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Record for a UF thesis. Title & abstract won't display until thesis is accessible after 2008-02-29.

Permanent Link: http://ufdc.ufl.edu/UFE0020109/00001

Material Information

Title: Record for a UF thesis. Title & abstract won't display until thesis is accessible after 2008-02-29.
Physical Description: Book
Language: english
Creator: Erie, Jean-Marie George
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2007

Subjects

Subjects / Keywords: 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

Statement of Responsibility: by Jean-Marie George Erie.
Thesis: Thesis (Ph.D.)--University of Florida, 2007.
Local: Adviser: Norton, David P.
Electronic Access: INACCESSIBLE UNTIL 2008-02-29

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2007
System ID: UFE0020109:00001

Permanent Link: http://ufdc.ufl.edu/UFE0020109/00001

Material Information

Title: Record for a UF thesis. Title & abstract won't display until thesis is accessible after 2008-02-29.
Physical Description: Book
Language: english
Creator: Erie, Jean-Marie George
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2007

Subjects

Subjects / Keywords: 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

Statement of Responsibility: by Jean-Marie George Erie.
Thesis: Thesis (Ph.D.)--University of Florida, 2007.
Local: Adviser: Norton, David P.
Electronic Access: INACCESSIBLE UNTIL 2008-02-29

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2007
System ID: UFE0020109:00001


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1 PULSED LASER DEPOSITION OF DOPED ZnO AND (Mg, Zn)O FILMS FOR OPTOELECTRONIC APPLICATIONS By JEAN-MARIE GEORGE ERIE A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLOR IDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2007

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2 2007 Jean-Marie George Erie

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3 To my dearly beloved parents whose uncondition al love, endless encouragement, support and patience made this possible.

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4 ACKNOWLEDGMENTS My gratitude goes to my committee chair fo r persevering with me as my advisor throughout the time it took me to complete my resear ch and dissertation. I am also grateful to my mentor for her support and enthusiasm. I thank the members of my dissertation committee for their contributions and their good-natured support. They have generously given their time and expertise to better my work. I must acknowledge the Army Research O ffice, the Naval Office of Research and Scientific Development, the National Science F oundation and the College of Engineering at The University of Florida for their financial support and fellowships. I need to express my gratitude and deep appreciation to my closest associates whose friendship, hospitality, knowledge and wisdom have supported, en lightened, and entertained me over the many years of our friendship. They have consistently helped me maintain a healthy perspective on life reality. I also thank my brothers and my extended family for their support and care.

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5 TABLE OF CONTENTS page ACKNOWLEDGMENTS...............................................................................................................4 LIST OF TABLES................................................................................................................. ..........8 LIST OF FIGURES................................................................................................................ .......11 ABSTRACT....................................................................................................................... ............20 1 INTRODUCTION..................................................................................................................21 2 MATERIALS AND METHODS...........................................................................................23 The Approach................................................................................................................... ......23 Substrate Preparation.......................................................................................................23 Experimental Set-up........................................................................................................23 Pulsed Laser Deposition........................................................................................................ .24 Characterization............................................................................................................... .......25 Hall Measurement...........................................................................................................26 Variable Field Hall Measurement...................................................................................27 Temperature-Dependent Hall Measurements..................................................................29 Photoluminescence..........................................................................................................31 3 LITERATURE REVIEW.......................................................................................................35 ZnO and Related Alloys as Wide Band Gap Materials..........................................................35 Applications................................................................................................................... .........35 ZnO Crystal Structure......................................................................................................36 ZnO Band Structure.........................................................................................................37 Doping ZnO..................................................................................................................... .......37 Doping Rules...................................................................................................................37 Doping ZnO n-type..........................................................................................................38 Doping ZnO p-Type.............................................................................................................. .39 Compensation..................................................................................................................39 Native Defects in ZnO.....................................................................................................39 Hydrogen Doping............................................................................................................41 State of the Art for p-Type Doping.................................................................................42 Co-Doping Method..........................................................................................................44 Phosphorous, Arsenic and Antimony Doping.................................................................45 Epitaxial Growth of ZnO Films.......................................................................................46 Alloying ZnO with MgO........................................................................................................49 Crystal Structure and Stability.........................................................................................52 MgxZn(1-x)O Band Gap....................................................................................................53 Native Defects.................................................................................................................54 Applications and Advantages..........................................................................................55

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6 Epitaxial Growth of (Mg, Zn)O.......................................................................................56 Challenges..................................................................................................................... ..57 4 ZINC OXIDE DOPED WITH ARESEN IC FROM ARSENIC TRI-OXIDE........................61 Results and Discussion......................................................................................................... ..61 Summary........................................................................................................................ .........67 5 ZINC OXIDE DOPED WITH ARESENIC FROM Zn2As3 AS THE SOURCE OF As.......74 Results and Discussion......................................................................................................... ..74 ZnO Films Doped with 0.02 at % As on c-Sapphire.......................................................74 ZnO Films Doped With 0.2 at % As on c-Sapphire........................................................79 ZnO films Doped with 2 at % As grown on ZnO and MgO buffer.................................86 Effects of Target Concentration......................................................................................94 Effects of Persistent Photoconductivity...........................................................................95 Effects of MgO Alloying on ZnO:As0.002........................................................................96 Mg.10Zn.90O Doped with 0.02 at % As............................................................................99 Summary........................................................................................................................ .......100 6 STABILITY OF THE ARSE NIC DOPED ZnO FILMS.....................................................153 Results and Discussion.........................................................................................................153 Low Temperature RTA study...............................................................................................153 RTA Temperature Study...............................................................................................153 RTA Time Study...........................................................................................................155 High Temperature Annealing...............................................................................................157 Nitrogen RTA at 800C.................................................................................................157 Effects of Growth Temperatur e on Post-Annealing Properties.....................................158 Effects of Growth Pressure on Film Properties After Ar Annealing.............................160 Effects of Target Processing..........................................................................................161 Effects of Arsenic Source..............................................................................................164 Effects of Aging............................................................................................................164 Effects of Photoresist Coa ting on Transport Properties................................................165 Summary........................................................................................................................ .......166 7 OPTICAL CHARACTERIZATI ON OF OXIDIZED ZnNxO(1-x) FILMS...........................206 Results and Discussion.........................................................................................................206 ZnO Films Doped with 0.2 and 2 at % N on c-Sapphire...............................................206 Thermal Oxidation of ZnNxO(1-x) Films.........................................................................208 Summary........................................................................................................................ .......211

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7 8 GROWTH AND CHARACTERIZATION OF Nb AND Ta DOPED ZnO........................224 Results and Discussion.........................................................................................................224 ZnO films Doped with 0.1 at % Nb and Ta...................................................................225 ZnO films Doped 1 at % Ta..........................................................................................231 Thermal Stability of the 1 at % Ta Doped ZnO films...................................................234 Summary........................................................................................................................ .......240 9 FUTURE WORK..................................................................................................................268 APPENDIX: CHARACTERIAZATI ON ON UNDOPED ZnO FILM.......................................269 Photoluminescence of Single Crystal ZnO...........................................................................269 Photoluminescence of Undoped ZnO film...........................................................................269 Hall measurements of Undoped ZnO films..........................................................................270 LIST OF REFERENCES.............................................................................................................275 BIOGRAPHICAL SKETCH.......................................................................................................285

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8 LIST OF TABLES Table page 3-1. Properties of some compound semiconductors..................................................................58 3-2. n-type dopant atoms...................................................................................................... .....59 3-3. p-Type dopant candidates with th eir valence and ionic radius..........................................59 3-4. Literature survey for p-type, N doped ZnO.......................................................................60 3-5. Properties of MgO and ZnO..............................................................................................60 4-1. Room temperature Hall data calculated from....................................................................68 5-1. Hall data for the films doped with 0.02 at % of As.........................................................103 5-2. Hall data for the films doped with 0.02 at % of As and grown at 30 mTorr of oxygen..103 5-3. Hall data for the films doped with 0.2 at % of As...........................................................103 5-4. Room temperature Hall data fo r the films doped with 0.2 at % of As............................103 5-5. Summary of the structural, optical and transport properties............................................104 5-6. Summary of the averag e Hall measurement data............................................................104 5-7. Hall measurement data for samples dope d with 2 at % As and grown at 500C............104 5-8. Surface roughness, in nm, of several films doped with 2 at % As..................................105 5-9. Transport properties of the ZnO films doped with 0.2 and 2.0 at % As..........................105 5-10. Transport properties of the ZnO films doped with 0.2 at % As.......................................105 5-11. Transport properties of the undoped Mg0.05 Zn0.95O........................................................105 5-12. Transport properties of the Mg0.05Zn0.95O:As0.002 films grown on sapphire....................106 6-1. Transport properties of th e films doped with 2 at % As..................................................170 6-2. Resistivity of the films doped w ith 2 at % as from zinc arsenide....................................170 6-3. Before and after nitrogen rapid thermal anneal at 200C................................................170 6-4. Hall data for the as-grown and RTA in nitrogen at 800C Films....................................171 6-5. Transport properties of f ilms grown from Targets 1 and 2..............................................171

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9 6-6. Hall measurements for the samples A Ire RTA for 5 s in oxygen...................................171 6-7. Hall data for the film doped with 0.02 at % As on bare sapphire....................................172 6-8. Hall measurements for the film dope d with 0.02 at % As on bare sapphire....................172 6-9. Hall data for the film doped with 0.2 at % As.................................................................172 6-10. Hall measurements for the film doped with 0.2 at % As on bare sapphire......................173 6-11. Transport properties of the films doped with 2 at % As and grown 500C.....................173 6-12. 800C oxygen RTA of the film 500 C and 5 mTorr of oxygen for 2 hrs..........................173 6-13. Effects of an 800C, O2, 5 seconds RTA on the transport properties..............................173 6-14. Transport properties of the films gr own from different targets at 500C........................174 6-15. Transport properties of 10 mi n Ar annealed films at 600C............................................174 6-16. Transport properties of as-grown and aged 2 at % As doped samples............................174 6-17. Transport properties of as-grown and PR covered films.................................................175 7-1. Room temperature Hall data for the 0.2 and 2 % doped films........................................213 7-2. Room temperature Hall data for the as-grown and thermally..........................................213 8-1. Room temperature Hall measurements data for the 0.1 at % Nb ZnO films...................242 8-2. Room temperature Hall measurem ents for the 0.1 at % Nb............................................242 8-3. Room temperature Hall measuremen ts data for the 0.1 at % Ta ....................................242 8-4. Room temperature Hall measurem ents data for the 1.0 at % Ta.....................................242 8-5. Temperature dependent Hall data for the 1.0% Ta doped film........................................243 8-6. FWHMs of the High resolution -rocking curves shown in Figure 8-18.......................243 8-7. Hall measurements for the 1.0 at % Ta doped ZnO films...............................................243 8-8. Hall measurements for the ZnO:Ta films........................................................................244 8-9. Hall data for the films grown at 90 mTorr and annealed in air at 1000C......................244 8-10. Resistivity data for the films annealed in 1 atm of flowing oxygen................................244 8-11. Hall data for the films ann ealed in 1 atm of flowing Ar..................................................244

