semiconductor the Fermi level resides in the

conduction band. The free electrons in the semiconductor fill the lowest states in the

conduction band and the valence electrons can no longer be optically excited into these

filled states. This results in an apparent increase in the onset of absorption and the

gap shifts to higher energy [79]. This shift requires an increase in the electron density.

However, there is no systematic increase in the carrier density of the measured films with

additional cobalt.

As discussed earlier in Chapter 3, band tails can arise from perturbations in the band

structure caused by impurities and disorder. States introduced by impurities overlap at









high concentrations and evolve into an impurity band. As seen in figure 7-8, the tails'

intensity rises with increased cobalt. Therefore, the tails are most likely related to the

increase of impurity states. Kittilstved and coworkers saw a large MCD peak at -25,000

cm 1 (3.10 eV), which they assigned as a valence band-to-metal charge transfer (CT)

transition in ZnCoO [32]. The level was approximately 2,600 cm1 (0.322 eV) below the

conduction band. The onset of absorption in figure 7-8 starts around 2.9 eV. This is below

the value found by Kittilstved; however, their MCD peak had some breadth with the

onset of the peak beginning -23,000 cm1 (2.85 eV), which is in good agreement with

the absorption onset in figure 7-8. Therefore the low energy absorption onset is likely an

electron transition from the valence band to the cobalt impurity states.

7.3.3.2 Photoluminescence

Photoluminescence (PL) was performed on the series of films grown at 400'C in

vacuum. The PL was used primarily to verify the band-gap shift found in the optical

absorption data. Figure 7-9 shows the PL spectra taken at room temperature. For the

undoped film, a broad luminescence band is visible across most of the spectrum. Broad

green-yellow bands are typically attributed to defects in ZnO, including oxygen vacancies

[136]. The films are grown in low oxygen pressures and are non-stoichiometric which could

give rise to the broad defect bands. A more detailed study of the low temperature PL

properties would have to be done to say more about the bands. The dip in the intensity

at -517 nm is an artifact due to the blaze angle of the diffraction grating and not from

the film properties. The band-edge is visible in the undoped and 2%Co doped films, but

is quenched with higher cobalt doping. Higher resolution scans around the band-edge

emissions of the undoped and 2%Co samples are shown in the inset of figure 7-9. The

band-edge is blue-shifted from 3.25 eV to 3.31 eV in the 2%Co sample. This is consistent

with the band-gap determined from optical absorption.









7.3.4 Magnetic Characterization

The volume magnetization of the films was measured using a Quantum Design

superconducting quantum interference device (SQUID) magnetometer. Before measuring,

the backs and sides of the samples were etched in nitric acid (50% nitric/ 50% deionized

water) to remove excess silver paint and contaminants that could contribute a spurious

magnetic signal as measured by the SQUID. Each film surface was first coated in

photoresist and baked for 20min. at 50'C to help protect the film during etching and

then floated on top of the nitric acid for 3 min. The photoresist was removed by rinsing in

acetone.

The SQUID magnetization data is normalized using two common methods:

normalizing by the film volume and normalizing by the number of bohr magnetons

(pB) per Co atom. To convert the raw magnetization data from emu to PB/Co, all the

cobalt atoms are assumed to substitute on Zn sites and a cation density of 4.18x1022 cm3

is used for the conversion. The data is normalized using the cobalt concentrations from the

EDS data given in figure 7-1.

Magnetization data for a series of films grown at 400'C in vacuum with different

amounts of cobalt are displayed in figure 7-10. These films do not show cobalt-induced

secondary phases by XRD. The films with a nominal concentration of 2% and 5% cobalt

are ferromagnetic with hysteresis at 10K; however, the saturation magnetization gradually

decreases with increased temperature and the hysteresis disappears. The film with 30%Co

shows ferromagnetism up to 300K with a clear hysteretic shape. The sample has a room

temperature magnetization approaching 0.08 PB/Co.

The magnetization data for films with and without secondary phase formation are

compared in figure 7-11. As stated previously, the main difference between the films is the

growth pressure used during deposition. The 30%Co film with the secondary phase has

a larger magnetization than the film without. This would be expected if the secondary

phase is cobalt metal. Notice the film deposited in oxygen shows very little magnetization.









This is also true for films that were grown in higher oxygen pressures (up to 20mTorr

pO2). The film with cobalt precipitation has a room temperature magnetization that

saturates -0.22 ps/Co at higher fields (>1 Tesla). The film without precipitation has

a room temperature magnetization approaching 0.08 pB/Co. Both these values are

well below the 3ps moment of the (d7) high-spin configuration of Co+2. They are also

below the 1.72 pB/Co value of pure metallic cobalt. However, the values are in good

agreement with other values reported in the literature for ZnCoO (0.08 to 0.4 u/P/Co)

[48, 67, 137]. Recent LSDA+U calculations predict that the nearest-neighbor exchange

couplings of cobalt in ZnO should be antiferromagnetic, irrespective of the geometrical

nearest-neighbor arrangement [138]. Therefore at high cobalt concentrations, where

statistically a greater number of nearest-neighbors will develop, one would expect a lower

moment per Co than that given by lower Co doping concentrations. However, this is not

what is observed in the present case.

7.3.5 Electrical Transport

7.3.5.1 Hall effect

The origins of ferromagnetism in cobalt-doped ZnO are still not fully understood.

Whether the ferromagnetism is truly intrinsic as a result from the interaction of carriers

with magnetic dopants, or if the ferromagnetism is extrinsic and arises from secondary

phases or nanoclusters is an important consideration. The usefulness of a DMS rests on its

ability to produce and manipulate spin polarized currents. If the ferromagnetism in these

materials is solely localized in secondary phases and does not polarize the free carriers,

then the DMS is of limited utility in spintronic device applications.

A DMS that contains an asymmetry in the carrier spin-density (a spin-polarized

current) should exhibit an anomalous Hall effect (AHE) in transport measurements. For

ferromagnetic materials the Hall equation is given by:

p, = RoB +RsM









where py is the Hall resistivity, Ro is the ordinary Hall coefficient, Rs the anomalous

Hall coefficient, and B and M are the magnetic flux and magnetic field vectors normal

to the film surface. Ro is caused by the Lorentz force acting on moving charge carriers.

The anomalous term, Rs, is usually ascribed to spin-dependent scattering of carriers at

local atomic moments. Carriers of opposite spin are scattered in different directions at

each moment. This modifies the charge accumulation at each end of the sample. At higher

fields, the magnetization, M, will saturate causing the anomalous Hall component, RsM, to

become constant. At this point, changes in the Hall curve are determined by the ordinary

component and should be linear in field. The ordinary Hall coefficient can be extracted

from the high-field, linear region and the electrical properties of the material determined.

Variable-field Hall effect measurements were performed on Hall bridge patterned films.

The films were patterned during growth by depositing through a stainless steel shadow

mask. Electrical contacts were soldered to the sample using indium metal. The films were

mounted inside a Quantum Design Physical Property Measurement System (PPMS) to

control the ambient temperature and applied magnetic field, and the electrical data was

collected using a Keithley high-impedance Hall effect system (the components of this

set-up are described in the appendix).

Both the transverse (py) and longitudinal (p.) resistivity were measured by applying

a longitudinal current (I..). Measured values of py will have a contribution from p if

the voltage leads are slightly misaligned. Therefore the resistance data was geometrically

and field averaged to remove any asymmetries and to account for thermally induced

voltages. Since py is antisymmetric with applied magnetic field, the data was averaged

by pxy,odd(H) = I[pzy(+H) pzy(-H)] to help remove any parts of the signal from the

longitudinal component [139]. Conversely, p., is symmetric with respect to applied

magnetic field and can be averaged using p,evn,,,(H) = 1[p.=(+H) + p.(-H)].

For the ZnO films doped with 30% cobalt, an anomalous Hall signature is observed in

samples grown at 400'C and in vacuum as given in Figure 7-12. Included in the figure is a









film with and without cobalt precipitation. Both films show AHE at room temperatures,

but the effect is much more pronounced in the film with cobalt precipitation. Derivatives

are shown in the insets of each graph to show the change of slope created by the AHE.

Interestingly, there is no evidence of hysteresis in either Hall curve within the resolution

of the instrumentation. A weak AHE is also observed in the film doped with 15%Co. For

films doped with less than 15% Co, no AHE is observed at the chosen growth conditions.

Although the AHE is typically attributed to scattering of carriers by local magnetic

moments, it has been suggested that nonmagnetic materials with ferromagnetic precipitates

can also exhibit an AHE. This has been observed in Co-doped TiO2 films by Shinde et

al. [140]. They report an AHE in superparamagnetic, highly reduced cobalt-doped rutile

TiO2 films that contain cobalt metal clustering close to the film/substrate interface.

Superparamagnetic granular metal composites are also known to show anomalous Hall

behavior [141, 142]. While the larger AHE is suspect in the films with precipitation, the

origin of the AHE in the non-precipitated films is not obvious. The AHE could arise

from the cobalt atoms that are substitutional in the ZnO lattice or possibly from small

nanoparticles of cobalt metal not observable with XRD. Evidence for a true AHE from

substitutional cobalt atoms is inferred from annealing experiments, which is discussed

shortly.

It is important to examine the possibility of cobalt particles as the source of

ferromagnetism. The films with 30%Co grown at 400'C show hysteresis with a finite

coercive field and remanence at 300K from the SQUID magnetometry data. Superparamagnetic

particle systems will not show hysteresis above their blocking temperature, where the

temperature is high enough for the superparamagnetic moments to fluctuate faster than

experimental measuring times. This suggests that if the cobalt precipitates are behaving

as superparamagnets, then their blocking temperature should be higher than 300K.

The blocking temperature may be estimated using [140, 143]:

TB 25
B -25k,









where K is the magnetic anisotropy constant (K = 4.1x105 J/m3 for hexagonal cobalt

[144], V is the particle volume, and kB is Boltzmann's constant. Using this equation, the

estimated diameter for a spherical cobalt particle with a blocking temperature of 300K

is -7.8nm. This is close to the particle size seen in the 30%Co film from TEM, which is

about 5nm. The smaller particle size should show lower blocking temperatures, suggesting

the blocking temperature in the films should be less than 300K, which is inconsistent with

the magnetization curves at the same temperature. This suggests that if the particles are

superparamagnetic, they do not contribute to the ferromagnetic moment seen in the films

at room temperature. However, since TB scales with the cube of the particle radius, a

small deviation in the particle size will have a large effect on the blocking temperature

and makes such an argument hard to justify based solely on the blocking temperature

equation.

7.3.5.2 Magnetoresistance

The magnetoresistance (MR) of the films was measured simultaneously with the

Hall data. The MR can provide some useful insights into the transport properties of

semiconductors, including the potential landscape of the impurity distribution and

lattice disorder. Aluminum was added as a codopant to some of the films to study the

effects of electron concentration on the transport properties. Aluminum is a well know

n-type dopant in ZnO; Al+3 substitutes on a Zn+2 site and dopes an electron into the

lattice. Changes in the MR behavior at low temperature as a function of the electron

concentration are found to be consistent with the notion of a critical carrier concentration

at the Metal-to-Insulator Transition (MIT). The cobalt doped films transition from

positive to negative MR as the electron concentration crosses the MIT.

In order to understand the MR results, a brief discussion of the metal-to-insulator

transition is in order. Doped semiconductors will become metallic when sufficient overlap

overcomes the localizing effects of electron-electron correlation and disorder. At a certain









concentration, electrons delocalize and their wavefunctions extend throughout the lattice.

This is a well known metal-to-insulator transition in semiconductor materials.

At dilute limits, impurity states are isolated and electrons are confined to a

hydrogenic orbital around their associated impurity. The radius of this orbital if given

by r=H (m/m*)aH, where c is the relative dielectric constant of the host material, m* is

the carrier effective mass, and aH is the hydrogenic Bohr radius (aH= 53pm) [31, 145].

As the impurity concentration is increased, the impurity orbitals begin to overlap to form

an impurity band. The amount of overlap required to overcome the localizing effects of

correlation is given by the Mott criterion,

(n,) rH 0.25

where n, is the critical concentration. The critical concentration for ZnO is about

4x1019 cm 3. Below the critical concentration carriers remain bound to localized sites, but

can move by hopping between occupied and empty states. This is the insulating regime of

the MIT. Above the critical concentration, the impurity band states become delocalized

and the carriers are itinerant.

As a reference, the MR behavior of an undoped ZnO film deposited in vacuum at

600'C is shown in figure 7-13. The MR is negative over the entire field range and is

dependent on temperature. This is consistent with other reports found in the literature

[78, 146, 147]. The negative MR in these reports was observed for highly n-type ZnO films

with electron concentrations exceeding 1020. Negative MR in highly doped semiconductors

is thought to be caused by the weak localization correction to the conductivity [148]. Hall

measurements on the film presented in figure 7-13 give carrier concentrations on the order

of 1018. This is well below the critical concentration of the MIT and demonstrates that the

MR of nonmagnetic ZnO is negative on both sides of the MIT.

Figure 7-14 shows the MR behavior at various temperatures for three films doped

with 5%Co. These films were deposited under slightly different conditions to impart

different electron concentrations. The growth conditions are indicated in the figure for









each film. The electron concentrations as measured by the Hall effect are also provided

in the figure. 2 at% Al was added to the third film to substantially increase the electron

concentration. The MR for the three films is negative and qualitatively similar above

100K. The MR at high temperature is rather small being less than -0.5%. However, at low

temperature the MR is substantially different between the films. At 10K, a progression

from positive to negative MR occurs as the electron concentration is increased. Below

the critical concentration of the metal-to-insulator transition, the MR is positive over the

entire magnetic field. Near the critical concentration, a small negative MR component

appears at low field. The MR value decreases from -10% to -1.5% as the MIT is

approached. At concentrations much larger than the critical concentration (in the metallic

conduction regime) the MR is negative over the entire field range, similar to the MR

behavior of the undoped ZnO film.

The MR behavior was also studied for films heavily doped with cobalt. Figure 7-15

shows the MR data collected for two films, one doped with 30% Co and the other codoped

with 30% Co and 2% Al. Again the Al is added to increase the electron concentration

of the doped film. A similar progression from positive to negative MR across the MIT

boundary is seen at low temperature. At higher temperature, the film doped only with

cobalt, develops a kink in the MR curve near 10,000 Oe. To study the kink formation, the

shape of the MR curve was tracked at temperatures between 10K to 100K and is shown

in figure 7-16. The relative magnitude of the negative MR peak at zero-field gradually

increases with temperature. Between 50K and 75K, this peak increases and the MR

component at high-field becomes negative. The MR of the Co- and Al- codoped films

is remarkably similar to the 5%Co, 2%A1 codoped film, and qualitatively similar to the

undoped film. This suggests that electronic transport is not greatly affected by the cobalt

concentration above the MIT.

Figure 7-17 shows the temperature dependence of the resistivity for the 5% and

30%Co films. In the insulating regime, the resistivity decreases as the temperature is









increased. This is typical behavior for semiconductor films. Above the MIT, the resistivity

increases with increased temperature, typical of metallic behavior. The temperature

dependence is consistent with the notion that the Al doped films have crossed over into

the metallic conduction regime.

The progression from positive to negative MR in figures 7-14 and 7-15 is commensurate

with the metal-to-insulator transition. On the insulating side of the MIT, the MR

is positive, and on the metallic side of the MIT, the MR becomes negative. Near

the transition point, negative MR is seen at low field and positive MR at high field.

However, as the temperature is increased, the positive MR component vanishes and only

negative MR is seen. A similar dependence of MR on carrier concentrations around the

MIT in cobalt-doped ZnO have been reported by Xu et al. and Kim et al. [51, 149].

Clearly, a model including the interplay of carrier concentration, doped cobalt spins, and

temperature dependence is needed to explain the peculiar MR in these samples.

7.3.6 Effects of Annealing

In an effort to better understand the nature of these films and the origin of

ferromagnetism, the 30%Co samples were annealed in both oxidizing and reducing

atmospheres. The films were deposited in vacuum and it is assumed that most of the

electrons are associated with oxygen vacancies in the lattice. Annealing the films in

oxidizing conditions should fill the vacancies and reduce the electron concentration. A

film was annealed at 500'C for 1 hour in 1 atm of oxygen. Table 7-2 shows the change in

electrical properties after annealing. The electron concentration slightly decreased while

the resistivity slightly increased. Most notably, the AHE in the oxygen annealed film

was significantly smaller than that seen in the as-grown film. The Hall resistivity and

magnetoresistance are shown in Figure 7-18. Since the change in electron concentration

is small, one may require an alternative explanation to the diminished AHE coefficient.

One possibility is that some of the cobalt metal dissolves into the ZnO lattice. This seems

unlikely given that the high cobalt concentration is already metastable. It is also possible









that the cobalt metal reacts with oxygen to form cobalt oxide. However, XRD after

annealing shows that the peak tentatively assigned to cobalt metal is still present in the

film and that no cobalt oxides have formed (Figure 7-18(c)). Annealing may cause some

change in microstructure, such as a change in size of the precipitates which can alter the

transport behavior.

It is also possible that the oxygen anneal at 500'C is sufficient to drive out any

hydrogen that resides in the lattice. Recent theoretical calculations suggest that hydrogen

might mediate ferromagnetic interactions between cobalt atoms in a ZnO matrix [150].

To further explore this possibility, we have also annealed 30% Co-doped ZnO films in

forming gas (4%H2/Ar) at 500'C for 1 hour. Unfortunately, this anneal in hydrogen led

to a decomposition of the ZnO:Co film, suggesting that the Co is indeed substitutional

and metastable in the ZnO matrix. The SEM micrograph in Figure 7-19 shows the film

decomposition after annealing in hydrogen. An XRD comparison in Figure 7-20 shows a

large (almost 10 fold) increase in the intensity of the (111) Co metal peak as compared

to the as grown film, indicating an increase in Co metal clustering. High-resolution

four-circle XRD indicates the cobalt phase is aligned with the ZnO lattice. Cobalt metal

can exist in either the hexagonal or fcc structure. The fcc phase is stable above 425'C but

is often observed as a metastable phase at room temperature. Since the hexagonal and fcc

structures differ only in their stacking sequence (hexagonal ABAB and fcc ABCABC) the

fcc (111) planes have the same d-spacing as the (0001) hexagonal planes. Since standard

0-20 XRD measures planes parallel to the sample surface, the hexagonal (0001) and fcc

(111) orientations are indistinguishable [151]. However, since the periodicities of the planes

are different, off-axis peaks (planes not parallel to the surface) can be used to identify

the stacking sequence. To distinguish between the two phases, an off-axis diffraction scan

through the Co (1 0 L) reciprocal lattice points was performed as shown in figure 7-21(a).

The (1 0 L) scan should show peaks at integral L positions for hexagonal stacking and at

particular integer/3 positions for cubic stacking [56]. The L-scan shows that the packing is









mostly cubic ABC stacking. However, a scan through the hexagonal cobalt (101) position,

which cuts through the rod at L=1 in the L-scan, shows there is a small intensity at (101)

as shown in the inset of the figure. This suggest some stacking faults are present or small

regions of hexagonal cobalt. Additionally, a phi-scan through the cubic Co (200) at x=35'

shows in-plane alignment of the cobalt grains as shown in figure 7-21(b). Clustering should

lead to an increase in the observed ferromagnetism since a larger volume of cobalt metal

will have a larger magnetization, which is verified in Figure 7-22.

While the presence of Co metal precipitates in the films provides a possible

explanation to the magnetic behavior, an examination of the magnetic behavior of the

Co-doped ZnO films grown at different conditions provides circumstantial evidence that

the cobalt precipitates are not the origin of the magnetic behavior. First, varying the

growth conditions of the films has a large effect on the observed magnetization. Films

grown at higher temperatures (500'C and 600'C) in vacuum show very little, if any,

magnetization. Given that these large concentrations of cobalt are highly metastable,

one would expect a stronger tendency to form more segregated cobalt metal at the

higher temperature, but thermodynamically the antiferromagnetic CoO phase is stable.

SQUID measurements show no evidence for ferromagnetism for ZnO:Co films grown at

the elevated temperatures. XRD scans for these films were given in figure 7-5. SQUID

characterization for these films is shown in figure 7-23. The coexistence of a cobalt metal

XRD peak and the absence of ferromagnetism in the film grown at 500'C in vacuum is an

interesting result. One would expect, from the sensitivity of the SQUID magnetometer,

that the film would show magnetization in the presence of cobalt metal. The presence of

CoO could be responsible for reducing the magnetization by removing cobalt from the

lattice and precipitates. However the reduction in magnetization is over 100 times smaller

as compared to the film grown at 400'C and may suggest that the small particles of cobalt

metal do not make a large contribution to the magnetization.




































5 10 15 20 25 30

cobalt conc. in targets (mol%)


EDS results for a select number of films grown under different conditions. The
data shows the cobalt concentration in the films as compared to the cobalt
concentration in the target.


Table 7-1: Possible cobalt-induced secondary phases.


Structure
Cubic (111)
Hex (0002)
Cubic (200)
Spinel (400)
Spinel (400)


CoAl204 Spinel (400)


20 (deg) Coupling
44.216 Ferromagnetic
44.762
42.401 Antiferromagnetic
44.808 Antiferromagnetic
44.738 Antiferromagnetic (n-type)
Ferromagnetic (p-type)
44.692 Antiferromagnetic


T*(K)
Tc = 1373

Tn = 291
Tn =30
N/A Ref[152]

Tn < 40


I I I *
- 4000C, vacuum
- 400C, 0.02mTorr
-A 5000C, vacuum
-- 6000C, vacuum


Figure 7-1.


Phase
Co

CoO
Co304
ZnCo204












Sapphire
ZnO (006)
(002)


U)


d
4-,



->10
U)
C
c
C


Sapphire


0 10 20 30 40 50 60 70 80 90 100


29 (degrees)



Figure 7-2. XRD scans for a series of films grown in vacuum at 400'C. The films are
predominately c-axis oriented ZnO. At high cobalt concentrations, a cobalt
induced secondary phase appears near 20=44.4. TEM and High-resolution
XRD suggest this phase is a mixture of cubic and hexagonal cobalt with
stacking faults.
















































Figure 7 4. Convergent bTeam TEMII diffraction patterns of ZnO film doped with 30~%Co
grown at 400 C in vacuum. (a) ZnO (hexagonal) film at [120] zone axis,
(b) ZnO film nanmo particles, (c) ZnO film a nano particles sapphire
substrate, (d) NTano particle at [120] zone axis with doo2 =0.20nm and
dio10 .13nm, which is consistent with metallic cobalt.









S11




















ZnO Sapphire
(002) (006)


35 40 45 50 55 60


26 (degrees)


ZnO


35 40 45 50 55 60

2e (degrees)


Figure 7-5. XRD scans for ZnO films doped with 30%Co. (a) Films deposited at 400C
with different pressure conditions. (b) Films deposited at 500'C and 600'C in
vacuum.


108

10'
107
- 16
E 106

105

S104

10s
c













































400 600 800 1000


1200


Temperature (oC)


Figure 7-6. Thermodynamic predominance diagram for cobalt oxides.






















Undoped 10%Co


1 0 2%Co
0 8- 30%Co
E 15%Co
0 8 216

I--
x T 200
-, 187

E06
0 6 15 16 17 18 19 2 0 21 22 23 24
C Photon Energy (eV)
0o

E 04-
Undoped
-0 2% Co
02- 10% Co
--- 15% Co
-*- 30% Co
00

2 3 4 5 6

Photon Energy (eV)



Figure 7-7. UV-Vis transmission of Co-doped ZnO films deposited in vacuum at 400'C.
The inset shows a close up view of the absorption levels which correspond to
the A2 2 E(G) (1.9 eV), A2 -4 Ti(P) (2.0 eV), and A2 12 A1(G) (2.2
eV).

















S-n- Undoped
50 2% Co
10% Co
---15% Co
40 30% Co
30% Co


low energy
"1 ,1 I ', ,, :


0---
26 28 30 32 34 36
Photon Energy


38 40 42 44
(eV)


Optical band-gaps of Co-doped ZnO films. (a) (achv)2 plots for Co-doped
ZnO films deposited in vacuum at 400C. Straight line fits through the linear
regions of the plot are extrapolated to hv = 0 to find the band-gap. (b) Plot
of the band-gap values as a function of nominal cobalt concentration. Filled
circles are the band-gap and hollow triangles are the onset of absorption at low
energy.


38-


36-


3 34-
,.
m
32-
C
30-


28-


o High energy fits 0-'
A Low energy fits



-00
G .- (b)






A A
A

0 5 10 15 20 25 30 35
Cobalt Concentration (at.%)


Figure 7-8.
























Room Temperature PL

024

0 22 Band-edge
emission


Blaze


016 3 25 eV

1 ,Undoped
012 -

oio 331 ev
2%Co
008
37 36 35 34 33 32 31 30 29 28
Photon Energy (eV)


Undoped



2%Co


5%Co

S15%Co

....-----. 30%Co


I I


800


900
900


Wavelength (nm)


Figure 7-9. PL results for Co-doped ZnO films deposited in vacuum at 400CC. The dip
in the PL spectra at -517 nm is an artifact from the diffraction grading blaze
angle. The PL intensity decreases with higher cobalt doping. The inset shows
higher resolution scans around the band-edge peak for the undoped and 2%Co
samples. The values correspond well with the absorption data.

































0 00


-0 06


-2000 -1000


02


01-


00-


-0 1


-0 2-1


-2000 -1000


0 1000

H (Oe)


, 30% Co
- 5% Co
-- 2% Co


0 1000 2000


-1000 0 1000
H (Oe)


-1000 0 1000


Figure 7-10.


SQUID magnetization curves for Co-doped ZnO films deposited at 400 C in
vacuum. The films do not show any secondary phases by XRD measurement.


2000








































-2000 -1000 0 1000 2000
H (Oe)


300K





WitsBB


ipitates
lo precip
rates


-1000 0 1000 2000
H (Oe)


-2000


a E


Pf ff
P ^ --^ --


--a~B44-0>t

/1


pi"-+


3-30% Co (vac) Precipitates
-- 30% Co (vac) No Precipitates
30% Co (0 02mTorr) No Precip
15% Co (vac) Precipitates


-2000 -1000 0 1000 2000
H (Oe)


Figure 7-11. SQUID magnetization curves for Co-doped ZnO films deposited at 400'C.
The growth pressure and whether the film contains secondary phase
precipitation is indicated in the legend.


-2000 -1000


0 1000 2000
H (Oe)





















400K



... 1ii


10k 0 10k 20k
H (Oe)


5t SG moothlng


H(Oe)

-40k -20k 0
H (Oe)


Ik 4I
20k 40k


300K


No Precipitation


80k 0k k 20k 0 20k k 0k 80k
H (Oe)


-60k -40k -20k


0
H (Oe)


20k 40k 60k 80k


Figure 7-12. Anomalous Hall effect in 30%Co-doped ZnO. Films were grown at 400C
in vacuum. (a) film with cobalt precipitation. The upper inset shows the
Hall curves at different temperatures. The lower inset shows the derivative
of the Hall curve. (b) film without cobalt precipitation. The inset shows the
derivative of the Hall curve.


0 0010-


0 0005 -


0 0000 -


0003-


-0 002-


-0003--
-80k





































300K o o01
0035504 n=5 4010'" cm
-0 0 00
0 035500-
0 0 -001
0 035496 0
-0 02
0
0035492 0- 0 003


0035488- 0 004
80k 60k 40k -20k 0 20k 40k 60k 80k
00414-


00413 n=437x10 cm 02

0 o
E 00412 04

00411
08 01
00410 0



0k 60k 40k -20k 0 20k 40k 60k 80k
00490 -

00485 10K 0

00480 n=375x108 cm3

00475- 0 1 ~~

00470 2


0 0
00460 0
0 0
0 0 -6
00455 -

0k -60k -40k -20k 20k 40k 860k 80k


Applied Field (Oe)




Figure 7-13. The magnetoresistance of an undoped ZnO film. The film was deposited at

600'C in vacuum. The measurement temperature and carrier concentration

are indicated in each figure.






























n

o


oo

300K
n835x10' cm


100K
n 25xl0' cm


n~n


06
00



300K "0
300K
n=475x109 cm











100K












10K
n 420x10 cm



Applied Field (Oe)


n>n


Snm22 c2210m
300K
n=3 25x1 0 cm -















10K
n=32x100 cm


Figure 7-14. The magnetoresistance of 5%Co-doped ZnO films. The film with n
grown at 600'C in 0.02mTorr oxygen. The film with n-nc was deposited at

400'C in vacuum. The film with n>nc was deposited at 400'c in vacuum and

is codoped with 2%A1 to impart a high electron concentration.



















129


o
o
0 0



.. 10K
n=724x1001cm




























n

... 300K
n=1 82x10'" cm
D566 IO 0
o c


100K
n=1 88xs10'cm'


10K
n=1 33x10'" m


n > n


...... 300K
n =1 94x10m cm'


O o


...,,,, 100K
n1 8E910 cm3


10K
n lj n=1 86xlO cmm


Applied Field (Oe)





Figure 7-15. Magnetoresistance of 30%Co-doped ZnO films. Both films were deposited

at 400'C in vacuum. Neither films show signs of a secondary phase by XRD

measurement. The film with n>nc is codoped with 2%A1 to impart a high

electron concentration.















100K


7
75K










50K


I I I I I I I I I
-80k -60k -40k -20k 0 20k 40k 60k 80k


Applied Field (Oe)


Figure 7-16. Magnetoresistance of 30%Co-doped ZnO film at temperatures between 10K
to 100K. The development of the negative MR kink is established at low
temperature followed by a transition to negative MR across the entire field
range with increased temperature.


,,












5%Co


0 024


30%Co


n

50 100 150 200 250 300 350

E
0 n -n

o


0 50 100 150 200 250 300 35


0
0
2
0
o 0 n > n


S s50 100 150


0
a


200 250 300 350


nn
c


Ooo0ooo


0 0 0 0 0 0 0 0 0 0


0 50 100 150 200 250 300
o



0 5 0 100 150
0 50 100 150


n>n
c
200 250 300


Temperature (K)


Temperature (K)


Figure 7-17. Temperature dependent resistivity measurements for 5% and 30% Co-doped
ZnO films.



Table 7-2. Transport data for a 30%Co-doped ZnO film with cobalt precipitation. Data is
for before and after annealing in 02 at 500'C.


(cm3 C 1)


As deposited
02 annealed


Carriers Mobility
(cm 3) (cm2 v 1 s1)


-1.685 3.70x1018
-2.998 2.08x1018


Zero-field p.,
(Q-cm)
0.2367
0.3344


0 01SL-
0 008


0007


0 006


0 00136
























As Deposited


00K











20k 10k 0 10k 20k 30
H (Oe)


02 Annealed


30k 20k 10k 0 10k 20k
H (Oe)


4o
0
0 0
2 300K
0
o 0
o 0

o 0
0
0 0


H (Oe)


26 (deg)


26 (deg)


Figure 7-18. Hall resistivity and magnetoresistance for a 30%Co-doped ZnO film with
cobalt precipitation. Data is shown before and after annealing in 02 at 500.
Insets show the derivates of the Hall resistivity to show AHE.


3


300K


(a) 00008

0 0004


0 0008

30k



(b) 0236

0 2364


o 0 2360


a 0 2356


0 0
0 300K
0 0


o o
o o


30k



000

0 05

010

0 15

0 20


Co (002)


/















































Figure 7-19 SEM micrographs at different magnifications of the surface of a Co-doped
ZnO film that has been annealed in H2/Ar at 500C for 60min Magnification
and scale bars are given in each micrograph The arrow in (a) shows the
region of film degradation This area was brown in color compared to the
green color of the film









134



















C As
H2


Co metal
(002)


deposited
Annealed


r- f


44 46 48 50


29 (deg)


Figure 7-20. 0-20 XRD for a 30%Co-doped ZnO film before and after annealing in forming
gas at 500'C.


42


1



















(10L)_scan#23

(a (10 L) Scan













0 05 1 15 2 25 3 35
L (Hex coords)


phiCoscan#28


-90 -60 -30 0 30 60 90
| (degs)


Figure 7-21. High-resolution XRD scans for 30%Co-doped ZnO film that has been
annealed in forming gas at 500'C. (a) L-scan along the Co (10L). The inset
shows an off-normal 0-20 scan through the hexagonal Co (101). (b) phi-scan
through the Co (200) at x=35.


















































20
2500 2000 1500 1000 500 0 500 1000 1500 2000 2500


2500 2000 1500 1000 500 0 500 1000 1500 2000 2500


H (Oe) H (Oe)




Figure 7-22. SQUID magnetometry for a 30%Co-doped ZnO film before and after
annealing in forming gas at 500'C.

























137





















10




05-


IN


E -



-0 5-
400C, 0.02mTorr
500C, vac
S- 600C, vac
-1 0-
-1500 -1000 -500 0 500 1000 1500
H (Oe)


Figure 7-23. SQUID magnetometry for 30%Co-doped ZnO films deposited under
differerent conditions.









CHAPTER 8
CONCLUSION

The advancement of spintronics as a practical technology depends upon the

development and understanding of semiconductors that can support spin-polarized carrier

operation above room temperature. The research presented in this dissertation explored

the possibility of using wide band-gap oxide semiconductors as spintronic materials. Both

transition-metal doped Cu20 and ZnO were investigated. Mn and Co were used as the

transition-metal dopants to provide localized spins in the host semiconductor lattice. Thin

films of these materials were deposited using pulsed laser deposition. Various material

properties including the structure, magnetic, and electronic transport properties were then

characterized to gain a better understanding of the materials.

8.1 Mn-doped Cu20

The magnetic properties of Cu20 films doped with 1 at.% Mn were investigated

under different growth conditions. This research was stimulated by Dietl's theoretical

prediction that carrier-mediated ferromagnetism is favored in Mn-doped, highly p-type,

wide band-gap semiconductors, such as ZnMnO and GaMnN. Cu20 is a naturally p-type

semiconductor with a wide (direct) band-gap and therefore holds interest in exploring spin

behavior in oxides.

The Mn solubility in Cu20 was found to be small and the precipitation of Mn-oxides

was favored at high growth temperatures. However, metastable incorporation of Mn in

the Cu20 could be achieved at low temperature (400'C). These phase pure samples were

found to be non-ferromagnetic. Ferromagnetism with a Tc-50K was observed in the films

deposited at higher temperatures, but appears to be associated with a Mn304 secondary

phase which has a Tc near 50K.