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10 8-12. Summary of the effects of a nnealing on the electrical properties....................................245 8-13. Surface roughness....................................................................................................... .....245 8-14. Hall data for an undoped ZnO film grown on a thin MgO buffer ..................................245 8-15. Effects of cooling pressure on the electrical properties of 1.0 at % Ta ..........................245 8-16. Effects of 5 sec oxygen rapid thermal a nnealing on the electrical properties ................245

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11 LIST OF FIGURES Figure page 2-1. Our PLD system set-up.................................................................................................... ..32 2-2. Example of variations of Hall voltage with applied magnetic field .................................32 2-3. Energy level structure of a compensated p-type semiconductor.......................................33 2-4. Experimental set-up for colle ction of photoluminescence spectra....................................33 2-5. Simplified electron-holes recombin ation diagram inside the band gap.............................34 4-1. Powder XRD pattern for films grown at 400 C in 0.3, 3 and 30 mTorr of O3/O2..............69 4-2. Powder XRD pattern for films grown at 600 C in 15 and 30 mTorr of O3/O2...................69 4-3. Changes in surface roughness and c-axis length with growth ozone/oxygen....................70 4-4. Room temperature photoluminescence spectra.................................................................71 4-5. Low temperature photolu minescence spectra at 16K........................................................71 4-6. Room temperature variable magnetic fiel d Hall coefficient. The slope of the linear........72 4-7. Changes of transport properties with growth ozone/oxygen pre ssure for the films..........73 5-1. Powder X-ray diffraction of: A600C in 3 mTorr of O3/O2; B300C.........................107 5-2. Room temperature photoluminescence of the 0.02 at % As ZnO films.........................107 5-3. Photoluminescence at 300Kof the 0.02 at % As doped ZnO films grown on sapphire...108 5-4. Photoluminescence spectrum at 16K for an undoped film grown at 700C....................108 5-5. Low temperature photoluminescence spectra for the 0.02 at % As doped ZnO film......109 5-6. Low temperature PL spectra for the 0.02 at % As doped ZnO film................................109 5-7. Photoluminescence at 16K spectrum fo r the 0.02 at % As doped ZnO film grown........110 5-8. Photoluminescence at 16K for the 0.02 at % As doped ZnO film grown at 500C.........110 5-9. Photoluminescence at 16K for the 0.02 at % ZnO:As grown at 600C and 30 mTorr ...111 5-10. Resistivity, carrier concentra tion and mobility of the films doped..................................111 5-11. Variable magnetic field Hall measur ement of 0.02 at % As doped ZnO film.................112

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12 5-12. Changes in resistivity, carrier concen tration and mobility for the films doped...............112 5-13. Variable magnetic field Hall measurement for the film doped with 0.02 at % As..........113 5-14. Effects of growth temperature on the caxis length for the 0.02 at % As doped ZnO.....113 5-15. Powder X-ray diffraction of the films doped with 0.2 at % As.......................................114 5-16. X-ray diffraction of the ZnO films dope d with 0.2 at % As and grown on sapphire......114 5-17. C-axis length for the ZnO films doped with 0.2 at % As as a function of growth..........115 5-18. Photoluminescence spectra for the ZnO film s doped with 0.2 at % As and grown at 3..115 5-19. Photoluminescence of the ZnO films doped with 0.2 at % As and grown at 30 mTorr..116 5-20. 16K photoluminescence spectra for the ZnO film doped with 0.2 at % As....................116 5-21. PL at 16K for the ZnO film doped with 0.2 at % As and grown at 500C......................117 5-22. PL at 16K for the ZnO film doped with 0.2 at % As and grown at 600C......................117 5-23. Transport properties of the ZnO films doped with 0.2 at % As.......................................118 5-24. Transport properties of the ZnO film s doped with 0.2 at % As and grown at 30............118 5-25. Effects of growth oxygen pressure on the c-axis spacing................................................119 5-26. High resolution rocking curves fo r ZnO films doped with 2 at % As.............................119 5-27. Effects of growth pressure on the crys talline qualities (FWHM) of the ZnO films........120 5-28. High resolution rocking curves for ZnO films doped with 2 at % As and grown...........120 5-29. Rocking curves for ZnO films doped with 2 at % As and grown at 500C.....................121 5-30. 300K PL spectra for ZnO films doped with 2 at % As and grown on ZnO.....................121 5-31. PL at 300K spectra for ZnO films do ped with 2 at % As and grown on MgO................122 5-32. Room temperature PL for ZnO films dope d with 2 at % As and grown in 150 mTorr...122 5-33. Low temperature Photo Luminescence sp ectra of selected 2 at % As doped ZnO.........126 5-34. Dependence of Acceptor-bound exciton (AX), Donor-Acceptor Pair transition...........126 5-35. Dependence of Acceptor-bound exciton (AX) emission energy on growth oxygen......127 5-36. Dependency of Acceptor binding energy (Eb A) on acceptor-bound exciton (AX).........127

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13 5-37. Dependency of acceptor binding energy (Eb A) on growth pressure (a) and hole............128 5-38. Resistivity as a func tion of growth Press. and grown on ZnO buffer layer.....................129 5-39. Resistivity as a function of gr owth temperature for the ZnO films.................................129 5-40. Resistivity as a function of growth pressure for the ZnO films doped with 2 at %.........130 5-41. Hall voltages against applied magnetic fi eld for ZnO films doped with 2 at % As........130 5-42. Carrier density and resistivity Vs reci procal temperature for the ZnO films doped........131 5-43. Charge balance equation analysis of the hole concentration for the ZnO samples.........132 5-44. Plots of surface roughness as a function of growth temperature for the ZnO film..........133 5-45. Changes in carrier concen tration, resistivity, Hall coeffi cient and Hall mobility as a....133 5-46. Dependency of hole concentration on Acceptor binding energy (Eb A)...........................134 5-47. C-axis length of the ZnO films doped w ith 2 at % As and grown at 500C at 5, 50.......134 5-48. Dependency of Acceptor binding energy (Eb A) on Arsenic concentration......................135 5-49. Hall voltage of ZnO:As0.02 films doped with 0.2 and 2.0 at % As grown at 500C........135 5-50. Transport properties of the ZnO films doped with 0.2 and 2 at % As grown at 500C...136 5-51. PL at 300K spectra of the ZnO films doped with 0.2 and 2.0 at % As and grown at......136 5-52. Hall coefficient for the ZnO films doped with 0.2 at % As and grown at 500C and.....137 5-53. Effects of Persistent Photoconductivity on the resistivity when the ZnO films..............137 5-54. Hall voltage the ZnO sample doped with 2 at As and grown at 500C in 5 mTorr.........138 5-55. Variation of Hall voltage with applie d magnetic field after the ZnO sample doped.......138 5-56. Variations of Hall voltage after the ZnO sample doped with 2 at As and grown............139 5-57. Resistivity as a functi on of time in the dark and measured in AIR for the ZnO.............139 5-58. Variation of Hall voltage with applied magnetic field for the ZnO sample doped with.140 5-59. Hall voltage with applie d magnetic field for the ZnO sample doped with 2 at % As.....140 5-60. Hall voltage tor the ZnO sample doped with 2 at % As grow n at 500C and 100...........141 5-61. Powder XRD of the Mg0.05Zn0.95O:As0.002 films grown in O2 at 500C..........................142

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14 5-62. Powder XRD of the Mg0.05Zn0.95O:As0.002 films grown at 600C and 60 mTorr of O2...142 5-63. c-axis spacing Mg0.05Zn0.95O:As0.002 films grown at 500C on sapphire.........................143 5-64. c-axis spacing as a function of growth temperature for the Mg0.05Zn0.95O:As0.002..........143 5-65. -rocking curves of the Mg0.05Zn0.95O:As0.002 films grown at 500C on in 30, 60..........144 5-66. FWHMs of the -rocking curves of the Mg0.05Zn0.95O:As0.002 films grown at 500C....144 5-67. Room temperature luminescence of the Mg0.05Zn0.95O:As0.002 films grown...................145 5-68. PL at 25K of an undoped Mg0.05Zn0.95O films grown in 60 mTorr of O2........................145 5-69. Photoluminescence at 25K of the Mg0.05Zn0.95O:As0.002 films grown 500C and 60.......146 5-70. Photoluminescence at 25K of the Mg0.05Zn0.95O:As0.002 films grown 500C and 90.......146 5-71. PL of the Mg0.05Zn0.95O:As0.002 films grown at 500 C and 120 mTorr of O2..................147 5-72. Resistivity against O2 growth pressure for the Mg0.05Zn0.95O:As0.002 films grown........147 5-73. Dependence of Hall voltage on applied magnetic field for the Mg0.05Zn0.95O:As0.002 ...148 5-74. Hall voltage Vs applied magnetic filed for the Mg0.05Zn0.95O:As0.002 films...................148 5-75. Powder XRD patterns of the Mg0.1Zn0.9O:As0.002 films grown in 0.1 mTorr of O2.........149 5-76. High resolution w-rocking curve of the Mg0.1Zn0.9O:As0.002 deposited at 400C...........149 5-77. Backscattered electron image of the Mg0.1Zn0.9O:As0.002 deposited at 400C.................150 5-78. Backscattered electron image of the Mg0.1Zn0.9O:As0.002 deposited at 700C.................150 5-79. Variations of Mg content with substrate temperature......................................................151 5-80. Variations of c-axis of the films and estimated band gap with Mg content....................151 5-81. Variations the band gap with c-axis of the films.............................................................152 5-82. Transport properties of the films as a function of Mg content and growth.....................152 6-1. High-resolution omega rocking curves for the films doped with 2 at % As : A)............176 6-2. Photoluminescence spectra of the films doped with 2 at % As from zinc arsenide........176 6-3. Photoluminescence at 30K of the films dope d with 2 at % As from zinc arsenide.........177 6-4. Resistivity of the films doped with 2 at % As and grown at 500C in 50 mTorr............177

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15 6-5. Field dependent Hall voltage with annea ling temperature of the films doped with 2.....178 6-6. Effects of annealing temperature on carrier density of the films doped..........................179 6-7. Hall mobility as a function of the 60 seconds N2 RTA temperature...............................179 6-8. Photoluminescence spectra of the fi lms doped with 2 at % As and RTA in N2..............180 6-9. PL at 30K of the films dope d with 2 at % As and RTA in N2 at 200C...........................181 6-10. Temperature dependence of the photoluminescence from 30-350K...............................182 6-11. Resistivity as a function of annealin g time for the films doped with 2 at % As..............182 6-12. Changes in resistivity and carrier concentration as a function of RTA...........................183 6-13. Changes in carrier mobility as a function of RTA time for the 2 at % As.......................183 6-14. As grown room temperature photoluminescence of the film 2 at % As..........................184 6-15. Room temperature photoluminescence spectra of the annealed 2 at % As.....................184 6-16. 300K photoluminescence spectra for as-grown samples A and B...................................185 6-17. Resistivity as a function of O2 RTA temperature for sample A......................................185 6-18. Changes in resistivity as a function of growth temperature for the film doped with.......186 6-19. Room temperature photoluminescence of the as-grown films doped with 0.02 at %.....186 6-20. Photoluminescence of the 0.02 at % ZnO:As films grown in 3 mTorr of O3/O2............187 6-21. Photoluminescence at 300K of the 0.02 at % ZnO:As films grown on sapphire............187 6-22. Photoluminescence spectra at 20 K of the as-grown and annealed films doped.............188 6-23. Transport properties for the as-grown and the 10 min 800C O2 annealed films............189 6-24. PL at 300K spectra of the 0.2 at % ZnO:As films grown in 3 mTorr of O3/O2 and........189 6-25. PL at 16K spectrum of the 0.2 at % Zn O:As films grown on sapphire and at 600C......190 6-26. PL at 300K spectrum of the 0.2 at % ZnO:As films in 30 mTorr of O2..........................190 6-27. Transport properties of: (a) as-grown a nd (b) annealed in flowing oxygen at 800C....191 6-28. Plot of resistivity as a function of growth Pressure for the films doped..........................192 6-29. Hall coefficient of the films doped w ith 2 at % As and grown at 500C and..................192

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16 6-30. Hall voltage of the film s doped with 2.0 at % of As and grown at 500C and 50...........193 6-31. Hall coefficient for the films doped with 2.0 at % of As and 150 mTorr........................193 6-32. PL at 300K spectra of films doped w ith 2 at % As and grown on ZnO buffer................194 6-33. PL at 300K spectra of films deposited at 500C from the ZnO target doped..................194 6-34. Transport properties of the films grown at 500C on MgO buffer..................................195 6-35. Transport properties of the films gr own from a target doped with 2 at % As.................195 6-36. Resistivity of films O2 RTA at 800C.............................................................................196 6-37. Carrier concentration and mobility of the films rta at 800C in oxygen..........................196 6-38. Effects of aging on: (a ) Variable Hall measurements and (b) low temperature..............197 6-39. Hall measurements of an as-grown and aged (D in Table 6-16) film doped...................198 6-40. Variable Hall measurements of an as-grown and aged (F in Table 6-16).......................198 6-41. Hall measurements of the as-grown Sample G................................................................199 6-42. Variable Hall measurements of Sample G right after curing the PR...............................199 6-43. Hall Variable Hall measurements for Sample G 48 hrs after curing the PR...................200 6-44. Variable Hall measurements for Sample H right after curing the PR..............................200 6-45. Variable Hall measurements for Sample H 24 hrs after curing the PR...........................201 6-46. Hall measurements of Sample H 48 hrs after curing the PR. The Hall data....................201 6-47. Variable Hall measurements of as Sample I after PR cure..............................................202 6-48. Variable Hall measurements of Sample I 48 hrs after PR cured.....................................202 6-49. Variable Hall measurements of as-grown Sample J........................................................203 6-50. Variable Hall measurements of Sample J right after curing the PR................................203 6-51. Variable Hall measurements of Sample J 48 hrs after curing the PR..............................204 6-52. Hall measurements of as grown Sample K......................................................................204 6-53. Variable Hall measurements of Sample K right after curing the PR...............................205 6-54. Variable Hall measurements of Sample K 48 hrs after curing the PR.............................205

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17 7.1. Powder X-ray diffraction pattern for th e film grow at 500C and 120 mTorr of N2.......214 7.2. Photoluminescence spectrum at 300K of an undoped single crystal ZnO substrate.......215 7.3. Photoluminescence spectra at 300K for the films grown in oxygen...............................215 7.4. Photoluminescence spectra at 300K fo r the film doped with 2 % nitrogen.....................215 7.5. Photoluminescence spectrum at 16 K for the film doped with 0.2 % N..........................216 7.6. Photoluminescence spectrum at 16 K for the film doped 2% N and grown at 500C.....216 7.7. Photoluminescence at 16 K for the film doped with 2 % N and grown at 500C............217 7.8. Photoluminescence at 16 K for the film doped with 2 % N and grown at 700C............217 7.9. Photoluminescence spectrum at 16K for th e film grown at 500 C in 120 mTorr of N2..218 7.10. Photoluminescence spectrum at 16K for an undoped film grown at 700C....................218 7.11. As grown and for the film grown fr om the target doped with 10 at % N........................219 7.12. X-ray diffraction pattern of the film grown from the ta rget doped with 10 at % N........220 7.13. X-ray diffraction pattern of the film grown from the ta rget doped with 10 at % N2.......220 7.14. Photoluminescence at 300K of the before and after 35 minutes annealed at 700C.......221 7.15. Photoluminescence before and after 35 minutes annealed at 700C for the film............221 7.16. Photoluminescence at 20K of the f ilm grown at 500C and 1 mTorr of N2....................222 7.17. Photoluminescence at 20K of the f ilm grown at 500C and 120 mTorr of N2................222 7.18. Photoluminescence of the film grown at 500C and 120 mTorr of N2 from the 10 %....223 7.19 Changes in Hall voltage for the film doped with 10 at % N at 500C.............................223 8-1. C-axis lengths of the ZnO films dop ed with 0.1 at % Nb deposited at 500C.................246 8-2. Effects of growth temper ature on the room temperature.................................................247 8-3. Room temperature photoluminescence spectra of the 0.1 at % Nb doped ZnO..............247 8-4. 20K photoluminescence spectra for selected films doped with 0.1% at Nb....................248 8-5. PL at 300K spectra for selected films doped with 0.1% at Ta.........................................249 8-6. PL at 16K of the film doped with 0.1% at Ta deposited at 500C and 30 mTorr of O2..249

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18 8-7. Changes in resistivity [opaque s quares and solid line], Carrier density..........................250 8-8. Resistivity [opaque squa res and solid lin e in (a)],...........................................................251 8-9. 20K photoluminescence spectra of. undoped ZnO film (dotted line); 0.1 at % Ta.........252 8-10. Powder X-ray diffraction of the 1.0 at % Ta doped ZnO grown on thin MgO...............252 8-11. Surface roughness as a function of growth pressure for the 1.0 at % Ta doped..............253 8-12. PL at 300K of selected films dope d with 1 at % Ta on MgO buffer layer......................253 8-13. Resistivity (a), Hall coefficien t (b), carrier concentration (c)..........................................254 8-14. Logarithmic plots of resistivit y versus reciproc al temperature........................................255 8-15. Hall coefficient versus temperatur e for the film grown at 500C and 30........................255 8-16. Arrhenius plot of carrier concentr ation versus reciprocal temperature...........................256 8-17. Carrier mobility as a function of te mperature for the film grown at 500C and.............257 8-18. High resolution rocking curves: (A) As grown; 500C and 1mTorr of oxygen; (B).......257 8-19. Powder XRD pattern of the ZnO:Ta0.001 film annealed in air at 1000C for 60 min.......258 8-20. Powder XRD of: (1) annealed in air at 900C for 5 min, (3)..........................................258 8-21. PL of the ZnO:Ta0.001 films grown at 500C and 1 mTorr of O2 and annealed in air......259 8-22. PL for the ZnO:Ta0.001 films grown at 500C and 90 mTorr of O2..................................259 8-23. Resistivity as a functi on of air annealing temperatur e for the films grown at 1..............260 8-24. Resistivities as a function of air 1000C annealing time for the ZnO:Ta0.001..................260 8-25. Resistivities as a function of oxyge n annealing temperature for the ZnO:Ta0.001 ...........261 8-26. PL spectra of the films annealed in flowing oxygen (1 atm) for 5 min...........................261 8-27. Plots of Carrier density (uppe r), Carrier Mobility (middle)............................................262 8-28. PL spectra of the films annealed in flowing argon (1 atm) for 5 min..............................263 8-29. Effects of annealing on the films surfaces.......................................................................263 8-30. Effects of cooling pressure on the room temperature PL of 1.0 at % Ta.........................264 8-31. Effects of 5 sec oxygen rapid thermal annealing on the electrical properties.................264

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19 8-32. Comparison between the low temp erature photoluminescence spectra..........................265 8-33. As grown film AFM image..............................................................................................266 8-34. AFM image for the film annealed at 1000C for 5 min...................................................266 8-35. AFM image of the film annealed at 1000C for 30 min..................................................266 8-36. AFM image of the film annealed at 1000C for 60 min..................................................267 A-1. Photoluminescence spectrum at 300K of an undoped single crystal ZnO substrate.......271 A-2. Photoluminescence spectrum at 25K of an undoped single crystal ZnO substrate.........271 A-3. Photoluminescence spectrum at 300K of an undoped film.............................................272 A-4. Photoluminescence spectrum at 16K for an undoped film..............................................272 A-5. Room temperature Hall propertie s of undoped ZnO films grown at 500C....................273 A-6. Room temperature Hall properties of undoped ZnO films..............................................274

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20 Abstract of Dissertation Pres ented to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy PULSED LASER DEPOSITION OF DOPED ZnO AND (Mg, Zn)O FILMS FOR OPTOELECTRONIC APPLICATIONS By Jean-Marie George Erie August 2007 Chair: David P. Norton Major: Materials Science and Engineering I analyzed the effects of doping ZnO films with As, N, Nb and Ta and (Mg, Zn)O films by pulsed laser deposition. For the As doped films, photoluminescence and Hall measurements revealed the films were compensated and compensation depended on dopant concentration. The As related acceptor-bound exciton, acceptor binding energy and thermal activation energy was dependent of dopant content and O2 growth pressure. Binding energy of the As related acceptor varied from 190 meV for the ZnO films doped with 0.02 atomic percent of (at %) As to 90 meV for a films doped with 2 at % As. The plot of acceptor optical binding energy against p1/3 suggests that the binding energy at infinite dilution to be approximately 160 meV. The ZnO films doped with 0.2 at % As doped on MgO buffer layer showed the lowest degree of compensation with resistivity, carrier density and mobility on the order of 71 .cm, 2 x 1016 cm-3 and 2 cm2/(V.s), respectively. N doped films showed acceptor bound emission and N-acceptor binding energy of 160 meV and N doped ZnO optical binding energy did not show any dependence on film N concentration. The donor-bound exciton emission for the Nd and Ta doped films Ire around 3.31 eV and 3.33 eV, respectively. The Mg0.05Zn 0.95O:As0.002 film grown at 500C and 60 mTorr showed p-type behavior, where as, the As doped films with higher Mg content were n-type regardless of growth conditions.

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21 CHAPTER 1 INTRODUCTION The increasing demand for high-power, high-frequ ency, more advanced, reliable and faster electronic devices has pushed the scientif ic community to look into non-conventional semiconducting materials to meet such needs. Over 40 years ago, Rose1 wrote: while germanium and silicon have been purified to approximate the idealized models with which physicists are constrained to work, the magnitude of effort required to effect similar purification and control in compound semiconductors is onl y now beginning to loom in its real and impressive proportions. The abil ity to engineer the band structur es and physical properties of compound semiconductors has enabled us to accomplish what had, otherwise, seemed unachievable. Breakthroughs in high-quality growth of galliu m nitride and its alloys made large-scale manufacturing and process integr ation possible and led to the successful commercialization of blue light emitting diodes (LEDs), blue laser di odes with lifetime in excess of 10000 hours, fieldeffect transistors with poIr outputs of 2.6 W/mm at 10 GHz and ultraviolet photodetectors.2-5 One major problem arises when fabricating these structures; since their st able phase is hexagonal crystal structure, these material s are grown epitaxially via latti ce or domain matching on 6H-SiC (0001), ZnO(0001), sapphire ( -Al2O3) and Si(111).6 Because of the lack of native substrates, devi ces are prone to high defects and dislocations density caused by the misfit between the films and the substrates. Defects adversely degrade the lifetime of the devices in a variety of ways.3 In detectors they cause excess dark current, noise and reduce sensitivity. In light emitting diodes and lasers, they reduce radiative efficiency and operation lifetime. Moreover, they are the main cau se for instability in devices such as fieldeffect transistor.