Spintronic concepts based on ferromagnetic semiconductors require the distribution

of charge carriers in the semiconductor to be spin-polarized. Magnetism derived from

localized magnetic precipitates is of little utility for semiconductor-based spintronics if









the carriers are not polarized. Therefore, based on the collected data, it is concluded that

Mn-doped Cu20 has limited use as a ferromagnetic semiconductor under the studied

growth conditions. The low solubility of Mn in the Cu20 lattice limits experimentation

using larger concentrations of Mn. Perhaps larger hole concentrations or alternative

fabrication techniques could be used to induce ferromagnetism.

8.2 Mn-doped ZnO

Several of the proposed models for ferromagnetism in ZnO DMS emphasize the

importance of holes mediating the exchange interaction between doped Mn spins. This

research investigated the trends in the magnetization as a function of carrier concentration

in order to elucidate the role of charge carriers on the ferromagnetism in these materials.

The carrier concentration was varied using Sn and P as electronic codopants. Codoping

allowed independent control over the magnetic and electronic properties by doping each

separately. This provided a platform to study the effects of carrier concentration on the

observed magnetic properties. Sn acts as an n-type dopant providing extra electrons to

the ZnO. P acts a p-type dopant that supplies holes to compensate the native electron

concentration in ZnO. Thin films doped with 3 at.%Mn, and either Sn or P as the

codopant, were deposited in an oxygen atmosphere of 20mTorr in a temperature range of

400-6000C.

The electron concentration in the ZnMnO:Sn films was controlled by varying the

Sn content. Initially, the magnetization increased with minimal Sn doping resulting in a

maximum magnetization of 0.5 pB/Mn atom at 300K. However, with increased Sn doping

there was an inverse correlation between the Sn content and the saturation magnetization.

As the electron density increased with Sn doping, the magnetization decreased.

The trend in magnetization showed a similar carrier dependence using P as a

codopant. Under conditions where the acceptor dopants were activated, the magnetization

was enhanced. The resistivity of the as-deposited film was -4,000 -cm at 300K with an









electron density of -3x1015 cm and became too insulating to measure after annealing.

The saturation magnetization after annealing was around 0.15 PB/Mn atom.

The correlation of saturation magnetization with electron density is interesting

and provides some insight into the mechanism for ferromagnetism in ZnMnO. Overlap

of the Mn d-states with the valence band suggests holes are necessary in order to

induce ferromagnetic ordering. Magnetism in low free-electron density material is

consistent with the bound magnetic polaron model in which bound acceptors mediate

the ferromagnetic ordering. The holes may be delocalized, but with low mobility, thus

yielding low conductivity. In this case, the carrier mediated mechanism may suffice

without the need to invoke bound polarons as inherent to the ferromagnetic ordering. In

either case, the addition of electrons to the system will move the Fermi energy level up in

the band-gap, resulting in a decrease in hole density and a reduction in magnetization.

While the trend of magnetization versus carrier density was similar in the Sn and

P codoped films, the measured magnetization was slightly different. The magnetization

was anticipated to be larger for the annealed P-doped films, as compared to the Sn-doped

films, since the Fermi level should reside closer to the valence band. However, this was

not observed; films with minimal Sn doping exhibited a higher magnetization per Mn

atom. The reason for this is unclear, but could be a result of other defects formed after

the annealing process. The XPS data also suggested a change in valence of some of

the Mn atoms after annealing. This may have caused a reduction in magnetic moment

that competed with the suppressed electron density. Nevertheless, the overall trend of

magnetization versus carrier density was consistent between the experiments.

The results of this work are consistent with the observations of Kittilstved and

coworkers [32]. The authors concluded that ferromagnetism in ZnO is closely tied

to the charge transfer electronic structure of the transition-metal dopant. For Mn,

ferromagnetism is induced when the holes from the acceptor states hybridize with the









Mn ions. Ferromagnetism was observed when the ZnO was locally doped p-type, but no

ferromagnetism was observed when doped n-type.

8.3 Co-doped ZnO

The effect of cobalt-doping on the magnetic and magneto-transport behavior in ZnO

was also investigated. Ferromagnetism was found in films deposited at low temperature

(400'C) in vacuum, while films deposited in oxygen or at higher temperatures were

non-magnetic. Films deposited under vacuum had rather high electron concentrations

and are presumably doped with oxygen vacancies. Segregation of cobalt metal occurred in

films doped with 15% or greater Co concentrations when deposited in low base pressure

(<10 5 Torr) vacuum conditions. The precipitates were small (-5nm) and aligned

with the ZnO lattice. The segregation could be suppressed by depositing in higher base

pressures (>10-5 Torr), but the process was not fully reproducible. These films also show

ferromagnetism albeit with lower magnetization than the films with metallic cobalt. This

suggests that the cobalt metal particles could be responsible for some of the observed

magnetic properties in the films with precipitation. However, the role and the extent of

magnetization from the particles is uncertain. The distribution of particles that behave

ferromagnetically, or behave superparamagnetically because of their small size, and to

what extent the film properties have on the ferromagnetism in the precipitated films, is

unclear.

The cobalt-doped films also exhibited peculiar magnetoresistance that had a strong

dependence on the carrier concentration. The MR was studied in films with 5%Co and

30%Co doping, along with an undoped ZnO film as a reference. At low temperature, the

cobalt-doped films undergo a progression from positive to negative MR as the electron

concentration is increased and the films cross over the metal-to-insulator transition (MIT).

At large carrier concentrations, well into the metallic regime, the MR behavior for the

undoped, 5%Co, and 30%Co films is quite similar. This suggests that the electronic

transport is not greatly affected by the cobalt concentration above the MIT. One possible









explanation is that since the Fermi energy has moved fully into the conduction band,

the transport is dominated by the character of the conduction band. In an applied

magnetic field, the cobalt d-states and the donor impurity states are split into spin-up

and spin-down sub-bands. If the Fermi energy lies in the conduction band, there is likely

no overlap of the Fermi energy with the spin-polarized d-bands of the cobalt, and the

conduction is characteristic of undoped ZnO.

A possible future experiment that may prove interesting in understanding the MR

behavior could be executed through the fabrication of field-gated Hall bars. Voltage

applied to a top gate electrode could be used to vary the electron concentration near the

vicinity of the gate. Magnetoresistance and Hall measurements could then be performed

at varying gate voltages to examine their behavior with different electron concentrations.

Ideally, one could smoothly study the MR behavior across the MIT and closely examine

the changes near the transition point.









APPENDIX A
HALL EFFECT SYSTEM AND EQUIPMENT

A.1 Introduction

I've added this information for readers interested in detailed information regarding

the resistivity and Hall effect methods used in this work. I first wrote this section as a

reference for my research group to introduce them to the equipment. However, after a few

changes, I decided that this section also made a valuable addition to the dissertation. I've

included operational information about our current equipment, simple Hall effect theory,

and explanations for the advantages and disadvantages of certain measurement techniques.

I hope this section will prove beneficial to those who come across it.

A.2 Hall Effect Equipment

The Hall effect was discovered by Edwin Hall in 1879 and has become a powerful tool

for the characterization of electronic materials. The charge carrier type, concentration, and

mobility can all be determined from accurate Hall measurements.

Most of the resistivity and Hall effect measurements in this thesis were performed

using a Keithley high-impedance Hall effect system. Control over the temperature and

magnetic field was accomplished by placing samples inside a Quantum Design (QD)

Physical Property Measurement System (PPMS). The PPMS is a versatile tool for

probing various material properties. Equipment options are available for measuring a

sample's magnetic, electric, and thermal properties in the same piece of equipment.

The PPMS's open architecture also allows room for user customization and other tools

may be integrated into the system. For instance, customizable probe heads allow the

introduction of extra signal cables or optical feedthrus into the sample arena. The

PPMS acts as a platform, or a "measurement environment", to control the ambient

temperature and field around the sample. The temperature of our system can be adjusted

from 1.4K-400K, and there is a longitudinal superconducting magnet that can be swept

to a maximum of 7 Tesla. The PPMS is a helium filled system. The helium is used









for cooling the sample space to low temperature and also keeps the magnet below its

superconducting critical temperature. The helium bath is surrounded by a nitrogen filled

jacket that reduces the helium boil-off. At low temperatures, the sample chamber must

be pumped to prevent freezing of atmospheric gases in the system. This is controlled

through an integrated gas handling system. There are packaged options for both DC

resistance and AC transport measurements available from QD. However, some of the

samples produced in our lab are highly resistive, >100 MQ, and our current PPMS

resistivity options can not support samples of such large resistance. Measuring transport

properties on these materials required the use of separate electronics. A high-impedance

Hall effect configuration using Keithley components was used, including a 236 Source

Measurement Unit (SMU), two 6514 Electrometers, and a 7152 4x5 matrix switch card

housed in a 7001 switching mainframe. A diagram of the equipment is give in Figure

A-1. For consistency and convenience, the external electronics were used for most of the

measurements in this work, even for samples that could be easily measured by the internal

PPMS electronics. Integration and control over the PPMS and Keithley electronics was

accomplished using the LabView programming language. LabView was used to program

the automated resistivity and Hall routines, which included programs for Van der Pauw

and Hall bridge samples, magnetic field and temperature sweeps, and time dependent

measurements. If necessary, LabView also provides flexibility for creating new routines for

future experiments.

Samples are mounted to pucks that can accommodate sample sizes up to 1cm x lcm

square. There are 12 available signal wires fed to the sample space that may be used for

measurements. For example, the standard QD resistivity puck allows up to three samples

configured for 4pt. measurements to be measured simultaneously through the 12 signal

leads (4 leads are used for each of the three samples). Access to these wires from outside

the system is granted through a Lemo connector.









The Keithley high-impedance Hall effect system is shown in Figure A-1. This section

will provide a brief description of the equipment used and the reasons behind their

selection. The basic premise of an electrical measurement is to provide a known stimulus

to a sample and to measure the sample's response to that stimulus. For instance, a basic

resistivity measurement may be made by applying a known current to a sample, measuring

the subsequent potential difference, and calculating the sample's resistance by using Ohm's

law. As with most scientific measurements, accuracy is of prime concern. Various internal

and external factors can effect the accuracy of electrical measurements. These include

error sources from the experimental set-up, including leakage currents, ground loops,

equipment offset voltages and currents, and environmental noise. The sample itself will

also provide sources of error, such as thermal Johnson noise, non-ohmic contacts, and

photovoltaic effects. Some of these errors can be significantly reduced by the proper choice

of equipment and measurement procedure.

The Keithley 236 SMU contains both source and measurement capabilities in one

unit. The SMU is used to drive current into the sample and measure the voltage difference

created. The 236 provides a current range of 100fA to 100mA providing flexibility for

samples of various resistance. Both the source and sense leads pass through the 7152

switch card. The 7152 provides automated switching of the signal cables between the 4

sample contacts through a grid of interconnected relays. This can be seen in the switch

card diagram included in Figure A-1. Each crossing point of the grid represents a relay

that can be opened or closed to redirect the cabling to each sample contact point. These

relays also switch the guard conductors. Switch cards eliminate the need to physically

reconfigure the sample contacts for each leg of the Van der Pauw routine. The high and

low signal voltages generated in the sample are sent to the Keithley 6514 electrometers.

These electrometers have a huge input impedance (>200 TQ) which reduces input loading

errors on high resistance samples. In our set-up, the two electrometers act as unity gain

buffers between the sample and the SMU voltmeter. The unity gain buffer is an amplifier









circuit that duplicates the input voltage at the output terminal of the amplifier. The

voltage gain is unity ( =1). Since the buffer's input impedance is large it can replicate

the voltage without drawing any current from the source. Ideally, voltage measurements

draw zero current [153]. In fact, any current at the voltmeter input represents an error,

called the input offset current. This is the advantage of placing a buffer amplifier (the

electrometers) in the circuit. The output from the electrometers is then read by the SMU's

voltmeter for measurement.

A.3 Sample Geometry and Measurement Technique

The two common sample geometries for Hall measurements are Hall bridge and

Van der Pauw samples. While the exact dimensions of Hall bridge samples are needed

for accurate data, the Van der Pauw technique allows samples of arbitrary shape

and size. Both measurement techniques are described in the ASTM standard F76-86

"Standard Test Methods for Measuring Resistivity and Hall Coefficient and Determining

Hall Mobility in Single-Crystal Semiconductors". Included in the ASTM standard are

specifications for sample geometries and the mathematical equations used to calculate

the material's electrical properties. Both types of sample geometries were used in this

work. The LabView programs used to drive the measurement routines were based on the

information provided in the ASTM standard. This section will give a brief explanation of

the advantages and disadvantages of each measurement.

Resistivity measurements can be performed using a 2pt. method, where only 2 leads

are connected to the sample, by shorting a voltmeter across the same leads used to drive

current through the sample. A diagram for the 2pt. configuration is provided in Figure

A-2. The drawback of 2pt. measurements is that, in addition to the sample resistance, the

voltage drop across the contact and lead resistances are also measured by the voltmeter.

This can cause serious loading errors if the contact and cable resistance is larger than

the sample resistance. A more accurate approach requires 4 contacts to the sample.

Current is driven into the sample through two leads and the voltage is measured across









the other two. Ideally, no current flows across the voltmeter, but this depends on the input

resistance of the meter. Quality voltmeters have a large input resistance (10MQ-10GQ) to

reduce the input offset current. Hall bridge samples are measured using the 4pt. contact

method. Current is sourced between the two longitudinal contacts and the voltage is

measured between two contacts situated parallel to the current direction. The hall voltage

is measured across two contacts that are placed perpendicular to the current direction.

The Van der Pauw technique was first proposed by L.J. Van der Pauw in 1958. The

technique is widely used since it provides a convenient method for probing the transport

properties of materials. The Van der Pauw technique's primary advantage is that contacts

may be placed arbitrarily around the sample perimeter. This enables samples of arbitrary

shape and size to be measured directly. The need to fabricate complex sample geometries

or detailed knowledge of the sample dimensions is eliminated. Only knowledge of the

thickness is needed to calculate the volume resistivity and carrier density from their

respective 2-dimensional sheet values. This can be highly beneficial in certain cases, such

as when the sample material is difficult to process or when the facilities and/or the time

for processing samples is not available.

There are several disadvantages to the Van der Pauw technique. There are errors

associated with contact size effects and shorting of the Hall voltage [154]. Ideally, contacts

should be kept as small as possible. As mentioned above, the Van der Pauw technique

allows arbitrary placement of contacts to determine the resistivity of the sample. On

the contrary, this is not true for the Hall voltage. Misalignment of the voltage leads

will cause additional error. Deviations from perpendicular alignment with the applied

current can cause mixing of the perpendicular and parallel components of the created

voltage. This distorts the true Hall voltage being measured. Note that this is also true for

Hall bridge samples; however, the fixed geometry of the bridge limits the misalignment.

Current-reversal averaging techniques (described later) can help alleviate misalignment

error. Another drawback of the technique is measurement time. Van der Pauw requires a









total of 8 voltage measurements around the sample-each side of the sample is measured

under positive and negative current polarities-whereas a hall bridge patterned sample

only requires 2 measurements. This can be especially problematic when long delay times

are needed for the circuit to settle, such as when measuring high resistance samples.

Settling time is a function of the RC time constant. It represents the time needed to

saturate any shunt capacitance in the circuit to obtain an accurate measurement of the

circuit voltage. This idea is represented in figure A-3. The Van der Pauw technique

also requires a switching system if leads are hard wired to the sample. This system

can be manual or automatic, and is used to reconfigure the current and voltage leads

around the sample without having to remove and replace each sample contact during the

measurement.

Small signal currents are needed to make measurements on high resistance sources.

During Hall and resistivity measurements, ohmic heating in the sample should be kept

to a minimum to retain measurement accuracy and reduce thermal noise. Since power

is equal to resistance multiplied by the square of the current (P=I2R), low currents

are needed to keep the power dissipation low. Typically, keeping the power in the

sample below 1-5 mW is suggested for general transport measurements. Leakage

currents-currents that 'leak' into or out of the circuit through unwanted resistive

pathways-can degrade the accuracy of low current measurements [153]. This can be

a common error in low-level measurements using standard cables, where portions of

the signal current can leak through the cable's insulation resistance. Use of a guarding

technique with triaxial cables can reduce leakage currents. Triaxial cables contain an extra

conductor that surrounds the signal cable (separated by an insulator) and lies underneath

the ground sheath of the cable. A potential is driven along the guard that matches the

potential on the signal cable. This creates a space of zero potential between the guard and

signal and reduces the driving force for current leakage between them. The leakage current

combined with the input offset current represents the total error current of the circuit









[153]. The open-loop gain of the guard also reduces the effects of cable capacitance by

reducing the shunt capacitance charging time [153]. The Keithley electronics support the

use of guarding and the equipment is connected using triaxial cables. However, the PPMS

does not use triaxial cabling which poses a limitation of using guarded measurements with

the current set-up (see the following sections for more details).

A.4 Hall Effect Method

The physics of the Hall effect can be found in almost any solid-state physics or

electronic materials science book, as well as plenty of sources on the world wide web. It's

not my intention of rehashing this information since it is easily found, but only to provide

the basic ideas and equations for the measurement. The driving force behind the Hall

effect is the Lorentz force acting between moving charge carriers and a magnetic-field

applied perpendicular to the carrier's velocity. The Lorentz force is given by:



F = q(vxB) (A-l)

where q is the carrier charge (1.602x10 19 for electrons and holes), v is the carrier

velocity, and B is the magnetic induction.

A current is applied to one end of the sample and carriers are driven through the

length of the sample. The magnetic-field is applied perpendicular to the carrier velocity

and pushes the carriers towards the edge of the sample as dictated by the Lorentz force.

This results in a slight charge imbalance between the two sample edges. To maintain the

flow of current, the charge imbalance creates a potential drop across the sample to balance

the Lorentz force. This is the Hall voltage, VH, and is given by:



V. = B (A-2)
nxqxt
where n is the carrier concentration, t is the sample thickness, and I is the applied

current. If the sample resistance is known, the carrier mobility, p, can be calculated:











n (A3)

where RHl=nq is defined as the Hall coefficient.

However, if the sample under consideration is composed of two carriers, such as

a semiconductor with a sufficient number of electrons and holes in the material, a two

carrier model for the Hall coefficient is more appropriate [154]:



RH = (A-4)
e (pp, + n~,)
Taking Hall voltages at three magnetic-field points (+Hi, H=0, and -Hi, where H,

is some value of magnetic field) is usually sufficient to determine the Hall coefficient.

The two measurements at H1 are used to average over both field polarities and the

measurement at H=0 can be used calculate the carrier mobility. However, I have found

it extremely beneficial to measure Hall voltages at many points over a full magnetic-field

sweep. Since the Hall coefficient is proportional to the quotient of VH/B, the coefficient

can be calculated from the slope of a line fit through the data. This provides a better

statistical average, averaging over many points rather than just three, and the accuracy

of the data is easily examined based on how well the points follow a linear path. This

proved particularly useful for difficult samples which exhibit weak Hall voltages that

stagger between n- and p-type (positive and negative slope). Taking Hall measurements at

varying-fields is also necessary for studying the anomalous Hall effect (AHE) in magnetic

samples. In fact, sensitive magnetic properties can even be determined in samples that

show a large AHE component.

A.5 Limitations and Tips for Better Measurements

There are several limitations of the current set-up that could use refinement. The

main difficulty we have experienced is measuring Hall voltages on our p-type ZnO films.

The doped holes in these materials are highly compensated by the intrinsic electron

concentration, and both carrier types are present in the material. The Hall coefficient









for this condition was given above. The mobility of holes in these materials is small,

and the condition for p-type material, pp2 > nP is difficult to fulfill. These samples

produce a very weak Hall voltage, even at fields of 7T, and are extremely difficult to

measure accurately. The electrometers have a resolution of 10 pV, which does not provide

enough sensitivity for measuring the weak Hall voltage. Typically measurements are

noisy and teeter between n- and p-type conductivity. Replacing the electrometer with a

more sensitive voltmeter, such as a nanovoltmeter with nanovolt sensitivity, may provide

enhanced measurement accuracy for these difficult samples.

The 236 SMU is a convenient tool because it provides sourcing and measuring

capabilities in one cost-effective unit. However, this also provides a slight limitation when

measuring highly-resistive samples. The SMU is not capable of performing measurements

in 4 terminal mode when the potential difference between the source and sense leads

exceeds 4 volts. This is easily seen during a measurement because every Van der Pauw

voltage measurement around the sample converges to the same value. Whenever this

was encountered, a digital multimeter was borrowed from another piece of equipment

to perform the measurements. Since the other DMM was isolated from the SMU, this

problem could be alleviated.

An additional limitation comes from interfacing the external Keithley electronics

to the sample inside the PPMS. The wires inside the PPMS that connect the external

port on the PPMS to the sample fixture are not triaxial. Therefore a discontinuity

exists between the guarded triaxial cables from the Keithley electronics to the sample

fixture-the guard is not carried all the way to the sample. Currently, the discontinuity in

cable type occurs at a custom-made junction box. The guards terminate at the junction

box, where only the center conductor (the portion carrying the signal) is carried to the

PPMS cabling. Creating a PPMS probe that uses triaxial cables and carries the guard all

the way to the sample should provide further measurement accuracy for samples that are

difficult to measure.









Misalignment in the sample contacts can add spurious values to the measured

voltages due to mixing of the parallel and perpendicular voltage components. A simple

way to help separate the two components is by field averaging the collected data. The

magnetoresistance of a sample will be an even function of the applied field. So the curve

produced should have mirror symmetry reflected across the y-axis on a resistivity vs. field

plot. Therefore, the magnetoresistance can be averaged by RPx = [R (+H) + Ra^(-H)].

On the contrary, the Hall resistivity is an odd function of field, which produces symmetry

through the origin. Odd symmetry produces graphs that are unchanged after a 180 degree

rotation around the origin. The Hall data can be averaged by Ry = 2[RH(+H) RH(-H)]

These averaging techniques can be used for both single field sweeps and hysteretic sweeps,

where the field is swept in one direction and then swept back in the reverse direction.

When averaging the hysteretic sweeps it is advisable to average the initial downward

sweep with the ascending return sweep and vice-versa [155]. While these simple averaging

schemes can aid in separating the two components, the data may still be skewed if one

component is much larger than the other.




















7152 Switch Card


Figure A-1. The resistivity and Hall measurement system. The sample is placed inside
the PPMS dewar to control the ambient temperature and magnetic field. The
electrical data is collected with external Keithley electronics.



























4 point


I---------i
SSample
I AAAA I


Figure A-2.


r--------------
Sample
I ai a


Circuit diagrams for 2-point and 4-point resistivity measurements. In a 2pt.
measurement, the contact and lead resistance are measured by the voltmeter.
In a 4pt. measurement, the current through the voltmeter approaches zero, so
only the sample voltage is measured.


2 point















































00 10 20 30 40 50
T= RC
S


Figure A-3.


Circuit shunt capacitance and settling time. The circuit shunt capacitance
that must be fully charged before an accurate measurement of the circuit
voltage, Vm, is possible. The graph shows the settling time necessary to
reach a certain percentage of the final voltage. The settling time is given as
multiples of the circuit's RC time constant. Adapted from [153].


(VMNs)

Percent of
Final Value (%)









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BIOGRAPHICAL SKETCH

Mathew Ivill was born in Ft. Myers, Florida. He was named after his grandfather,

who also only had one 't' in his name. He grew up and attended school in the city of Cape

Coral, Florida, where he lived with his parents, Richard and Linda, and older brother

David.

Mathew began his undergraduate study at the University of Florida in 1996 (the year

of UF's first national championship football title) and found his way into the Department

of Materials Science and Engineering his sophomore year. In 2000, during his final

year of undergraduate study, he began research in Dr. David Norton's group, growing

epitaxial CeO2 films on InP substrates. He later earned his B.S. in Materials Science and

Engineering in the spring of 2001. He spent the following summer at Sandia National Lab

in Livermore, CA, fabricating and testing samples to study the interface adhesion between

PMMA and silicon (while also enjoying the beautiful national parks around California).

Upon his return as a graduate student to the University of Florida, he began studying

the effects of transition-metal doping in oxides for the field of spintronics. He was also

fortunate enough to take some time off from his studies to live in London, England for 5

months. He has presented at 7 professional conferences, and was awarded the Grand Prize

in Materials Science at the Florida Chapter of the American Vacuum Society, Orlando,

two years in a row for his posters on transition-metal doped ZnO.

In his free time during graduate school, he was a member of the Gator Kendo Club.

He has also enjoyed playing intramural softball with the Braised Cabbages.





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DEVELOPMENTOFTRANSITION-METALDOPEDCu2OANDZnODILUTEMAGNETICSEMICONDUCTORSByMATHEWP.IVILLADISSERTATIONPRESENTEDTOTHEGRADUATESCHOOLOFTHEUNIVERSITYOFFLORIDAINPARTIALFULFILLMENTOFTHEREQUIREMENTSFORTHEDEGREEOFDOCTOROFPHILOSOPHYUNIVERSITYOFFLORIDA2007 1

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Tomyfamilyandfriends,especiallytomyparentsandKathryn,fortheirloveandsupport. 3

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ACKNOWLEDGMENTSGraduateschoolhasproventobequiteawildride,lledwithmomentsofgreataccomplishment,balancedwithmomentsofinsanefrustration.Luckily,I'vehadtheprivilegeofcollaboratingwithamazingpeoplewhohavehelpedmewhileIworkedonthisdissertation.Itwouldbeimpossibletolisteveryonewhohashelped.Nevertheless,IwilldomybesttonamethosewhomademyyearsandallthemanyhoursandlatenightsinthelabattheUniversityofFloridaaneducational,memorable,andenjoyableexperience.Firstofall,Ithankmyadvisor,Dr.DavidNorton,fortakingmeintohisresearchgroup.Iamdeeplygratefulforhisconstantencouragementandsupportthroughoutmygraduatestudies.Hehashelpedmegrowbothacademicallyandpersonally,andhasbeenanexceptionalrolemodel.Withouthishelpthisdissertationwouldnotbepossible.Ialsothanktheothermembersofmycommittee:Dr.StephenPearton,Dr.CammyAbernathy,Dr.ArtHebard,andDr.SusanSinnottfortheirhelpandguidance.IamgratefultoDr.BrentGila,apersonwhopossessesavastamountofpracticalknowledge.Ilearnedsomethingneweverytimewechattedinlab.I'vespentmanyenjoyablehoursworkingwithmyfellowgroupmembers.Iamhappytothankthemall,bothpastandpresent,fortheirsupportandfriendshipinandoutofthelab,includingDr.Seh-JinPark,Dr.Beong-SeongJeong,Dr.Hyung-jinJohnnyBae,Dr.KyunghoonKim,Dr.JenniferSigman,Dr.GeorgeErie,Dr.YaunjieLi,Dr.SeemantRawal,MiteshPatel,VijayramVaradarajan,MichealJones,HyunsikKim,Li-ChiaTien,PatrickSadik,CharleeCallender,Lii-CherngDanielLeu,JoeCianfrone,FernandoLugo,RyanPateandZivinPark.IamespeciallygratefultobothDr.Young-WooHeoandDr.YongwookKwon,bothofwhomtookmeundertheirwingwhenIstartedgraduateschool,andwhosekindnessandpersonalityhavebeenagreatinspiration.Young-WootaughtmemanyofthelabtechniquesIusedthroughoutgradschool,includinghowtodepositlmsusingPLD,andwasalwayskindenoughtolendahandwithmyresearch,evenwhenhewasextremelybusywithhisown.Yongwookalsotaughtmemanythings;heintroduced 4

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metoLabViewprogramming,electricalcharacterization,andalsotaughtmesomeKendoalongtheway!Iwishthemallgreatsuccessandhappiness.IthankthemembersofDr.Abernathy'sresearchgroup,includingDr.RachelFrazier,Dr.JenniferHite,andDr.GeraldThalerfordiscussionspertainingtomagneticsemiconductors.Ialsothankmycollaboratorsandfriendsfromthephysicsdepartment,RiteshDas,Dr.JoshKelly,andDr.RyanRairighforprovidingtheSQUIDmeasurementspresentedinthisdissertationandtheirinsightfuldiscussionsinterpretingthedata.AlsoRajivMisrawhooeredsomeveryusefuladviceregardingelectricaltransportmeasurements.IthankDr.JohnBudaiandDr.MatthewChisholmforcollaborationwithhigh-resolutionXRDandTEMonselectedsamples.IalsothankDr.ValentinCraciunandEricLambersfromtheMajorAnalyticalInstrumentationCenterMAICfortheirhelpwithmaterialcharacterizationusingXRDandXPS.IthankDr.SarahRussellGonzalez,oneofthefriendliestpeopleIknow,forintroducingmetotheLaTeXtypesettingprograminwhichthisdissertationwaswritten.Ithankmyparentsandmybrotherfortheirloveandunwaveringsupport.They'vealwaysbeenaroundwhenIneededthemandtheycontinuetosupportmeinmydecisions.IwouldhavenevermadeitthisfarwithoutthemandIamveryfortunatetohavetheminmylife.Finally,Ithankmyancee,KathrynKennedy,forherendearingloveandsupportandhoursofproofreading.Shewasalwaystheretocheermeupwhenexperimentswentwrong,toinspiremewhenlifebecameoverlyfrustrating,andtocelebratewithmewhenthingsnallywentright. 5

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TABLEOFCONTENTS page ACKNOWLEDGMENTS ................................. 4 LISTOFTABLES ..................................... 9 LISTOFFIGURES .................................... 10 ABSTRACT ........................................ 13 CHAPTER 1INTRODUCTION .................................. 15 1.1ChargeandSpin ................................ 15 1.2WhatisSpintronics? .............................. 16 1.3DiluteMagneticSemiconductors ........................ 17 2REVIEWOFDILUTEMAGNETICSEMICONDUCTORS ........... 19 2.1MagneticSemiconductingMaterials:AShortHistory ............ 19 2.2DMSTheory:ThePhysicalOriginsofFerromagnetisminDMS ...... 22 2.2.1Dietl'sMean-eldTheory ........................ 22 2.2.2First-principlesDesign:DFTCalculations .............. 23 2.2.3FerromagnetisminDisorderedAlloys ................. 24 2.2.4FerromagnetisminaSpin-splitConductionBand ........... 25 2.3ExperimentalProgressinZnODMS ..................... 26 2.3.1Mn-dopedZnO ............................. 27 2.3.2Co-dopedZnO .............................. 28 3THINFILMDEPOSITIONANDEXPERIMENTATION ............ 34 3.1PLDasaToolforThinFilmOxides ..................... 34 3.2PLDSystemUsedforThisWork ....................... 35 3.2.1TheGrowthEnvironment ....................... 35 3.2.2TheLaserSource ............................ 36 3.3TypicalGrowthProcedure ........................... 37 3.3.1SubstratePreparation .......................... 37 3.3.2ThinFilmGrowth ............................ 38 3.4FabricatingPLDAblationTargets ....................... 39 3.5ThinFilmCharacterization .......................... 40 3.5.1X-rayDiraction ............................ 40 3.5.2MagnetoresistanceandHallEectMeasurements ........... 41 3.5.3ElectronDispersiveSpectroscopyEDS ................ 42 3.5.4OpticalAbsorption ........................... 43 3.5.5SQUID .................................. 45 6

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4PROPERTIESOFMn-DOPEDCu2ODMS .................... 49 4.1Introduction ................................... 49 4.2Cu2O:AWide-bandgapP-typeSemiconductor ................ 49 4.3Experimental .................................. 51 4.4ResultsandDiscussion ............................. 52 5PROPERTIESOFZnOCODOPEDWITHMnANDSn ............. 70 5.1Introduction ................................... 70 5.2Experimental .................................. 71 5.3ResultsandDiscussion ............................. 71 6PROPERTIESOFZnOCODOPEDWITHMnANDP ............. 82 6.1Introduction ................................... 82 6.2Experimental .................................. 82 6.3ResultsandDiscussion ............................. 83 7PROPERTIESOFCOBALT-DOPEDZnO .................... 100 7.1Introduction ................................... 100 7.2Experimental .................................. 100 7.3ResultsandDiscussion ............................. 101 7.3.1ChemicalComposition ......................... 101 7.3.2StructureandPhaseAnalysis ..................... 101 7.3.2.1Filmswithprecipitation ................... 101 7.3.2.2Filmswithoutprecipitation ................. 102 7.3.3OpticalProperties ............................ 103 7.3.3.1Opticalabsorption ...................... 103 7.3.3.2Photoluminescence ...................... 105 7.3.4MagneticCharacterization ....................... 106 7.3.5ElectricalTransport ........................... 107 7.3.5.1Halleect ........................... 107 7.3.5.2Magnetoresistance ...................... 110 7.3.6EectsofAnnealing ........................... 113 8CONCLUSION .................................... 139 8.1Mn-dopedCu2O ................................ 139 8.2Mn-dopedZnO ................................. 140 8.3Co-dopedZnO ................................. 142 APPENDIX AHALLEFFECTSYSTEMANDEQUIPMENT .................. 144 A.1Introduction ................................... 144 A.2HallEectEquipment ............................. 144 7

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A.3SampleGeometryandMeasurementTechnique ............... 147 A.4HallEectMethod ............................... 150 A.5LimitationsandTipsforBetterMeasurements ................ 151 REFERENCES ....................................... 157 BIOGRAPHICALSKETCH ................................ 171 8

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LISTOFTABLES Table page 2-1ListofZnO-basedDMSexperimentalresults .................... 33 5-1ResistivityasafunctionofSncontentincodopedZnO:3%Mnlms. ....... 76 7-1Possiblecobalt-inducedsecondaryphases. ..................... 116 7-2Transportdatafora30%Co-dopedZnOlmwithcobaltprecipitation. ..... 132 9

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LISTOFFIGURES Figure page 2-1SchematicrepresentationofmagneticexchangebetweentwoMnionsmediatedbyadelocalizedhole. ................................. 30 2-2PredictedCurietemperaturesbasedonDietl'scalculations ............ 31 2-3Illustrationofboundmagneticpolarons. ...................... 32 3-1Schematicofpulsedlaserdepositionsystem. .................... 47 3-2VisualizationofBragg'slawforx-raydiraction. ................. 48 4-1CubicunitcellofCu2O. ............................... 57 4-2PhaseStabilitycurvesfortheCu-Cu2O-CuOsystem. ............... 58 4-3X-raydiractiondataforepitaxialCu2Oon01MgO .............. 58 4-4X-raydiractiondataforMn-dopedCu2Olmsgrownon01MgAl2O4inanoxygenpressureof1mTorr. ............................. 59 4-5X-raydiractiondataforMn-dopedCu2Olmsgrownon01MgAl2O4inanoxygenpressureof0.1mTorr. ............................ 60 4-6X-raydiractiondataforMn-dopedCu2Olmsgrownon01MgAl2O4invacuum. ........................................ 61 4-7Phaseassemblageforlmsgrownunderdierentconditions. ........... 62 4-8MagneticbehaviorforanepitaxialMn-dopedCu2Olmgrownat300Cand1mTorrofoxygen. .................................. 63 4-9MagneticbehaviorforMgAl2O4substrate. ..................... 64 4-10MagneticbehaviorforanepitaxialMn-dopedCu2Olm ............. 65 4-11LowtemperaturephotoluminescencespectraforMn-dopedCu2Olms. ..... 66 4-12Transportdatafor1%Mn-dopedCu2Olms. ................... 67 4-13Temperature-dependenttransportdatafor1%Mn-dopedCu2Olm. ...... 68 4-14Field-varyingtransportmeasurementsfora1%Mn-dopedCu2Olm. ...... 69 5-1X-raydiractionofZnOlmscodopedwithMnandSn .............. 77 5-2X-raydiractionofanepitaxialZnOlmdopedwith3%Mnand0.1%Sn .... 78 5-3XPSspectraforZnO:3%Mnlmcodopedwith0.01%Sn. ............. 79 10

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5-4PlotshowingthedependenceofthecoerciveeldonSnconcentrationatdierentSQUIDmeasurementtemperatures. ......................... 80 5-5Magnetizationmeasuredat300KforepitaxialZnO:3%Mnlmsthatarecodopedwith0.001%Sn,0.01%Sn,0.1%Sn,andnoSn. .................. 81 6-1X-raydiractionofZnOlmscodopedwithMnandPbothbeforeandafterannealing. ....................................... 90 6-2High-resolution!-rockingcurvesonZnOlmscodopedwithMnandPbeforeandafterannealing. ................................. 91 6-3AFMscansonepitaxialZnO:3%Mn,2%Plmsbeforeandafterannealing. ... 92 6-4ResistivityandcarrierconcentrationbehaviorofP-dopedZnOlms. ...... 93 6-5Transportdatafortheas-depositedZnO:3%Mn,2%Plmat300K. ....... 94 6-6XPSspectraforZnO:3%Mnlmcodopedwith2%P. ............... 95 6-7RoomtemperatureopticaltransmissionforZnO:3%Mn,2%Plms ....... 96 6-8RoomtemperatureSQUIDmeasurementsforepitaxialZnO:3%Mn,2%Plms 97 6-9SQUIDmeasurementat10KforepitaxialZnO:3%Mn,2%Plmsbeforeandafterannealing. .................................... 98 6-10Field-cooledandzeroeld-cooledmagnetizationmeasurementsforaZnO:3%Mn,2%Plmannealedat600CinO2. .......................... 99 7-1EDSresultsforaselectnumberoflmsgrownunderdierentconditions. .... 116 7-2XRDscansforaseriesoflmsgrowninvacuumat400C. ............ 117 7-3TEMmicrographsofasampledopedwith30%Co ................. 118 7-4ConvergentbeamTEMdiractionpatternsofZnOlmdopedwith30%Co ... 119 7-5XRDscansforZnOlmsdopedwith30%Co. ................... 120 7-6Thermodynamicpredominancediagramforcobaltoxides. ............. 121 7-7UV-VistransmissionofCo-dopedZnOlms .................... 122 7-8Opticalband-gapsofCo-dopedZnOlms. ..................... 123 7-9PLresultsforCo-dopedZnOlms ......................... 124 7-10SQUIDmagnetizationcurvesforCo-dopedZnOlmsdepositedat400Cinvacuum. ........................................ 125 7-11SQUIDmagnetizationcurvesforCo-dopedZnOlms ............... 126 11

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7-12AnomalousHalleectin30%Co-dopedZnO. .................... 127 7-13ThemagnetoresistanceofanundopedZnOlm. .................. 128 7-14Themagnetoresistanceof5%Co-dopedZnOlms. ................. 129 7-15Magnetoresistanceof30%Co-dopedZnOlms. ................... 130 7-16Magnetoresistanceof30%Co-dopedZnOlmattemperaturesbetween10Kto100K. ......................................... 131 7-17Temperaturedependentresistivitymeasurementsfor5%and30%Co-dopedZnOlms. ....................................... 132 7-18Hallresistivityandmagnetoresistancefora30%Co-dopedZnOlmwithcobaltprecipitation. ..................................... 133 7-19SEMmicrographsatdierentmagnicationsofthesurfaceofaCo-dopedZnOlmthathasbeenannealedinH2/Arat500Cfor60min. ............ 134 7-20-2XRDfora30%Co-dopedZnOlmbeforeandafterannealinginforminggasat500C. ..................................... 135 7-21High-resolutionXRDscansfor30%Co-dopedZnOlmthathasbeenannealedinforminggasat500C. ............................... 136 7-22SQUIDmagnetometryfora30%Co-dopedZnOlmbeforeandafterannealinginforminggasat500C. ............................... 137 7-23SQUIDmagnetometryfor30%Co-dopedZnOlmsdepositedunderdiererentconditions. ....................................... 138 A-1TheresistivityandHallmeasurementsystem. ................... 154 A-2Circuitdiagramsfor2-pointand4-pointresistivitymeasurements. ........ 155 A-3Circuitshuntcapacitanceandsettlingtime. .................... 156 12

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AbstractofDissertationPresentedtotheGraduateSchooloftheUniversityofFloridainPartialFulllmentoftheRequirementsfortheDegreeofDoctorofPhilosophyDEVELOPMENTOFTRANSITION-METALDOPEDCu2OANDZnODILUTEMAGNETICSEMICONDUCTORSByMathewP.IvillAugust2007Chair:Dr.DavidNortonMajor:MaterialsScienceandEngineeringTheeldofspintronicshasrecentlyattractedmuchattentionbecauseofitspotentialtoprovidenewfunctionalitiesandenhancedperformanceinconventionalelectronicdevices.Oxidematerialsprovideaconvenientplatformtostudythespin-basedfunctionalityinhostsemiconductingmaterial.Recenttheoreticaltreatmentspredictthatwideband-gapsemiconductors,includingZnO,canexhibithightemperatureferromagneticorderingwhendopedwithtransitionmetals.Thisworkfocusedonthepossibilityofusingwideband-gapoxidesemiconductorsaspotentialspintronicmaterials.Thestructure,magnetic,andelectronictransportpropertiesoftransition-metaldopedZnOandCu2Owereinvestigated.MnandCowereusedastransitionmetaldopants.ThinlmsofthesematerialswerefabricatedusingpulsedlaserdepositionPLD.TheMnsolubilityinCu2OwasfoundtobesmallandtheprecipitationofMn-oxideswasfavoredathighgrowthtemperatures.PhasepureMn-dopedCu2Osampleswerefoundtobenon-magnetic.Sampleswerep-typewithcarrierconcentrationsontheorderof1014{1016cm)]TJ/F19 7.97 Tf 6.586 0 Td[(3.TheeectsofcarrierconcentrationonthemagneticpropertiesofMn-dopedZnOwerestudiedusingSnandPaselectroniccodopants.Snactsasann-typedopantprovidingextraelectronstotheZnO.Pactsasap-typedopantthatsuppliesexcessholestocompensatethenativeelectronconcentrationinZnO.TheelectronconcentrationwasdecreasedusingP,butthelmsremainedn-type.Aninversecorrelationwasfound 13

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betweentheferromagnetismandtheelectronconcentration;theferromagneticcouplingbetweenMnspinsincreasedwithdecreasingelectronconcentration.ThenatureofferromagnetisminCo-dopedZnOwasalsoinvestigated.Ferromagnetismwasfoundinlmsdepositedat400Cinvacuum,whilelmsdepositedinoxygenorathighertemperatureswerenon-magnetic.Filmsdepositedundervacuumhadratherhighelectronconcentrationsandwerepresumablydopedwithoxygenvacancies.TheCo-dopedlmsalsoexhibitedpeculiarmagnetoresistanceMRthathadastrongdependenceonthecarrierconcentration.Atlowtemperatures,aprogressionfrompositivetonegativeMRwasobservedwithincreasedelectronconcentrationasthelmscrossedoverthemetal-to-insulatortransitionMIT. 14

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CHAPTER1INTRODUCTION1.1ChargeandSpinThemicroelectronicsindustryhasbeenattheforefrontofinformationtechnology,providingthedevicesnecessaryforfastandecientinformationprocessingandstorage.Theintegratedcircuit,packedwithtransistorarraysfabricatedfromsemiconductingsiliconmaterial,hasremainedtheworkhorseoftheinformationprocessingindustryoverthepast50years.Thesedevicesutilizethechargepropertiesofelectronsorholestocontroltheowofcurrentthroughthecircuit.Bytheproperorganizationoftransistorsonachip,informationcanbecomputedbythedigitallogicofelectroniccharge.Increaseintherawcomputationalpowerandspeedofthesetransistorsoverthepast50yearshasbeenpropelledbyoneimportanttrend|miniaturization.SteadilyfollowingtheearlypredictionbyGordonMoorecommonlycalledMoore'sLawminiaturizationhasledtothenumberoftransistorsonachiptodoubleaboutevery18months.Intel'scutting-edgeprocessorsarecurrentlyfabricatedwith65nmline-widths,andnextgeneration45nmarchitecturesarehittingthemarketsoon[ 1 ].Ontheotherhand,thetechnologyofinformationstoragereliesonanotherfundamentalpropertyofelectrons:thequantummechanicalelectronspin.Magnetisminsolidsisadirectconsequenceofthespinpropertyofelectrons.Electronshavetwoavailablespinstates,spin-upandspin-down.Permanentmagneticmaterialscontainanimbalanceinthenumberofspin-upandspin-downelectrons.Binaryinformationmaybeencodedintheformofnon-volatilemagneticdomainswithinthegrainsofferromagneticmaterial.Asthesizeofthesedomainsshrink,moreinformationcanbestoredperunitareaofmaterial.Increasesinmagneticstoragedensityhaveoccurredatratesfasterthananyotherindustryinhistory,withstoragedensityincreasingover50milliontimessincethecreationoftherstharddiskdrivein1957[ 2 ].Thearealdensityofdrivescontinuestoincreaseatarateofabout100%peryearandcontemporarystate-of-the-artdiskshavebeendevelopedwith 15

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densitiesover100Gbits/in2[ 3 ].Advancesinperpendicularrecordingareprojectedtopusharealdensitiesevenhigherbydeterringthesuperparamagneticlimitto1Tb/in2,andresearchintonovelrecordingschemescouldpushdensitiesevenhigher[ 4 5 ].Whilethechargeandspinpropertiesoftheelectronhaveseparatelyspawnedtwoofthemostrapidlyimprovingtechnologiesofourtime,littlehasbeendonetocombinethetwotofunctionsimultaneouslyinthesamematerial,anideathatmayleadtofurtherenhancementsininformationtechnologybeyondthelimitsofminiaturization.1.2WhatisSpintronics?Thenascenteldofspinelectronicsorspintronicsstandstoassimilatethesetwofundamentalpropertiesoftheelectron|chargeandspin|toformthebasisforanewclassofdevicedesign[ 6 { 10 ].Operatingbythemanipulation,transport,anddetectionofchargecarrierspins,spintronicsisexpectedtoimproveupontraditionalelectronicandphotonicdevices,allowingforenhancementsintheformofreducedpowerconsumption,fasterdeviceoperation,andnewformsofinformationcomputation.Spintronicsmayleadtodevicessuchasspin-polarizedLEDs,spin-FETs,andspin-basedqubitsforquantumcomputers.Increasedfunctionalitiesarealsoexpected,suchasintegratedmagnetic/electronicoperationsonthesamechip.Currently,veryfewspintronicdeviceshaveappearedonthemarketbuthavealreadymadeanastoundingimpactontechnology.Forexample,metallic-multilayeredstructuresdisplayinglargeamountsofmagnetoresistanceso-calledGiantMagnetoresistancehavereplacedconventionalhard-diskreadheads,leadingtohugeincreasesinhard-diskstorage[ 6 ].Thesedevicesconsistofasandwichstructurewherelayersofferromagneticandnon-ferromagneticmaterialarealternatelystacked.Theresistancethroughthedevicedependsontherelativemagneticorientationoftheferromagneticlayers.Whenthelayersaremagneticallyalignedsothattheirdirectionsofmagnetizationpointoppositetooneanother0degreesmisalignment,theresistancethroughthedeviceislarge.Conversely,whenthemagneticlayersarealignedparallel,theresistanceisreduced.Thesesensors 16

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havebecometheprominentformofsolid-statemagneticsensingtechnology.Asthesizeofmagnetizedbitsshrinkandarealdensitiescontinuetoincrease,sensorsmustbecomeevenmoresensitivetothesmallchangesinmagnetization.TheadventofGMRtechnologyhasbecomeamulti-billiondollarindustryandrevolutionizedthereadheadsusedinharddisktechnology.1.3DiluteMagneticSemiconductorsOfparticularinterestisthecreationandcontrolofspin-polarizedcurrentsinsemiconductingmaterial.Ferromagneticsemiconductorsprovideeasierintegrationofspintronicsintoexistingsemiconductordevices.Forinstance,highlyecientspininjectionispossiblebetweensemiconductor/semiconductorinterfaces,whereasonlyspinpolarizationsofafewpercentarepossiblebetweenaferromagneticmetal/semiconductorinterfaceduetotheconductivitymismatch[ 11 ].Ferromagneticsemiconductorsallowthe`tools'ofconventionalsemiconductortechnologytobeutilized.Thesetoolsincludep-njunctionsandheterostructures,whichprovideaconvenientplatformforawidevarietyofdevicesthatallowforelectronicgainandlightemission.However,inordertofullyrealizesemiconductor-basedspintronics,signicantchallengesrelatedtothelifetime,control,anddetectionofspinpolarizedcarriersinsemiconductorsmustbeaddressed.Materialsthatcanretaintheirferromagnetismcomfortablyaboveroomtemperaturearecrucialtothepracticalapplicationofspintronicdevices.ThesereasonshavecreatedinterestindevelopingaclassofmaterialsknownasdilutemagneticsemiconductorsDMS.DMSsaresemiconductorsdopedwithafewpercentofmagneticatoms.Themagneticatomsoccupylatticesitesandinduceferromagnetismintheotherwisenon-magneticsemiconductorhost.DMSstypicallyhaveorderingtemperaturesmuchlowerthanroomtemperature.Therehasbeensomesuccessinattainingroomtemperatureferromagnetisminvarioussemiconductors,buttheresultsarenon-reproducibleamongresearchgroupsandcontradictorytosometheoreticalpredictions.Thusthereisagreatresearchopportunityforthestudyofferromagnetism 17

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insemiconductorsforbothreal-worldapplicationsandfortheaddedknowledgetofundamentalphysics. 18

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CHAPTER2REVIEWOFDILUTEMAGNETICSEMICONDUCTORSTheideaofmagnetisminsemiconductorsisnotnew.Thequestionofwhetherchargeandspincancoexistinthesamematerialtoenhancematerialpropertieshasbeenaddressedformanyyears.Whilemagnetisminmetallicandinsulatingmaterialswaswellknown,thepossibilityofmagneticorderinginsemiconductorswasnotdiscovereduntilthemid1900s.Fromratherdicultbeginnings,theeldofmagneticsemiconductorshasseensignicantprogress,especiallyintheadvancementofmagneticallydopedsemiconductors.Signicantchallengesstillremainrelatedtothepreparationandgrowthofmaterials,understandingthephysicaloriginsofferromagnetisminsemiconductors,andraisingthemagneticorderingtemperatures,justtonameafew.Thissectionwillrstprovideaglimpseintothehistoryofexperimentalprogresssurroundingmagneticsemiconductors.Next,areviewoftheleadingtheoriesthatdelveintothephysicaloriginsofmagneticcouplinginDMSisgiven.Lastly,abriefreviewoftherecentprogressinmagneticallydopedZnOispresented.2.1MagneticSemiconductingMaterials:AShortHistoryThestoryofmagneticsemiconductorsoriginatesfromhumblebeginnings.Inthe1960sand1970s,semiconductingbehaviorinferromagneticmaterialwasuncoveredwiththediscoveryoftheeuropiumchalcogenidesEuOandchromiumspinelsCdCr2S4,CdCr2Se4.Thesematerialsaretrueferromagneticsemiconductorsinthesensethattheyhavemagneticatomsbuilt-into"thecrystalsublattice.Theelegantinterplaybetweenbandelectronsandlocalizedmagneticionsinthesematerialsbroughtaboutextensiveresearchandscienticinterestintotheeld.However,thesematerialshavenotprogressedbeyondtheeldofacademicresearchforseveralreasons[ 12 13 ].Firstofall,theircrystalstructuresareincompatiblewithconventionalsemiconductors,likeSiandGaAs,makingtheirintegrationwithcontemporaryelectronicsdicult.Thesynthesisofthesematerialsisalsocumbersomeandhardtoreproduce,makingindustrialproductionofthecrystals 19

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expensive.Andlastly,lowferromagneticorderingtemperaturesTc<100Kmakethemlessattractiveforpracticalapplications.Movinginaslightlydierentdirection,otherworkhasfocusedonmakingnon-magneticsemiconductorsmagneticbydopingthemwithsmallamountstypicallyafewpercentofmagneticatoms.Thisclassofmaterialshasattractedrenewedinterestinthedevelopmentofmagneticsemiconductors.Suchcompoundsareknownasdilutedmagneticsemiconductors,orDMS,becauseofthediluteconcentrationsofmagneticimpurities.Notice,thesematerialsarefundamentallydierentthantheEuchalcogenidesandCrspinelssincethemagneticatomsarearticiallyaddedintothelattice;themagneticatomsarenotapartoftheperiodiccrystalstructureoftheparentmaterial.EarlystudiesofDMSmaterialsbeganwithMn-dopedII-VIalloysoftheformAII1)]TJ/F21 7.97 Tf 6.586 0 Td[(xMnxBVIwhereAII=Zn,Cd,HgandBVI=S,Se,Te.Thesematerialswereheavilystudiedinthe1980sandcomprehensivelyreviewedbyFurdyna[ 14 ].ItisworthwhiletoreviewsomeaspectsofthesematerialssinceZnOalsobelongstotheII-VIfamilyofsemiconductors.Theternarynatureofthesecompoundsmakesthemamenabletotuningthelatticeandbandparametersbyvaryingalloycomposition,makingthemanattractivecandidateforthepreparationofheterostructuredevices.Thealloyscrystallizeintoeitherthezinc-blendeorwurtzitestructureandareformedbysp3tetrahedralbonding,incorporatingthevalences-electronsfromthegroupIImetalandthep-electronsfromthegroupVIelement.ElementalMnhasahalf-lled3d-shellandtwovalences2electrons.ManganeseatomsmaysubstituteonthegroupIIsitesasMn+2bygivingupthesetwovalenceelectrons.HighsolubilitiesofMninthehostmaterialswhilemaintainingthezinc-blendeorwurtzitestructuresarepossible,whichisthoughttoarisefromthechemicalsimilarityofMn+2tothegroupIIelement.The3d-shellofMnisexactlyhalf-lledandrequiressubstantialenergytoaddanelectron;thismakesthe3d5orbitactchemicallysimilartoa3d10orbit.Themagneticpropertiesofthesealloysaredictatedbytheexchangeinteractionsbetweenlocalatomicmomentsprovidedbythe 20

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Mnandthesp-bandelectrons,andhavedramaticeectsontheopticalandelectricalpropertiesofthematerial,suchasgiantFaradayrotationandboundmagneticpolaronformation.Drivenmostlybysuperexchangemechanisms|anindirectexchangeinteractionmediatedthroughtheanion|thesesystemsexhibithightemperatureparamagnetism,lowtemperaturespin-glassphase,andtypeIIIantiferromagneticordering[ 14 ].Neutrondiractionstudiesshowthattheantiferromagneticorderingofthesestructuresislimitedtoshortranges,implyingthatthemagneticorderingisconnedtotheformationofsmallclusterregions[ 14 ].Recently,however,ferromagneticorderinghasbeenachievedinlow-dimensionalquantumwellsdrivenbyhole-mediatedexchange,butwithlowCurietemperatureTc<2K[ 15 ].AnadditionalobstacletothepracticalapplicabilityofII-VImaterialisthecapabilityofdopingthematerialbothn-typeandp-typebipolardoping.Again,thesematerialsarenotpracticalsincethematerialsthatdidshowferromagnetismwererestrictedtoloworderingtemperaturesTcjustafewdegreesabove0K.Intheearly90s,atechnologicalbreakthroughintheadvancementofDMSoccurredwiththediscoveryofferromagnetismupto35KinMn-dopedInAs[ 16 { 18 ].InAsisanestablishedIII-Vcompoundsemiconductormaterial.TransitionmetalspeciesareknowntohaveverylowsolubilityinhostIII-Vmaterials,buttheproblemwasovercomebynon-equilibriumepitaxialgrowthusinglowtemperaturemolecularbeamepitaxyLTMBE.III-Vmaterialsndwidespreadapplicationintheelectronicsandoptoelectronicsindustriesashigh-speeddigitaldevices,visibleandinfra-redlight-emittingdiodesandlasers,andmagneticsensors.ThedemonstrationofferromagnetisminInMnAsoeredtheintriguingopportunitytostudyspin-basedphenomenainthesewellestablishedsemiconductordevices.Eventually,thesuccessofLTMBEgrowthofInMnAsledtothedevelopmentMn-dopedGaAsDMS.SegregationofMnsecondaryphases,namelytheMnAsphase,wassuppressedusinglowtemperaturegrowthTg=250C.If,however,thetemperaturewasraisedortheMnuxwastoohigh,phasesegregationcouldoccur.Mnactsasanacceptor 21

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dopantwhensubstitutedonthegroupIIIsitesleadingtohighholeconcentrationswhich,asexplainedlater,isanecessityforferromagnetisminthematerial.Thecouplingofthechargeandspin-basedprocesseshavebeenrepeatedlyproveninGaMnAs,includingtherealizationofspin-polarizedlightemission[ 19 ]andelectricalandopticalcontrolovertheferromagnetism[ 20 ].Unfortunately,despiteeortstoraiseit,GaMnAsislimitedbyitslowCurietemperatureof170Kwellbelowroomtemperature.RaisingtheCurietemperaturehasbeenthebiggestchallengeforGaMnAsbasedDMS.2.2DMSTheory:ThePhysicalOriginsofFerromagnetisminDMSUnderstandingthephysicalmechanismbehindmagneticorderinginDMSmaterialsisanessentialingredienttotheirfurtherdevelopment.Indeed,ifbothaconceptualandquantitativefoundationtotheoriginofferromagnetisminthesematerialsisdeveloped,theymayprovidethedirectionnecessarytoasuccessfulrecipeforthefabricationofhigherTcmaterials.Atthepresenttime,however,thereisanincompleteunderstandingoftheoriginofferromagnetisminDMSmaterialandthesubjectremainsanissueofactivedebate.Thissectionwilldiscussthecontemporarytheoriesonthesubjectandthepaththeyprovideforsubsequentresearch.2.2.1Dietl'sMean-eldTheoryThemotivationforstudyingZnOforspintronicsbeganwiththeworkofDietletal.[ 21 ].Dietlandcoworkersemployedamean-eldmodelofferromagnetism,asoriginallydescribedbyZener,tothecaseofIII-VandII-IVcompoundsemiconductors.Zener'smodelproposedthatferromagnetismisdrivenbytheexchangeinteractionbetweencarriersandlocalizedmoments.Zener'searlymodelwasabandonedbecauseitwasfoundinappropriatetodescribethemagnetismoftransitionmetals.However,DietlfoundthatitcouldbeusedtoaccuratelypredicttheferromagneticCurietemperaturesofGaMnAsandZnMnTe.Themodelassumesthattheferromagneticexchangeinteractionsoccurbetweenlocalizedspinsdopedintothesemiconductormatrixandaremediatedbychargecarriers. 22

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Thesespinsareassumedtoberandomlydistributedthroughoutthehostsemiconductorlattice.Specically,thedopedMnionsresideongroupIIorIIIsitesandprovidethelocalizedspins.Conceptually,theaectmaybeenvisagedasafeedbackbetweenthemagneticspinsandthecarriersbyaspin-spinprocess;thelocalizedmagneticspininducescarrierpolarizationwhichtheninducesmagneticpolarizationandsoonFigure2-1[ 22 23 ].ThemodelsuggeststhathighvaluesofTcareobtainableinp-typematerial,whiletheTcofn-typematerialshouldbeconstrainedtolowertemperatures.Thiscanbeattributedtoboththelargep-dexchangeintegralNoanddensityofstatesofthevalenceband,whiletheconductionband'ss-dexchangeNoanddensityofstatesaresignicantlysmaller[ 24 ].NotethatinthecaseofIII-Vsemiconductors,MnalsoactsasanacceptordopantwhereasitsubstitutesisovalentlyintheII-VIsemiconductors.DietlextendedhiscalculationtopredicttheCurietemperaturesofothersemiconductorsystemsandoxides.ThepredictionswerebasedonaMnconcentrationof5%andaholeconcentrationof3.5x1020cm)]TJ/F19 7.97 Tf 6.587 0 Td[(3.Theresultsaresummarizedingure2-2asafunctionofbandgap.Ofparticularimportance,arethepredictedCurietemperaturesinexcessofroomtemperatureforGaNandZnO.Diet'stheoryhasprovenusefulinunderstandingtheexperimentalresultsforGaMnAs.However,itdoesnotappeartobeconsistentwiththeexperimentalresultsoftransitionmetaldopedwideband-gapsemiconductors,suchasthepredictionsforGaNandZnO.Thisstemsfromseveralreasons,includingthedicultyinexperimentallypreparingp-typeZnOmaterialandthemanyobservationsofferromagnetisminn-typeZnODMS.Nevertheless,Dietl'soriginaltheoryhasledtomultipleexperimentalandcomputationalstudiesoftransitionmetaldopinginZnO.2.2.2First-principlesDesign:DFTCalculationsSatoandKatayama-Yoshidahaveemployedrst-principlesdesigntoinvestigateferromagnetisminbothsemiconductorandoxidespintronics[ 25 { 27 ].ThemagneticstabilityoftransitionmetaldopedZnOwascalculatedusingdensityfunctionaltheory 23

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DFTwithintheframeworkofthelocaldensityapproximationLDA.TherandomdistributionoftransitionalmetalionsoverthelatticecreatingdisorderinthealloywasinherentlyincludedinthecalculationsbythecoherentpotentialapproximationCPA.Magneticstabilitywascalculatedbycomparingthetotalenergydierencebetweentheferromagneticandspin-glassstate,thelowerofthetworepresentingthegroundstateofthesystem.InthecaseofMn,theirresultsareconsistentwithDietl'stheorythattheferromagneticstateisstabilizedwiththeadditionofholedoping,andwithoutholesthespin-glassstateisfavored.However,V,Cr,Fe,Co,andNiimpuritieswerepredictedtobeferromagneticwithouttheneedofadditionalchargecarriers.Electrondopingfurtherstabilizedtheferromagneticstateinthesealloys.Theirworkalsopointstoacontributionofd-statesattheFermilevel,hintingatsomedelocalizationofd-states.Itwassuggestedthatthiscouldleadtoferromagneticorderingthroughadouble-exchangeinteraction,inwhichferromagneticalignmentisstabilizedbythehoppingof3delectronsbetweenneighboringTMsites.Thismechanismisdrivenbypartiallyunoccupiedup-spinordown-spinstatesinthe3d-bandandisthereforenotpossibleinthecaseofMn,whichexhibitsahalf-lled3dband.Inthecaseofp-typedoping,however,thetransferenceofweakly-bound3delectronsbetweenMnionsmaybemediatedbythepresenceofholes.Thevalencebandp-stateshybridizewiththe3d-statesofMnanditinerantholescanretaintheird-likecharacter.ThisstabilizestheferromagneticphaseforMndoping.2.2.3FerromagnetisminDisorderedAlloysAnadditionaltheoreticalapproachconsiderswhetherferromagneticorderingbetweenthelocalizedspinscanoriginatefromlocalizedcarriers.Again,themodelisdevelopedinthemean-eldtreatmentbutaccountsforpositionaldisorderinthealloy.Numericalstudieswithinthemean-eldtreatmentshowthatthenatureofferromagnetismisstronglyaectedbythisdisorderandthattheTccanbepushedtohighertemperatureswithincreasingrandomnessinthepositionofMnions[ 28 ]. 24

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FerromagnetisminthislocalizedcarrierregimecanbeexplainedthroughtheformationofboundmagneticpolaronsBMPs[ 29 30 ].ABMPisaquasi-particlecomprisedofthelocalizedcarrierandthemagneticatomsencompassedwithinitsradiusFigure2-3.Thelocalizedcarrierisboundtoitsassociateddefectsuchasadonoratomifthecarrierisanelectroninahydrogenicorbitalofradius,rh=m/m*ao,whereisthehighfrequencydielectricconstant,m*istheeectivemass,andaoistheBohrradius3pm[ 31 ].Thisradiuscanbelarge8AinZnOextendingoverseverallatticeconstants,andcanencompassanumberofmagneticdopantsdependingontheirconcentration.TheexchangeinteractionbetweentheboundcarrierandthemagneticmomentstendstoalignthemomentsparalleltooneanotherinsidetheBMP.Athightemperatures,theBMPsmaybeisolatedfromoneanother.However,asthetemperatureislowered,theBMPradiusgrowsandtheindividualBMPsbegintooverlap.OverlappingBMPsbecomecorrelatedandtheirspinsalign,producinglongrangeferromagneticinteractions[ 29 ].Atacriticaltemperature,theoverlappingBMPsarepercolatedthroughoutthesampleandthetransitiontoferromagnetismoccurs.TheBMPmodelisequallyapplicableton-typeorp-typematerial[ 30 ].TheBMPmodelallowsforferromagnetisminaninsulatingorsemi-insulatingregime.ThisisespeciallyattractiveinthecaseofZnOwherehighp-typedoping,asrequiredbyDietl'smodel,isinherentlydicult.2.2.4FerromagnetisminaSpin-splitConductionBandCoeyetal.haveproposedanothermodelforferromagnetisminDMSmaterialsbasedonaspin-splitdonorimpurityband[ 31 ].Themodelisconsistentwiththeobservedmagnetizationforn-typetransition-metaldopedZnO.Inthismodel,donordefectswhichcouldarisefromeitheroxygenvacanciesorzincinterstitialsinthecaseofZnOoverlapatlargeconcentrationstoformanimpurityband.TheimpuritybandcaninteractwithlocalmagneticmomentsthroughtheformationofboundmagneticpolaronsBMP.WithineachBMP,theboundcarrierinteractswiththemagneticdopantsinsideitsradiusand 25

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canalignthespinsofthemagneticdopantsparalleltooneanother.FerromagnetismisachievedwhentheBMPsstarttooverlaptoformacontinuouschainthroughoutthematerial,thuspercolatingferromagnetismintheDMS.However,Coeyshowedthatinthismodel,toachieveahighTc,afractionofthepolaronicchargemustdelocalizeorhybridizeontoeachmagneticdopant.Inabandscheme,thisoccurswhentheimpuritybandoverlapswithunoccupiedd-levelsofthemagneticdopant.ItwasshownthatforSc,Ti,andV,thespin-upstatesofthe3dTMmetalarealignedwiththeimpuritylevels,resultingissignicantalignment.SimilarlyforFe,Co,andNidoping,thespin-downstatesperformthesamefunction.Interestingly,MnandCrdopingwouldnotleadtostrongmagnetizationduetosmallhybridization.WithintheframeworkofCoey'smodel,Kittilstvedetal.haveperformeddetailedspectroscopicexperimentsoncobalt-dopedZnO[ 32 ].TheirresultsshowthatthesinglyionizedCo+stateliesclosetotheconductionband,similarinenergytoashallowdonorstate.Sincetheenergiesaresimilar,chargetransfercanoccurbetweenthecobaltatomsandthedonorimpurities,thusleadingtothehybridizationnecessaryforferromagnetism.Kittilstvedetal.hasalsoshownthatthisleadstoaninherentpolaritydierenceforferromagnetismincobaltandmanganese-dopedZnO.Whereasferromagnetismincobalt-dopedZnOiscloselytiedtothepresenceofshallowdonors,manganese-dopedZnOiscloselytiedtothepresenceofshallowacceptors.ThedierenceliesinthelocationofthesinglyionizedMn+3state,whichsitsclosetothevalencebandinZnO.Thisideaisdescribedfurtherattheendofthenextsection.2.3ExperimentalProgressinZnODMSOntheexperimentalfront,therehasbeenawidedistributioninthemagneticpropertiesreportedfortransitionmetaldopedZnO.Experimentshavenowcoveredabroadrangeofparameters,includingvarioustransition-metaldopantseveryelementintherstrowofthetransitionmetalserieshasnowbeensurveyed,compositionalvariations,preparationtechniquesandgrowthconditions,andpost-growthprocessing. 26