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22 Zinc oxide has attracted a lot of attention la tely as a possible replacement to the nitrides currently used for high powered electronics, light emitting diodes (LEDs), detectors, sensors, UV and photonic applications. Among zinc oxides ch ief advantages are: high bulk free exciton energy of about 60 meV, availability of large area native substrate ma terials, high radiation resistance and ease of processing. The majo r disadvantage to ZnO is achieving p-type conductivity. The difficulty in attain ing highly p-type has been mainly attributed to the presence of compensating defects such as oxygen vacancies (Vo) and zinc interstitials (Zni). Recently, several groups have reported p-type ZnO via N, As, P doping a nd the N codoping with Ga, In, Al or As. There have been few reports of light emitting diodes.7, 8 9, 10 Despite our advances, p-type ZnO is still not fully underst ood and the lack of knowledge is impeding the progress and use of ZnO as a semiconducting material. My study addressed the electrical structural and optical prop erties of ZnO and (Mg, Zn)O thin films grown by pulsed laser deposition (PLD). As, N, Ta and Nb Ire used as dopants to understand the consequences of their incorporation on the electrical, optical and structural of zinc oxide thin films. In addition to varying impurity concentrations and growth conditions, I used different impurity sources such as arsenic oxid e, zinc arsenide, and zinc di-arsenide to understand the dopant incorporation. I accessed the th ermal stability of the film by rapid thermal and tube furnace anneals.

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23 CHAPTER 2 MATERIALS AND METHODS The Approach Using pulsed-laser depositi on (PLD), epitaxial films Ire deposited in different environments, on different types of substrates, and with different dopant elements including As, N, Ga, Mg, Ta and Nb in order to understand how the electrical, optical and structural properties are affected. Hall measurement using a Van der Pauw set-up was used to extract electronic transport parameters such as resistivity, carri er density and mobility. Temperature-dependent Hall measurements Ire performed to extract carrier thermal activation energies. Photoluminescence and optical transmission spectra Ire collected to underst and the effects of the dopants on the optical properties. Scanning El ectron Microscopy, powder and high resolution xray diffraction Ire used to investigate the mi crostructural properties Surface morphology and chemical composition Ire investigated by means of atomic force microscopy, electron dispersive spectroscopy and X-ray phot oelectron spectroscopy. Substrate Preparation Epi-grade substrates Ire cleaned using succes sive ultrasonic baths of ultrahigh purity trichloroethylene, acetone and methanol. The subs trates Ire then mounted on the substrate holder using silver paint. Before film deposition, thin buffer layers of either ZnO (20-50 nm), MgO (520 nm) or a combination of both Ire grown on the sapphire substrates using a KrF (248 nm) excimer pulsed laser at 1 Hz and energy density of 1-3 J/cm2. The total films thickness ranged from 500 to 700 nm. Experimental Set-up To understand the effects of doping the film with As, N, Mg, Nb and Ta, I investigated the effects of several parameters such as dopa nt concentration, background gas, substrate

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24 temperature, dopant source and th e stability of the doped films. Growing films at high growth pressure adversely affects films quality. As back ground gas pressure increases collisions between the gas molecules and the plume generated by the la ser/target interaction. As a result of the increase collision frequency, the plume broadens and the particles arrive at the surface with much lower kinetic energy. At moderate subs trate temperature, low species kinetic energy translates into low surface mobility and, therefore, islands form at the substrates or growing film interface. Both defects in density and non-uniform ity increase with increasing growth pressure. Because arsenic is very volatile and Zn3N2 is unstable at temperat ure as low as 400K, I am limited to grow the films at 300-400C lower than half the melting temperature of ZnO. Film growth at low temperatures further reduces surfa ce mobility of the already low energy particles, thereby causing further dete rioration in film quality. Pulsed Laser Deposition My PLD system set-up, also called laser molecu lar beam epitaxy, consists of a high power UV KrF (248 nm) excimer laser, optics for sca nning and focusing the beam, an ultrahigh vacuum deposition chamber with various gases available for growth (figure 2.1). This system allowed me to to grow multi-components, complex, stoich iometric or non-stoichiometric films at temperatures ranging from 25 to 1000C in almo st any desirable background gas. The KrF excimer laser has a characteristic wavelength of 248 nm, pulse duration of 30 ns with adjustable pulse rate and pulse energy densit y. Using fused silica lenses, the laser beam was focused on the rotating target at a 45 angle of incidence while maintaini ng a constant energy density in the vicinity of 1.5 J/ cm2 at the target surface. For my experiments, the target-substrate distance was maintained at about 6.0 cm.

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25 The laser-target intera ction is a very complex physic al phenomenon. Once the high energy ultraviolet beam from the excimer laser reaches the target surface, it is absorbed by a thin surface layer resulting in both equilibrium and non-equili brium evaporation of the target surface and producing the depositing flux or plume. Depending on the laser energy density and background gas pressure, the ablation plume can consist of a mixture of highl y energetic (0.1-100 eV) species including molecules, electrons, atoms, ions clusters, micron-sized solid particulates and molten globules. For properly chosen ablation parameters, the plume is mostly atomic and diatomic species. Ablated species travel towards the substrate and deposit in the form of a thin film. The quality of the film depends on the co mposition, optical and top ographical properties of the target, nature and partial pressure of ambien t gas, as Ill as the pr operties of the substrate.11 The deposition chamber was equipped with a loading compartment, allowing me to maintain the chamber at base pressure of 10-8 Torr when not in use. It also had a retractable multiple target rotation system. The substr ate was heated by two quartz lamps and its temperature was controlled electr onically. Before each growth cy cle, the chamber was baked at 200C overnight and then filled with the desired b ackground gas to the desired pressure. Characterization The microstructural properties of the f ilms Ire mostly investigated by X-ray powder diffraction and high resolution X-ray analysis. I used atomic force microscopy to measure the films surface roughness. Composition and chemi cal bonding Ire determined by means of X-ray photoelectron spectroscopy and wavelength disper sive X-ray spectroscopy. Hall measurements and photoluminescence, described below, were used to measure the electrical and optical properties of the films.

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26 Hall Measurement The Hall effect is among the most popular met hods to characterize the transport properties of semiconducting materials. It allows one to de termine the carrier density mobility and majority carrier type. Moreover, temperat ure-dependent Hall measurements yield valuable information on impurities, defects, and scattering mechanisms. Fo r a system with a dominant single carrier type, Hall measurements at a single field and its revers al, to remove magneto resistance, are sufficient to evaluate the transport properties of the ma terial. When the carriers are subjected to low magnetic fields, their energy distribution is un affected. Therefore scattering mechanisms and rates which determine carrier mobility are not altered by the presence of the applied field. The measured Hall coefficient ( RH) and the resistivity ( ) are given by12 ne H r B H R ) ( (2-1) And H H H hR ne r (2-2) The conductivity ( ) tensor is given by 2 21 B nexx (2-3) and 21 B B nexy (2-4) where n is the carrier concentration, H is the Hall mobility, rH is the Hall scattering factor, which depends on the scattering mechanisms, e is the charge of an electron, and B is the

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27 magnetic field. For a single crystal, rH is estimated to be around un ity. Since the Hall coefficient and resistivity are independent of the applied magnetic field, measur ements at the +/ fields are sufficient as long as sample inhomogeneity can be ignored. Carrier mobility, which is governed by the sca ttering mechanism, is re lated to the average relaxation time { } by { m e (2-5) The relaxation time is affected by the different scattering mechanisms that are present in the system and is given by Matthiessens rule: i 1 1 (2-6) where T is the total relaxation time and i the independent scattering mechanism processes present. The major scattering mechanisms present in ZnO can be classified as: Ionized impurity scattering caused by long-ra nge Coulomb interactio ns and unintentional impurities, which create local perturbation in the band edge. Polar LO-phonon, which is caused by the lattice vibrations due the electric field of the moving charge in the polar semiconductor. Piezoelectric scattering aris ing from the electric fiel ds produced with phonons. Native defects and dislocati on scattering due to structur al and compensation defects.12-15 Variable Field Hall Measurement To dope a semiconductor, one introduces impurity atoms with different valence states than what they are substituting for in the lattice. Si nce charge neutrality must be maintained, charge compensation in the form of point defects (Vo and Zni) arise spontaneously. Both dopants and

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28 compensation introduce levels into the band ga p creating conduction in more than one band. Several authors have addresse d this problem by computing RH (B) and (B) from sophisticated models based on the solution to the Boltzmann e quation with an energy -dependent relaxation time. However, the final expressions are not only complex but also can be fitted by more than one models.16, 17 Dave Look13 proposed the use of a two-band mixed conduction model, as shown in equations 1-7 and 1-8: ) (n n p n p ep (2-7) and 2 2 2) (n p n p Hn p e n p R (2-8) where p and n are the densities of holes and electrons, respectively. The film has a positive Hall coefficient (p-type) when 2 2 n pn p and a negative Hall coefficient (n-type) when the opposite occurs. Four-point Van der Pauw Hall measurements zero out the adverse effects of sample geometry and contacts potentials. The DC excita tion currents used must be Ill within the linear ranges of the I/V curves for th e contacts pairs. Sample shape and contacts inhomogeneity are compensated for by employing both positive and negative currents and magnetic fields.18, 19 Different excitation currents Ire used to verify the Hall voltages. The effects of persistent photoconductivity relaxation effects can be minimi zed in the transport measurements by keeping the samples in the dark for at least 24 hours before taking Hall measurements. Upon exposure to light, electronholes pairs are crea ted in transparent semiconductors from photon with energy wider than the band gap of the materials. If the light source is maintained,

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29 both the electron and holes generated by illuminati on become available for conduction. When the light source is removed, the photogenerate d electrons and holes recombine causing photorelaxation. The materials return to its dark conductivity state. The rate of recombination depends on the density of preferred recombina tion sites such disloca tions and density of electronic traps present in the materials. For ZnO, in addition to the electron-holes pairs created by exposure to ultraviolet light, oxygen at the surface and crystallite interfaces can be photodesorbed. The desorbtion of oxygen from the surface causes conductivity to increase. When ZnO oxide is placed in the dark, the re verse occurs. Oxygen is adsorbed at the surface causing conductivity to decrease.20-22 In compensated semiconductors, conduction occurs in two bands, making it very difficult to access the transport properties of the material s since it creates scatte r in the data of the variations of Hall voltage with the applied filed.15, 19 I extracted the Hall coefficients (RH) from the plot of the measured Hall voltage as a func tion of the applied magnetic fields. The slope of the linear fit is the Hall coefficient (RH), as illustrated Figure 2-2, and is used to calculate the transport properties using equations (1) and (2). If the scatter in the data is too large for a reasonable linear fit, I are then unable to access th e electrical properties of the films accurately and it is then labeled indeterminate. This pr ocedure is necessary to unambiguously determine the carrier density, mobility and conduction type Films showing p-type behavior Ire further characterized using a Quantum Design Physical Properties Measurement System allowing measurements at magnetic field up to 7 Tesla a nd temperature down to 4 K using DC currents. Temperature-Dependent Hall Measurements Temperature-dependent Hall measurements allow one to estimate dopant thermal activation energy using the ChargeBalance equa tion. Figure 2-3 is an energy diagram of a compensated p-type semiconductor with Na acceptor impurity with energy level Ea above the