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Theobservedresultsareoftenconictingandnon-reproduciblebetweenresearchgroups.Thediscrepancyintheobservedpropertiesandtheirinterpretationlikelystemsfromdierentgrowthtechniquesandconditions,andinsucientcharacterization.MostofthedicultiesariseindeterminingifthematerialisatrueDMSTMatomsrandomlysubstitutingZnlatticesitesorifferromagnetismoriginatesfromTMclusteringordopant-inducedsecondaryphases.Inanycase,theresultsindicatethattheunderlyingmechanismsofferromagnetisminZnODMSarequitesensitivetogrowthconditionsandmustbeclearlydelineatedbycarefulanalysis.DescribingalltheexperimentaltrialsinZnODMSoverthepastseveralyearswouldbetediousandoverwhelming.Therearealreadyseveralreviewscoveringthesubject[ 33 { 36 ].AcompilationofsomeresultsislistedinTable2-1[ 36 ].Instead,toprovideaavoroftheexperimentalprogress,abriefsummaryoftheimportantachievementssurroundingtransition-metaldopedZnOisprovided.2.3.1Mn-dopedZnOByfar,thetwomoststudiedmagneticdopantsinZnOhavebeenMnandCo.FukumuraandcoworkerswerethersttostudyMn-dopedZnODMSusingPLD[ 37 ].Alargesolubilityof35%MnwasachievedwhileretainingthewurtzitestructureofZnOreminiscentoftheearlierstudiesonII-Mn-VIcompoundsdiscussedearlier.Thisisoverthethermodynamicsolid-solubilitylimitofMninZnOandisatestamenttothenon-equilibriumconditionsobtainablebythinlmgrowth.Theylatershowedtheheavilydopedalloytoexhibitspin-glassbehaviorwithaspin-freezingtemperatureof13KduetostrongantiferromagneticexchangecouplingbetweenneighboringMnatoms[ 38 ].ThehighsolubilityofMnachievedinZnOmotivatedotherexperimentaleortsintothesynthesisofZnMnO.Whilesomegroupsreportedferromagnetism,othersobservedantiferromagnetic,spin-glass,orparamagneticbehaviorforexample,refertoreference[ 36 ]. 27

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SharmaandcoworkerswerethersttoreportferromagnetismaboveroomtemperatureindiluteMn-dopedZnObulkandthinlmsamples[ 39 ].Bulkpelletswithanominalconcentrationof2at%MnEDSshowedtheactualconcentrationtobemuchlower:0.3at%sinteredbelow700CwerefoundthehaveaCurietemperatureof420K.Additionally,thinlmsdepositedbyPLDwith2.2at%Mnwereshowntoexhibitferromagnetismatroomtemperature.However,usingsimilarpreparationtechniquestoSharmaetal.,Kundaliyaandcoworkers[ 40 ]convincinglydemonstratedthattheobservedhigh-temperatureferromagnetismresultedfromametastablephaseoxygen-vacancy-stabilizedMn2)]TJ/F21 7.97 Tf 6.586 0 Td[(xZnxO3)]TJ/F21 7.97 Tf 6.587 0 Td[(,andnotfromtheproposedcarrier-mediatedferromagnetismbetweenMnatoms.ThereisalsodiscrepancyinthereportedoveralldistributionofMnatoms.Forexample,ahomogenousdistributionofMnwasobservedbyChengandChien[ 41 ],whileJinetal.[ 42 ]foundclusteringofMnatoms.Clearly,thoroughcharacterizationisneededtofullyappreciateandunderstandtheoriginofferromagnetisminthesematerials.2.3.2Co-dopedZnOOneoftheearlyworksoncobalt-dopedZnODMSwasbyUedaetal.[ 43 ].Theyfoundthematerialtobeferromagneticabove280Kwith5-25%Coand1%Aladdedasann-typedopantwithouttheadditionofsecondaryphases.Dierencesinthemagnetizationwereattributedtodierencesintheconductivity;lmswithhighercarrierconcentrations1020cm)]TJ/F19 7.97 Tf 6.586 0 Td[(3showedferromagneticfeatureswithhigherMsandTc.Sincethen,additionalexperimentalstudieshaveinvestigatedthepropertiesandoriginofferromagnetismincobalt-dopedZnO.Again,theresultsareconictingwithreportsofferromagnetisminphasepurelms[ 44 45 ],ferromagnetismfromclusters[ 46 ],andnoobservedferromagnetism[ 47 ].TherstreportofreversiblecontrolledswitchingofferromagnetisminanyDMSwasdemonstratedbySchwartzandGamelinincobalt-dopedZnO[ 48 ].ThereversibilitywasmediatedbytheincorporationandremovalofZninterstitials.TheZninterstitialZniisaknownn-typedopantthatproducesashallowdonorlevelbelow 28

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theconductionband.DiusingZniintothelatticelowerstheconductivityandactivatesroomtemperatureferromagnetism.RemovingZni,byheatinginair,returnedthelmstoaninsulatingstateandsubsequentlyquenchedtheferromagnetism.Theprocesswasreversibleovermanycycles.Thisreversibilityisevidencethatfreecarriersactivateferromagnetismincobalt-dopedZnO.TheprocesswasobservedinbothMOCVDgrownlmsandZnO:Conanoparticlelmspreparedbyspincoating.StronghybridizationofZnidonorstateswithCo+2statesneartheconductionbandwhich,asexplainedearlier,istheoreticallybelievedtocauseferromagnetismwasusedtoexplainthemagneticordering.Conductionelectrons,derivedfromtheZnidonors,delocalizeoverseveralCo+2ionsandferromagneticallyaligntheirspinsthroughadoubleexchangeinteraction.Importantly,fromthesamegroup,KittilstvedwasabletodemonstrateachemicalpolaritydierencebetweentheferromagnetisminZnCoOandZnMnO[ 49 ].Specically,p-typeZnMnOledtoferromagnetism,whileferromagnetisminZnCoOwasactivatedbyn-typedoping.DopingoftheZnMnOwasdoneonalocallevelbyN-cappingZnMnOnanoparticleswithamines.ZnCoOnanoparticlelmsweremaden-typebycappingwithoxygen.Reversingthecappinglayers,ZnCoO:NandZnMnO:O,ledtothedisappearanceofferromagnetisminbothsetsoflms.Opticalabsorption,MCD,andphotoconductivitymeasurementswereemployedtounderstandthisinherentpolaritydierence[ 32 ].Forn-typeZnCoO,theauthorsshowedthataresonanceinthechargetransferCo+1!Co+2+e)]TJ/F21 7.97 Tf 0 -8.189 Td[(CB,E0.27eVanddonorstateenergiescanleadtoalargehybridizationnecessaryforferromagnetism.ForZnMnO,asimilarresonancewasobservedbutderivedfromtheMn+3stateclosetothevalencebandMn+3!Mn+2+h+VB,E0.22eVwithacceptorstateenergies. 29

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Figure2-1.SchematicrepresentationofmagneticexchangebetweentwoMnionsmediatedbyadelocalizedhole.Adaptedfrom[ 22 ]. 30

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Figure2-2.PredictedCurietemperaturesbasedonDietl'scalculations[ 21 ]for3%Mnandaholeconcentrationof3.5x1020cm)]TJ/F19 7.97 Tf 6.587 0 Td[(3After[ 50 ]. 31

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Figure2-3.Illustrationofboundmagneticpolarons.AnelectronboundwithinitshydrogenicorbitalcouplestomagneticimpuritiescausingthemtoalignparalleltooneanotherinsidetheorbitradiusAdaptedfrom[ 31 ]. 32

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Table2-1:ListofZnO-basedDMSexperimentalresultsAfter[ 36 ] CompoundTMSubstrateFabricationGrowthO2pressurePost-TcKNotesRefscontentmethodtempCTorrannealing ZnO:Mn<0.35c-sapphirePLD6005x10)]TJ/F6 4.981 Tf 5.397 0 Td[(5N/A[ 37 ]ZnO:Mn0.36c-sapphirePLD6005x10)]TJ/F6 4.981 Tf 5.397 0 Td[(5N/ASpin-glass[ 38 ]Zn1)]TJ/F9 4.981 Tf 5.396 0 Td[(xTMxOc-sapphirePLD500-60010)]TJ/F6 4.981 Tf 5.397 0 Td[(9to10)]TJ/F6 4.981 Tf 5.397 0 Td[(6N/A[ 47 ]ZnO:Co0.02-0.5c-sapphirePLD300-70010)]TJ/F6 4.981 Tf 5.397 0 Td[(6to10)]TJ/F6 4.981 Tf 5.397 0 Td[(1Spin-glass[ 51 ]ZnO:Mn.01-0.36c-sapphirePLD6105x10)]TJ/F6 4.981 Tf 5.397 0 Td[(5Paramagnetic[ 41 ]ZnO:Mn0.07a-sapphiresputtering4000.06Paramagnetic[ 41 ]ZnO:Mn0.03-0.2GaAssputtering500-6008x10-4[ 52 ]ZnO:Co,Mn,0.05-0.25r-sapphirePLD350-6002-4x10)]TJ/F6 4.981 Tf 5.397 0 Td[(5280-2B/Co[ 43 ]Cr,orNi300ZnO:Ni0.01-0.25c-sapphirePLD300-7001x10)]TJ/F6 4.981 Tf 5.397 0 Td[(5Superpara-or[ 53 ]ferromagneticZnO:V0.05-0.15r-sapphirePLD30010)]TJ/F6 4.981 Tf 5.397 0 Td[(5to10)]TJ/F6 4.981 Tf 5.397 0 Td[(3>3500.5B/V[ 54 ]ZnO:Co,Fe<0.15SiO2/SiMagnetron6002x10)]TJ/F6 4.981 Tf 5.397 0 Td[(3600C,10min,>30012-15emu/cm3[ 55 ]sputtering1.0x10)]TJ/F6 4.981 Tf 5.396 0 Td[(5TorrZnO:Co0.03-0.05BulkZnOIon700C,5min,>300OrientedCo[ 56 ]implantationunderO2precipitatesZnO:Co0-0.25c-sapphireSol|gel<350700C,1min>3500.56B/Co[ 57 ]ZnO:Mn0-0.3c-sapphirePLD>30-450.15-0.17B/Mn[ 58 ]ZnO:Mn<0.04Sintered500-700Air,atm.>4250.006emu/gm[ 39 ]pelletspressuresinglephaseZnO:Mn0.02FusedPLD400>4250.05emu/gm,[ 39 59 ]quartzsinglephaseZnO:Fe,Cu0-0.1Solid-state8975500.75B/Fe[ 60 ]reactionZnO:Co0.015PLD6505x10)]TJ/F6 4.981 Tf 5.397 0 Td[(5>300Ferromagnetic[ 61 ]ZnO:Co,Al0.04-0.12GlassRF1x10)]TJ/F6 4.981 Tf 5.397 0 Td[(2>3500.21B/Co[ 62 ]sputteringinArZnO:Mn,Sn0-0.3Implantation700C,5min250Ferromagnetic[ 63 ]ZnO:Mn,SnMn=0.03,c-sapphirePLD400-6000.02>300Ferromagnetic[ 64 ]Sn<0.1%ZnO:MnandCo0.05-0.15CrystallineAntiferromagnetic[ 65 ]precursorZnO:MnandCo<0.05BulkMeltgrowth1000Paramagnetic[ 66 ]ZnO:Co0.1O-faceZnOPLDAntiferromagnetic[ 67 ]ZnO:Co<0.35r-sapphireMOCVD350-60040500C,20min>350Ferromagnetic[ 67 ]invacuumZnO:CoandFe<0.15SiO2/SiMagnetron6002x10)]TJ/F6 4.981 Tf 5.397 0 Td[(3600,10min>300Ferromagnetic[ 68 ]sputtering10)]TJ/F6 4.981 Tf 5.397 0 Td[(5Torr12-15emu/cm3ZnO:Mn0.1r-sapphirePLD6500.1>3000.075B/Mn[ 69 ]ZnO:MnandCu0.05-0.1r-sapphirePLD6500.14000.1B/Mn[ 69 70 ]ZnO:Sc,Ti,V,0.05r-sapphirePLD6000.1-750>3000.5B/Ti,5.9B/Co[ 60 ]Fe,CoorNi0.3B/ScZnO:Mn0.02BulkPowder,500>3000.16B/Mn[ 39 59 ]pelletspelletsandlaser-ablatedlmsZnO:CrSTOPLD>400Ferromagnetic[ 71 ]ZnO:Mn0.08TetrapodsEvaporation60043Zn,MnMn204phases[ 72 ]ZnO:Mn0.05ZnOsubPLD200-600250HighTclowerTG[ 73 ]ZnO:Mn0.02PolypelletsPowder400-8000.4>3000.18B/Mn[ 40 ]pelletsandandPLDthinlmsZnO:Mn0.02-0.1PelletsPowder400500-800C>300Interfacialphase[ 74 ] 33

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CHAPTER3THINFILMDEPOSITIONANDEXPERIMENTATIONPulsedLaserDepositionPLDwasusedforallthelmdepositionpresentedinthisdissertation.Thissectionwillcovertheexperimentalapparatusandmethodsfortheselmdepositions.AbriefdiscussionofthebasicPLDprocesswillbegivenrst.ThisisfollowedbyadescriptionofthespecicPLDsystemusedfordeposition,includingthegrowthchamberandUVexcimerlaser.Thentheexactexperimentalstepsusedtodepositlmsinthisworkwillbedescribed,includingthepreparationofoxidetargetsthesourcematerialsandthetypicalgrowthprocedureforobtainingepitaxialthinlms.3.1PLDasaToolforThinFilmOxidesPLDhasevolvedintoasuccessfulandwidely-usedresearchtoolforthefabricationofthin-lmcomplexoxides[ 75 76 ].PLDisaphysicaldepositionmethod.High-intensitylaserpulsesupwardsof100sMWcm)]TJ/F19 7.97 Tf 6.586 0 Td[(2areusedtovaporizeatomicspeciesfromatargetofdesiredchemicalcomposition.Thetargetmaterialistypicallyasolidsource,butliquidsources,suchasorganicliquids,arealsofeasible[ 77 ].Thelaserenergyisabsorbedbythetargetandarapidlyexpandingplume,containingelectronsandground-andexcited-stateneutralatomsandions,isejectedfromthetargetsurface[ 76 ].Thisplumeishighlydirectional.Itisemittedperpendiculartothetargetsurfaceandhasanangulardistributiongivenbyf=cosn,wherefisthedistributionofablateduxandn5-25[ 76 ].Theablatedmaterialiscollectedontoaheatedsubstratelocatedseveralcentimetersfromthetarget.PLDgrowthoersseveraladvantagesintherealmofthinlmdeposition[ 76 ]: 1. Congruentstoichiometrictransferofthetargetmaterialtothedepositedlm.Thevariouschemicalcomponentsofcomplexoxides,orothermulticomponentmaterials,areevaporatedsimultaneously.Thisallowsthecontroloflmcompositionbysimplypreparingtargetsofthedesiredcomposition.Typically,thecompositionalstoichiometryofthetargetisaccuratelyreproducedinthelm. 2. Almostanythingcanbeablated,includinghighmeltingtemperatureoxides. 34

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3. Thelaserbeamiscapableofimpartingconsiderableamountsofkineticenergytotheablatedchemicalspecies.Theadditionalkineticenergycanleadtolargerstickingcoecientsandenhancedadatommobilityonthegrowthsurface,permittingepitaxyatreducedgrowthtemperatures. 4. Depositioninabackgroundgasispermissible.Thehighlyenergeticplumecanreadilyundergogas-phasereactionswithambientgases,suchasO2,H2O,andNO2,providinganadditiondegreeoffreedominthegrowthparametersofthinlmoxides.3.2PLDSystemUsedforThisWork3.2.1TheGrowthEnvironmentAcommercialPLDsystembuiltbyNeocerawasusedforthethinlmdepositionsdescribedinthisdissertation.AdiagramofthesystemisprovidedinFigure3-1.Thesystemiscomposedofonemainvacuumchamberwhichisusedforgrowth.Thischamberisevacuatedbyapairofpumpsthataresituatedunderneaththechamberandcanbesealed-obythegatevalve.Thelaserusedfortargetablationisseparatefromthesystem.Theothercomponentsofthesystemarecontainedwithinthegrowthchamber.TheNeocerasystemusesaPfeierturbopumptocreateandmaintainthelowpressuresneededforgrowth.Theturbopumpconsistsofaseriesofrapidlyspinningbladestotransfergasoutofthechamber.Essentially,gasmoleculescollidewiththebladesandarekineticallydriventowardthepumpexhaust.Thisisaccomplishedthroughmultiplestagesuntiltheexhaustgasiscompressedtothefore-linepressureofthebackingpump.Theturbopumpiswatercooledbyarecirculatingchillertosuppressoverheating.Thesystem'sturbopumpisbackedbyanoil-lessfour-chamberdiaphragmpump.Thisbackingpumpservestwopurposes.Beforetheturbopumpcanbeswitchedon,thebackingpumproughpumpsthegrowthchamberdownfromatmospheretoaninletpressuremanageablebytheturbopump.Thegrowthchamberisevacuatedtoapressureofatleast240Torrbeforetheturbopumpisswitchedon.Thebackingpumpservesitssecondpurposebymaintainingthefore-linepressurefortheturbopumptooperateecientlyandremovingtheexhaustgasintotheatmosphere.Sincethereisnoload-lock, 35

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thechamberhastobeevacuatedfromatmosphereforeverygrowthrun.Theultimatebasepressureofthesystemis2x10)]TJ/F19 7.97 Tf 6.586 0 Td[(6Torr.Thepressureinsidethegrowthchamberismonitoredbyapairofvacuumgauges.Pressuresabove1mTorraremonitoredwithaGranvilleconvectrongauge.Pressuresbelow1mTorraremonitoredusingacold-cathodegauge.Theablationtargetsaremountedontoarotatingcarousel.Atotalofsixtargetscanbemountedatonetimeandeachtargetcanbeaccessedduringgrowthbyrotatingthecarouseltothepositionofthedesiredtarget.Thisallowsforthegrowthoflmstacksthatcanbeusedfordevicefabrication,compositionallygradedlayers,superlattices,etc.Theablationtargetisalsorotatedarounditscenteraxisduringdeposition.Targetrotationhelpsmaintainuniformtargetablationandlmdeposition.Thesubstratesaremountedontoasolidblock,resistiveheater.Theheaterismountedtoavacuumangethatcanberemovedtomountthesamplesandprovideaccesstothegrowthchambertoexchangetargets.LaseraccessintothegrowthchamberisgrantedthroughanopticalwindowmadefromSUPRASILfusedsilica,whichhasalargeopticaltransparencytoUVradiationinordertominimizeattenuationofthelaserenergy.3.2.2TheLaserSourceALambda-PhysikCompex205KrFpulsedlaserisusedastheablationsource.Thislaserproducesacoherentbeamwitha248nmwavelength.Thelaserenergyandpulserepetitioncanbevariedtosuitaparticularexperiment.Mostofthegrowthinthisdissertationwasdoneatenergiesaround300mJwithrepetitionratesrangingfrom1to10Hz.ThelaseroutputisdirectedintothePLDchamberbyaseriesoffourmirrors.Oneofthemirrorsismountedtoamicroprocessordrivenrotationalstage.Thisstagecanbeprogrammedtorepeatedlyscanthelaseracrossthesurfaceoftheablationtarget.Thereisanapproximate10%lossofbeamenergyateachmirror.Thebeamisfocusedusingtwoplanefocusinglensestofocusboththehorizontalandverticalwidthofthe 36

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beam.Typically,thebeamwasfocusedtoaspotsizeof4-6mm2.TheCompex205isachemicalexcimerlaserthatemitsultravioletUVlightatawavelengthof248nm.Theterm`excimer'isderivedfromEXCIteddiMER.TheCompexusesaKrFexcimer,derivedfromultra-highpurityKrandF2gases,asthesourceoflightemission.Theexcimerisformedbysupplyingelectricalenergytothegasmix.SinceKrisanoblegas,thebondedKrFcomplexisahighlyunstableexcitedstateandquicklydecomposestotheunbondedgroundstate.Thestoredchemicalenergyisreleasedintheformofaphotonwiththecharacteristicdecompositionenergythelasingenergy:2KrFg=2Krg+F2g+energy48nm3.3TypicalGrowthProcedureAtypicalgrowthrunusingthePLDsystemisdescribednext.Describingthedepositionprocessindetailcanalleviateambiguityateachstepandprovideameansoftrackingdowndierencesinlmpreparationbetweenusers.Theprocessisdescribedintwosteps,substratepreparationandlmgrowth.3.3.1SubstratePreparationThesapphiresubstratesusedforZnOlmgrowthtypicallycameintheformof2"diameterwafers.Squaresubstrates6.5mmx6.5mmwerecleavedfromthelargerwafersusingadiamondscribe.TheMgAl2O4andMgOsubstratesusedfortheCu2Olmscamepre-cutin1cmx1cmsquaresandwerecutintofour5mmx5mmsquaresforlmgrowth.BeforeeachdepositionthesubstrateswerechemicallydegreasedinsequentialbathsofTCE,acetone,andmethanol.Thedegreasingmethodwasthesameforallsubstrates.First,about15mLofeachsolventwaspouredintoseparate50mLbeakers.Thebottomsofthebeakerswerethenplacedintoanultrasonicbathoftapwatertoaidthesolventinremovinganycontaminationorresidueinsidethebeaker.Thebeakerswereultrasonicatedforatleast5minutes.Next,eachbeakerwasdumpedoutandrelledwithabout20mLoffreshsolvent.ThesubstrateswererstplacedintothebeakerofTCEandultrasonicallyrinsedforatleast5minutes.Tokeepthesubstratesothebottomofthebeaker,the 37

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substrateswereheldwithplastictweezerswhichwerepinchedclosedusingabinderclip.Thesubstrateswerethensubsequentlycleanedintheacetonebeakerandthenthemethanolbeaker.Inordertokeepthesubstratescleanofdriedsolventresidue,theywereimmediatelyblowndrywithN2afterbeingremovedfromthemethanol.Aftercleaningthesubstratestheyweremountedtotheheaterblock.Theheaterblockwascleanedbeforeattachingthesubstrates.Theblockwasscrapedcleanofdrysilverpaintleftfromthepreviousdepositionusingastainlesssteelrazorblade,etchedin1:1Nitricacid/DiH2O,andwipedwithacetone.ThesubstrateswereattachedtotheblockwithsilverpaintTedPella`Leitsilber200'silverpaintandallowedtodryfor20mininsideafumehood.Aglasspetridishwasplacedupside-downoverthesubstratestoprotectthemfromdustduringthedryingtime.3.3.2ThinFilmGrowthAfterdrying,thesamplesandheaterassemblyaremountedinsidethechamberbyreattachingtheheaterangetothesystem.Thegatevalveisopenedandthebackingpumpisswitchedontoroughpumpthegrowthchamberfromatmosphere.Whenthepressuredropsbelow250Torr,theturbopumpisthenturnedon.Thesystemisallowedtopumpforafewhoursuntilapressurebelow1x10)]TJ/F19 7.97 Tf 6.586 0 Td[(5Torrisreached.Thetemperaturecontrollerisrampedtothedesiredsubstratetemperatureatarateof10C/min.Oxygenisowedintollthechamberwith10mTorrofO2.Thisisthoughttohelpdriveoanyhydrocarboncontaminationfromthesubstrateduringheating.Whenthesettemperatureisreached,thesampleshutterisclosedtoblockthesubstratesfromthetarget,theoxygenowisstopped,andthetargetispre-ablatedat300mJforatleast1000laserpulses.Pre-ablationcleansthetargetsurfacebeforeactuallmgrowth.Sincethechamberpressurerisesfromvaporizedspeciesothechamberwallduringheating,thechamberisleftattemperaturein10mTorrO2untilthebasepressuredropsbelow1x10)]TJ/F19 7.97 Tf 6.586 0 Td[(5Torr.Tostartthelmdeposition,theshutterisippedopen,oxygenisintroducedintothechamberatthedesiredgrowthpressure,thelaserenergyandrepratearesetandthelaser 38

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isswitchedon.Duringthegrowth,thelaserbeamisscannedandthetargetisrotatedtoprovideuniformablationacrossthetargetsurface.Oncethegrowthiscompleted,thetemperatureisrampedbacktoroomtemperature.Theoxygenoverpressureisusuallymaintainedatthegrowthpressureduringcooldown.Thesamplesarethenremovedfromthesystemandstoredinadesiccatoruntilneeded.3.4FabricatingPLDAblationTargetsCeramictargetsofknowncompositionwereusedasthesourcematerialforthePLDlmgrowth.Filmsweredepositedbyablatingthesetargetsinsidethegrowthchamber.Theablationtargetswerefabricatedusingthesolid-statereactionofmixedoxidepowders.Thehighpuritysourcepowders>99.99%purewerepurchasedfromAlfa-Aesar.Targetswerepreparedbasedonatomicpercentconcentrationsofthedopantspeciesincontrasttoweightpercent.Forexample,aZn0:97Mn0:03Otargetcontains3at%MndopedintotheZnO.Standarddimensionalanalysisusingatomicweightsfromtheperiodictablewereusedtoconvertat%toweightpercentsotheappropriateamountsofpowdercouldbeweighed.Thepowderswereweighedinaplasticweighingdishusinganelectronicscale.Afterweighing,thepowdersweretransferredtoanaluminamortar.Methanolwasaddedtothepowderstoaidwithmixingandthepowdersweregroundandthoroughlymixedusinganaluminapestle.Thepowderswerethenallowedtodryinair.Auniaxialpresswasusedtocompactthepowdersintoa1"diameterdiscusingastainlesssteeldie.Theappliedpressurewastypically2{4metrictons.Eachdiscwasplacedonathinpieceofaluminaandplacedinsideahightemperatureboxfurnace.Thepelletswerethencoveredwithanaluminacrucible.Theboxfurnacewasraisedto1000Catarateof10C/min.Thetargetsweresinteredat1000Cinairfor12hoursandallowedtocooltoroomtemperature. 39

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3.5ThinFilmCharacterization3.5.1X-rayDiractionX-raydiractionXRDisaversatiletoolforstudyingthecrystallinenatureofmaterials.Analysisofgrainsize,lmquality,andphaseidenticationarepossiblethroughXRDtechniques.IusedpowderXRDtoinvestigatethestructuralinformationofthedepositedlms.Specically,crystallinephasesofmaterialweredeterminedbytheta-2thetadiractionscans.ThelmsweremeasuredwithaPhilipsAPD3720diractometer.ThesystemusesaCux-raysourcethatemitsprimarilyCuK1photonswithawavelengthof1.5405A.ANilterabsorbsmostoftheothercharacteristicwavelengthsfromthesource,althoughsomeCuK2andKphotonsescapetowardthesample.Thesewavelengthsareimportantbecausetheygenerateextradiractionpeaksinthecollecteddata.InXRD,theincidentx-raysinteractwiththematerialthroughconstructiveanddestructiveinterferencefromtheperiodiclatticeplanesofthecrystal.TheconditionforconstructiveinterferenceisgovernedbyBragg'sLaw:n=2dsinwhereisthex-raywavelength,disthedistancebetweenatomicplanes,andistheanglebetweentheseplanesandthex-raysourceFigure3-2.Thereforeonlyspecicd-spacingssatisfytheBraggconditionatacertainangle.ThePhilipssystemisconguredinBragg-Brentanogeometry.Thex-raysourceispointedatthesampleandheldinaxedposition.Thesamplerotatesaroundtheangle,andthedetectorrotatesaround2.Inthis-2geometry,onlyplanesparalleltothesamplesurfacecansatisfytheBraggcondition.Thisprovidescrystallineorientationinformationofthedepositedlm.Thex-rayintensityisplottedasafunctionof2torevealthediractionpatternforasample.Phaseidenticationispossiblebycomparingthediractionpeaksto 40

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knownmaterialstandardsintheJCPDScatalog.Withthismethod,recognitionofanydopant-inducedsecondaryphaseswithinthesemiconductorhostwasdetermined.Additionalx-raycharacterizationforselectsampleswascarriedoutonaPhilipsX'perthigh-resolutiondiractometer.!-rockingcurveswereusedtoassessthelatticeparameterandcrystallinequalityofthedepositedlms.Theepitaxialqualitybetweenlmswasquantiedbythe!-rockingcurveFWHMfull-widthathalfmaximum.-scanswereusedtomeasurethein-planeorientationandlmtexturetoconrmepitaxialregistrywiththesubstrate.3.5.2MagnetoresistanceandHallEectMeasurementsTransportmeasurementsprovidevaluableinsightintotheelectronicstructureofmaterials.Electricalpropertiesaresensitivetomaterialimperfections,includingelectricaldopants,atomicimpurities,latticestrain,andstructuraldefects.ThetransportpropertiesoflmsgrowninthisworkwereprobedbymagnetoresistanceandHalleectmeasurements.SamplesweremeasuredusingeitherVanderPauworHallbridgegeometries.VanderPauwsampleswereapproximately6.5mmx6.5mmsquare.Hallbridgespecimenswerepatternedduringlmgrowthbydepositingthroughastainlesssteelshadowmask.Contactsweresolderedtothesamplesusingindiummetal.ThelmsweremountedinsideaQuantumDesignPhysicalPropertyMeasurementSystemPPMStocontroltheambienttemperatureandmagneticeldnearthesample,andtheelectricaldatawascollectedusingaKeithleyhigh-impedanceHalleectsystem.Detailsofthecongurationareprovidedintheappendix.TheHalleectprovidesvaluableinformationaboutthechargecarriersinamaterial.TheHalleectiscausedbytheinteractionofmovingchargecarrierswithanappliedmagnetic.ThisinteractionisdictatedbytheLorentzforceandleadstoanaccumulationofchargeatthesampleedge.TheaccumulatedchargedistributioninducesapotentialdroptheHallvoltagebetweenthesampleedgesuntiltheLorentzforceisbalanced.The 41

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carriertype,mobility,anddensitycanbedeterminedfrommeasurementsoftheinducedHallvoltage.Halleectmeasurementsbecomeespeciallyvaluableinthestudyofmagneticmaterials.Inaferromagneticmaterial,theHallresistivityisdescribedbytheempiricalexpression:xy=RsM+RoBRoistermedtheordinaryHallcoecientandiscausedbytheLorentzforce.RsistheanomalousHallcoecientthatresultsfromanitesamplemagnetization.Whiletheordinarytermiscausedbyclassicalphysics,theanomaloustermderivesfromquantummechanicaleects.Theanomaloustermresultsfromspin-dependentscatteringbetweenchargecarriersandmomentcarryingcenters.Thephysicaloriginofthisinteractionisthespin-orbitcouplingofthechargecarrierasitpassesbyamagneticimpurity.Thestrengthofthistermisgivenbythestrengthofthespin-orbitcouplingandtherelativedensityofspin-upandspin-downelectrons.Magnetoresistancecanprovidekeyinsightsintothetransportpropertiesofmagneticsemiconductors,includingthepotentiallandscapeofimpurities,latticedisorder,andelectron-electroninteractions[ 78 ].MagnetoresistancewastypicallymeasuredsimultaneouslywiththeHalldata.3.5.3ElectronDispersiveSpectroscopyEDSEDSwasusedtodeterminechemicalcompositions.TheEDSanalysiswasperformedinsideaJEOL6335Feld-emissionscanningmicroscopettedwithaliquidnitrogencooledEDSdetector.EDSmeasurestheemittedx-rayspectrumofamaterialwhenbombardedbyabeamofenergeticelectrons.Thehighenergyelectronsarefocusedontoasamplewheretheytransfertheirkineticenergyintothelatticethroughaseriesofcollisionalevents.Asmallfractionoftheseeventsarecapableofionizinganatombyejectinganinnershellelectron.Theatomthenrevertsbacktoitsgroundstatebyllingthevacancywithanelectronfromahigherenergyorbital,andeitheraphotonorAugerelectronisemitted. 42

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Theemittedphotonshaveacharacteristicenergydependingontheatomthatemitsit.Therefore,elementalinformationcanbedetermined.TherelativeamountofeachelementcanbecalculatedbycomparingthepeakheightsandapplyingthecorrespondingZAFcorrections,whereZistheatomicnumber,Aisabsorption,andFisthex-rayuorescence.3.5.4OpticalAbsorptionInformationrelatedtotheelectronicbandstructureofsemiconductorscanbeinferredfromtheirinteractionwithlight.Theabsorptionofanopticalphotonnecessarilyinvolvesthetransitionofanelectronfromanoccupiedenergystatetoavacantstateofhigherenergy.Energyandmomentummustbeconservedintheprocess.Considerasemiconductorwithalledvalencebandandemptystatesatthebottomoftheconductionband.Ifthephotonenergyisbelowthelowestallowedelectronictransition,inthiscasetheband-gapenergy,thephotonisnotabsorbedandpassesthroughthematerial;thematerialistransparenttotheparticularphotonenergy.If,however,thephotonenergyisgreaterthanthegapenergy,thephotonisabsorbedbyanelectronictransitionfromthevalencetotheconductionband.Therefore,theonsetofopticalabsorptionishEg.Thesemiconductoractsasalow-passopticallter.Theabsorptioncoecient,,fordirectinterbandtransitionsisgivenbytherelation[ 79 ]:h=Aoh-Eg1 2where,histhephotonenergy,Egistheband-gap,andAoisaparameterassociatedwiththetransitionprobabilityandrefractiveindex.iscalculatedfromtheabsorbancedatausing=2:3A t,whereAistheabsorbanceandtisthelmthickness.Ifaplotofh2vs.hrevealsastraightline,thesamplehasadirectgapabsorption.Theband-gapisdeterminedbyextrapolatingthelinearportiontozero.Transitionswherethewavevectorisnotconserved,k6=0,arealsopossible.Theseareindirecttransitionsacrossthegap.Inordertoconservethetotalmomentum,these 43

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transitionsrequireexchangingmomentumwiththelatticethroughtheabsorptionorcreationofaphonon.Comparedtodirecttransitions,therequiredelectron-phononinteractionmakesthesetransitionslesslikelytooccur.Theabsorptioncoecientforindirecttransitionsis[ 79 ]:h=A1h-Eg2Besidesband-to-bandabsorptions,othertransitionsarealsopossible,includingimpurity-to-banddiscussedbelow,impurity-to-impurity,excitonic,andintrabandtransitions[ 79 ].Perturbationsinthebandstructureofthematerialcanleadtotheformationofbandtails[ 79 ].Theperturbationsarecausedbyimperfectionsinthecrystallattice,suchasimpurities,structuraldefects,andlatticedisorder.Thebandtailingextendstheconductionandvalencebandstatesintotheenergygap.ElectronictransitionsbetweenthetailscauseanexponentiallyincreasingabsorptioncoecientknownasUrbach'srule:dln/dh=1 kBT.Intheabsorptionspectrum,theseUrbachtailsappearasabroadeningofthelowestenergytransitions.IonizedimpuritieswillinteractwithbandcarriersthroughtheCoulombinteraction.Apositivelychargeddonor,forexample,willattractconductionbandelectronsandrepelvalencebandholescausingalocaldistortioninthebandstructure.Thiseectcanbemoreorlessstrongdependingonhowmanyimpuritiesareclusteredtogetherataparticularspotinthematerial.Additionally,impuritiescanalterthedensityofstates.Whendopedinlargeconcentrations,thediscreteimpuritystatescanevolveintoimpuritybands.Opticaltransitionsbetweenthesebandsarepermissible.Fortheparticularcaseoftransitionmetaloxides,intra-iontransitionsbetweend-states,ord-dtransitions,provideanadditionalavenueforopticalexcitation.Theelectroniccongurationofthemetalionisperturbedbythelocalchemicalenvironment,theso-calledcrystaleldsplitting.Theorbitaldegeneracyofthed-statesisliftedthestatesaresplitintodierentenergylevels.Thesplittingisstronglyinuencedby 44

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coordinationgeometry.Forexample,octahedralsitesresultinacompletelydierentarrangementofenergystatesthan,say,atetrahedralsymmetrywouldproduce.Thesplittingsarecharacteristicofthemetalionandthelattice.Therefore,transitionmetalsdopedintooxidelatticeswillproducecharacteristicabsorptionlinesthatcanbeusedtoidentifythecoordinationandvalencestateofthedopant.Thisbecomesparticularlyusefulathighdopingconcentrationswhichmaybenearthesolidsolubilitylimit.Inthiswork,aPerkin-ElmerLambda800UV/Visdouble-beamspectrometerwasusedforopticalabsorptionmeasurements.Samplesweremountedtorigidopaquepanelsthatcontainasmallhole,3mmindiameter,forlighttopassthrough.Thisensuresthatonlylightpassingthroughthesampleisdetected.Italsokeepsthetransmittedbeamsizeconstantbetweensamples.Thesampleswereheldnormaltotheincidentlight.Absorbancespectraweremeasuredusingunpolarizedlightwithwavelengthsbetween200nmand900nm.Thedatawasusedtodeterminethebandgapasafunctionofdoping.Theopticalsignatureofmagneticdopantswasalsousedtoconrmtheirsolubilityandvalencestateinthehostlattice.3.5.5SQUIDMagneticmeasurementswereperformedbymycollaboratorsinDr.ArtHebard'sresearchgroupintheUniversityofFlorida'sphysicsdepartment.ThemeasurementsweredoneusingaQuantumDesignMPMSSQUIDSuperconductingQuantumInterferenceDevicemagnetometer.Currently,SQUIDsprovidethemostsensitiveresolutionsformagneticeldmeasurements.Magnetizationversuseldloopsweretakenatvarioustemperatures.Ferromagneticmaterialswillproduceanitehysteresisintheseloops.Therefore,loophysteresiswasusedtoverifyferromagnetisminmysamples.Thesamplesweremountedinsideaplasticdrinkingstrawwhichhasnoferromagneticcomponentsandplacedperpendiculartotheappliedeld. 45

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Additionally,FieldCooled/ZeroFieldCooledFC/ZFCmeasurementswereusedtotrackthemagneticresponseasafunctionoftemperature.FCmeasurementswereperformedbymeasuringthemagnetizationasthesampleiscooleddownto10Kinasmallappliedeld.ZFCmeasurementswereperformedbycoolingthesampleto10Kinzeroeldandthenapplyingasmalleldasthesamplewasheatedbacktoroomtemperaturewhilemeasuringthemagnetization.ThediamagneticresponseofthesampleandsubstrateweresubtractedfromtheMvs.Hdata.Thiswasperformedforeachmeasurement.Themagnetwassweptto5Ttorevealthediamagneticandparamagneticresponseofthesampleandsubstrate.Theslopeofthehigheldresponseisthebackgroundmagneticsusceptibility,=M H.Thesusceptibilitymultipliedbytheappliedeldisthensubtractedfromeachdatapoint. 46

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Figure3-1.Schematicofpulsedlaserdepositionsystem. 47

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Figure3-2.VisualizationofBragg'slawforx-raydiraction.Adaptedfrom[ 80 ]. 48

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CHAPTER4PROPERTIESOFMn-DOPEDCu2ODMS4.1IntroductionSemiconductingoxidesoerthepotentialforexploringandunderstandingspin-basedfunctionalityinasemiconductinghostmaterial.Dietl'stheoreticalpredictionsuggeststhatcarrier-mediatedferromagnetismshouldbefavoredforheavily-dopedp-typeZnO.However,thisposesaseriouschallengeexperimentallysinceZnOisnaturallyann-typesemiconductorduetointrinsicdefectsanddiculttodopeevenmoderatelyp-type.Thisposestheopportunityofusingalternativewide-gapsemiconductorsthatarenaturallyp-typetofabricatehighCurietemperatureDMSmaterials.Inthissection,thepropertiesofMn-dopedCu2Oareexplored.Cu2Oisap-type,directwidebandgapsemiconductorthatmayholdinterestinexploringspinbehaviorinanoxideDMS.ActivitiesfocusedonunderstandingMnincorporationduringthin-lmsynthesis,aswellaselectricaltransportandmagneticcharacterization.FerromagnetismisobservedinselectMn-dopedCu2Olms,butappearstobeassociatedwithMn3O4secondaryphases.Inphase-pureMn-dopedCu2Olms,noevidenceforferromagnetismisobservedabovethatattributedtothesubstrate.4.2Cu2O:AWide-bandgapP-typeSemiconductorCu2OCupriteisoneoftheearliestknowsemiconductormaterials.Thedevelopmentofcopper-Cu2Orectiersdatesbacktotheearly1920's,decadesbeforesilicondeviceswoulddominatethesemiconductormarket.Therectierswereeasilyfabricatedbyoxidizingpurecopperathigh-temperatureinsideafurnace,andwereadvantageousoverearlierpointcontact"cat-whiskerrectierssincetheinterfaceremainedfreeofcontaminationandcouldbemadeuniformlyoverlargeareasofcopper[ 81 ].Cu2Oalsohasalargetheoreticalsolarcelleciency%andwasheavilystudiedforphotovoltaicpropertiessincethe1970's;howeverpracticalapplicationofCu2Osolarcellsislimitedduetodicultieswithimprovingthesemiconductor'selectricalproperties[ 82 83 ] 49

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Cu2Oisoneofthefewbinaryp-typesemiconductingoxides.Ithasadirectband-gapof2.0eV.Thestructureiscubicwitha=4.27AFigure4-1.OxygenatomssitonabcclatticeandareencagedbyatetrahedronofCuatoms.EachCuatomistwo-foldcoordinatedresultingintherareoccurrenceofO-Cu-Olinearbonding.EachCu-Obondlengthisabout1.84A.ThestatesatthetopoftheCu2OvalencebandarepredominantlyofCu3dcharacterwithO2pstateslyinglowerintheband[ 84 ].Thedominantmethodforp-typeconductionistypicallyattributedtotheformationofCuvacanciesinthelattice:Cu1+!Cu0+h+UsingbothDFTandDFT+Ucalculations,NolanandElliotthavecalculatedtheformationenergyoftheCuvacancytobeontheorderof0.4-1.7eV[ 84 ].Theyndthevacancycreatesanacceptorstatearound0.2eVabovetheFermilevel,andthattheholesaredelocalizedaroundthevacancy.Experimentally,theacceptorstateshavebeendeterminedtobe0.26eV[ 85 ],0.5eV[ 86 ],0.4eV[ 87 ]abovethevalenceband.ThestatesatthetopoftheCu2OvalencebandarepredominantlyofCu3dcharacterwithO2pstateslyinglowerintheband[ 84 ].OxidestypicallyhavesmallholemobilitiessincetheO2pvalencebandsarelocalized;however,thefully-occupied3d10statesinCu2Oaremobilewhenconvertedtoholes[ 84 ].Reasonableholemobilitiesontheorderof100cm2v)]TJ/F19 7.97 Tf 6.587 0 Td[(1s)]TJ/F19 7.97 Tf 6.586 0 Td[(1havebeenfoundexperimentally[ 82 88 ].Thenumberofreportsexploringthemagneticpropertiesoftransition-metaldopedCu2Oareratherlimited.Thereisalsovariationintheobservedferromagnetismbetweengroups,similartothecontemporarystateofresearchintootherDMSmaterials,likeTM-dopedZnO.WeiandcoworkersfoundbulkpelletsofCu2Odopedwithanominalconcentrationof1.7at%Mntobeferromagneticupto300K[ 89 ].Theyreportamagneticsaturationof0.4B/Mnat10Kwhichdroppedto0.05B/Mnatroomtemperature.Theseresultswerecalculatedusingthe1.7at%nominalconcentration,buttheirEDSresultsindicatedaMnconcentrationof0.3-0.5at%whichwouldresultinahighersaturationupto2.5B/Mnat10K.Thesamegroupalsofoundroomtemperature 50

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ferromagnetismwithasaturationmomentof0.6B/MninCu2O:Mnlmsdepositedbynear-roomtemperatureelectrodeposition[ 90 ].Inbothcases,theCu2OlmsdopedwithMnshowedlowerresistivitythanundopedlmsbyafactorof2.Mn-dopedC2OlmsweredepostiedbyPanetal.usingrfmagnetronsputtering[ 91 ].Thelmswereparamagneticabove5Kandshowedsomeweakferromagnetismnear5Kwithamomentof5.3B/Mn.Kaleetal.usedPLDtostudy5at%cobalt-dopedCu2Olmsthatwerecodopedwith0.5at%Al,V,andZn[ 92 ].FerromagetismwasonlyobservedinlmscodopedwithAl,butpersisteduptoroomtemperatturewithamomentof0.44B/Co.DopingZnorVdidnotshowferromagnetismbutcausedtheresistivitytodecreaseandincrease,respictively.Sincetherewasnocorrelationbetweentheferromagnetismandresistivity,theysuggestedtheferromagnetismmayberealtedtoorbitaldefectsintroducedbytheAlwhichhassandpvalenceorbitals.Incontrast,Antonyetal.didobserveferromagnetismat400Kin5%Co-dopedCu2Onanoclusters[ 93 ].DFTcaluculations,employedbySiebererandcoworkers,showthattheferromagneticpropertiesofCu2OdopedwithCoorMnisdependantuponthepresenceofdefectscopperoroxygenvacancies[ 94 ].Thefounddefect-freeMn-dopedCu2Otobeanitferromagnetic,whilelong-rangeferromagnetismcouldoccurwhendefectsarepresent.ForthecaseofCu2OdopedwithCo,theexchangeconstantsaremostlyferromagneticindefectfreematerialonlynearest-neighborsiteswerefoundtohaveantiferromagneticcouplingwhentheHubbardUisincreased.CuvacancieswerefoundtoincreasetheTc,whileoxygendeciencyintroducedstrongoscillationsinthemagneticexchange.4.3ExperimentalPulsed-laserdepositionwasusedforlmgrowth.Manganese-dopedCuOtargetswerefabricatedusinghigh-purityCuO.995%,withMnO29.999%servingasthedopingagent.Thetargetswerepressedandsinteredat860Cfor12hr,followedby950Cfor2hrinair.TargetswerefabricatedwithacompositionofCu1:9Mn0:1O.AKrFexcimerlaserwasusedastheablationsource.Alaserrepetitionrateof5Hzwasused,witha 51

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targettosubstratedistanceof4cmandalaserpulseenergydensityof1-3J/cm2.Thegrowthchamberexhibitsabasepressureof10)]TJ/F19 7.97 Tf 6.586 0 Td[(6Torr.Filmthicknesswasapproximately300nm.Singlecrystal01MgAl2O4andMgOwereusedasthesubstratematerialinthisstudy.Four-circleX-raydiractionwasusedtodeterminecrystallinityanddopantsolubility.SQUIDmagnetometrywasusedtocharacterizethemagneticpropertiesofthedepositedlms.Inparticular,itwasusedtodeterminethepresenceofferromagnetismandmeasuretheCurietemperature.Photoluminescencewasusedtocharacterizetheopticalpropertiesofthematerial.4.4ResultsandDiscussionTheepitaxialgrowthofCu2OthinlmsisdependentuponachievingoxidationconditionsinwhichtheCuionassumesa1+valence.Figure4-2showsthethermodynamicstabilitycurvesforCu-Cu2O-CuOasafunctionofoxygenpartialpressureandtemperature[ 95 ].Usingthis,wendthatepitaxialCu2Olmscanbegrownfromhigh-purityCuOorCutargetsusingpulsed-laserdepositioninanoxygenambient.Figure4-3showstheX-raydiractiondatafora01Cu2Olmgrownon001MgOwithPO2=4x10)]TJ/F19 7.97 Tf 6.586 0 Td[(4Torr,T=750C,usingaCutarget.Similarresultswereobtainedonperovskitesubstrates,yieldingp-typelmswithacarrierdensityof1015cm3andaroomtemperaturemobilityof26cm2v)]TJ/F19 7.97 Tf 6.586 0 Td[(1s)]TJ/F19 7.97 Tf 6.587 0 Td[(1.TheprimaryfocusofthisworkwastoinvestigatethesynthesisandpropertiesofepitaxialCu2OdopedwithMn.TheselectionofMnasthetransitionmetaldopantisbasedonthebestavailableevidenceindeterminingwhichmagneticimpuritiesarelikelytoyieldferromagnetism.ManganesedopinghasresultedininterestingmagneticphenomenoninII-VIandIII-Vsemiconductors,includingspinglassorantiferromagneticbehaviorforanumberofsystems,andispredictedtoyieldahighCurietemperatureinZnOaspreviouslydiscussed.IntheCu2Ostructure,theCu1+cationistwofoldcoordinatedwithanionicradiusof0.46A.Mn2+doesnotnormallyexhibitatwofoldcoordinationinbulkmaterials.However,theradiusforfourfoldcoordinatedMn2+is0.66A,whichisclosetothatforthefourfoldcoordinatedCu1+.60A.Asa2+cationon 52

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a1+site,Mnwouldbeexpectedtocompensateforthep-typeconductivityobservedinCu2O.TheresistivityofMn-dopedCu2Olmswassignicantlyhigherthatthatforundopedlms.ThephasestabilityandsolidsolubilityofMninCu2Owasinvestigatedasafunctionofdepositiontemperatureandoxygenpressure.X-raydiractionwasusedtodetermineconditionsthatlimitsegregationofsecondaryphases.Filmgrowthwascarriedoutoveratemperaturerangeof300-700Candanoxygenpressureof10)]TJ/F19 7.97 Tf 6.586 0 Td[(3,10)]TJ/F19 7.97 Tf 6.587 0 Td[(4Torr,orinvacuum.Thebasepressureof10)]TJ/F19 7.97 Tf 6.586 0 Td[(5Torrconsistsmostlyofwatervapor.Figures4-4to4-6showtheX-raydiractiondataforMn-dopedlmsgrownundertheseconditions.Severalitemsshouldbenoted.First,theCu2Ophaseisdominantovermostofthisrange.ThisissurprisingasthethermodynamicstabilitybehaviorsuggeststhatCuOshouldbethestablephaseforT<600C,PO2=10)]TJ/F19 7.97 Tf 6.587 0 Td[(3Torr,andT<550C,PO2=10)]TJ/F19 7.97 Tf 6.586 0 Td[(4Torr.Twopossibilitiesexistinexplainingthisdiscrepancy.First,thedopingofCu2OwithMnmayshiftthephasestabilityline.Second,epitaxyatlowtemperaturemaybesucienttostabilizetheCu2Ophase.AsegregationofMnoxidesintheMn-dopedlmswasalsoexaminedforthevariouslm-growthconditions.TheX-raydiractionresultsindicatethepresenceofantiferromagneticMn2O3asanimpurityCu2OphaseforT500C.Forlmsgrowninvacuum,aweakpeakthatcouldbeassociatedwitheitherMn2O3orferromagneticMn3O4isobserved.Phase-pureCu2OlmswereobtainedatT400C,indicatingthemetastableincorporationofMnintheCu2Omatrix.Figure4-7showsthephaseassemblageasafunctionofgrowthconditions.ThemagneticpropertiesofMn-dopedsamplesweremeasuredusingaQuantumDesignSQUIDmagnetometer.Measurementsweremadeonlmsgrownatlowtemperatures,inwhichnoMn3O4impurityphasepeakswereevidentintheX-raydiractionpatterns,aswellaslmsgrownatelevatedtemperatures.InordertocharacterizethemagneticpropertiesoftheMn-dopedsamples,eld-cooledandzeroeld-cooledmagnetizationmeasurementswereperformedfrom4.2to300K.BytakingthedierenceMbetween 53

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thesetwoquantities,thepara-anddiamagneticcontributionstothemagnetizationcanbesubtracted,leavingonlyameasureofferromagneticbehavior.Figure4-8showstheMdierenceasafunctionoftemperatureforaMn-dopedsamplegrownat300Cin1mTorrofO2.Atlowtemperature,asmall,butniteeld-cooledminuszeroeld-cooledmagneticsignalpersistsupto250Kasseeningure4-8.However,themagnetizationsignalissmall,andcanbeattributedtoabackgroundmagneticsignalfromtheMgAl2O4substrate.Figure4-9showsthetemperaturedependentMandMvs.HbehaviorforthesubstratematerialwithoutaCu2Olm.Forthisandotherphase-puresamples,noferromagnetismcouldbedetectedabovethatattributedtothesubstrate.WealsoinvestigatedthemagneticpropertiesofMn-dopedCu2Ogrownat700C.TheseepitaxialMn-dopedCu2OlmsareclearlyferromagneticwithaTcof48Kasseeningure4-10.Akeyrequirementinunderstandingferromagnetismintransitionmetaldopedsemiconductorsistodelineatewhetherthemagnetismoriginatesfromsubstitutionaldopantsoncationsites,orfromtheformationofasecondaryphasethatisferromagnetic.Theimportanceofthisissuecannotbeunderstated.Theconceptofspintronicsbasedonferromagneticsemiconductorsassumesthatthespinpolarizationexistsinthedistributionofsemiconductorcarriers.Localizedmagneticprecipitatesmightbeofinterestinnanomagnetics,butisoflittleutilityforsemiconductor-basedspintronics.Thequestionofprecipitatesvs.carrier-mediatedferromagnetismiscomplex,andisacentraltopicofdiscussionforothersemiconductingoxidesthatexhibitferromagnetism,inparticulartheCo-dopedTiO2system[ 96 97 ].Severalissuesmustbeaddressedinordertogaininsightintothepossibleroleofsecondaryphaseprecipitatesinthemagneticpropertiesoftransitionmetaldopedsemiconductors,specicallyforCu2Olms.First,oneshouldidentifyallcandidatemagneticphasespossiblefromtheassemblageofelements.ThecoincidenceofTcwithaknowncandidatesecondaryferromagneticphaseindicatesalikelysourceofatleastpartofthemagneticsignature.Forthepresentmaterial,metallicMnisantiferromagnetic,withaNeeltemperatureof100K.Inaddition,nearlyallof 54

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thepossibleMn-basedbinaryandternaryoxidecandidatesareantiferromagnetic.TheexceptiontothisisMn3O4,whichisferromagneticwithaCurietemperatureof46K[ 98 99 ].X-raydiractionmeasurementsonthesampleconsideredingure4-10showevidencefortheMn3O4phase.Obviously,thesimplestexplanationforferromagneticbehaviorinthismaterialisMn3O4precipitates.Inadditiontomagnetization,theopticalpropertieswerealsoexamined.ThephotoluminescencepropertiesofthelmsweremeasuredatroomtemperatureusingaHe-Cdlaser5nmandtakenoverawavelengthrangeof350-800nm.Thepowerdensitywas1W/cm2.A0.3nmscanninggratingmonochromatorwithaPeltier-cooledGaAsphotomultiplierwasutilized.Theplotingure4-11showsphotoluminescencespectraforMn-dopedCu2Olmsdepositedat400and700C.Thepeakat610nmcorrespondstothe1sexciton[ 100 { 102 ].Thispeakisratherweakandbroad,butismostevidentinthelmgrownat700C.Notealsothepeakat735nm,whichhaspreviouslybeenassociatedwithextrinsicdefectsintheCu2Omaterial[ 103 ].Mostoftheadditionalbroadpeakscouldbeattributedtothebackgroundphotoluminescencefromthesubstrate.Theemergenceofintrinsicphotoluminescenceinthelmgrownat700CisconsistentwiththesegregationofMnfromtheCu2Omatrix.ItisalsopossiblethatluminescencefromeitherMn2+orMn4+alsocontributestothespectraobserved.Withtheadventofnewequipmentavailableinthelabformeasurements,resistivityandHalldatawerecollectedonlmsthatwereaged3{4years.Therewasnovisualdecompositionofthelms,butadetailedstudyofthemicrostructurewasnotconducted.MeasurementswereperformedinVanderPauwcongurationusingavaryingmagneticeldwithamaximumstrengthof1T.Solderedindiumcontactswereplacedonthecornersofthesample.Figure4-12showsthecollectionofmeasurementsonvarioussamples.Thelmsappeartobecomeslightlymoreresistantwithhighergrowthtemperature,butthetrendissmallandmaybeinsignicant.TheHallmobilityincreaseswithhighergrowthtemperature,buttheredoesnotseemtobeasignicantrelationshipwithgrowthpressure. 55

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Thismaybeanindicationofbettercrystallinequalitywithhighergrowthtemperatures,whichleadstoenhancedmobilityofthecarriers.Whilethemobilityincreases,thecarrierconcentrationdropsowithhighertemperature.Itisalsoimportanttorememberthatdierentsecondaryphasesareevolvinginthemicrostructureasthegrowthconditionsarevariedwhichpresumablyhaveanimpactontheobservedelectricalproperties.TemperaturedependentHalleectdatawasstudiedonanewlydepositedlmgrownat400Cin1mTorrO2onaMgAl2O4substrate.HalldatawastakeninaVanderPauwcongurationwithfour80nmthickplatinumcontactssputteredonthecornersofthesample.Ptwaschosenforit'shighbarrierheight5.6eVtomakeanohmiccontacttothep-typelm.Thesamplewasmeasuredinatemperaturerangefrom200Kto400K.Theroomtemperatureresistivitywas885ohm-cm,butquicklybecametooresistive>106ohm-cmtomeasureaccuratelybelow200K.TheacceptorstateactivationenergycanbecalculatedassuminganArrhenius-typeactivationoftheform:log=logA+)]TJ/F21 7.97 Tf 6.587 0 Td[(Ea 2:303kT,whereistheresistivityinohm-cm,Aisaconstant,Eaistheactivationenergy,kisBoltzmann'sconstant,andTistemperatureindegreesKelvin.Theslopewascalculatedbyalinearleast-squarestthroughthedatainsetofgure4-13.Eawasdeterminedtobe0.25eV.Thisvalueissimilartothevalue.2eVcalculatedforacceptorstatesarisingfromCuvacanciesusingrst-principlesDFTandDFT+Utheories[ 84 ].Itisalsoconsistentwiththe0.22-0.25eVvaluesfoundforCu2OlmsdopedwithNications[ 104 ].Thetemperaturedependentresistivity,mobility,andcarrierconcentrationisgiveningure4-13withthelogvs.1/Tactivationt.Thevalueoftheresistivityandhallresistivityat300Kasafunctionofeldisgiveningure4-14.Thepositiveslopeofthehallresistivityclearlyindicatesp-typebehaviorandislinearthroughouttheeldsweep.NoindicationofanomalousHalleectisobserved.TheresistivityhasaslightnegativeMR-0.4%at7Tesla.Thelmbecametooresistivetomeasureatlowertemperature. 56

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Figure4-1.CubicunitcellofCu2O.OxygenatomsoccupybcclatticepositionsthataresurroundedbyatetrahedronofCuatoms.Cuatomarethereforelinearlycoordinatedtotwooxygenatoms.ThisformsarareoccurrenceoflinealO-Cu-Obonding. 57

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Figure4-2.PhaseStabilitycurvesfortheCu-Cu2O-CuOsystem. Figure4-3.X-raydiractiondataforepitaxialCu2Oon01MgO,showingbothaout-of-planeandbin-planeorientation. 58

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Figure4-4.X-raydiractiondataforMn-dopedCu2Olmsgrownon01MgAl2O4inanoxygenpressureof1mTorr. 59

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Figure4-5.X-raydiractiondataforMn-dopedCu2Olmsgrownon01MgAl2O4inanoxygenpressureof0.1mTorr. 60

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Figure4-6.X-raydiractiondataforMn-dopedCu2Olmsgrownon01MgAl2O4invacuum. 61

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Figure4-7.Phaseassemblageforlmsgrownunderdierentconditions. 62

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Figure4-8.MagneticbehaviorforanepitaxialMn-dopedCu2Olmgrownat300Cand1mTorrofoxygen. 63

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Figure4-9.MagneticbehaviorforMgAl2O4substrate. 64

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Figure4-10.MagneticbehaviorforanepitaxialMn-dopedCu2Olmgrownat700Cinvacuum. 65

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Figure4-11.LowtemperaturephotoluminescencespectraforMn-dopedCu2Olms. 66

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Figure4-12.Transportdatafor1%Mn-dopedCu2Olms.Theresistiviya,mobilityb,andholeconcentrationcareplottedasafunctionofgrowthtemperaturefordierentoxygenpressures. 67

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Figure4-13.Temperature-dependenttransportdatafor1%Mn-dopedCu2Olm.aTheresistiviyvs.temperature.Theinsetshowsresistivityvs.1/Twithalinearttocalculatetheactivationenergyofacceptorstatesabovethevalenceband.bThemobility[hollowcircles]andtheholeconcentration[lledtriangles]. 68

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Figure4-14.Field-varyingtransportmeasurementsfora1%Mn-dopedCu2Olm.aHallresistivity.DottedlineisalineartusedtocalculatetheHallcoecient.Theinsetshowsthederivativeofthedatatoemphasizethereisnocurvaturefromtheanomaloushalleect.bMagnetoresistance. 69

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CHAPTER5PROPERTIESOFZnOCODOPEDWITHMnANDSn5.1IntroductionAsdiscussedinChapter2,severalrecenttheoriesregardingtheoriginofferromagnetisminZnODMSemphasizetheimportanceofholesinmediatingtheexchangeinteractionbetweendopedMnatoms.Dietl'smean-eldcalculationspredictthatroomtemperatureferromagnetismispossibleinMn-dopedZnOthatisheavilydopedwithholes,whilecarrier-mediatedferromagnetisminn-typematerialshouldbelimitedtolowertemperatures.TherecentworkbyKittilstvedandcoworkersdemonstratethattheMn+3chargetransferenergyliesclosetothevalenceband,similarinenergytoZnOacceptorstates.Itisthoughtthiscanleadtolargehybridizationnecessarytosupportferromagneticordering.Theadvancementofspintronicsasatechnologydependsuponthedevelopmentandunderstandingofsemiconductorsthatcansupportspin-polarizedcarrieroperationatoraboveroomtemperature.Inthischapter,thesynthesisandmagneticpropertiesofMn-dopedZnOepitaxiallmscodopedwithSnareexamined.Codopingallowsindependentcontroloverthemagneticandelectronicpropertiesofthematerialbydopingforeachseparately.InII-VImaterials,Mn+2isisovalentanddoesnotintroducecarriers.BycodopingII-VIsemiconductors,Mnprovidesthelocalizedspinswhileanadditionaldopantcanbeusedtocontrolthecarrierconcentration.ThisprovidesaconvenientplatformtostudytheeectsofcarrierconcentrationontheobservedmagneticpropertiesinZnODMS.AsagroupIVcation,Sncanexistineitherthe4+or2+valence,althoughthe4+valenceismostcommon.Assuch,itcanserveeitherasadoublyionizeddonororisoelectronicimpurity.FortheZnOlmsdepositedinthiswork,Snbehavesasadonor.Themagnetizationdependenceonthecarrierdensityisinvestigated.Superconductingquantuminterferencedevicemagnetometrymeasurementsindicatethatthelmsareferromagneticwithaninversecorrelationbetweenmagnetizationand 70

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electrondensityascontrolledbySndoping.Magnetisminlowfree-electrondensitymaterialisconsistentwiththeboundmagneticpolaronmodel,inwhichboundacceptorsmediatetheferromagneticordering.Increasingtheelectrondensitydecreasestheacceptorconcentration,thusquenchingtheferromagneticexchange.Thisresultisimportantinunderstandingferromagnetismintransition-metal-dopedsemiconductorsforspintronicdevices.5.2ExperimentalEpitaxialMn,Sn-dopedZnOlmsweregrownbyconventionalpulsed-laserdeposition.Laserablationtargetswerepreparedfromhigh-puritypowdersofZnO9.999%,withMnO29.999%andSnO29.95%servingasthedopingagents.Thepressedtargetsweresinteredat1000Cfor12hinair.Thetargetswerefabricatedwithanominalcompositionof3at.%Mnand0,0.1,0.01,and0.001at.%Sn.ALambdaPhysikKrFexcimerlaserwasusedastheablationsource.Thelaserenergydensitywas1-3J/cm2withalaserrepetitionrateof1Hzandtarget-to-substratedistanceof6cm.Thegrowthchamberexhibitsabasepressureof10)]TJ/F19 7.97 Tf 6.587 0 Td[(5Torr.Filmsweredepositedontosingle-crystal,c-planeorientedsapphiresubstrates.Filmgrowthwasconductedoveratemperaturerangeof400-600C.Anoxygenpressureof20mTorrwasusedforalllmgrowthexperiments.Filmthicknesseswereapproximately300-400nm.X-raydiractionwasusedtodeterminethecrystallinityandsecondaryphaseformation.SuperconductingquantuminterferencedeviceSQUIDmagnetometrywasusedtocharacterizetheferromagneticbehaviorofthedopedlms,focusingonthelmsgrownat400C.5.3ResultsandDiscussionThephasestabilityandsolidsolubilityofMnintheZnOlmswereinvestigatedasafunctionofgrowthtemperatureforlmswithvaryingSnconcentrations.Figure5-1showsthex-raydiractionscansforlmsdepositedunderthegivengrowthconditions.Inallcases,thedominantlmpeakscorrespondtoc-axisperpendicularZnO.Notethat, 71

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forsomeofthelms,apeaklocatedaround64.55appears.Atrstglance,thispeakwasassignedtothe0reectionofMn3O4.The00diractionpeakofMn3O4hasa2valueof64.65whichcloselymatchestheobservedpeak.However,theformationofMn3O4at400Candnotathighertemperaturesin0.1%Sndopedsamplesispeculiarandshouldbequestioned.NotethatpreviousreportsfromFukumuraindicatedthatepitaxialZnOlmswithaMnconcentrationashighas35%couldbeachievedwhilemaintainingthewurtzitestructureusingpulsed-laserdeposition[ 37 ].Theobservedpeaksaresmallafewhundredcountsabovebackgroundandrathersharp.AsimilarpeakisobservedinundopedZnOlms,whichisincludedinFigure5-1.Therefore,itisbelievedthispeakisassociatedwiththeZnOandrepresentstheKartifactfromtheZnO04peak.TheXRDusedintheseexperimentshasaNilterforattenuatingCuKradiation,buttheKlineisclearlypenetratingbecausetheKpeakfromtheZnO002peakispresentat31.TheKpeakscanbecheckedbytakingthed-spacingoftheZnO02planesandcalculatingtherespective2valueforKradiation=1.3922AusingtheBraggequationnl=2dsin.The2valuefortheZnO04Kpeakisdeterminedtobearound64.6,whichiscommensuratewiththeobservedpeak.NotethatevenifthepeakdoesrepresenttheformationofMn3O4inthelm,thephasewouldnotcontributetoahightemperatureferromagneticsignalsincetheCurietemperatureisonly50Kasmentionedpreviously.AlsonoticethattheprecipitationofSn-containingphasesisnotobservedinthediractionscan,norisitexpectedevenifpresentasthenominalconcentrationofSninthelmsis0.1%.TheepitaxialnatureoftheZnOlmswasdeterminedusingfour-circlehigh-resolutionx-raydiraction.Figure5-2ashowsan!rockingcurveabouttheZnO002peakforthelmgrownonc-planesapphiresubstrateatagrowthtemperatureof500CandSnconcentrationof0.1%.A1 2divergenceslitwasplacedoverthex-raysourceanda1mmreceivingslitwasplacedinfrontofthedetector.TheZnO02rockingcurvedisplaysafullwidthathalfmaximumFWHMof0.47.Thein-planealignmentisevidentin 72

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thephiscanandpolegureoftheZnO01planeshowningure5-2b.The60peakintervalsareconsistentwiththehexagonalsymmetryoftheepitaxialZnOwurtzitestructure.TheMnvalencestateintheZnOlatticewasinvestigatedusingx-rayphotoemissionspectroscopyXPS.Figure5-3showsthecorelevelXPSspectraforaZnOlmdopedwith3%Mnand0.01%Sn.ThedatawaschargecorrectedbyshiftingtheO1speakto530.1eV.ThelmwassputteredwithAr+for4min.toremovesurfacecontamination.TheMn2p3=2bindingenergyis640.7eV.ThisisconsistentwithvaluesassignedtoMn+2inZnO[ 105 106 ].Thereisa6eVenergydierencebetweentheMn2p3=2anditshigherbindingenergysatellitepeak.ThebindingenergyandsatellitepeakareconsistenttothereportedforsinglecrystalMnObyLangelletal.[ 107 ].Theyfoundthesatellitestructuretobeparticularlysensitivetotheoxidestoichiometry.InthecaseofMn2O3theMn+3valencethebindingenergywashigherat641.1eVandthesatellitestructuretendedtodecreaseforthehigheroxidephases[ 107 ].ThebindingenergiesformetallicMnandMn+4sitclosetotheMn+2value,buttheenergyforMnhasbeenseenat637.7eVandthatforMn+4at642.4eVinZnO[ 105 ].ThebindingenergyandsatellitestructureforthemeasuredlmsuggeststhatmostoftheMndopedintotheZnOisinthe+2valencestate.Theroom-temperatureresistivityfortheMn-dopedZnOlmswithvaryingconcentrationsofSnwasdeterminedusingafour-pointvanderPauwgeometry.DefectchemistrycalculationsforMn-dopedZnOindicatethatMn+2formsadonorlevel2.0eVbelowtheconduction-bandedge[ 108 ].ThisdeepdonorstatewithMnsubstitutionontheZnsiteinZnOhasnodirecteectontheelectronconcentrationatroomtemperature.However,defectchemistrycalculationsalsoindicateareductioninZninterstitialswithMndoping.ZninterstitialsaregenerallyacceptedastheprimaryshallowdonordefectsinnominallyundopedZnO.ThiswillyieldanincreaseinresistivityforMn-dopedlmsascomparedtoundopedmaterial[ 108 { 110 ].TheMn-dopedZnOlmswithnoSnexhibita 73