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30 valence band and Nd compensating acceptors. Assuming a nondegenerate, single-donor / singleacceptor model, the hole concentration at any temperature is given by15, 19: i Ai Ai Dn N N p / 1 (2-9) Where i denotes a particular acceptor, NAi is the density of acceptor I, ND is the intrinsic donor density, and Ai is given by the following relationship: Ai = ( g1/ g0) N 'V exp( A/ k ) T3/2exp( EA0/ kT ) (2-10) N 'V = 2(2 mp* k )3/2/ h3, (2-11) and EA = EA0 A T (2-12) The term g1/ g0 is a degeneracy factor referri ng to electrons, not holes and EA is the acceptor thermal activation energy. From equation (9), it can be shown that the hole concentration is: 1 ) ( ) ( 4 1 ) ( 2 1 ) (2 1 2 D A D A A D AN N N N T p (2-13) Expanding equation (13) a llows us to calculate p (T) leads us to: ) 2 / exp( ) (1kT E N N N Nv T pA A A D (2-14) The slope (-Ea/k) of the semi logarithmic plot of p(T) T-3/2 against T-1 allows me to derive the activation energy (Ea). Note that p (T) is the experimental hole concentration at temperature (T) derived from RH at the specific temperature.13, 15, 23

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31 Photoluminescence Photoluminescence is one of the most popul ar methods for optically characterizing semiconductors. It is a non-equilibrium proce ss by which electrons within the material, upon absorption of light, are photo-excited to the cond uction band or other empt y states with alloid transitions. Upon return to the ground state, th e excess energy is released as photons and/or phonons radiation in a direct semiconductor depending on the nature of the defects present in the lattice. The period between abso rption and emission is dependent on the nature of the transition and the lifetime of the state. For band-to-band tr ansition, lifetime is typically extremely short, on the order of 10 nanoseconds. Under special circum stances, however, this period can be extended into minutes or hours. Figure 2-5 shows simplifie d electron-holes recombination diagrams inside the band gap of a direct semiconductor: (a) undo ped and high crystalline quality semiconductor where free electrons in the conduction band reco mbine with free holes in the valence band (Xa); (b) and (c) are lightly n-type and p-type doped semiconductor, respectively showing Xa recombination and impurity band transition a nd recombination, donor-bound exciton (DX) and acceptor-bound exciton (AX); (d) is for a compensated semiconductor where donor and acceptor pair transition (DAP) occurs in addition to Xa, DX and AX. Both DX and AX radiations can be due to levels introduced into the band gap by the departure from stoicheometry such as vacancies and interstials or intentional impurities dopant. The photoluminescence properties of the films, set-up pictured in Figure 2-4, Ire measured using a He-Cd laser (325 nm) with 1 W/cm2 power density as excitation source. Spectra were taken over a wavelength range of 340 to 700 nm. The system employs a 0.3 m scanning grating monochromator with a Peltier-cooled GaAs photomultiplier.

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32 Figure 2-1. Our PLD system set-up. -1.5x104-1.0x104-5.0x1030.0 5.0x1031.0x1041.5x104-3.0x106-2.5x106-2.0x106-1.5x106-1.0x106-5.0x1050.0 5.0x105 Linear fit; RH = slope = 81.436C/cm-3Applied Magnetic Field (G)VH Figure 2-2. Example of variations of Hall voltage with applied magnetic field in a compensated ZnO film. The slope of the linear fit is the Hall coefficient and is used to calculate carrier concentration, ty pe and Hall mobility.

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33 Figure 2-3. Energy level structure of a compensated p-type semiconductor. The band edges, Ev and Ec, and the Fermi Level are defined with reference to an arbitrary Zero. Er is the energy to the rth excited states of acceptor atom below the acceptor energy Ea and Above Ev. Figure 2-4. Experimental set-up for co llection of photoluminescence spectra. Er Ev + Ea } Nd E Conduction Band Valence Band Ev Ec Ea EF = ( Ev+ ) Ec Na 0 } Nd Conduction Band Valence Band Ev Ec Ev + Ea Ea EF = ( Ev+ ) Na 0

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34 Figure 2-5. Simplified electron-hole s recombination diagram inside the band gap of an direct semiconductor: (a) undoped semiconductor where band to band, Xa, recombination occurs; (b) and (c) are ligh tly n-type and p-type dope d semiconductor, respectively showing Xa recombination and impurity band transition and recombination, DX and AX; (d) is for a compensated semiconducto r where DAP occurs in addition to Xa, DX and AX. = E(AX) Xa Ec Ev ED = E(DX) = EDEA = E(DAP) d = Ec Ev = Eg = Xa Ec E v a Xa Ec E v ED = ED Ev = E(DX) b = Ec EA= E(AX) E A Xa Ec E v c

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35 CHAPTER 3 LITERATURE REVIEW ZnO and Related Alloys as Wide Band Gap Materials There is an increasing interest in zinc oxide as an altern ative semiconductor to the III-V nitrides. ZnO has several advantages over Ga N and other wide band gap materials. Among its advantages: (1) its free exciton is bound with hi gher energy (60meV); (2) availability of large area native substrate; (3) it can be easily chemi cally It processed; (4) it has higher radiation damage resistance.3 Its direct band gap of 3.32 eV can be engineered by substituting Zn cations by Cd for a reduction toward the visible or by Mg to extend it to the UV range while maintaining the wurtzite crystal structure. Moreover, epitax ial ZnO films with good qua lity can be deposited at loIr growth temperature than GaN, allowing the possibility of tran sparent junctions on low cost substrates such as glass. 3, 4, 24-26 Table 3-1 is a summary of ZnO properties compared to other semiconductors. Applications ZnO oxide has been used in a wide range va riety of applications due to its chemical, structural, surface and electrical properties. The possibility of engineering its band gap makes it an ideal candidate for blue lasers, light emitti ng diodes, and other photonic application. A. Ohtomo et al. reported stimulated emission due to excitonic recombination at room temperature by optically pumping ZnO nanocryst alline epitaxial thin films.4 Laser action with very low threshold intensity (24 kW/cm2) took place using naturally occurring grain boundaries as cavity mirrors.25 A. Tsukazaki27 have fabricated a ZnO p-i-n hom o-junction grown on lattice match ScAlMgO4 substrate by molecular-beam epitaxy us ing nitrogen as the dopant for the p-type layer. ZnO and ZnMgO have also been deposit ed on p-AlGaN/p-GaN/cplane sapphire by radio-

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36 frequency plasma-assisted molecular beam ep itaxy to produce UV light-emitting diodes based on a pn hetero-junction.28, 29 Other applications include its use in gas and chemical sensors both as thin film and as nanostructure because of its surface ability to selec tively react with or absorb certain gases. Nitin Kumar and contributors engineered nanoscale zinc oxide struct ures, which were effectively used for the identification of the biothreat agent, Bacillus anthracis by successfully discriminating DNA sequence from other genetically related species.30 Furthermore, ZnO properties make it well suited for phosphors, piezoelectric transduc ers, varistors, transp arent conducting films, transparent high power electronic, surface acoustic wave devices and as contacts for solar cells. While many of these devices can be fabricated using other wide band gap semiconductors, zinc oxide is environmentally friendly and is found in abundance in nature. In addition, processing temperatures are typically lower for ZnO than its rivals making it economically more attractive. 24, 31-33 ZnO Crystal Structure ZnO exists in three different crystal structures: a rock salt (NaCl) structure which is formed under moderate hydrostatic pressure and in some instances still remain s metastable at zero pressure, a zinc-blende configuration, and its mo st common structure, the wurtzite (hexagonal) crystal structure.34 The wurtzite crystal structure is cons idered a closed pack ed hexagonal and the basis vector (a = 3.25A) is orient ed along the axis of the hexagons and has a length 3/8 times the separation between the two hexagonal faces (c = 5.12A). Each Zn atom is tetrahedrally surrounded by four O atoms. The tetrahedrons are or iented so that the atoms fit in to the same axis but are offset to form two interpenetrating closed-packed hexagonal latt ices. This tetrahedral coordination is typical of sp3 covalent bondi ng, but has a substantia l ionic character.

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37 ZnO Band Structure As with other materials with the wurtzite cr ystal structure, all th ree valence band maxima are separated from each other and each is split by negligible spin-orbit interactions. The E-K relationship is35 E = A(k2 x + K2 y) + Bk2 z C(k2 x + K2 y)1/2 (3-1) where the z direction is chosen parallel to the c axis. The constant energy surfaces are, therefore, spheroidal with the c-axis as the principal axis for both the Conduction Band Minimum and Valence Band Maximum. There have been numerous experimental and theoretical studies of ZnO band structure since the initial calculations proposed by Russler36 in 1969. The nature of the valence band ordering remains unclear. The dcation electrons make theoretic al calculations based on the Local Density Approximation of the band struct ure demanding and difficult. Several authors disregard the 3d-electron contribut ions, but it was shown that their explicit consideration as valence electrons is a necessary step towards a qu antitative description of the band structure and other electronic properties.37-39 Doping ZnO Doping Rules Generally, doping a semiconductor is limited by40: local bonding effects, change in chemical potentials due to the presence of impurity atoms Fermi-level induced compensation effects The enthalpy energy change of the system due to the presence of a dopant D with charge q can be expressed as:

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38 H(D, q)( EF) = q EF + nD ( D H) + Eb (3-2) where D and H are the dopant and host chemical potentials, EF, the electro-chemical potential, nD is the dopant density, and Eb is the excess energy ge nerated by the local bonding around the dopant as defined as: Eb = E(host + defects) E (host) (3-3) E is the total energy with respec t to free-atoms. These rules must be taken into account to enhance dopant solubility and avoid pinning of the Fermi-Level due to compensating defects. Doping ZnO n-type ZnO oxide is normally n-type due to defects and exhibits carr ier concentration raging from 1016 cm-3 for high quality single crystals 41 to 1021 cm-3 for intentionally doped samples. The intrinsic n-type conductivity have been initially attribut ed to the presence of native defects such as oxygen vacancies (Vo), zinc interstitials (Zni) induced by the materials tendency to deviate from stoichiometry42, 43, unintentional hydrogen doping43-46 and other structural defects such as dislocations or grain boundary induced conductivity47. Extrinsically, ZnO is easily doped n-type by cation substitution. Table 3-2 presents a list of el ements that are used to n-type doped ZnO. Al, Ga and In are the most frequently used dopants. The M3+ metal ions most likely substitute for Zn2+ allowing an electron to be release to the conduction band according to the following equation 48 written in standard Krger-Vink notations. M2O3 ZnO M. Zn + e + O2 (g) (3-4)

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39 Doping ZnO p-Type Compensation While n-type conductivity is achieved in a straightforward manner, p-type doping remains a challenge. This doping asymmetry is a common occurrence in wide band gap semiconductors. This phenomenon is associated to low formation energy of compensating defects such as anion vacancies, cation interstitials and unintentional hydrogen doping. While the pinning level for ntype doping is inside the conduc tion band, for ZnO, the pinning le vel for p-type doping is high above the valence band maxima inside the ba nd gap. This causes degeneracy due to the spontaneous formation of compensating defects when ZnO is p-doped as the Fermi level gets closer to the valence band. T hus the pinning level of the Fermi energy is independent of the dopant and is determine by the formation en ergy and energy levels of the compensating defects.42 Native Defects in ZnO The predominant type of native defect in ZnO remains a controversy. Many suggest zinc interstitials based on ionic size cons iderations or ionic diffusion data49, 50 while others support oxygen vacancies based on reaction rate experiment s, annealing studies, photoluminescence and Hall measurements.50, 51 Van De Walle et al recently calculated the form ation energies of point defects and hydrogen in ZnO using first-principles, plane-wave soft-pseudopotential technique together with supercell approach. According to their calculations, the concentration of a defect, Cd, in a crystal depends on the free energy of formation: Cd = Nsitese -( G /kBT) (3-5) where Nsites is the number of sites in the la ttice where defects can occur, kB is the Boltzmann constant, T is the temperature in Kelvin and G is given by

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40 G = E T S + P V (3-6) Here E is the change in the total energy of the system; S and V are the changes in vibrational entropy and volume due to the presence of the defect in the lattice, respectively. Since the change in volume is negligible and th e change in entropy is of comparable order for defects, the charged defect formation energy can be articulated in terms of chemical potentials and the Fermi level: E = E(NZn, NO) NZnZn NZnO + q F (3-7) where E(NZn, NO) is the total energy of the system, NZn and NZn are the number of Zn and O sites in the system, Zn and O are the chemical potential of zinc and oxygen, q and F are the electronic charge and the Fermi level. To avoid the formation of Zn metal and lo ss of oxygen, the boundaries for the chemical potentials are defined as: ZnO Zn < O Zn (3-8) and ZnO O < O O (3-9) Here ZnO Zn, ZnO O, O O and O Zn are the chemical potentials of Zn in ZnO, O in ZnO, pure Zn metal and O2 gas, respectively.52 Theoretical calculations suggest VO is the dominant donor in zinc rich conditions and VZn for low zinc partial pressure. Howe ver, recent studies identified VO as a deep acceptor and Zni as a shallow donor with activation ener gy around 30 meV 5 meV. Hence, Zni is the dominant hole killer.41, 53, 54 It is worth noting that the complexi ng of vacancies and interstitials with dopants can significantly change th e formation energy of the def ects by altering their potential

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41 energy and the Fermi level. Moreover, large atom ic displacements near impurities can play an important and perhaps a signifi cant role in dop ing passivation.55 Hydrogen Doping From first-principle calculati ons, the concentration of hydrogen that can be incorporated in the lattice is determined by its forma tion energy E through the expression: CH = Nsites e (E / kBT) (3-10) where Nsites is the number of sites in the lat tice the dopant can be incorporated; kB is the Boltzmann constant and T is th e absolute temperature. The formation energy of H impurity with charge q is given by H q ( EF) = q (EF + EVBM) [ H + 0.5 E(H2) + E(host + Hq) E (host) (3-10) where H represents hydrogen chemical potential and EVBM is the relative position of the Valence Band Maximum, respectively. Corrected LDA computations demonstrate that only H+ is stable in ZnO for any position of the Fermi energy.43, 46, 56 The hydrogen pinning level, (+/-), is situated well inside the conduction band of ZnO. This translates into hydrogen acting exclusively as a donor with sufficiently low formation energy to allow large solubility in ZnO. Even though H is a shallow donor in ZnO, it may be helpful in p-type doping. Several researchers proposed the use of hydrogen as a co-dopa nt with p-type dopants such as N, P, As and Sb. As an n-type dopant, H in ZnO ha s the beneficial effect of activating VO, thereby facilitating oxygen substitution a nd increasing dopant solubility. In addition, its low formation energy can allow it to overtake other compensating defects such as Vzn and Zni. Subsequent removal via low temperature anneal will uncompensated acceptors.44, 56-61

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42 State of the Art p-Type Doping As mentioned earlier, p-doping remains a challenge because spontaneous formation of compensating defects; unintentional H doping; low dopant solubility; acceptor-defect complexing; deep impurity levels causing significant resi stance to the formation of shallow acceptor levels; Group V antisites. Table 3-3 lists the known agen ts to cause acceptor levels into ZnO. Lithium has long been used as a dopant for ZnO.62, 63 However, Group I elements tend to migrate to the interstitial sites rather than substitutional sites, in part becau se of their smaller ionic radii; therefore, they act mainly as shallow donors. Lithium, however, act s as a deep acceptor and yields to semiinsulating films. Both experimental data and first-principal LDA calculations reveal only insulating ZnO is achieved by group I doping due to the formation of meta l interstitials (Mi), which are shallow donors, and MZn-Mi defect complexes. For K, oxygen vacancies are most likely the hole killer, not Ki .64, 65 AgZn and CuZn could also act as acceptors in ZnO. Experiments by Fan et al. earlier pointed toward Ag acting as an amphoteric dopant, existing on both substitutional Zn sites and the interstitial sites.66, 67 Experiments on ZnO: Cu shows decr ease in conductivity proportional to the square of the Cu concentration, whic h was explained by the passivation of CuZn acceptors by O vacancies and/or the formation of acceptor-type Cu2+Cu2+ pairs.68-70 Kanai reported that Ag and Cu behaves as deep acceptors level lying about 0.23 and 0.17 eV below the conduction band,

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43 respectively.71, 72 Therefore, it is still uncertain whet her Cu and Ag are actually incorporated on substitutional Zn sites.66, 73, 74 In terms of standard Krger-Vink notations, the passivation of group I and Ib elements in zinc oxide can be written as: M2O ZnO MZn + h + M (s) (3-11) And M (s) ZnO M. i + e (3-12) So far, Group V elements remain th e most promising dopants in ZnO. Over the past couple of years, the bulk of research on ZnO p-doping has focused on N as a dopant. C. H. Park and contributors55, 64, table 3-3, calculated the bond length and energy of formation of NO. It appears that N is the best dopant for ZnO based on the similar bond lengths, low ionization energy and lack of antisite NZn.64 Various sources of dopants including N2, NO, N2O, NH3 and Zn2N3 have been used since Menegishi75 first reported successful p-type doping using NH3. D. C. Look et al,76 reported p-type doping using N2 through a plasma source by MBE on Li doped semi-insulating ZnO substrates. Nitrogen, with a concentration as high as 1019 cm-3, was incorporated into the film. This was c onfirmed by secondary ion mass spectroscopy. Carrier density, however, was on the order of 9 x 1016 cm-3 with a mobility of 2 cm2/(V.s). Although exact acceptor peak locations for P and N are st ill disputed, D. C. Look observed that by low temperature photoluminescence of phosphorous a nd nitrogen doped ZnO films, acceptor related peaks at 3.32 eV and 3.357 eV due to neutral-ac ceptor-bound excitons. The estimated acceptor level, based on Haynes rule is around 170 and 200 meV. Wang 77and Li78 both reported p-type behavior of ZnO with carrier density above 1017 cm-3 and mobility of about 1 cm2/(V-s) by simple oxida tion annealing of Zn3N2 on amorphous

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44 quartz and fused silica in air and oxygen. The films were polycrystal line with [0 1 0] and [0 0 2] as axis of majority grain alignment as revealed by powder XRD. Although p-type ZnO using N doping can be ach ieved using different growth apparati and doping sources, only about 0.1% of the N incorporated in the latt ice seem to be electrically active. A large number of electrically inactive dopants are suspected of acting as neutral scattering centers and, therefore, may be accountable for the low carrier m obility of the films. Moreover, Barnes79 reported conversion from pto ntype of their films in a matter of days. The metastability of N doping may, therefore, be the s ource of the scattering in the reported data and the difficulties in reproducibility. Co-Doping Method While p-type ZnO is possible by doping with nitrogen alone, the films are generally resistive. To overcome the low so lubility of nitrogen in ZnO, codoping with a group III such as Ga, Al, or In element was proposed to solve th e unipolarity in ZnO based on Yamamotos ab initio LDA calculations. 58, 80, 81 Solubility is enhanced through the formation of ion pairs between donor and acceptor ions. This contributes to: 1) reduci ng the Madelung energies of the lattice, 2) enhancement of acceptor incorpora tion because of stronger acceptor-donor attraction, 3) formation of an acceptor-donor-acceptor comple x that causes a decrease in the acceptor level toward the valence band and an increase in th e donor level towards th e conduction band, and 4) increasing carrier mobility due to short range dipole-like scattering mechanism. Experimentally, codoping was fi rst demonstrated by M. Joseph82 in 2001 by PLD using GaN as the co-dopant and N2O as background gas. They reported film resistivity as low as 0.5 ohm-cm and carrier conc entration as high as 1019 cm-3. However, carrier mobility was on the order of 0.07 cm2/(V-s). Z. -Z. Ye83 also reported fabricating ptype ZnO by codoping. Using Al and N, they achieved carri er concentration around 1017 cm-3 with mobility of about 0.3 cm2/(V-

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45 s). Codoping involving In experime nts were also performed. The unusually low resistivity, high carrier density and high mobility that was reporte d brings the validity of these results into question.84 Although the co-doping method demonstrated a huge increase in N solubility, it has its drawbacks. The limitations are the low growth temperature, low success in reproducibility and the presence of GaN or AlN pr ecipitates in the films, both of which could be responsible for the low mobility of carriers due to in crease number of scattering centers. Phosphorous, Arsenic and Antimony Doping Since p-type doping of ZnO using other elemen ts has yielded either semi-insulating ZnO or appeared to be a metastable in the case of nitrogen, attention turned to P, As and Sb. Ab initio electronic band structure calculations based on the local density approximation indicate increased Madelung energy with group V anion substitution. Consequently, these acceptor states are expected to be significantly localized.85 The large difference in si ze and bond length for these impurities increases the probability of vacancy fo rmation to compensate for the large lattice strain induced by the impurity inco rporation. Moreover P, As and Sb are expected to have levels deep inside the band gap once they substitute for O because of the increasing p-orbital energy as one moves down group V column.86 Work by K. Kim and contributors on Pdoped ZnO films shows that post-activation processing can yield phosphorus doped p-type and semi-insulating ZnO grown by sputter deposition. 87 Previous results on phosphorus-doped (Zn, Mg)O, in particular C-V and I-V characteristics, indicate that phosphorus yiel ds an acceptor state and p-type behavior.88, 89 The conversion from highly conductive n-type to low carrier density p-type behavior by thermal annealing suggests that th e shallow donor states in the as-dep osited films are relatively unstable, and that perhaps phosphorous may form a deep acceptor in ZnO as predicted.89