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resistivityontheorderof102-cmwithacarrierdensityofmid-1016/cm3.Thiscarrierdensityissubstantiallylowerthanthatseenforundopedepitaxiallmsandisconsistentwiththereductionofshallowdonors.LimitedresultsonthedopingbehaviorofSninZnOindicatethatitintroducesadonorstate[ 111 { 115 ],althoughinsomeII-VIcompoundsemiconductors,Snisanamphotericdopant,substitutingoneithertheIIorVIsite[ 116 117 ].ForZnO,theexpectationisthatSnwillsubstituteontheZnsiteduetoaclosematchinionicradiibetweenZn+2.074nmandSn+4.069nm.Fortheepitaxiallmsconsideredinthiswork,Snbehavesasadonor.TheresistivityofZnO:MnlmswithvariousSncontentisshowninTable5-1.TheresistivityofthelmsdropsrapidlywithSndopingwithaminimumof0.185cmforaSnconcentrationof0.1%.ThemostcommonvalencestateofSnis+4,yieldingadoublyionizeddonorifdopedsubstitutionallyontheZnsite.Hallmeasurementsindicatethatthelmsareincreasinglyn-typewithSndopingupto0.1at.%.ItshouldalsobenotedthatotherworkhasshownthattheadditionofSntoZnOceramicsalsoyieldsanenhancementincrystallinity[ 111 112 ].ThemagneticpropertiesofthelmsweremeasuredusingaQuantumDesignSQUIDmagnetometer.Thediamagneticresponsesofthesubstrateandhostsemiconductorweresubtractedfromthemagnetizationplots.TheprimaryfocusofthemeasurementswastodeterminehowthemagneticpropertiesofthelmschangedasafunctionofelectrondensityascontrolledbySnconcentration.SamplesthatshowedminimalamountsofpossibleMn3O4precipitationviax-raydiractionwereusedfortheSQUIDmeasurements.AlltheMvs.Hloopsshowhystereticbehaviorwithanitecoercivityandloopclosure.Fromthehysteresiscurves,anincreaseinloopwidthisobservedwithincreasingSnconcentration.Figure5-4showsthecoerciveeldasafunctionofSnconcentration,suggestingdomainpinningastheSndopingisincreased.ItisunclearwhytheadditionofSnenhancesthehystereticmagnetizationresponseintheMn-dopedlms.ItmightindicateenhancedpinningofdomainsduetotheSndopants. 74

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MostinterestingisthesaturationmagnetizationbehaviorasSncontentisincreased.Asnotedearlier,increasingSnconcentrationincreaseselectrondensityandconductivity.Figure5-5showstheroom-temperaturemagnetizationversuseldbehaviorfortheZnOsamplescontaining3%MnandSncontentsof0,0.1,0.01,and0.001%.MagnetizationisgivenasthemagneticmomentperMndopantion.Initially,thereisanincreaseinmagnetizationwithminimalSndoping.However,withincreasingSndoping,thereisaninversecorrelationbetweentheSncontentandsaturationmagnetization.AstheelectrondensityincreaseswithSndoping,themagnetizationdecreases.ThemaximummagnetizationcorrespondstoamagneticmomentperMnionof0.5B/Mn.ThisisconsistentwiththeboundmagneticpolaronmodelinwhichonlyafractionoftheMnionsareexpectedtoorderferromagneticallyduetocompetingsuperexchangeantiferromagneticinteractionsbetweenneighboringMnions[ 118 ].TheinversecorrelationofsaturationmagnetizationwithelectrondensityisinterestingandprovidessomeinsightintothemechanismforferromagnetisminMn-dopedZnO.OverlapoftheMnd-stateswiththevalencebandsuggeststhatholesarenecessaryinordertoinduceferromagneticorder.Forsemi-insulatinglmstoexhibitferromagnetism,theboundmagneticpolaronmodelprovidesamechanismwherebyholesthatarelocalizedatorneartheMnionsareresponsibleformediatingferromagnetism.TheadditionofelectronstothesystemwillmovetheFermienergylevelupinthebandgap,resultinginadecreaseinholedensityandareductioninmagnetization.ThisisconsistentwithKittilstvedandcoworker'sobservationwhereferromagnetismwasinducedwhentheholesfromtheacceptorstateshybridizewiththechargetransferstateMn.FerromagnetismwasobservedwhentheZnOwaslocallydopedp-type,butnoferromagnetismwasobservedwhendopedn-type.Thisappearsconsistentwithearlyworkontrivalent-dopedZn,MnOwherenoferromagnetismwasobservedforheavilyn-typelms.ItmayalsoexplainthediscrepancyfromotherstudiesofMn-dopedZnOlmsinwhichtheintrinsicdefect-mediateddonorstatesarehighindensity.ItisimportanttonotetheneedtomaintainaMnconcentrationlow 75

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enoughtoavoidMnMnantiferromagneticinteractions,whicharelikelytodominatehighMn-dopedZnOlms.Inconclusion,themagneticandtransportpropertiesofMn-dopedZnOthinlmscodopedwithSnwereexamined.ResultsindicatethatthelmsareferromagneticwithaninversecorrelationbetweenmagnetizationandelectrondensityascontrolledbySndoping.Theresultsaremostconsistentwiththeboundmagneticpolaronmodelinwhichboundacceptorsmediatetheferromagneticordering.Increasingtheelectrondensitydecreasestheacceptorconcentration,thusquenchingtheferromagneticexchange.Thisresultisrelevanttounderstandingferromagnetismintransition-metaldopedsemiconductors. Table5-1:ResistivityasafunctionofSncontentincodopedZnO:3%Mnlms. Snconcentration0.0%0.001%0.01%0.1% Resistivity195320170.185-cm 76

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Figure5-1.X-raydiractionofZnOlmscodopedwithMnandSngrownatoxygenpartialpressureof20mTorrandgrowthtemperatureof400,500,and600C.ThediractionpatternforanundopedZnOlmgrownat400Cinvacuumisalsoshown.Thepeakat2=64.55isclearlyevidentintheundopedlmandattributedtotheKartifactfromtheZnO04reection. 77

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Figure5-2.X-raydiractionofanepitaxialZnOlmdopedwith3%Mnand0.1%Snthatwasdepositedat500CandpO2=20mTorr.aan!-rockingcurveoftheZnO2peakwithaFWHMof0.47.bin-plane-scanandpolegureoftheZnO101planes. 78

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Figure5-3.XPSspectraforZnO:3%Mnlmcodopedwith0.01%Sn.aZn2p3/2spectrum,bMn2pspectrum,andcO1sspectrum.ThelmwassputteredinArfor4min.andchargecorrectedtotheO1speak. 79

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Figure5-4.PlotshowingthedependenceofthecoerciveeldonSnconcentrationatdierentSQUIDmeasurementtemperatures. 80

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Figure5-5.Magnetizationmeasuredat300KforepitaxialZnO:3%Mnlmsthatarecodopedwith0.001%Sn,0.01%Sn,0.1%Sn,andnoSn.ThereappearstobeaninversecorrelationoftheSncontentwiththesaturationmagnetization. 81

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CHAPTER6PROPERTIESOFZnOCODOPEDWITHMnANDP6.1IntroductionInthepreviouschapter,themagnetizationdependenceoncarrierconcentrationwasinvestigatedusingSnasadonordopant.Themagnetizationhadaninversecorrelationwithelectrondensity,suggestingthathighermagnetizationwasfavoredastheFermilevelmoveddowninthebandgap.Ifthemagnetizationdecreaseswithelectronconcentration,itmightbeexpectedthatdopingwithholeswillincreasethemagnetization.ThischapterinvestigatesthemagneticpropertiesofMn-dopedZnOcodopedwithP,aknownp-typedopantinZnO.SuperconductingQuantumInterferenceDeviceSQUIDmagnetometrymeasurementsindicatethatthelmsareferromagneticwithaninversecorrelationbetweenmagnetizationandelectrondensityascontrolledbyPdoping.Inparticular,underconditionswheretheacceptordopantsareactivated,leadingtoadecreaseinfreeelectrondensity,magnetizationisenhanced.Theresultisconsistentwithhole-mediatedferromagnetisminMn-dopedZnO,inwhichboundacceptorsmediatetheferromagneticordering.Increasingtheelectrondensitydecreasestheacceptorconcentration,thusquenchingtheferromagneticexchange.Thisresultisimportantinunderstandingferromagnetismintransitionmetaldopedsemiconductorsforspintronicdevices.6.2ExperimentalEpitaxialMn,PdopedZnOlmsweregrownbyconventionalpulsed-laserdeposition.LaserablationtargetswerepreparedfromhighpuritypowdersofZnO9.999%,withMnO29.999%andP2O59.95%servingasthedopingagents.Thepressedtargetsweresinteredat1000Cfor12hrinair.Thetargetswerefabricatedwithanominalcompositionof3at.%Mnand2at.%P.ALambdaPhysikKrFexcimerlaserwasusedastheablationsource.Thelaserenergydensitywas1-3J/cm2withalaserrepetitionrateof1Hzandtarget-to-substratedistanceof6cm.Thegrowthchamberexhibitsabase 82

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pressureof10)]TJ/F19 7.97 Tf 6.587 0 Td[(5Torr.Filmsweredepositedontosinglecrystal,c-planeorientedsapphiresubstrates.Thegrowthtemperaturewas400C.Anoxygenpressureof20mTorrwasusedforalllmgrowthexperiments.Filmthicknesseswereapproximately300to400nm.X-raydiractionwasusedtodeterminethecrystallinityandsecondaryphaseformation.SQUIDmagnetometrywasusedtocharacterizetheferromagneticbehaviorofthedopedlms.TheimpurityconcentrationsweremeasuredbyEDSanddeterminedtobewithin10%ofthenominalconcentrations.6.3ResultsandDiscussionThephasestabilityandsolidsolubilityofMnintheZnOlmswereinvestigatedbeforeandafterannealingforlmswithPcodoping.Figure6-1showsthex-raydiractionscansforlmsdepositedunderthegivengrowthconditions.Inallcases,thedominantlmpeakscorrespondtoc-axisperpendicularZnO.Notethat,fortheselms,segregationoftheMn3O4phaseisnotobservedinthediractiondata.Asmentionedearlier,theMn3O4phaseisferromagneticwithaCurietemperaturelessthan50K.PreviousreportsfromFukumuraetal.indicatedthatepitaxialZnOlmswithaMnconcentrationashighas35%couldbeachievedwhilemaintainingthewurtzitestructureusingpulsedlaserdeposition[ 37 ].Uponannealing,ashiftinthed-spacingfortheZnOisobserved.ThismayindicateasegregationofeitherPorMninthelmswiththermalprocessing.Figure6-2showshigh-resolution!-rockingcurvestakenaroundtheZnO002onaseparatebutsimilarlygrownlm.Thesimilarshiftind-spacingappearsafterannealing.Thescansweretakenusinga1 2divergenceslitoverthex-raysourceanda1mmbrassslitoverthedetectoroptics.TheFWHMbeforeandafterannealingbarelychanges,suggestingthecrystallinitydoesnotchangesucientlywiththeanneal.Furthermore,in-planealignmentismeasuredwith-scansthatshowthesix-foldsymmetryoftheZnOwurtzitestructureandconrmepitaxialregistrywiththesubstrate.Thesurfaceroughnessofthelmslightlydecreasesasmeasuredbytapping-modeAtomicforcemicroscopy 83

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AFMFigure6-3.Thegrainstructurehasalsonotablycoalescedandcoarsenedintolargergrains.DefectchemistrycalculationsforMn-dopedZnOindicatethatMn+2formsadonorlevel2.0eVbelowtheconductionbandedge[ 108 ].ThisdeepdonorstatewithMnsubstitutionontheZnsiteinZnOhasnodirecteectontheelectronconcentrationatroomtemperature.However,defectchemistrycalculationsalsoindicateareductioninZninterstitialswithMndoping.ZninterstitialsaregenerallyacceptedastheprimaryshallowdonordefectsinnominallyundopedZnO.ThiswillyieldanincreaseinresistivityforMn-dopedlmsascomparedtoundopedmaterial[ 108 { 110 ].TheMn-dopedZnOlmswithnoPexhibitaresistivityontheorderof102-cmwithacarrierdensityofmid-1016/cm3.Thiscarrierdensityissubstantiallylowerthanthatseenforundopedepitaxiallms,andisconsistentwiththereductionofshallowdonors.LimitedresultsonthedopingbehaviorofPinZnOindicatethatitintroducesadonorstateforas-depositedlms,anacceptorstatewhenannealed.ThebehaviorofphosphorusinZnOepitaxiallmsbothas-depositedanduponannealinghasbeenreportedindetailelsewhere[ 119 ].Fortheas-depositedlms,theinclusionofphosphorusyieldsasignicantincreaseinelectrondensity,resultinginZnOthatishighlyconductiveandn-type.Theshallowdonorbehaviorintheas-depositedlmsisinconsistentwithPsubstitutionontheOsite,andpresumablyoriginatesfromeithersubstitutionontheZnsiteortheformationofaphosphorus-bearingcomplex.Previousworkhasshownthatthedefect-relatedcarrierdensityinnominallyundopedZnOcanbereducedviahightemperatureannealinginoxygenorair.Inthecaseofundopedmaterial,thereductionindonordensityispresumedduetoeitherareductioninoxygenvacancies,Zninterstitials,orperhapsout-diusionofhydrogenthatisincorporatedintheZnOlatticeduringsynthesis.Inordertoreduceelectrondensityannealinginoxygencanbeperformed.Figure6-4showstheresistivityoflmsannealedatvarioustemperatures.Notethattheresistivityoftheas-depositedphosphorusdopedlmsissignicantlylower 84

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thanthatforthenominallyundopedlm.Foras-depositedlms,ashallowdonorstatedominatestransport.Asthelmsareannealedatincreasingtemperatures,theresistivityofthephosphorus-dopedlmsincreasesrapidly.Thisisparticularlyevidentforlmssubjectedtoannealingtemperaturesof600Corhigher.Whenannealedat700C,thephosphorusdopedZnOlmsbecomesemi-insulatingwitharesistivityapproaching104-cm.Theconversionoftransportbehaviorfromhighlyconductingtosemi-insultingwithannealingshouldbeattributedtoatleasttwofactors.First,thedefectassociatedwiththeshallowdonorstateinas-depositedlmsappearstoberelativelyunstable.Thiswouldexplaintheincreaseinresistivity,butwouldalonepredictasaturationofresistivityatthevaluegivenbytheundopedmaterial.Thedependenceofpost-annealedresistivityonphosphoruscontentsuggeststhatadeeplevelassociatedwithphosphorusdopantispresent.Thisis,infact,consistentwiththeexpectedresultsthatPsubstitutionontheoxygensiteyieldsadeepacceptor.Figure6-4alsoshowsthecarrierconcentrationforP-dopedlms.ForallofthedatashowninFig.6-4,theHallsignwasnegative,indicatingn-typematerial.BothcarrierdensityandHallmobilitydataforsomeannealedsamplesareabsentintheplots.FromthemeasurementsyieldingunambiguousHallvoltage,boththecarrierdensityandmobilityinphosphorus-dopedlmsareobservedtodecreasewithannealing.Thisisconsistentwithareductionintheshallowdonorstatedensityandactivationofadeepacceptorlevelinthegap.Figure6-5showsthelongitudinalxxandtransverseHallxyresistivitymeasuredforaZnOlmdopedwith3%Mnand2%Patroomtemperature.ThelmwasmeasuredinsideaQuantumDesignPhysicalPropertyMeasurementSystemPPMSmodiedwithanexternalhigh-impedanceHalleectset-upfromKeithleyelectronics.ThelmwasmeasuredintheVanderPauwcongurationwithfourindiumcontactssolderedtothecornersofasquaresample.Theasdepositedlmisn-typewithacarrierconcentrationof2.8x1015cm)]TJ/F19 7.97 Tf 6.586 0 Td[(3andmobilityof0.56cm2v)]TJ/F19 7.97 Tf 6.586 0 Td[(1s)]TJ/F19 7.97 Tf 6.586 0 Td[(1.Thexxinzero-eldis3,975-cm.Thereisasmallnegativemagneto-resistanceMRthatreachesavalueofabout-0.35%at7 85

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Tesla.ThissmallnegativeMRisalsoseeninundopedZnOlmsatroomtemperature.Theresistivityisastrongfunctionoftemperatureandquicklybecomestooresistive,>105-cm,tomeasureatlowertemperatures.Afterannealingthesamplefor60mininatubefurnaceat600in1atmO2,thesampleresistancewasmeasuredtobe>109,whichwastoohightomeasureaccuratelywithpracticalsettlingtimes.Phosphorusisbelievedtocreateacceptorstatesafterbeingthermallyactivatedbyannealing[ 119 ].ThisrequiresthesubstitutionofP)]TJ/F19 7.97 Tf 6.586 0 Td[(3ontoanO)]TJ/F19 7.97 Tf 6.586 0 Td[(2latticesitetoformtheacceptorstate.However,phosphorusexistsinavarietyofvalencestates,including+3,+5,and-3.ToinvestigatethePincorporationintothelm,XPSwasusedtoinvestigatethechargestateofthePion.Figure6-6showstheXPSspectraforaZnO:3%Mn,2%Plmbeforeandafterannealing.Thereisabroadstructurebetween128and132eV,andthenanotherpeakcenteredaround133.6eV.TheP+5valencestateinP2O5hasabindingenergyof135.6eV[ 120 ].ThetheP)]TJ/F19 7.97 Tf 6.587 0 Td[(3ionofZn3P2hasalowerbindingenergyof128.7eV[ 121 ].Notethatpurephosphorushasabindingenergyof130.5eV[ 121 ].Therefore,thehigherbindingenergypeakofFigure6-6islikelyrelatedtothe+5valence.ThebroadstructureatlowerenergycouldbetheexistenceofP)]TJ/F19 7.97 Tf 6.586 0 Td[(3bondedtoZncations.Therefore,thedatasuggestsacoexistenceofphosphoruschargestates.Theredoesnotappeartobeasignicantshiftinthebindingenergiesafterannealing.Beforeannealing,thebindingenergyoftheMn2p3=2peakis641eVandthereisadenitesatellitestructureathigherenergy.ThespectrumisconsistentwithMninthe+2valencestate.Afterannealing,thereisaslightpeakshifttohigherenergyandthesatellitestructuresignicantlydecreases.Langelletal.havereportedsimilarbehaviorforsinglecrystalMnOafterannealinginoxygen,whichwasattributedtotheformationofhigheroxidizedMnphases[ 107 ].It'sprobablethatthereisamixtureofMnvalenceafterannealing.OpticalabsorptionwasmeasuredusingaPerkin-ElmerLambda800UV/Visdual-beamspectrometer.Theabsorbanceofeachsamplewasmeasuredusingunpolarized 86

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lightwithwavelengthsrangingbetween200nmto900nm.Figure6-7displaystheopticaltransmissionfortheZnO:3%Mn,2%Plmbeforeandafterannealing,alongwithanundopedZnOreferencesample.Theundopedsampleshowsasharpexcitonabsorption.Thisabsorptionissmearedoutinthedopedlmsandabroadabsorptionisseenaround3eVinthedopedlms.TherearetwoabsorptionpeaksassignedtoMn2+inZnOat2.95and3.26eVfromthe6A1S!4T2Gand6A1S!4A1,4EGtransitions,respectively[ 122 ].Thesearespin-forbiddend-dtransitionsandshouldbeweak.Therefore,thein-gapabsorptionaround3eVismostlikelyattributedtheMn2+6A1S!4T2Gtransition.DirectinterbandtransitionsfollowtheTaucrelationh=Aoh-Eg1 2,whereistheabsorptioncoecient,hthephotonenergy,andAoisaparameterassociatedwiththetransitionprobablilityandrefractiveindex.h2vs.hisplottedintheinsetofgure6-7.Straightlinesaretthroughthelinearregionsoftheplotstoextractthebandgapofeachsample.ThebandgapwidenswithMn-doping.Boththein-gapabsorptionandblue-shifthavebeenreportedpreviouslyinZn1)]TJ/F21 7.97 Tf 6.587 0 Td[(xMnxOlms[ 37 ].Afterannealing,thebandgapdecreasesslightly.ThiscouldbecausedbyMnsegregatingoutofthelatticeathightemperature,whichwouldbeconsistentwiththepeakshiftintheXRDdataandthedisappearanceofthesatellitepeakinXPS.ThemagneticpropertiesofthelmsweremeasuredusingaQuantumDesignSQUIDmagnetometer.Thediamagneticresponsesofthesubstrateandhostsemiconductorweresubtractedfromthemagnetizationplots.TheprimaryfocusofthemeasurementswastodeterminehowthemagneticpropertiesofthelmschangedasafunctionofelectrondensityascontrolledbyPdoping.SamplesthatshowedminimalamountsofMn3O4precipitationviax-raydiractionwereusedfortheSQUIDmeasurements.Figure6-8showstheroomtemperaturemagnetizationasafunctionofappliedmagneticeldforepitaxialZnO:3%MnlmsbothwithoutandwithPco-doping.FortheMn-dopedlmwithnoP,saturationinthemagnetizationisobserved,butwithlittleevidenceforhysteresisintheMvs.Hcurves.Theas-depositedZnOlmdopedwithbothMnandP 87

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showedareductioninmagnetization/Mnion.ThisisconsistentwiththeproposedmodelsforMn-dopedZnO,whereferromagneticorderingisnotfavoredbyelectrondoping.MostinterestingisthesaturationmagnetizationbehaviorasthePdopedsamplesareannealed.Asnotedearlier,increasingPconcentrationinasdepositedlmsinitiallyincreaseselectrondensityandconductivity.Figure6-8showstheroomtemperaturemagnetizationversuseldbehaviorfortheZnOsamplescontaining3%Mnand2%Pannealedinoxygen.Magnetizationisgivenasthemagneticmoment/Mndopantion.Initially,thereisadecreaseinmagnetizationwithPdoping.However,withannealing,thereisaninversecorrelationbetweentheelectroncarrierconcentrationandsaturationmagnetization.Similarresultsareseenat10KasshowninFigure6-9.Initially,astheelectrondensityincreaseswithPdoping,themagnetizationdecreases.TheinversecorrelationofsaturationmagnetizationwithelectrondensityisinterestingandprovidessomeinsightintothemechanismforferromagnetisminMn-dopedZnO.OverlapoftheMnd-stateswiththevalencebandsuggeststhatholesarenecessaryinordertoinduceferromagneticorder.Forsemi-insulatinglmstoexhibitferromagnetism,theboundmagneticpolaronmodelprovidesamechanismwherebyholesthatarelocalizedatorneartheMnionsareresponsibleformediatingferromagnetism.However,themostimportantobservationisthattheactivationofacceptorstatesforholeformationisnecessaryinordertoachieveferromagnetism.Theholesmaybedelocalized,butwithlowmobility,thusyieldinglowconductivity.Inthiscase,thecarriermediatedmechanismmaysucewithouttheneedofinvokingboundpolaronsasinherenttotheferromagneticordering.Ineithercase,theadditionofelectronstothesystemwillmovetheFermienergylevelupinthebandgap,resultinginadecreaseinholedensityandareductioninmagnetization.ThisappearsconsistentwithearlyworkontrivalentdopedZn,MnOwherenoferromagnetismwasobservedforheavilyn-typelms.ItmayalsoexplainthediscrepancyfromotherstudiesofMn-dopedZnOlmsinwhichtheintrinsicdefect-mediateddonorstatesarehighindensity.Itshouldbenotedthattheamountofmagnetizationinthematerialremains 88

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relativelylowatalltemperaturesasseeninthetemperaturedependentmagnetizationdataingure6-10.ItisimportanttonotetheneedtomaintainaMnconcentrationlowenoughtoavoidMn-Mnantiferromagneticinteractions,whicharelikelytodominatehighMndopedZnOlms.TheresultsofthisstudyareconsistentwithpreviousstudiesonthecarriertypedependenceinCo-andMn-dopedZnOnanocrystallinelms[ 32 49 ].Inthatcase,ferromagnetismwasobservedinMn-dopedZnOnanocrystalsonlywhennitrogen,agroupVacceptordopant,wasintroducedduringthesynthesisprocess.Incontrast,noferromagnetismwasobservedinCo-dopedZnOnanocrystalswhenprocessedwithnitrogen.Basedonthisandotherproperties,theauthorsconcludethatferromagnetisminZnOiscloselytiedtothechargetransferelectronicstructureofthetransitionmetaldopant.ForMn,ferromagnetismisinducedwhentheholesfromtheacceptoriondelocalizeontoMn2+.Again,ourresultsareconsistentwiththisconclusion. 89

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Figure6-1.X-raydiractionofZnOlmscodopedwithMnandPbothbeforeandafterannealing.ThetargetwasZnOdopedwith3%Mnand2%P 90

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Figure6-2.High-resolution!-rockingcurvesonZnOlmscodopedwithMnandPbeforeandafterannealing.ThescansaretakenaroundtheZnO2reection.Theinsetsshow-scansaroundtheindicatedpeaks,showingthelm'sin-planealignmentwiththesubstrate. 91

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Figure6-3.AFMscansonepitaxialZnO:3%Mn,2%Plmsbeforeandafterannealing.Scansweretakenintapping-mode.Thesurfaceroughnessdecreasedandthegrainshavecoalescedintolargergrainswithannealing.AFMimagingsoftwareprovidedby[ 123 ]. 92

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Figure6-4.ResistivityandcarrierconcentrationbehaviorofP-dopedZnOlmsbothas-depositedandannealedinoxygen.Thedatashownisfor1at.%PdopedZnO 93

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Figure6-5.Transportdatafortheas-depositedZnO:3%Mn,2%Plmat300K.aTransverseHallxyandblongitudinalxxresistivity. 94

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Figure6-6.XPSspectraforZnO:3%Mnlmcodopedwith2%P.aO1sspectrum,bZn2p3 2spectrum,cP2pspectrum,dMn2pspectrum.ThedatawaschargecorrectedtotheO1speakat530.1eV. 95

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Figure6-7.RoomtemperatureopticaltransmissionforZnO:3%Mn,2%Plmsbothas-depositedandafterannealingat600Cinoxygen.DataforanundopedZnOlmisalsoincluded.Inset:Taucplotswithlineartstodeterminetheopticalband-gapforeachlm. 96

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Figure6-8.RoomtemperatureSQUIDmeasurementsforepitaxialZnO:3%Mn,2%Plmsbeforeandafteranneal.AlsoshownisaZnO:3%MnlmwithnoP. 97

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Figure6-9.SQUIDmeasurementat10KforepitaxialZnO:3%Mn,2%Plmsbeforeandafterannealing. 98

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Figure6-10.Field-cooledandzeroeld-cooledmagnetizationmeasurementsforaZnO:3%Mn,2%Plmannealedat600CinO2. 99

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CHAPTER7PROPERTIESOFCOBALT-DOPEDZnO7.1IntroductionThetwopreviouschaptersinvestigatedthepropertiesofMn-dopedZnOasapotentialmaterialforspintronicapplications.Inthischapter,cobaltisinvestigatedasamagneticdopantinZnO.Themagneticandmagneto-transportbehaviorofdopedlmswereexamined.Cobaltconcentrationisvariedoverawiderange,from0to30at.%Co.Acombinationofx-raydiraction,opticalabsorption,andtransmissionelectronmicroscopywereusedtoexaminethesolubilityofcobaltintheZnOlatticeandphasesegregationofcobaltmetal.Filmsdepositedat400Cinvacuumwerefoundtobeferromagnetic,whilelmsdepositedinoxygenorathighertemperatureswerefoundtobenonmagnetic.Segregationofcobaltmetaloccursinlmsdopedwith15at.%orgreaterCoconcentrationswhendepositedinvacuum,andtheprecipitatesarefoundtobeorientedwithinthelattice.Thesegregationcanbesuppressedbydepositinginhigherbasepressures,buttheprocessisnotfullyreproducible.PeculiarMRisobservedinthelmsandtheMRchangesasthecarrierconcentrationcrossesthemetal-to-insulatortransition.7.2ExperimentalCobalt-dopedZnOlmsweredepositedviapulsedlaserdepositionPLDontoc-planeorientedsapphiresubstrates.Theablationtargetswerepreparedthroughthesolid-statereactionofmixedoxidepowders.AppropriateamountsofZnOAlfaAesar,Puratronic,99.9995%andCo3O4AlfaAesar,Puratronic,99.9985%powdersweregroundandmixedinmethanol,driedinair,pressedintopellets,andsinteredat1000Cfor12hoursinair.ThetargetsweremixedtogiveproportionsofZn1)]TJ/F21 7.97 Tf 6.587 0 Td[(xCoxOwithx=0.00,0.02,0.05,0.15,and0.30.AKrFexcimerlaser8nmwavelengthwasusedfortargetablationusingarepetitionrateof1Hzandalaserenergydensityof1{3J/cm2.Atemperaturerangeof400C-600Candoxygenpressuresupto2x10)]TJ/F19 7.97 Tf 6.586 0 Td[(5Torrwereusedin 100

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theexperiments.Thevacuumbasepressureofthechamberwas7.0x10)]TJ/F19 7.97 Tf 6.586 0 Td[(6Torr.Filmthicknesseswere200{300nmasmeasuredbymechanicalprolometry.7.3ResultsandDiscussion7.3.1ChemicalCompositionEnergyDispersiveSpectroscopyEDSwasusedtomeasurethepercentageofcobaltinthelms.Figure7-1showsthecobaltconcentrationinthelmasafunctionoftargetcompositionforlmsdepositedunderdierentconditions.Thecobaltconcentrationisgenerallylargerinthelmsthanthepreparedtargets.ThislikelyoccursbecauseZnhasahighervaporpressurethanCo,thuslessZnisincorporatedintothelmduringdeposition.Thecobaltconcentrationisalsohigherwithincreasedsubstratetemperature.7.3.2StructureandPhaseAnalysis7.3.2.1FilmswithprecipitationCrystalstructureandphaseanalysiswerecharacterizedusingX-raydiractionXRDinBragg-Brentanogeometry.Figure7-2showsthe-2x-raydiractionpatternsforaseriesoflmsgrownat400Cinvacuumbasepressure7x10)]TJ/F19 7.97 Tf 6.586 0 Td[(6Torrwithvaryingcobaltconcentration.TheprimarypeakscorrespondtothewurtziteZnO02indicatinggoodtexturewiththec-planeofthesapphiresubstrate.Asthecobaltconcentrationisincreasedabove10%,theappearanceofanewpeakbeginstodeveloparounda2valueof44.4degrees.ThepeakissmallonlyafewhundredcountsabovebackgroundanddoesnotcorrespondtoanyZnOorsubstratepeaks.Boththesmallintensityand2positionmakeidenticationofthepeakdicultusingXRDasthereareseveralcobaltcontainingphaseswithsimilar2valuesofaround44.4degrees,includingthespinelfamilyofcobaltoxidesandcobaltmetal.However,exactdeterminationofthepeakiscriticalsincethepresenceofferromagneticcobaltmetalcouldcontributetothemagneticsignatureofthelms.Thecubicandspinelcobaltoxidesareantiferromagnetic,thoughsomepapersreportthatsmallnanoclusterpowdersofcobaltoxidesareferromagneticduetouncompensated 101

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surfacespins[ 124 125 ].Table7-1isalistofthepossiblephaseswiththeirrespective2diractionvaluesandmagneticcharacter.HighresolutionXRDandTEMwereusedtocharacterizethesecondaryphaseobservedinthepowderXRDscans.Cross-sectionalTEMwasusedtomorepreciselydelineatethenatureandlocationoftheextraphaseasshowningure7-3.Mostoftheprecipitationoccursnearthelm/substrateinterfaceintheformofsmall5nmparticulates.However,formostofthelm,thecobaltdopantappearstoresideintheZnOlatticewithoutprecipitation.ConvergentbeamTEMdiractionpatternsofthelmandnano-precipitatesareshowningure7-4.Theparticlesappeartobeorientedwiththelatticeandhaved-spacingsofd002=0.20nmandd)]TJ/F19 7.97 Tf 6.587 0 Td[(210=0.13nm.Thisisconsistentwithmetalliccobaltandtheparticlesaretentativelyassignedassuch.Thepeakintensitywastoolowtocollectinformationusinghigh-resolutionXRDontheasgrownlm.However,afterannealingthesamplesinhydrogen%H2/Arbalanceat500C,theamountofcobaltincreasesattheexpenseofZnOcausingpartialdegradationofthelm,andhigh-resolutionXRDcouldbeemployedwhichisdiscussedinalatersection.ThechangeinmicrostructureatlowannealingtemperaturesuggeststhatthecobaltthatissubstitutionalintheZnOlatticeisnotstableatmoderatelyhightemperatures.7.3.2.2FilmswithoutprecipitationBymodifyingthegrowthconditions,thesecondaryphasecanbesuppressed.Figure7-5ashowsXRDscansforsamplesgrownat400Catdierentpressureconditions.Thecobaltphaseformsatlowbasepressures.Bydepositinginhigherbasepressures>10)]TJ/F19 7.97 Tf 6.586 0 Td[(5Torroraddingasmallamountofoxygen,theCophasedisappearsfromtheXRDscans.Itshouldbenotedthattheprocessisnotfullyrepeatable.ThecobaltphaseisstillpresentbyXRDforsomedepositionsbutseemslessprobableatthehigherbasepressures.Asaroughestimate,thephaseissuppressedin80{90%ofthesamplesgrowninhigherbasepressures.Itisassumedthepressureatvacuumiscomposedmostlyofwatervapor. 102