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46 P-type behavior has also been reported in As-doped ZnO films grown by pulsed-laser deposition10, 26, 90, 91 with carrier density on the order of mid-1017 cm-3. Ryu et al,90 demonstrated p-type behavior of ZnO films doped with As grown by Hybrid Beam Deposition. Carrier concentration as high as 4 x 1017 cm-3 with mobility as high as 35 cm2/(V-s) and film resistivity of 2 .cm from an arsenic concentration of 3 x 1018 cm-3. Temperature-dependent Hall measurements showed a thermal activation ener gy of 120 meV; an ionization activation energy was estimated to be 130 meV.90 In other work, p-type behavior in As-doped ZnO films is observed only after moderate temperature (200C) annealing after deposition.92, 93 The films were grown on c-plane sapphire at 600C in an O2 pressure of 50 mTorr. The as-grown films were n-type with a resistivity above 100 -cm and an n-type carrier density on the order of 1016/cm3. Upon annealing in nitrogen at 200C, the films showed a conversion to p-type with carrier density of 1017-1018 cm-3. It is unclear what mechanism yielded this activation of the acceptor conduction upon annealing. Theoretical work by Limpijumnong94 suggests that recent observations of p-type behavior in Asand Sb-doped ZnO may not be due to substitution on the oxygen site, but instead from the formation of co mplexes associated with As or Sb on the Zn site. Evidence of the involvem ent of complex associated w ith the p-type ZnO doped with antimony was pointed out in work by Lopatiuk-Tirpak95, 96 by the change of acceptor activation energy with dopant concentration. Activati on energies of about 212, 175, 158, and 135 meV were obtained for samples with carrier concentra tions of 1.3 x 1017, 6.0 x 1017, 8.2 x 1017, and 1.3 x 1017 cm-3, respectively. Epitaxial Growth of ZnO Films ZnO films have been grown on various substr ates including sapphire silica, diamond, SiC, Si, GaAs, GaP, GaN, ZnO single crystal, glass and ScAlMgO4, Mg Al2O4, CaF2.97-100 Single crystal ZnO would be the ideal substrate for the growth of ep itaxial films. However, its

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47 conductive nature and the tendency fo r its conductivity to dramatica lly increase when heated in a non-oxidative environment make it challenging to measure the electric al property of the overlaying film. ScAlMgO4 is a highly attractive substrate due to its small lattice mismatch to ZnO (0.09%). However, ScAlMgO4 is extrem ely expensive and remains commercially unavailable. Sapphire is by far the most frequently used substrate due to the availability of low cost, large area, epi-grade substrates. For ZnO f ilms grown on ror csapphire, the epitaxial relationship is as follows: ZnO / (0 1 1 2) Al2O3 (1 1 2 0) ZnO (01 12) Al2O3 [0 0 0 1] ZnO (0 1 11) Al2O3 ZnO / (0 0 0 1) Al2O3 (0 0 0 1) ZnO (0 0 0 1) Al2O3 with the following in-plane orientation relationships: [1 0 1 1] Al2O3 [1 1 2 0] ZnO [1 1 2 0] Al2O3 [0 1 1 0] ZnO Despite their large lattice mismatch, epitaxial film s have been grown at substrate temperature as low as 300 C. The crystalline quality of the films increases with temperature up to 800 C. For pulsed laser deposited films, the films are usually textured and c-axis oriented with a ZnO [1 0 1 0] Al2O3 [1 0 1 0] ZnO in plane relationship. This rela tionship is a direct superposition of the hexagons of the basal plane of ZnO and sapphire with a 31.8% lattice mismatch. With increasing temperature, the lat tice rearranges itself to a thermodynamically more stable [1 1 2 0] Al2O3 [0 1 1 0] ZnO relationship. In this configura tion the ZnO hexagons are aligned to the underlying sapphires oxygen sublatti ce causing 30% rotation with a reduced lattice mismatch of

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48 about 18%. Therefore, at low temp eratures and fast deposition ra tes, orientation is governed by kinetics based on local interface energy whereas thermodynamics dictate the relationship at high temperature and low growth rate.61, 101 Other properties are also affected by the use of sapphire as a substrat e. The large lattice mismatch between the two material s induces structural defects su ch as dislocations and other interfacial defects that are detrimental to the films mobility. As mentioned in Ch. 2, carrier mobility can be written as: sds iis acp pop i i Total 1 1 1 1 1 1 (3-13) Where pop, acp, iis, dis are the change in mobility due to polar optical phonon, acoustic phonon scattering, ionized impurity scattering and structural defects induced scattering.19 V. Srikant102 found that the optical absorption coeffi cient, band edge, and near band edge characteristics are affected by the nature of the substrate us ed. The band edge of films on c [(0001)] and R -plane [(1102)] sapphire were found to be 3.29 and 3.32 eV, respectively. The deviations in energy from single crystal value of ZnO (3.3 eV) were attr ibuted to the thermal mismatch strain using the known deformation potentials of ZnO a nd sapphire. In contrast, films grown on fused silica consistently ex hibit a band edge ~ 0.1 eV lower than that predicted using the known deformation potential and the thermal mismatch strains. The band edge of the films grown on fused silica are determined by an inte rplay between shifts due to the deformation potential and band bending at gr ain boundaries. Additionally, the spread in the tail ( E0) of the band edge for the different films is found to be very sensitive to the defect structure in the films. 102, 103 To alleviate some of the problems accompanie d by the growth of ZnO on sapphire, several groups have used MgO buffer laye rs grown at low temperature, low and high temperature ZnO

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49 buffers, annealed MgO or ZnO, GaN, combinations of thin and/or thin MgO/ZnO and annealed MgO/ZnO buffer layers. Cubic MgO grows along the MgO [111] on (0001) sapphire with 8% lattice mismatch and the following relationship: 0 1 1 MgO 0 1 1 23 2O Al And 2 1 1 MgO 0 1 013 2O Al Consequently, the use of buffer layers relieved so me of the stress that was induced by the larger lattice mismatch between ZnO and sapphire and also decreased dislocation density.104-107 Alloying ZnO with MgO To tune the direct band gap of zinc oxide to shorter wavelengths, researchers such as A. Ohtomo108. and Sharma109 alloyed zinc oxide and magnesium oxide to form a novel II-VI oxide semiconductor system, MgxZn1-xO (MZO) with x up to 0.33. This new system, having excellent optical properties, has similar lattice constants to those of ZnO but with a larger band gap of up to 4.0 eV (for x = 0.33) and can be used to construct excellent ZnO/(Mg, Zn)O quantum wells and superlattices as long as film surface and in terface flatness are controlled. According to the phase diagram110 the solubility limit of MgO in ZnO is only 4 mol %. Nevertheless, they were able to prepare metastable films by pulsed laser deposition with concentrations as large as 33 mol % with only 1% lattice constant increased in the a-axis and decreased the c-axis direction as x was increased. Diffraction patterns from brig ht-field Transmission El ectron Microscope of cross section of Mg0.34Zn0.66O/-Al2O3 samples support the formation of a single-phase MZO with decreasing dislocations with increasing film thickness.109 Beyond this concentration, MgO segregates at first to form a small phase field co mposed of mixture of wurtzite and cubic phases. As content of the film increases, the c ubic phase becomes the only phase present.6, 109

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50 The nonequilibrium nature of pulsed laser deposition facilitates the formation of the metastable materials.11 The laser-ablated species have high kine tic energies relative to a thermal evaporation source.11, 109 This excess energy results in an e nhancement of the species mobility at the substrates surface and the solid solubility of the two oxides producing a nonequilibrium extension of the wurtzite phase.11 Furthermore, the fairly large Madelung constant of wurtzite allows an increase of io nicity without any change in coordination number.111, 112 Thermodynamically, wurtzite phase films with x up to 0.22 mol % appear to be stable after being annealed at temperatures up to 850C. While the films with x = 0.15 mol % remained unchanged after an annealing at 1000C for an hour, the crystallinity of the sample with x = 0.22 was destroyed due to segregation of MgO particles. 113 Although most of the MgxZn1-xO films analyzed had the wurt zite crystal structure and were grown by either PLD or molecular-beam epitaxy (MBE), several research groups have demonstrated the capability of ot her techniques to grow the metast able phase. Using horizontal type Metal Organic Vapor-Phase Epitaxy, W.I. Pa rk produced high quality MZO films up to x = 0.49 mol % as thick as 0.7 nm on sapphire usin g a thin ZnO buffer layer with the purpose of improving crystallinity of the films. He re ported a broadening and blueshift of the NBE emission peak for the films from 3.394 eV for x = 0.0 to 4.05 eV for x = 0.49.114 Sol-gel deposition technique, which does not require vacuum apparatus, was used by D. Zhao115 to generate single phased epitaxial films of Mg concentrations up to 36 mol %. T. Terasakog also reported successful growth of polycrystalline films by means of Chemical Spray Pyrolysis on shapphire substrates.116 T. Minemotos group has also prepared single crystal MgxZn1-xO films by means of two-source ZnO and MgO RF magn etron co-sputtering. Phase segregation was observed at x = 0.58 and the band gap of MZO at x = 0.46, with a crystallographic structure

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51 similar to ZnO, was found to be 4.20 eV. Their results confirm that this method makes it possible to reduce the substrate temperature.117 Thus, one can deduce that the solubility limit of the Mg in ZnO is strongly affected by the growing mechanism. A study by J. H. Kang shows that films with x 0.21 are transparent and demonstrate excellent c-axis orientation without any MgO-rela ted peak in their spectra. However, increasing the substrate temperature cause a shift to higher angles due to the rising Mg composition. The band edges of the alloys shift gr adually to higher energies but retain absorption edges similar to ZnO with increasing x The pseudodielectric functions, at room temperature, indicate strong excitonic contribution to the fundamental band-ga p transition of the alloys despite alloying disorder.118 Takayuki119 reports observing absorption edge-sin gularity at room temperature, making MgxZn(1-x)O the second semiconductor besides ZnO for which this phenomenon has been observed. According to his study, the lineshape spectra of the absorption onset exhibit a powerlaw singularity of the form (E-E1)following Mahan. Here, E is the phonon energy and E1 is the threshold energy where the optical abso rption sets in. The coupling parameter, provides a measure of the total screening of valence holes carried out by the photocreated electron generated in the absorption process. In other words, is the strength of the singularities. The value for for the MgxZn1-xO alloy, for x = 0.21, was found to be higher than in ZnO grown under identical conditions. Possible reasons for the increase of the coupling factor can be explained by the following: decrease of electro n density, stronger hole localization, coulomb interactions and enhanced biding energy of excitons. It is known that the coupling parameter is a decreasing function of the electron density. Since Mg (1.2) is less electronegative than Zn (1 .7), it is expected that MgO is more ionic and

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52 has a wider band gap than ZnO. Therefore, the resulting ternar y alloy should produce a semiconductor with lower electron density than ZnO. Moreover, optic scattering is produced by the optic vibrations through polarization due to di fferences in electronegativity associated with the constituent cations. The effects of the excitonic binding energy on should be minimal because when x is in the vicinity of 29 mol %, the modulus of the en ergy is close to the binding energy in ZnO. The exciton binding energy is strongly affected by Mg concentration. From about 60 meV (pure ZnO) to a minimum value of 50 meV for x 0.17 and then increases to 58 meV for x = 0.29.120 This seems to be in very good agreement w ith the theoretical calculation made by M. A. Kanehisa et al for mixed alloys semiconductors.121 Crystal Structure and Stability It is clear that increasing the Mg concentration of MgxZn(1-x)O raises the ionicity of the alloy. However, since the Madelung constant for wurtzite crystal is fairly large, in the vicinity of rock salts, it favors the extension of the wMgxZn(1-x)O phase at room temperature. The continuous incorporation of Mg in ZnO beyond about x = 36 mol % results in a two phases field composed of a mixture of both wMgxZn(1-x)O and cMgxZn(1-x)O. This two phase mixture region persists until a magnesium concentra tion of up to 44 mol %. At high magnesium concentration, coulomb interactions become power ful enough to cause a change in the crystal structure as strong ionicity fa vors high coordination numbers.111 Moreover, S. Limpijumnong122 demonstrated in a recent study that the transition from wurtzite to the rocksalt structure involves no bond breaking but rather the formation of an ex tra bond which is can be facilitated by an inplane strain or compression of the c/a.