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Figure7-5alsoshowsXRDscansforlmsdepositedat500Cand600Cinvacuum.At500CtheCoOphasebeginstoformalongwiththecobaltmetal.At600CthecobaltmetalphasedisappearsandtheCoOphasebecomesmoreprominent.Thermodynamically,CoOisthemoststablephaseatthesetemperaturesandpressure.Thiscanbeseenfromthethermodynamicpredominancediagramforthecobaltoxidesgiveningure7-6.Interestingly,onewouldexpecttheformationofCo3O4atlowergrowthtemperaturesandnotthemetalliccobaltphasedeterminedfromTEM.ThissuggeststhecobaltmetalphaseisstabilizedbytheZnOlatticeandnotbythermodynamicconsiderations.ThisisconsistentwiththewellorientedparticlesobservedinTEM.7.3.3OpticalProperties7.3.3.1OpticalabsorptionEvidenceforCosubstitutionintheZnOlatticecanbeinferredfromopticalabsorptionmeasurements.Figure7-7showstransmissiondataforCo-dopedZnOlmsthatdonotshowprecipitationbyXRD.AnundopedZnOlmisincludedasreference.EachlmwasnormalizedbydividingbyitsmaximumobservedtransmissionT/Tmaxtocomparetheintensityofabsorptionpeaksbetweenlms.Threeadsorptionpeaksareapparentinthedopedlms.Thesepeaksarecharacteristicd-dtransitionlevelsattributedtoCo+2occupyingtetrahedrallatticepositions,andindicatethatcobaltissubstitutingasCo+2onZnlatticesitesinthelms[ 45 126 127 ].TheintensityoftheseabsorptionsinthedopedlmsalsoincreaseswithincreasedCoconcentrationsuggestingmostoftheCoissolubleinthelattice.Specically,thepeakslocatedatenergiesof1.9eV651nm,2.0eV8nm,and2.2eV64nmcorrespondtothe4A2!2EG,4A2!4T1P,and4A2!2A1G,respectively[ 128 ].Theband-gapofthealloyswascalculatedbyplottingh2vs.handextrapolatingthelinearportionoftheplottoh2=0.Theplotsaregiveningure7-8.Atlowcobaltdoping|theundopedlmandthelmdopedwith2%Co|theabsorptionedgeiswelldenedandcanbetreliably.Theexcitonpeakisclearlyvisibleintheundoped 103

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lmindicatinggoodqualityZnO.However,asthecobaltdopingisincreased,alowenergyabsorptiononsetappearsandthedopingsmearsoutthelinearregiongivingitamoreroundedshape.Lineartstoboththehighandlowenergyregionsaregiveningure7-8asafunctionofcobaltconcentration.Thehighenergyslopesgiveaband-gapenergythatlinearlyincreasesblue-shiftswithnominalcobaltconcentration.Thelowenergyslopesgiveanenergythatroughlyremainsconstantaround2.8{2.9eV.Somereportsintheliteratureobserveared-shiftinthebandgapenergyasthecobaltconcentrationisincreased[ 129 { 132 ].Thered-shiftistypicallyattributedtothesp-dexchangebetweentheZnObandelectronsandlocalizedd-electronsassociatedwiththedopedCo+2cations.Theinteractionleadstocorrectionsintheenergybands;theconductionbandisloweredandthevalencebandisraisedcausingthebandgaptoshrink[ 133 ].Ontheotherhand,otherpapershavereportedablue-shiftinthebandgapofZnOwithcobaltdoping.Pengetal.reportedablue-shiftintheband-gapofthematerialandared-shiftofthebandtails,whichissimilartoourobservations[ 134 ].Yooetal.alsoobservedablue-shiftinAlandCocodopedZnOlmswhichwereattributedtotheBurnstein-Mosseectfromanincreaseinthecarrierconcentration[ 135 ].Theblue-shiftinthegapenergyisprobablynotcausedbytheBurnstein-Mosseectintheselms.Inaheavilydopedn-typesemiconductortheFermilevelresidesintheconductionband.Thefreeelectronsinthesemiconductorlltheloweststatesintheconductionbandandthevalenceelectronscannolongerbeopticallyexcitedintotheselledstates.Thisresultsinanapparentincreaseintheonsetofabsorptionandthegapshiftstohigherenergy[ 79 ].Thisshiftrequiresanincreaseintheelectrondensity.However,thereisnosystematicincreaseinthecarrierdensityofthemeasuredlmswithadditionalcobalt.AsdiscussedearlierinChapter3,bandtailscanarisefromperturbationsinthebandstructurecausedbyimpuritiesanddisorder.Statesintroducedbyimpuritiesoverlapat 104

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highconcentrationsandevolveintoanimpurityband.Asseeningure7-8,thetails'intensityriseswithincreasedcobalt.Therefore,thetailsaremostlikelyrelatedtotheincreaseofimpuritystates.KittilstvedandcoworkerssawalargeMCDpeakat25,000cm)]TJ/F19 7.97 Tf 6.586 0 Td[(1.10eV,whichtheyassignedasavalenceband-to-metalchargetransferCTtransitioninZnCoO[ 32 ].Thelevelwasapproximately2,600cm)]TJ/F19 7.97 Tf 6.586 0 Td[(1.322eVbelowtheconductionband.Theonsetofabsorptioningure7-8startsaround2.9eV.ThisisbelowthevaluefoundbyKittilstved;however,theirMCDpeakhadsomebreadthwiththeonsetofthepeakbeginning23,000cm)]TJ/F19 7.97 Tf 6.587 0 Td[(1.85eV,whichisingoodagreementwiththeabsorptiononsetingure7-8.Thereforethelowenergyabsorptiononsetislikelyanelectrontransitionfromthevalencebandtothecobaltimpuritystates.7.3.3.2PhotoluminescencePhotoluminescencePLwasperformedontheseriesoflmsgrownat400Cinvacuum.ThePLwasusedprimarilytoverifytheband-gapshiftfoundintheopticalabsorptiondata.Figure7-9showsthePLspectratakenatroomtemperature.Fortheundopedlm,abroadluminescencebandisvisibleacrossmostofthespectrum.Broadgreen-yellowbandsaretypicallyattributedtodefectsinZnO,includingoxygenvacancies[ 136 ].Thelmsaregrowninlowoxygenpressuresandarenon-stoichiometricwhichcouldgiverisetothebroaddefectbands.AmoredetailedstudyofthelowtemperaturePLpropertieswouldhavetobedonetosaymoreaboutthebands.Thedipintheintensityat517nmisanartifactduetotheblazeangleofthediractiongratingandnotfromthelmproperties.Theband-edgeisvisibleintheundopedand2%Codopedlms,butisquenchedwithhighercobaltdoping.Higherresolutionscansaroundtheband-edgeemissionsoftheundopedand2%Cosamplesareshownintheinsetofgure7-9.Theband-edgeisblue-shiftedfrom3.25eVto3.31eVinthe2%Cosample.Thisisconsistentwiththeband-gapdeterminedfromopticalabsorption. 105

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7.3.4MagneticCharacterizationThevolumemagnetizationofthelmswasmeasuredusingaQuantumDesignsuperconductingquantuminterferencedeviceSQUIDmagnetometer.Beforemeasuring,thebacksandsidesofthesampleswereetchedinnitricacid0%nitric/50%deionizedwatertoremoveexcesssilverpaintandcontaminantsthatcouldcontributeaspuriousmagneticsignalasmeasuredbytheSQUID.Eachlmsurfacewasrstcoatedinphotoresistandbakedfor20min.at50Ctohelpprotectthelmduringetchingandthenoatedontopofthenitricacidfor3min.Thephotoresistwasremovedbyrinsinginacetone.TheSQUIDmagnetizationdataisnormalizedusingtwocommonmethods:normalizingbythelmvolumeandnormalizingbythenumberofbohrmagnetonsBperCoatom.ToconverttherawmagnetizationdatafromemutoB/Co,allthecobaltatomsareassumedtosubstituteonZnsitesandacationdensityof4.18x1022cm)]TJ/F19 7.97 Tf 6.587 0 Td[(3isusedfortheconversion.ThedataisnormalizedusingthecobaltconcentrationsfromtheEDSdatagiveningure7-1.Magnetizationdataforaseriesoflmsgrownat400Cinvacuumwithdierentamountsofcobaltaredisplayedingure7-10.Theselmsdonotshowcobalt-inducedsecondaryphasesbyXRD.Thelmswithanominalconcentrationof2%and5%cobaltareferromagneticwithhysteresisat10K;however,thesaturationmagnetizationgraduallydecreaseswithincreasedtemperatureandthehysteresisdisappears.Thelmwith30%Coshowsferromagnetismupto300Kwithaclearhystereticshape.Thesamplehasaroomtemperaturemagnetizationapproaching0.08B/Co.Themagnetizationdataforlmswithandwithoutsecondaryphaseformationarecomparedingure7-11.Asstatedpreviously,themaindierencebetweenthelmsisthegrowthpressureusedduringdeposition.The30%Colmwiththesecondaryphasehasalargermagnetizationthanthelmwithout.Thiswouldbeexpectedifthesecondaryphaseiscobaltmetal.Noticethelmdepositedinoxygenshowsverylittlemagnetization. 106

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Thisisalsotrueforlmsthatweregrowninhigheroxygenpressuresupto20mTorrpO2.Thelmwithcobaltprecipitationhasaroomtemperaturemagnetizationthatsaturates0.22B/Coathigherelds>1Tesla.Thelmwithoutprecipitationhasaroomtemperaturemagnetizationapproaching0.08B/Co.Boththesevaluesarewellbelowthe3Bmomentofthed7high-spincongurationofCo+2.Theyarealsobelowthe1.72B/Covalueofpuremetalliccobalt.However,thevaluesareingoodagreementwithothervaluesreportedintheliteratureforZnCoO.08to0.4uB/Co[ 48 67 137 ].RecentLSDA+Ucalculationspredictthatthenearest-neighborexchangecouplingsofcobaltinZnOshouldbeantiferromagnetic,irrespectiveofthegeometricalnearest-neighborarrangement[ 138 ].Thereforeathighcobaltconcentrations,wherestatisticallyagreaternumberofnearest-neighborswilldevelop,onewouldexpectalowermomentperCothanthatgivenbylowerCodopingconcentrations.However,thisisnotwhatisobservedinthepresentcase.7.3.5ElectricalTransport7.3.5.1HalleectTheoriginsofferromagnetismincobalt-dopedZnOarestillnotfullyunderstood.Whethertheferromagnetismistrulyintrinsicasaresultfromtheinteractionofcarrierswithmagneticdopants,oriftheferromagnetismisextrinsicandarisesfromsecondaryphasesornanoclustersisanimportantconsideration.TheusefulnessofaDMSrestsonitsabilitytoproduceandmanipulatespinpolarizedcurrents.Iftheferromagnetisminthesematerialsissolelylocalizedinsecondaryphasesanddoesnotpolarizethefreecarriers,thentheDMSisoflimitedutilityinspintronicdeviceapplications.ADMSthatcontainsanasymmetryinthecarrierspin-densityaspin-polarizedcurrentshouldexhibitananomalousHalleectAHEintransportmeasurements.ForferromagneticmaterialstheHallequationisgivenby:xy=RoB+RsM 107

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wherexyistheHallresistivity,RoistheordinaryHallcoecient,RstheanomalousHallcoecient,andBandMarethemagneticuxandmagneticeldvectorsnormaltothelmsurface.RoiscausedbytheLorentzforceactingonmovingchargecarriers.Theanomalousterm,Rs,isusuallyascribedtospin-dependentscatteringofcarriersatlocalatomicmoments.Carriersofoppositespinarescatteredindierentdirectionsateachmoment.Thismodiesthechargeaccumulationateachendofthesample.Athigherelds,themagnetization,M,willsaturatecausingtheanomalousHallcomponent,RsM,tobecomeconstant.Atthispoint,changesintheHallcurvearedeterminedbytheordinarycomponentandshouldbelinearineld.TheordinaryHallcoecientcanbeextractedfromthehigh-eld,linearregionandtheelectricalpropertiesofthematerialdetermined.Variable-eldHalleectmeasurementswereperformedonHallbridgepatternedlms.Thelmswerepatternedduringgrowthbydepositingthroughastainlesssteelshadowmask.Electricalcontactsweresolderedtothesampleusingindiummetal.ThelmsweremountedinsideaQuantumDesignPhysicalPropertyMeasurementSystemPPMStocontroltheambienttemperatureandappliedmagneticeld,andtheelectricaldatawascollectedusingaKeithleyhigh-impedanceHalleectsystemthecomponentsofthisset-uparedescribedintheappendix.BoththetransversexyandlongitudinalxxresistivityweremeasuredbyapplyingalongitudinalcurrentIxx.Measuredvaluesofxywillhaveacontributionfromxxifthevoltageleadsareslightlymisaligned.Thereforetheresistancedatawasgeometricallyandeldaveragedtoremoveanyasymmetriesandtoaccountforthermallyinducedvoltages.Sincexyisantisymmetricwithappliedmagneticeld,thedatawasaveragedbyxy;oddH=1 2[xy+H-xy-H]tohelpremoveanypartsofthesignalfromthelongitudinalcomponent[ 139 ].Conversely,xxissymmetricwithrespecttoappliedmagneticeldandcanbeaveragedusingxx;evenH=1 2[xx+H+xx-H].FortheZnOlmsdopedwith30%cobalt,ananomalousHallsignatureisobservedinsamplesgrownat400CandinvacuumasgiveninFigure7-12.Includedinthegureisa 108

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lmwithandwithoutcobaltprecipitation.BothlmsshowAHEatroomtemperatures,buttheeectismuchmorepronouncedinthelmwithcobaltprecipitation.DerivativesareshownintheinsetsofeachgraphtoshowthechangeofslopecreatedbytheAHE.Interestingly,thereisnoevidenceofhysteresisineitherHallcurvewithintheresolutionoftheinstrumentation.AweakAHEisalsoobservedinthelmdopedwith15%Co.Forlmsdopedwithlessthan15%Co,noAHEisobservedatthechosengrowthconditions.AlthoughtheAHEistypicallyattributedtoscatteringofcarriersbylocalmagneticmoments,ithasbeensuggestedthatnonmagneticmaterialswithferromagneticprecipitatescanalsoexhibitanAHE.ThishasbeenobservedinCo-dopedTiO2lmsbyShindeetal.[ 140 ].TheyreportanAHEinsuperparamagnetic,highlyreducedcobalt-dopedrutileTiO2lmsthatcontaincobaltmetalclusteringclosetothelm/substrateinterface.SuperparamagneticgranularmetalcompositesarealsoknowntoshowanomalousHallbehavior[ 141 142 ].WhilethelargerAHEissuspectinthelmswithprecipitation,theoriginoftheAHEinthenon-precipitatedlmsisnotobvious.TheAHEcouldarisefromthecobaltatomsthataresubstitutionalintheZnOlatticeorpossiblyfromsmallnanoparticlesofcobaltmetalnotobservablewithXRD.EvidenceforatrueAHEfromsubstitutionalcobaltatomsisinferredfromannealingexperiments,whichisdiscussedshortly.Itisimportanttoexaminethepossibilityofcobaltparticlesasthesourceofferromagnetism.Thelmswith30%Cogrownat400Cshowhysteresiswithanitecoerciveeldandremanenceat300KfromtheSQUIDmagnetometrydata.Superparamagneticparticlesystemswillnotshowhysteresisabovetheirblockingtemperature,wherethetemperatureishighenoughforthesuperparamagneticmomentstouctuatefasterthanexperimentalmeasuringtimes.Thissuggeststhatifthecobaltprecipitatesarebehavingassuperparamagnets,thentheirblockingtemperatureshouldbehigherthan300K.Theblockingtemperaturemaybeestimatedusing[ 140 143 ]:TB=KV 25kB 109

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whereKisthemagneticanisotropyconstantK=4.1x105J/m3forhexagonalcobalt[ 144 ],Vistheparticlevolume,andkBisBoltzmann'sconstant.Usingthisequation,theestimateddiameterforasphericalcobaltparticlewithablockingtemperatureof300Kis7.8nm.Thisisclosetotheparticlesizeseeninthe30%ColmfromTEM,whichisabout5nm.Thesmallerparticlesizeshouldshowlowerblockingtemperatures,suggestingtheblockingtemperatureinthelmsshouldbelessthan300K,whichisinconsistentwiththemagnetizationcurvesatthesametemperature.Thissuggeststhatiftheparticlesaresuperparamagnetic,theydonotcontributetotheferromagneticmomentseeninthelmsatroomtemperature.However,sinceTBscaleswiththecubeoftheparticleradius,asmalldeviationintheparticlesizewillhavealargeeectontheblockingtemperatureandmakessuchanargumenthardtojustifybasedsolelyontheblockingtemperatureequation.7.3.5.2MagnetoresistanceThemagnetoresistanceMRofthelmswasmeasuredsimultaneouslywiththeHalldata.TheMRcanprovidesomeusefulinsightsintothetransportpropertiesofsemiconductors,includingthepotentiallandscapeoftheimpuritydistributionandlatticedisorder.Aluminumwasaddedasacodopanttosomeofthelmstostudytheeectsofelectronconcentrationonthetransportproperties.Aluminumisawellknown-typedopantinZnO;Al+3substitutesonaZn+2siteanddopesanelectronintothelattice.ChangesintheMRbehavioratlowtemperatureasafunctionoftheelectronconcentrationarefoundtobeconsistentwiththenotionofacriticalcarrierconcentrationattheMetal-to-InsulatorTransitionMIT.ThecobaltdopedlmstransitionfrompositivetonegativeMRastheelectronconcentrationcrossestheMIT.InordertounderstandtheMRresults,abriefdiscussionofthemetal-to-insulatortransitionisinorder.Dopedsemiconductorswillbecomemetallicwhensucientoverlapovercomesthelocalizingeectsofelectron-electroncorrelationanddisorder.Atacertain 110

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concentration,electronsdelocalizeandtheirwavefunctionsextendthroughoutthelattice.Thisisawellknownmetal-to-insulatortransitioninsemiconductormaterials.Atdilutelimits,impuritystatesareisolatedandelectronsareconnedtoahydrogenicorbitalaroundtheirassociatedimpurity.TheradiusofthisorbitalifgivenbyrH=m/m*aH,whereistherelativedielectricconstantofthehostmaterial,m*isthecarriereectivemass,andaHisthehydrogenicBohrradiusaH=53pm[ 31 145 ].Astheimpurityconcentrationisincreased,theimpurityorbitalsbegintooverlaptoformanimpurityband.TheamountofoverlaprequiredtoovercomethelocalizingeectsofcorrelationisgivenbytheMottcriterion,nc1 3rH0.25wherencisthecriticalconcentration.ThecriticalconcentrationforZnOisabout4x1019cm)]TJ/F19 7.97 Tf 6.587 0 Td[(3.Belowthecriticalconcentrationcarriersremainboundtolocalizedsites,butcanmovebyhoppingbetweenoccupiedandemptystates.ThisistheinsulatingregimeoftheMIT.Abovethecriticalconcentration,theimpuritybandstatesbecomedelocalizedandthecarriersareitinerant.Asareference,theMRbehaviorofanundopedZnOlmdepositedinvacuumat600Cisshowningure7-13.TheMRisnegativeovertheentireeldrangeandisdependentontemperature.Thisisconsistentwithotherreportsfoundintheliterature[ 78 146 147 ].ThenegativeMRinthesereportswasobservedforhighlyn-typeZnOlmswithelectronconcentrationsexceeding1020.NegativeMRinhighlydopedsemiconductorsisthoughttobecausedbytheweaklocalizationcorrectiontotheconductivity[ 148 ].Hallmeasurementsonthelmpresentedingure7-13givecarrierconcentrationsontheorderof1018.ThisiswellbelowthecriticalconcentrationoftheMITanddemonstratesthattheMRofnonmagneticZnOisnegativeonbothsidesoftheMIT.Figure7-14showstheMRbehavioratvarioustemperaturesforthreelmsdopedwith5%Co.Theselmsweredepositedunderslightlydierentconditionstoimpartdierentelectronconcentrations.Thegrowthconditionsareindicatedinthegurefor 111

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eachlm.TheelectronconcentrationsasmeasuredbytheHalleectarealsoprovidedinthegure.2at%Alwasaddedtothethirdlmtosubstantiallyincreasetheelectronconcentration.TheMRforthethreelmsisnegativeandqualitativelysimilarabove100K.TheMRathightemperatureisrathersmallbeinglessthan-0.5%.However,atlowtemperaturetheMRissubstantiallydierentbetweenthelms.At10K,aprogressionfrompositivetonegativeMRoccursastheelectronconcentrationisincreased.Belowthecriticalconcentrationofthemetal-to-insulatortransition,theMRispositiveovertheentiremagneticeld.Nearthecriticalconcentration,asmallnegativeMRcomponentappearsatloweld.TheMRvaluedecreasesfrom10%to1.5%astheMITisapproached.AtconcentrationsmuchlargerthanthecriticalconcentrationinthemetallicconductionregimetheMRisnegativeovertheentireeldrange,similartotheMRbehavioroftheundopedZnOlm.TheMRbehaviorwasalsostudiedforlmsheavilydopedwithcobalt.Figure7-15showstheMRdatacollectedfortwolms,onedopedwith30%Coandtheothercodopedwith30%Coand2%Al.AgaintheAlisaddedtoincreasetheelectronconcentrationofthedopedlm.AsimilarprogressionfrompositivetonegativeMRacrosstheMITboundaryisseenatlowtemperature.Athighertemperature,thelmdopedonlywithcobalt,developsakinkintheMRcurvenear10,000Oe.Tostudythekinkformation,theshapeoftheMRcurvewastrackedattemperaturesbetween10Kto100Kandisshowningure7-16.TherelativemagnitudeofthenegativeMRpeakatzero-eldgraduallyincreaseswithtemperature.Between50Kand75K,thispeakincreasesandtheMRcomponentathigh-eldbecomesnegative.TheMRoftheCo-andAl-codopedlmsisremarkablysimilartothe5%Co,2%Alcodopedlm,andqualitativelysimilartotheundopedlm.ThissuggeststhatelectronictransportisnotgreatlyaectedbythecobaltconcentrationabovetheMIT.Figure7-17showsthetemperaturedependenceoftheresistivityforthe5%and30%Colms.Intheinsulatingregime,theresistivitydecreasesasthetemperatureis 112

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increased.Thisistypicalbehaviorforsemiconductorlms.AbovetheMIT,theresistivityincreaseswithincreasedtemperature,typicalofmetallicbehavior.ThetemperaturedependenceisconsistentwiththenotionthattheAldopedlmshavecrossedoverintothemetallicconductionregime.TheprogressionfrompositivetonegativeMRingures7-14and7-15iscommensuratewiththemetal-to-insulatortransition.OntheinsulatingsideoftheMIT,theMRispositive,andonthemetallicsideoftheMIT,theMRbecomesnegative.Nearthetransitionpoint,negativeMRisseenatloweldandpositiveMRathigheld.However,asthetemperatureisincreased,thepositiveMRcomponentvanishesandonlynegativeMRisseen.AsimilardependenceofMRoncarrierconcentrationsaroundtheMITincobalt-dopedZnOhavebeenreportedbyXuetal.andKimetal.[ 51 149 ].Clearly,amodelincludingtheinterplayofcarrierconcentration,dopedcobaltspins,andtemperaturedependenceisneededtoexplainthepeculiarMRinthesesamples.7.3.6EectsofAnnealingInaneorttobetterunderstandthenatureoftheselmsandtheoriginofferromagnetism,the30%Cosampleswereannealedinbothoxidizingandreducingatmospheres.Thelmsweredepositedinvacuumanditisassumedthatmostoftheelectronsareassociatedwithoxygenvacanciesinthelattice.Annealingthelmsinoxidizingconditionsshouldllthevacanciesandreducetheelectronconcentration.Almwasannealedat500Cfor1hourin1atmofoxygen.Table7-2showsthechangeinelectricalpropertiesafterannealing.Theelectronconcentrationslightlydecreasedwhiletheresistivityslightlyincreased.Mostnotably,theAHEintheoxygenannealedlmwassignicantlysmallerthanthatseenintheas-grownlm.TheHallresistivityandmagnetoresistanceareshowninFigure7-18.Sincethechangeinelectronconcentrationissmall,onemayrequireanalternativeexplanationtothediminishedAHEcoecient.OnepossibilityisthatsomeofthecobaltmetaldissolvesintotheZnOlattice.Thisseemsunlikelygiventhatthehighcobaltconcentrationisalreadymetastable.Itisalsopossible 113

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thatthecobaltmetalreactswithoxygentoformcobaltoxide.However,XRDafterannealingshowsthatthepeaktentativelyassignedtocobaltmetalisstillpresentinthelmandthatnocobaltoxideshaveformedFigure7-18c.Annealingmaycausesomechangeinmicrostructure,suchasachangeinsizeoftheprecipitateswhichcanalterthetransportbehavior.Itisalsopossiblethattheoxygenannealat500Cissucienttodriveoutanyhydrogenthatresidesinthelattice.RecenttheoreticalcalculationssuggestthathydrogenmightmediateferromagneticinteractionsbetweencobaltatomsinaZnOmatrix[ 150 ].Tofurtherexplorethispossibility,wehavealsoannealed30%Co-dopedZnOlmsinforminggas%H2/Arat500Cfor1hour.Unfortunately,thisannealinhydrogenledtoadecompositionoftheZnO:Colm,suggestingthattheCoisindeedsubstitutionalandmetastableintheZnOmatrix.TheSEMmicrographinFigure7-19showsthelmdecompositionafterannealinginhydrogen.AnXRDcomparisoninFigure7-20showsalargealmost10foldincreaseintheintensityofthe11Cometalpeakascomparedtotheasgrownlm,indicatinganincreaseinCometalclustering.High-resolutionfour-circleXRDindicatesthecobaltphaseisalignedwiththeZnOlattice.Cobaltmetalcanexistineitherthehexagonalorfccstructure.Thefccphaseisstableabove425Cbutisoftenobservedasametastablephaseatroomtemperature.SincethehexagonalandfccstructuresdieronlyintheirstackingsequencehexagonalABABandfccABCABCthefcc11planeshavethesamed-spacingasthe001hexagonalplanes.Sincestandard-2XRDmeasuresplanesparalleltothesamplesurface,thehexagonal01andfcc11orientationsareindistinguishable[ 151 ].However,sincetheperiodicitiesoftheplanesaredierent,o-axispeaksplanesnotparalleltothesurfacecanbeusedtoidentifythestackingsequence.Todistinguishbetweenthetwophases,ano-axisdiractionscanthroughtheCo10Lreciprocallatticepointswasperformedasshowningure7-21a.The0LscanshouldshowpeaksatintegralLpositionsforhexagonalstackingandatparticularinteger/3positionsforcubicstacking[ 56 ].TheL-scanshowsthatthepackingis 114

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mostlycubicABCstacking.However,ascanthroughthehexagonalcobalt01position,whichcutsthroughtherodatL=1intheL-scan,showsthereisasmallintensityat01asshownintheinsetofthegure.Thissuggestsomestackingfaultsarepresentorsmallregionsofhexagonalcobalt.Additionally,aphi-scanthroughthecubicCo0at=35showsin-planealignmentofthecobaltgrainsasshowningure7-21b.Clusteringshouldleadtoanincreaseintheobservedferromagnetismsincealargervolumeofcobaltmetalwillhavealargermagnetization,whichisveriedinFigure7-22.WhilethepresenceofCometalprecipitatesinthelmsprovidesapossibleexplanationtothemagneticbehavior,anexaminationofthemagneticbehavioroftheCo-dopedZnOlmsgrownatdierentconditionsprovidescircumstantialevidencethatthecobaltprecipitatesarenottheoriginofthemagneticbehavior.First,varyingthegrowthconditionsofthelmshasalargeeectontheobservedmagnetization.Filmsgrownathighertemperatures00Cand600Cinvacuumshowverylittle,ifany,magnetization.Giventhattheselargeconcentrationsofcobaltarehighlymetastable,onewouldexpectastrongertendencytoformmoresegregatedcobaltmetalatthehighertemperature,butthermodynamicallytheantiferromagneticCoOphaseisstable.SQUIDmeasurementsshownoevidenceforferromagnetismforZnO:Colmsgrownattheelevatedtemperatures.XRDscansfortheselmsweregiveningure7-5.SQUIDcharacterizationfortheselmsisshowningure7-23.ThecoexistenceofacobaltmetalXRDpeakandtheabsenceofferromagnetisminthelmgrownat500Cinvacuumisaninterestingresult.Onewouldexpect,fromthesensitivityoftheSQUIDmagnetometer,thatthelmwouldshowmagnetizationinthepresenceofcobaltmetal.ThepresenceofCoOcouldberesponsibleforreducingthemagnetizationbyremovingcobaltfromthelatticeandprecipitates.Howeverthereductioninmagnetizationisover100timessmallerascomparedtothelmgrownat400Candmaysuggestthatthesmallparticlesofcobaltmetaldonotmakealargecontributiontothemagnetization. 115

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Figure7-1.EDSresultsforaselectnumberoflmsgrownunderdierentconditions.Thedatashowsthecobaltconcentrationinthelmsascomparedtothecobaltconcentrationinthetarget. Table7-1:Possiblecobalt-inducedsecondaryphases. PhaseStructure2degCouplingT*K CoCubic1144.216FerromagneticTc=1373Hex00244.762CoOCubic0042.401AntiferromagneticTn=291Co3O4Spinel40044.808AntiferromagneticTn=30ZnCo2O4Spinel40044.738Antiferromagneticn-typeN/ARef[ 152 ]Ferromagneticp-typeCoAl2O4Spinel40044.692AntiferromagneticTn<40 116

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Figure7-2.XRDscansforaseriesoflmsgrowninvacuumat400C.Thelmsarepredominatelyc-axisorientedZnO.Athighcobaltconcentrations,acobaltinducedsecondaryphaseappearsnear2=44.4.TEMandHigh-resolutionXRDsuggestthisphaseisamixtureofcubicandhexagonalcobaltwithstackingfaults. 117

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Figure7-3.TEMmicrographsofasampledopedwith30%CothathascobaltprecipitationpresentintheXRDscan.Theprecipiationappearsmostlyatthesubstrate/lminterface. 118

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Figure7-4.ConvergentbeamTEMdiractionpatternsofZnOlmdopedwith30%Cogrownat400Cinvacuum.aZnOhexagonallmat[120]zoneaxis,bZnOim+nanoparticles,cZnOlm+nanoparticles+sapphiresubstrate,dNanoparticleat[120]zoneaxiswithd002=0.20nmandd210=0.13nm,whichisconsistentwithmetalliccobalt. 119

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Figure7-5.XRDscansforZnOlmsdopedwith30%Co.aFilmsdepositedat400Cwithdierentpressureconditions.bFilmsdepositedat500Cand600Cinvacuum. 120

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Figure7-6.Thermodynamicpredominancediagramforcobaltoxides. 121

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Figure7-7.UV-VistransmissionofCo-dopedZnOlmsdepositedinvacuumat400C.Theinsetshowsacloseupviewoftheabsorptionlevelswhichcorrespondtothe4A2!2EG.9eV,4A2!4T1P.0eV,and4A2!2A1G.2eV. 122

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Figure7-8.Opticalband-gapsofCo-dopedZnOlms.ah2plotsforCo-dopedZnOlmsdepositedinvacuumat400C.Straightlinetsthroughthelinearregionsoftheplotareextrapolatedtoh=0tondtheband-gap.bPlotoftheband-gapvaluesasafunctionofnominalcobaltconcentration.Filledcirclesaretheband-gapandhollowtrianglesaretheonsetofabsorptionatlowenergy. 123

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Figure7-9.PLresultsforCo-dopedZnOlmsdepositedinvacuumat400CC.ThedipinthePLspectraat517nmisanartifactfromthediractiongradingblazeangle.ThePLintensitydecreaseswithhighercobaltdoping.Theinsetshowshigherresolutionscansaroundtheband-edgepeakfortheundopedand2%Cosamples.Thevaluescorrespondwellwiththeabsorptiondata. 124

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Figure7-10.SQUIDmagnetizationcurvesforCo-dopedZnOlmsdepositedat400Cinvacuum.ThelmsdonotshowanysecondaryphasesbyXRDmeasurement. 125

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Figure7-11.SQUIDmagnetizationcurvesforCo-dopedZnOlmsdepositedat400C.Thegrowthpressureandwhetherthelmcontainssecondaryphaseprecipitationisindicatedinthelegend. 126

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Figure7-12.AnomalousHalleectin30%Co-dopedZnO.Filmsweregrownat400Cinvacuum.almwithcobaltprecipitation.TheupperinsetshowstheHallcurvesatdierenttemperatures.ThelowerinsetshowsthederivativeoftheHallcurve.blmwithoutcobaltprecipitation.TheinsetshowsthederivativeoftheHallcurve. 127

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Figure7-13.ThemagnetoresistanceofanundopedZnOlm.Thelmwasdepositedat600Cinvacuum.Themeasurementtemperatureandcarrierconcentrationareindicatedineachgure. 128

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Figure7-14.Themagnetoresistanceof5%Co-dopedZnOlms.Thelmwithnncwasdepositedat400cinvacuumandiscodopedwith2%Altoimpartahighelectronconcentration. 129

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Figure7-15.Magnetoresistanceof30%Co-dopedZnOlms.Bothlmsweredepositedat400Cinvacuum.NeitherlmsshowsignsofasecondaryphasebyXRDmeasurement.Thelmwithn>nciscodopedwith2%Altoimpartahighelectronconcentration. 130

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Figure7-16.Magnetoresistanceof30%Co-dopedZnOlmattemperaturesbetween10Kto100K.ThedevelopmentofthenegativeMRkinkisestablishedatlowtemperaturefollowedbyatransitiontonegativeMRacrosstheentireeldrangewithincreasedtemperature. 131

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Figure7-17.Temperaturedependentresistivitymeasurementsfor5%and30%Co-dopedZnOlms. Table7-2.Transportdatafora30%Co-dopedZnOlmwithcobaltprecipitation.DataisforbeforeandafterannealinginO2at500C. RoCarriersMobilityZero-eldxxcm3C)]TJ/F19 7.97 Tf 6.586 0 Td[(1cm)]TJ/F19 7.97 Tf 6.587 0 Td[(3cm2v)]TJ/F19 7.97 Tf 6.586 0 Td[(1s)]TJ/F19 7.97 Tf 6.587 0 Td[(1-cm Asdeposited-1.6853.70x10187.120.2367O2annealed-2.9982.08x10188.970.3344 132

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Figure7-18.Hallresistivityandmagnetoresistancefora30%Co-dopedZnOlmwithcobaltprecipitation.DataisshownbeforeandafterannealinginO2at500.InsetsshowthederivatesoftheHallresistivitytoshowAHE. 133

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Figure7-19.SEMmicrographsatdierentmagnicationsofthesurfaceofaCo-dopedZnOlmthathasbeenannealedinH2/Arat500Cfor60min.Magnicationandscalebarsaregivenineachmicrograph.Thearrowinashowstheregionoflmdegradation.Thisareawasbrownincolorcomparedtothegreencolorofthelm. 134

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Figure7-20.-2XRDfora30%Co-dopedZnOlmbeforeandafterannealinginforminggasat500C. 135

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Figure7-21.High-resolutionXRDscansfor30%Co-dopedZnOlmthathasbeenannealedinforminggasat500C.aL-scanalongtheCoL.Theinsetshowsano-normal-2scanthroughthehexagonalCo101.bphi-scanthroughtheCo200at=35. 136

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Figure7-22.SQUIDmagnetometryfora30%Co-dopedZnOlmbeforeandafterannealinginforminggasat500C. 137

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Figure7-23.SQUIDmagnetometryfor30%Co-dopedZnOlmsdepositedunderdiererentconditions. 138

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CHAPTER8CONCLUSIONTheadvancementofspintronicsasapracticaltechnologydependsuponthedevelopmentandunderstandingofsemiconductorsthatcansupportspin-polarizedcarrieroperationaboveroomtemperature.Theresearchpresentedinthisdissertationexploredthepossibilityofusingwideband-gapoxidesemiconductorsasspintronicmaterials.Bothtransition-metaldopedCu2OandZnOwereinvestigated.MnandCowereusedasthetransition-metaldopantstoprovidelocalizedspinsinthehostsemiconductorlattice.Thinlmsofthesematerialsweredepositedusingpulsedlaserdeposition.Variousmaterialpropertiesincludingthestructure,magnetic,andelectronictransportpropertieswerethencharacterizedtogainabetterunderstandingofthematerials.8.1Mn-dopedCu2OThemagneticpropertiesofCu2Olmsdopedwith1at.%Mnwereinvestigatedunderdierentgrowthconditions.ThisresearchwasstimulatedbyDietl'stheoreticalpredictionthatcarrier-mediatedferromagnetismisfavoredinMn-doped,highlyp-type,wideband-gapsemiconductors,suchasZnMnOandGaMnN.Cu2Oisanaturallyp-typesemiconductorwithawidedirectband-gapandthereforeholdsinterestinexploringspinbehaviorinoxides.TheMnsolubilityinCu2OwasfoundtobesmallandtheprecipitationofMn-oxideswasfavoredathighgrowthtemperatures.However,metastableincorporationofMnintheCu2Ocouldbeachievedatlowtemperature0C.Thesephasepuresampleswerefoundtobenon-ferromagnetic.FerromagnetismwithaTc50Kwasobservedinthelmsdepositedathighertemperatures,butappearstobeassociatedwithaMn3O4secondaryphasewhichhasaTcnear50K.Spintronicconceptsbasedonferromagneticsemiconductorsrequirethedistributionofchargecarriersinthesemiconductortobespin-polarized.Magnetismderivedfromlocalizedmagneticprecipitatesisoflittleutilityforsemiconductor-basedspintronicsif 139

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thecarriersarenotpolarized.Therefore,basedonthecollecteddata,itisconcludedthatMn-dopedCu2Ohaslimiteduseasaferromagneticsemiconductorunderthestudiedgrowthconditions.ThelowsolubilityofMnintheCu2OlatticelimitsexperimentationusinglargerconcentrationsofMn.Perhapslargerholeconcentrationsoralternativefabricationtechniquescouldbeusedtoinduceferromagnetism.8.2Mn-dopedZnOSeveraloftheproposedmodelsforferromagnetisminZnODMSemphasizetheimportanceofholesmediatingtheexchangeinteractionbetweendopedMnspins.Thisresearchinvestigatedthetrendsinthemagnetizationasafunctionofcarrierconcentrationinordertoelucidatetheroleofchargecarriersontheferromagnetisminthesematerials.ThecarrierconcentrationwasvariedusingSnandPaselectroniccodopants.Codopingallowedindependentcontroloverthemagneticandelectronicpropertiesbydopingeachseparately.Thisprovidedaplatformtostudytheeectsofcarrierconcentrationontheobservedmagneticproperties.Snactsasann-typedopantprovidingextraelectronstotheZnO.Pactsap-typedopantthatsuppliesholestocompensatethenativeelectronconcentrationinZnO.Thinlmsdopedwith3at.%Mn,andeitherSnorPasthecodopant,weredepositedinanoxygenatmosphereof20mTorrinatemperaturerangeof400{600C.TheelectronconcentrationintheZnMnO:SnlmswascontrolledbyvaryingtheSncontent.Initially,themagnetizationincreasedwithminimalSndopingresultinginamaximummagnetizationof0.5B/Mnatomat300K.However,withincreasedSndopingtherewasaninversecorrelationbetweentheSncontentandthesaturationmagnetization.AstheelectrondensityincreasedwithSndoping,themagnetizationdecreased.ThetrendinmagnetizationshowedasimilarcarrierdependenceusingPasacodopant.Underconditionswheretheacceptordopantswereactivated,themagnetizationwasenhanced.Theresistivityoftheas-depositedlmwas4,000-cmat300Kwithan 140

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electrondensityof3x1015cm)]TJ/F19 7.97 Tf 6.586 0 Td[(3,andbecametooinsulatingtomeasureafterannealing.Thesaturationmagnetizationafterannealingwasaround0.15B/Mnatom.ThecorrelationofsaturationmagnetizationwithelectrondensityisinterestingandprovidessomeinsightintothemechanismforferromagnetisminZnMnO.OverlapoftheMnd-stateswiththevalencebandsuggestsholesarenecessaryinordertoinduceferromagneticordering.Magnetisminlowfree-electrondensitymaterialisconsistentwiththeboundmagneticpolaronmodelinwhichboundacceptorsmediatetheferromagneticordering.Theholesmaybedelocalized,butwithlowmobility,thusyieldinglowconductivity.Inthiscase,thecarriermediatedmechanismmaysucewithouttheneedtoinvokeboundpolaronsasinherenttotheferromagneticordering.Ineithercase,theadditionofelectronstothesystemwillmovetheFermienergylevelupintheband-gap,resultinginadecreaseinholedensityandareductioninmagnetization.WhilethetrendofmagnetizationversuscarrierdensitywassimilarintheSnandPcodopedlms,themeasuredmagnetizationwasslightlydierent.ThemagnetizationwasanticipatedtobelargerfortheannealedP-dopedlms,ascomparedtotheSn-dopedlms,sincetheFermilevelshouldresideclosertothevalenceband.However,thiswasnotobserved;lmswithminimalSndopingexhibitedahighermagnetizationperMnatom.Thereasonforthisisunclear,butcouldbearesultofotherdefectsformedaftertheannealingprocess.TheXPSdataalsosuggestedachangeinvalenceofsomeoftheMnatomsafterannealing.Thismayhavecausedareductioninmagneticmomentthatcompetedwiththesuppressedelectrondensity.Nevertheless,theoveralltrendofmagnetizationversuscarrierdensitywasconsistentbetweentheexperiments.TheresultsofthisworkareconsistentwiththeobservationsofKittilstvedandcoworkers[ 32 ].TheauthorsconcludedthatferromagnetisminZnOiscloselytiedtothechargetransferelectronicstructureofthetransition-metaldopant.ForMn,ferromagnetismisinducedwhentheholesfromtheacceptorstateshybridizewiththe 141

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Mnions.FerromagnetismwasobservedwhentheZnOwaslocallydopedp-type,butnoferromagnetismwasobservedwhendopedn-type.8.3Co-dopedZnOTheeectofcobalt-dopingonthemagneticandmagneto-transportbehaviorinZnOwasalsoinvestigated.Ferromagnetismwasfoundinlmsdepositedatlowtemperature00Cinvacuum,whilelmsdepositedinoxygenorathighertemperatureswerenon-magnetic.Filmsdepositedundervacuumhadratherhighelectronconcentrationsandarepresumablydopedwithoxygenvacancies.Segregationofcobaltmetaloccurredinlmsdopedwith15%orgreaterCoconcentrationswhendepositedinlowbasepressure<10)]TJ/F19 7.97 Tf 6.586 0 Td[(5Torrvacuumconditions.Theprecipitatesweresmall5nmandalignedwiththeZnOlattice.Thesegregationcouldbesuppressedbydepositinginhigherbasepressures>10)]TJ/F19 7.97 Tf 6.586 0 Td[(5Torr,buttheprocesswasnotfullyreproducible.Theselmsalsoshowferromagnetismalbeitwithlowermagnetizationthanthelmswithmetalliccobalt.Thissuggeststhatthecobaltmetalparticlescouldberesponsibleforsomeoftheobservedmagneticpropertiesinthelmswithprecipitation.However,theroleandtheextentofmagnetizationfromtheparticlesisuncertain.Thedistributionofparticlesthatbehaveferromagnetically,orbehavesuperparamagneticallybecauseoftheirsmallsize,andtowhatextentthelmpropertieshaveontheferromagnetismintheprecipitatedlms,isunclear.Thecobalt-dopedlmsalsoexhibitedpeculiarmagnetoresistancethathadastrongdependenceonthecarrierconcentration.TheMRwasstudiedinlmswith5%Coand30%Codoping,alongwithanundopedZnOlmasareference.Atlowtemperature,thecobalt-dopedlmsundergoaprogressionfrompositivetonegativeMRastheelectronconcentrationisincreasedandthelmscrossoverthemetal-to-insulatortransitionMIT.Atlargecarrierconcentrations,wellintothemetallicregime,theMRbehaviorfortheundoped,5%Co,and30%Colmsisquitesimilar.ThissuggeststhattheelectronictransportisnotgreatlyaectedbythecobaltconcentrationabovetheMIT.Onepossible 142

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explanationisthatsincetheFermienergyhasmovedfullyintotheconductionband,thetransportisdominatedbythecharacteroftheconductionband.Inanappliedmagneticeld,thecobaltd-statesandthedonorimpuritystatesaresplitintospin-upandspin-downsub-bands.IftheFermienergyliesintheconductionband,thereislikelynooverlapoftheFermienergywiththespin-polarizedd-bandsofthecobalt,andtheconductionischaracteristicofundopedZnO.ApossiblefutureexperimentthatmayproveinterestinginunderstandingtheMRbehaviorcouldbeexecutedthroughthefabricationofeld-gatedHallbars.Voltageappliedtoatopgateelectrodecouldbeusedtovarytheelectronconcentrationnearthevicinityofthegate.MagnetoresistanceandHallmeasurementscouldthenbeperformedatvaryinggatevoltagestoexaminetheirbehaviorwithdierentelectronconcentrations.Ideally,onecouldsmoothlystudytheMRbehavioracrosstheMITandcloselyexaminethechangesnearthetransitionpoint. 143

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APPENDIXAHALLEFFECTSYSTEMANDEQUIPMENTA.1IntroductionI'veaddedthisinformationforreadersinterestedindetailedinformationregardingtheresistivityandHalleectmethodsusedinthiswork.Irstwrotethissectionasareferenceformyresearchgrouptointroducethemtotheequipment.However,afterafewchanges,Idecidedthatthissectionalsomadeavaluableadditiontothedissertation.I'veincludedoperationalinformationaboutourcurrentequipment,simpleHalleecttheory,andexplanationsfortheadvantagesanddisadvantagesofcertainmeasurementtechniques.Ihopethissectionwillprovebenecialtothosewhocomeacrossit.A.2HallEectEquipmentTheHalleectwasdiscoveredbyEdwinHallin1879andhasbecomeapowerfultoolforthecharacterizationofelectronicmaterials.Thechargecarriertype,concentration,andmobilitycanallbedeterminedfromaccurateHallmeasurements.MostoftheresistivityandHalleectmeasurementsinthisthesiswereperformedusingaKeithleyhigh-impedanceHalleectsystem.ControloverthetemperatureandmagneticeldwasaccomplishedbyplacingsamplesinsideaQuantumDesignQDPhysicalPropertyMeasurementSystemPPMS.ThePPMSisaversatiletoolforprobingvariousmaterialproperties.Equipmentoptionsareavailableformeasuringasample'smagnetic,electric,andthermalpropertiesinthesamepieceofequipment.ThePPMS'sopenarchitecturealsoallowsroomforusercustomizationandothertoolsmaybeintegratedintothesystem.Forinstance,customizableprobeheadsallowtheintroductionofextrasignalcablesoropticalfeedthrusintothesamplearena.ThePPMSactsasaplatform,orameasurementenvironment",tocontroltheambienttemperatureandeldaroundthesample.Thetemperatureofoursystemcanbeadjustedfrom1.4K{400K,andthereisalongitudinalsuperconductingmagnetthatcanbeswepttoamaximumof7Tesla.ThePPMSisaheliumlledsystem.Theheliumisused 144

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forcoolingthesamplespacetolowtemperatureandalsokeepsthemagnetbelowitssuperconductingcriticaltemperature.Theheliumbathissurroundedbyanitrogenlledjacketthatreducestheheliumboil-o.Atlowtemperatures,thesamplechambermustbepumpedtopreventfreezingofatmosphericgasesinthesystem.Thisiscontrolledthroughanintegratedgashandlingsystem.TherearepackagedoptionsforbothDCresistanceandACtransportmeasurementsavailablefromQD.However,someofthesamplesproducedinourlabarehighlyresistive,>100M,andourcurrentPPMSresistivityoptionscannotsupportsamplesofsuchlargeresistance.Measuringtransportpropertiesonthesematerialsrequiredtheuseofseparateelectronics.Ahigh-impedanceHalleectcongurationusingKeithleycomponentswasused,includinga236SourceMeasurementUnitSMU,two6514Electrometers,anda71524x5matrixswitchcardhousedina7001switchingmainframe.AdiagramoftheequipmentisgiveinFigureA-1.Forconsistencyandconvenience,theexternalelectronicswereusedformostofthemeasurementsinthiswork,evenforsamplesthatcouldbeeasilymeasuredbytheinternalPPMSelectronics.IntegrationandcontroloverthePPMSandKeithleyelectronicswasaccomplishedusingtheLabViewprogramminglanguage.LabViewwasusedtoprogramtheautomatedresistivityandHallroutines,whichincludedprogramsforVanderPauwandHallbridgesamples,magneticeldandtemperaturesweeps,andtimedependentmeasurements.Ifnecessary,LabViewalsoprovidesexibilityforcreatingnewroutinesforfutureexperiments.Samplesaremountedtopucksthatcanaccommodatesamplesizesupto1cmx1cmsquare.Thereare12availablesignalwiresfedtothesamplespacethatmaybeusedformeasurements.Forexample,thestandardQDresistivitypuckallowsuptothreesamplesconguredfor4pt.measurementstobemeasuredsimultaneouslythroughthe12signalleadsleadsareusedforeachofthethreesamples.AccesstothesewiresfromoutsidethesystemisgrantedthroughaLemoconnector. 145

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TheKeithleyhigh-impedanceHalleectsystemisshowninFigureA-1.Thissectionwillprovideabriefdescriptionoftheequipmentusedandthereasonsbehindtheirselection.Thebasicpremiseofanelectricalmeasurementistoprovideaknownstimulustoasampleandtomeasurethesample'sresponsetothatstimulus.Forinstance,abasicresistivitymeasurementmaybemadebyapplyingaknowncurrenttoasample,measuringthesubsequentpotentialdierence,andcalculatingthesample'sresistancebyusingOhm'slaw.Aswithmostscienticmeasurements,accuracyisofprimeconcern.Variousinternalandexternalfactorscaneecttheaccuracyofelectricalmeasurements.Theseincludeerrorsourcesfromtheexperimentalset-up,includingleakagecurrents,groundloops,equipmentosetvoltagesandcurrents,andenvironmentalnoise.Thesampleitselfwillalsoprovidesourcesoferror,suchasthermalJohnsonnoise,non-ohmiccontacts,andphotovoltaiceects.Someoftheseerrorscanbesignicantlyreducedbytheproperchoiceofequipmentandmeasurementprocedure.TheKeithley236SMUcontainsbothsourceandmeasurementcapabilitiesinoneunit.TheSMUisusedtodrivecurrentintothesampleandmeasurethevoltagedierencecreated.The236providesacurrentrangeof100fAto100mAprovidingexibilityforsamplesofvariousresistance.Boththesourceandsenseleadspassthroughthe7152switchcard.The7152providesautomatedswitchingofthesignalcablesbetweenthe4samplecontactsthroughagridofinterconnectedrelays.ThiscanbeseenintheswitchcarddiagramincludedinFigureA-1.Eachcrossingpointofthegridrepresentsarelaythatcanbeopenedorclosedtoredirectthecablingtoeachsamplecontactpoint.Theserelaysalsoswitchtheguardconductors.SwitchcardseliminatetheneedtophysicallyrecongurethesamplecontactsforeachlegoftheVanderPauwroutine.ThehighandlowsignalvoltagesgeneratedinthesamplearesenttotheKeithley6514electrometers.Theseelectrometershaveahugeinputimpedance>200Twhichreducesinputloadingerrorsonhighresistancesamples.Inourset-up,thetwoelectrometersactasunitygainbuersbetweenthesampleandtheSMUvoltmeter.Theunitygainbuerisanamplier 146

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circuitthatduplicatestheinputvoltageattheoutputterminaloftheamplier.ThevoltagegainisunityVo Vin=1.Sincethebuer'sinputimpedanceislargeitcanreplicatethevoltagewithoutdrawinganycurrentfromthesource.Ideally,voltagemeasurementsdrawzerocurrent[ 153 ].Infact,anycurrentatthevoltmeterinputrepresentsanerror,calledtheinputosetcurrent.Thisistheadvantageofplacingabuerampliertheelectrometersinthecircuit.TheoutputfromtheelectrometersisthenreadbytheSMU'svoltmeterformeasurement.A.3SampleGeometryandMeasurementTechniqueThetwocommonsamplegeometriesforHallmeasurementsareHallbridgeandVanderPauwsamples.WhiletheexactdimensionsofHallbridgesamplesareneededforaccuratedata,theVanderPauwtechniqueallowssamplesofarbitraryshapeandsize.BothmeasurementtechniquesaredescribedintheASTMstandardF76-86StandardTestMethodsforMeasuringResistivityandHallCoecientandDeterminingHallMobilityinSingle-CrystalSemiconductors".IncludedintheASTMstandardarespecicationsforsamplegeometriesandthemathematicalequationsusedtocalculatethematerial'selectricalproperties.Bothtypesofsamplegeometrieswereusedinthiswork.TheLabViewprogramsusedtodrivethemeasurementroutineswerebasedontheinformationprovidedintheASTMstandard.Thissectionwillgiveabriefexplanationoftheadvantagesanddisadvantagesofeachmeasurement.Resistivitymeasurementscanbeperformedusinga2pt.method,whereonly2leadsareconnectedtothesample,byshortingavoltmeteracrossthesameleadsusedtodrivecurrentthroughthesample.Adiagramforthe2pt.congurationisprovidedinFigureA-2.Thedrawbackof2pt.measurementsisthat,inadditiontothesampleresistance,thevoltagedropacrossthecontactandleadresistancesarealsomeasuredbythevoltmeter.Thiscancauseseriousloadingerrorsifthecontactandcableresistanceislargerthanthesampleresistance.Amoreaccurateapproachrequires4contactstothesample.Currentisdrivenintothesamplethroughtwoleadsandthevoltageismeasuredacross 147

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theothertwo.Ideally,nocurrentowsacrossthevoltmeter,butthisdependsontheinputresistanceofthemeter.QualityvoltmetershavealargeinputresistanceM{10Gtoreducetheinputosetcurrent.Hallbridgesamplesaremeasuredusingthe4pt.contactmethod.Currentissourcedbetweenthetwolongitudinalcontactsandthevoltageismeasuredbetweentwocontactssituatedparalleltothecurrentdirection.Thehallvoltageismeasuredacrosstwocontactsthatareplacedperpendiculartothecurrentdirection.TheVanderPauwtechniquewasrstproposedbyL.J.VanderPauwin1958.Thetechniqueiswidelyusedsinceitprovidesaconvenientmethodforprobingthetransportpropertiesofmaterials.TheVanderPauwtechnique'sprimaryadvantageisthatcontactsmaybeplacedarbitrarilyaroundthesampleperimeter.Thisenablessamplesofarbitraryshapeandsizetobemeasureddirectly.Theneedtofabricatecomplexsamplegeometriesordetailedknowledgeofthesampledimensionsiseliminated.Onlyknowledgeofthethicknessisneededtocalculatethevolumeresistivityandcarrierdensityfromtheirrespective2-dimensionalsheetvalues.Thiscanbehighlybenecialincertaincases,suchaswhenthesamplematerialisdiculttoprocessorwhenthefacilitiesand/orthetimeforprocessingsamplesisnotavailable.ThereareseveraldisadvantagestotheVanderPauwtechnique.ThereareerrorsassociatedwithcontactsizeeectsandshortingoftheHallvoltage[ 154 ].Ideally,contactsshouldbekeptassmallaspossible.Asmentionedabove,theVanderPauwtechniqueallowsarbitraryplacementofcontactstodeterminetheresistivityofthesample.Onthecontrary,thisisnottruefortheHallvoltage.Misalignmentofthevoltageleadswillcauseadditionalerror.Deviationsfromperpendicularalignmentwiththeappliedcurrentcancausemixingoftheperpendicularandparallelcomponentsofthecreatedvoltage.ThisdistortsthetrueHallvoltagebeingmeasured.NotethatthisisalsotrueforHallbridgesamples;however,thexedgeometryofthebridgelimitsthemisalignment.Current-reversalaveragingtechniquesdescribedlatercanhelpalleviatemisalignmenterror.Anotherdrawbackofthetechniqueismeasurementtime.VanderPauwrequiresa 148

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totalof8voltagemeasurementsaroundthesample|eachsideofthesampleismeasuredunderpositiveandnegativecurrentpolarities|whereasahallbridgepatternedsampleonlyrequires2measurements.Thiscanbeespeciallyproblematicwhenlongdelaytimesareneededforthecircuittosettle,suchaswhenmeasuringhighresistancesamples.SettlingtimeisafunctionoftheRCtimeconstant.Itrepresentsthetimeneededtosaturateanyshuntcapacitanceinthecircuittoobtainanaccuratemeasurementofthecircuitvoltage.ThisideaisrepresentedingureA-3.TheVanderPauwtechniquealsorequiresaswitchingsystemifleadsarehardwiredtothesample.Thissystemcanbemanualorautomatic,andisusedtorecongurethecurrentandvoltageleadsaroundthesamplewithouthavingtoremoveandreplaceeachsamplecontactduringthemeasurement.Smallsignalcurrentsareneededtomakemeasurementsonhighresistancesources.DuringHallandresistivitymeasurements,ohmicheatinginthesampleshouldbekepttoaminimumtoretainmeasurementaccuracyandreducethermalnoise.SincepowerisequaltoresistancemultipliedbythesquareofthecurrentP=I2R,lowcurrentsareneededtokeepthepowerdissipationlow.Typically,keepingthepowerinthesamplebelow1-5mWissuggestedforgeneraltransportmeasurements.Leakagecurrents|currentsthat`leak'intooroutofthecircuitthroughunwantedresistivepathways|candegradetheaccuracyoflowcurrentmeasurements[ 153 ].Thiscanbeacommonerrorinlow-levelmeasurementsusingstandardcables,whereportionsofthesignalcurrentcanleakthroughthecable'sinsulationresistance.Useofaguardingtechniquewithtriaxialcablescanreduceleakagecurrents.Triaxialcablescontainanextraconductorthatsurroundsthesignalcableseparatedbyaninsulatorandliesunderneaththegroundsheathofthecable.Apotentialisdrivenalongtheguardthatmatchesthepotentialonthesignalcable.Thiscreatesaspaceofzeropotentialbetweentheguardandsignalandreducesthedrivingforceforcurrentleakagebetweenthem.Theleakagecurrentcombinedwiththeinputosetcurrentrepresentsthetotalerrorcurrentofthecircuit 149

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[ 153 ].Theopen-loopgainoftheguardalsoreducestheeectsofcablecapacitancebyreducingtheshuntcapacitancechargingtime[ 153 ].TheKeithleyelectronicssupporttheuseofguardingandtheequipmentisconnectedusingtriaxialcables.However,thePPMSdoesnotusetriaxialcablingwhichposesalimitationofusingguardedmeasurementswiththecurrentset-upseethefollowingsectionsformoredetails.A.4HallEectMethodThephysicsoftheHalleectcanbefoundinalmostanysolid-statephysicsorelectronicmaterialssciencebook,aswellasplentyofsourcesontheworldwideweb.It'snotmyintentionofrehashingthisinformationsinceitiseasilyfound,butonlytoprovidethebasicideasandequationsforthemeasurement.ThedrivingforcebehindtheHalleectistheLorentzforceactingbetweenmovingchargecarriersandamagnetic-eldappliedperpendiculartothecarrier'svelocity.TheLorentzforceisgivenby:F=)]TJ/F20 11.955 Tf 9.298 0 Td[(q~vx~BA{1whereqisthecarriercharge.602x10)]TJ/F19 7.97 Tf 6.586 0 Td[(19forelectronsandholes,~visthecarriervelocity,and~Bisthemagneticinduction.Acurrentisappliedtooneendofthesampleandcarriersaredriventhroughthelengthofthesample.Themagnetic-eldisappliedperpendiculartothecarriervelocityandpushesthecarrierstowardstheedgeofthesampleasdictatedbytheLorentzforce.Thisresultsinaslightchargeimbalancebetweenthetwosampleedges.Tomaintaintheowofcurrent,thechargeimbalancecreatesapotentialdropacrossthesampletobalancetheLorentzforce.ThisistheHallvoltage,VH,andisgivenby:VH=IBz nqtA{2wherenisthecarrierconcentration,tisthesamplethickness,andIistheappliedcurrent.Ifthesampleresistanceisknown,thecarriermobility,,canbecalculated: 150

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=RH nA{3whereR)]TJ/F19 7.97 Tf 6.587 0 Td[(1H=nqisdenedastheHallcoecient.However,ifthesampleunderconsiderationiscomposedoftwocarriers,suchasasemiconductorwithasucientnumberofelectronsandholesinthematerial,atwocarriermodelfortheHallcoecientismoreappropriate[ 154 ]:RH=p2p)]TJ/F20 11.955 Tf 11.955 0 Td[(n2n epp+nn2A{4TakingHallvoltagesatthreemagnetic-eldpoints+H1,H=0,and-H1,whereH1issomevalueofmagneticeldisusuallysucienttodeterminetheHallcoecient.ThetwomeasurementsatH1areusedtoaverageoverbotheldpolaritiesandthemeasurementatH=0canbeusedcalculatethecarriermobility.However,IhavefounditextremelybenecialtomeasureHallvoltagesatmanypointsoverafullmagnetic-eldsweep.SincetheHallcoecientisproportionaltothequotientofVH/B,thecoecientcanbecalculatedfromtheslopeofalinetthroughthedata.Thisprovidesabetterstatisticalaverage,averagingovermanypointsratherthanjustthree,andtheaccuracyofthedataiseasilyexaminedbasedonhowwellthepointsfollowalinearpath.ThisprovedparticularlyusefulfordicultsampleswhichexhibitweakHallvoltagesthatstaggerbetweenn-andp-typepositiveandnegativeslope.TakingHallmeasurementsatvarying-eldsisalsonecessaryforstudyingtheanomalousHalleectAHEinmagneticsamples.Infact,sensitivemagneticpropertiescanevenbedeterminedinsamplesthatshowalargeAHEcomponent.A.5LimitationsandTipsforBetterMeasurementsThereareseverallimitationsofthecurrentset-upthatcoulduserenement.ThemaindicultywehaveexperiencedismeasuringHallvoltagesonourp-typeZnOlms.Thedopedholesinthesematerialsarehighlycompensatedbytheintrinsicelectronconcentration,andbothcarriertypesarepresentinthematerial.TheHallcoecient 151

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forthisconditionwasgivenabove.Themobilityofholesinthesematerialsissmall,andtheconditionforp-typematerial,p2p>n2n,isdiculttofulll.ThesesamplesproduceaveryweakHallvoltage,evenateldsof7T,andareextremelydiculttomeasureaccurately.Theelectrometershavearesolutionof10V,whichdoesnotprovideenoughsensitivityformeasuringtheweakHallvoltage.Typicallymeasurementsarenoisyandteeterbetweenn-andp-typeconductivity.Replacingtheelectrometerwithamoresensitivevoltmeter,suchasananovoltmeterwithnanovoltsensitivity,mayprovideenhancedmeasurementaccuracyforthesedicultsamples.The236SMUisaconvenienttoolbecauseitprovidessourcingandmeasuringcapabilitiesinonecost-eectiveunit.However,thisalsoprovidesaslightlimitationwhenmeasuringhighly-resistivesamples.TheSMUisnotcapableofperformingmeasurementsin4terminalmodewhenthepotentialdierencebetweenthesourceandsenseleadsexceeds4volts.ThisiseasilyseenduringameasurementbecauseeveryVanderPauwvoltagemeasurementaroundthesampleconvergestothesamevalue.Wheneverthiswasencountered,adigitalmultimeterwasborrowedfromanotherpieceofequipmenttoperformthemeasurements.SincetheotherDMMwasisolatedfromtheSMU,thisproblemcouldbealleviated.AnadditionallimitationcomesfrominterfacingtheexternalKeithleyelectronicstothesampleinsidethePPMS.ThewiresinsidethePPMSthatconnecttheexternalportonthePPMStothesamplexturearenottriaxial.ThereforeadiscontinuityexistsbetweentheguardedtriaxialcablesfromtheKeithleyelectronicstothesamplexture|theguardisnotcarriedallthewaytothesample.Currently,thediscontinuityincabletypeoccursatacustom-madejunctionbox.Theguardsterminateatthejunctionbox,whereonlythecenterconductortheportioncarryingthesignaliscarriedtothePPMScabling.CreatingaPPMSprobethatusestriaxialcablesandcarriestheguardallthewaytothesampleshouldprovidefurthermeasurementaccuracyforsamplesthatarediculttomeasure. 152

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Misalignmentinthesamplecontactscanaddspuriousvaluestothemeasuredvoltagesduetomixingoftheparallelandperpendicularvoltagecomponents.Asimplewaytohelpseparatethetwocomponentsisbyeldaveragingthecollecteddata.Themagnetoresistanceofasamplewillbeanevenfunctionoftheappliedeld.Sothecurveproducedshouldhavemirrorsymmetryreectedacrossthey-axisonaresistivityvs.eldplot.Therefore,themagnetoresistancecanbeaveragedbyRxx=1 2[Rxx+H+Rxx-H].Onthecontrary,theHallresistivityisanoddfunctionofeld,whichproducessymmetrythroughtheorigin.Oddsymmetryproducesgraphsthatareunchangedaftera180degreerotationaroundtheorigin.TheHalldatacanbeaveragedbyRxy=1 2[RH+H-RH-H].Theseaveragingtechniquescanbeusedforbothsingleeldsweepsandhystereticsweeps,wheretheeldissweptinonedirectionandthensweptbackinthereversedirection.Whenaveragingthehystereticsweepsitisadvisabletoaveragetheinitialdownwardsweepwiththeascendingreturnsweepandvice-versa[ 155 ].Whilethesesimpleaveragingschemescanaidinseparatingthetwocomponents,thedatamaystillbeskewedifonecomponentismuchlargerthantheother. 153

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FigureA-1.TheresistivityandHallmeasurementsystem.ThesampleisplacedinsidethePPMSdewartocontroltheambienttemperatureandmagneticeld.TheelectricaldataiscollectedwithexternalKeithleyelectronics. 154

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FigureA-2.Circuitdiagramsfor2-pointand4-pointresistivitymeasurements.Ina2pt.measurement,thecontactandleadresistancearemeasuredbythevoltmeter.Ina4pt.measurement,thecurrentthroughthevoltmeterapproacheszero,soonlythesamplevoltageismeasured. 155

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FigureA-3.Circuitshuntcapacitanceandsettlingtime.Thecircuitshuntcapacitancethatmustbefullychargedbeforeanaccuratemeasurementofthecircuitvoltage,Vm,ispossible.Thegraphshowsthesettlingtimenecessarytoreachacertainpercentageofthenalvoltage.Thesettlingtimeisgivenasmultiplesofthecircuit'sRCtimeconstant.Adaptedfrom[ 153 ]. 156

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BIOGRAPHICALSKETCHMathewIvillwasborninFt.Myers,Florida.Hewasnamedafterhisgrandfather,whoalsoonlyhadone`t'inhisname.HegrewupandattendedschoolinthecityofCapeCoral,Florida,wherehelivedwithhisparents,RichardandLinda,andolderbrotherDavid.MathewbeganhisundergraduatestudyattheUniversityofFloridain1996theyearofUF'srstnationalchampionshipfootballtitleandfoundhiswayintotheDepartmentofMaterialsScienceandEngineeringhissophomoreyear.In2000,duringhisnalyearofundergraduatestudy,hebeganresearchinDr.DavidNorton'sgroup,growingepitaxialCeO2lmsonInPsubstrates.HelaterearnedhisB.S.inMaterialsScienceandEngineeringinthespringof2001.HespentthefollowingsummeratSandiaNationalLabinLivermore,CA,fabricatingandtestingsamplestostudytheinterfaceadhesionbetweenPMMAandsiliconwhilealsoenjoyingthebeautifulnationalparksaroundCalifornia.UponhisreturnasagraduatestudenttotheUniversityofFlorida,hebeganstudyingtheeectsoftransition-metaldopinginoxidesfortheeldofspintronics.HewasalsofortunateenoughtotakesometimeofromhisstudiestoliveinLondon,Englandfor5months.Hehaspresentedat7professionalconferences,andwasawardedtheGrandPrizeinMaterialsScienceattheFloridaChapteroftheAmericanVacuumSociety,Orlando,twoyearsinarowforhispostersontransition-metaldopedZnO.Inhisfreetimeduringgraduateschool,hewasamemberoftheGatorKendoClub.HehasalsoenjoyedplayingintramuralsoftballwiththeBraisedCabbages. 171