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53 MgxZn(1-x)O Band Gap There have been few attempts of theoretical prediction of the MgxZn(1-x)O band structure and electronic properties. Usua lly, two compounds of related cr ystal structure are mixed; the band structure of the resulting compound has the same general features as those of the parent phases and can be considered as an interpolatio n between the structures of the two compounds. However, the lack of understa nding of disorder systems impe des the understanding of fine configurations of the bands.35 Disorder is introduced by statistical distribution of cation sites with the two different atoms (Mg and Zn) system, and can only be characterized in principle by the density of states D(E). Thus, the wave vector is no longer a good quantum number and the dispersion relation E(k) loses its meaning. In the case of MgxZn(1-x)O, the disagreements on the band properties of both end structures and the very fact that MgO and ZnO have dissim ilar crystal structures further complicates the problem. For ZnO, even though this material has been studied extensively in the past, there is long-standing controversy over the na ture of the valence band ordering.123 As for MgO, the assignment of the bands still remains uncerta in. Though, an alloy has no translational invariance,124 Lambrecht123 approached this problem by calculating the hypothetical forbidden band energy (Eg) of MgO in different struct ure employing Local Density Approximation and then applying Eg(x) = (1-x)Eg(ZnO) + xEg(MgO) bx(1 x) (3-14) where x is the relative MgO content in the alloy and b, the bowing parameter, is equal to 0.56. The small gap bowing and overall gap dependen ce they found are consistent with the data reported by Ohtomo et al. for the wurtzite pha se. However, experime ntal numbers show a discontinuity in the band gap, associated to the phase transition from wurtzite to cubic, in the vicinity of Mg concentrations of 0.36. Beyo nd this point, the band gag energy shows an upward

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54 shift and continues to increase linearly as a f unction of Mg content. The data for c-MgxZn1-xO agree with Virtual Crystal Appr oximation (eq: 3.16) assuming band gap for cubic system MgO and ZnO.125 Eg(MgxZn1-xO) = xEg(MgO) + (1 x)Eg(ZnO) (3-15) The estimates from VCA show linear dependenc y of the Eg of the cubic alloy on Mg content and agree well with their experimental findings. Moreover, they show the following relation: Eg (MgxZn1-xO) = 0.0400(x) + 3.20 (3-16) It appears that the fluctuations in the gap energy arise from the local compositional oscillations due to the different vapor pressures of the two oxides at elevated temperature. This compositional disorder results in a more ionic system in preferential fluctuation of the conduction band. Therefore, not only there be an increase in the strength of the localization making the effective hole mass heavier, but also additional scattering will be produced because of the disorders introduced by alloying effects. Native Defects It is well known that the optical, transport an d surface properties of an oxide are affected by the presence of defects and im purities, in particular point de fects. A detailed understanding and control of the nature and origin of the point defects is therefore of fundamental importance to synthesize Zinc oxide films with well defined prop erties. One of the most important defects in oxides is the oxygen vacancy. Oxygen vacancies ar e naturally present in ev ery oxide in the form of Schottky or Frenkel defects. In ZnO, it is still disputed whether hydrogen contamination, considered as an amphoteric impurity, or oxyge n vacancies and zinc interstitials give this material its natural n-type charac ter. Theoretical and experimental support both kinds of defects. On the other hand, in MgO, oxygen vacancies trigge r the formation of color centres, F, existing

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55 in three states: neutral, F+ and F+2. For the neutral and F+, the electrons associated with the oxygen vacancy remain trapped in the cavity formed by the removal of O2and O-.126 The F defects can degenerate to F+, and F+ can further liberate into F+2 either through luminescence or ionization. In the later case, the electron enters th e conduction band of MgO and may be trapped by defect states. One such defect is produced when an hydrogen impurity atom occupying a F+ or F2+ centre, Hor H-2 substitutional defects, is perturbed by a nearby impurity or vacancy.127 Due to the strong MgOs Madelung field, the oxygen vacancy does not result in a new bond and the electrons remain localized. Formation energy for the charged defects is relatively high co mpared to neutral vacancies.126, 128 Like all ionic materials, F is more stable than F+, and F+ is more stable than F2+. Moreover, it is suggested that defects are at lower energy states at low coordinated sites su ch as surfaces, terraces and corners. Cubic MgxZn1-xO has a lot of structural similarities to MgO. They have the same crystal structure and coordination nu mber, closely matched lattice constants and both have strong Madelung constants. Hence, it is expected th at defects will have st rong MgO character. The highly ionic nature of MZO implies that the formation of reactive defect centers will not be favored. In fact, the removal of an oxygen anion fr om a regular site might may create a cavity, a F center that resemble a regular O2of MgxZn1-xO, with the electron associated with the O2ions trapped inside due to electrostatic stabilization by the Madelung potential of the crystal. The trapping of the electron by the el ectrostatic field due to the re moval of the oxygen anion affects surface reactivity.126 Moreover, whereas defects may act as hopping centers for intrinsic films, in the case of the n-type films, they may also behave as scattering and trapping centers for the carriers. Applications and Advantages So far, most investigations on MgxZn1-xO have been centered on its wurtzite phase.

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56 Several growing techniques have been used to produce crystalline films with a magnesium concentration of up 46 mol %. It has been es tablished that band gap, exciton biding energy, photoluminescence peak can be tuned to a la rger value while the refractive index and birefringence decrease by increasing the Mg conc entration of the system. In many instances, researchers have classified MgxZn1-xO as an advantageous II-VI semiconductor.6, 109, 114, 118, 119, 129 The similarities in the struct ural and physical properties large exciton biding energy, high optical transparency and the larger forbi dden gap of the metastable hexagonal phase have assured the successful achievement of ZnO/ wMgZnO superlattices and quantum well devices. Moreover, the lower index of refraction of the te rnary alloy results in good optical confinement between the cladding and active region in heterost ructures that is key for the development of deep UV lasers. Other devices and applications where the electro-opt ical properties of this alloy system can be used to advantage include: spin tronics, biosensors, UV LEDs and sensors, white LEDs, optical communication, solar cells and as a protective and contact films for plasma display screens. Epitaxial Growth of (Mg, ZnO) As with ZnO, the growth of wMgxZn1-xO alloys are limited to lattice or domain matching epitaxy on hexagonal substrates, such as H6-SiC (0001), ZnO (0001), -sapphire (0001) or Si (111), which makes them inconvenient for manu facturing. The cubic phase, however, can be deposited on MgO substrate, GaAs or, as demonstrated by J. Narayan6, on Si (001) substrates either through the addition a TiN buffer layer or via domain-matching epitaxy where four lattice constants of the of the cubic phase match three lattice constant of Si.1,5 Such convenience is a big asset to large scale integration and reduces the manufacturing costs.

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57 Challenges Despite these attractive characte ristics, several challenges hinder the development of ZnO based devices. Though ZnO can be consistent ly produced as a highly n-type conductive by Al, Ga or In doping, there is still no reliable technology to consisten tly produce strong p-type conduction. Besides the improvements in thin films and bulk growth technology, defect formation mechanism that are responsible for zinc oxides departure from stoichiometry, root cause of intrinsic n-type cond uctivity and p-type doping compensa tion, remain uncertain. Semiinsulating materials result from group I and II el ements doping due to the formation of metal interstitial point defect s or oxygen vacancy and metal-defects complex. Despite the availability of 2" ZnO wafers, the bulk of the research remains on sapphire du e to ZnO electrical properties degradation in reducing atmosphere at temperature above 700C. Improvements in material quality via buffer layers are being pursued and will likely facilitate better epitaxy studies.

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58 Table 3-1. Properties of some compound semiconductors [www.semiconductors.co.uk] Properties ZnO AlN GaN ZnS ZnSe ZnTe GaAs Lattice Constants (A) Wurtzite WurtziteWurtzite WurtziteWurtzite Zinc Blende Zinc Blende a0 (nm) 0.3249 0.3111 0.3189 0.3811 0.398 0.427 0.5653 c0 (nm) 0.5207 0.4978 0.5185 0.6234 0.653 0.699 c0/a0 1.602 1.600 1.626 1.636 1.641 1.637 Density (g.cm-3) 5.606 3.255 6.095 3.98 Melting point (C) 1975 1850 1100 1240 1237 Coefficient of linear expansion (K-1) a0: 6.5 x 10-6 5.3x10-6 5.59x10-6 5.7x10-6 c0: 3.0 x 10-6 4.2x10-6 7.75x10-6 Dielectric constant 8.656 9.14 6.81 9.6 9.1 7.4 12.5 Refractive index 2.008, 2.029 2.15.05 at 3eV 2.67 at 3.38eV 2.356, 2.378 Band gap (eV) Type 3.34 Direct 6.2, Direct 3.45, Direct 3.68, Direct 2.82, Direct 2.39, Direct 1.4, Direct Intrinsic carrier concentration (cm-3) < 10-16 6.2 3.45 3.68 2.82 2.39 2.1x106 Exciton binding energy (meV) 60 25 40 22 7.5 Electron Mobility (cm2/v s) 200 1000 165 500 340 8500 Hole Mobility (cm2/v s) 20 1000 165 500 340 8500

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59 Table 3-2. n-type dopant atoms. Atom Concentration (at %) Resistivity (10-4 *cm) Carrier Density (1020 cm-3) Radius (pm) R/RZn Electronegativity Zn 60 0.00 1.70 Ga 1.7 6.1 1.2 14.5 47 -0.22 1.60 Al 1.2 3.2 1.3 15 39 -0.35 1.50 In 1.2 8.1 3.9 62 0.03 1.70 B 4.6 2 5.4 86 0.43 2.00 Y 2.2 7.9 5.8 180 2.00 1.20 Sc 2.5 3.1 6.7 162 1.70 1.30 Si 8 4.8 8.8 118 0.97 1.80 Ge 1.6 7.4 8.8 128 1.13 1.80 Ti 2 5.6 6.2 147 1.45 1.5 Zr 5.4 5.2 5.5 160 1.67 1.4 Hf 4.1 5.5 3.5 159 1.65 1.3 F 0.5 4 5 71.7 0.20 3.9 Table 3-3. p-Type dopant candidates with their valen ce and ionic radius. Values taken from literature cited throughout the text. Dopant Candidates Valence Bong Length Radius Predicted Energy level from VBM Cation pm eV LiZn +1 2.03 152 0.09 NaZn +1 2.10 186 0.17 AgZn +1 144 0.23 KZn +1 2.42 232 0.32 CuZn +1 128 0.17 VZn Anion NO -3 1.88 146 0.165 0.40 PO -3 2.18 212 0.95 AsO -3 2.23 222 0.935 1.15 SbO -3 145

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60 Table 3-4. Literature survey for p-type, N doped ZnO. Technique Source Resistivity ( .cm) Ref. CVD NH3 34 75 PLD N2O 4 130 PLD N2O, Ga 0.5 82 PLD N2O, 2 131 MOMBE MMHy 132 MBE N2 40 76 MOCVD NO 20 133 CVD Zn2N3 150 134 RF Diode. Sputt. N2, GaN 12 135 DC Mag. Sputt. NH3 1 136 Table 3-5. Properties of MgO and ZnO. MgO ZnO Crystal structure Cubic (NaCl) Wurtzite Density 3.6 g/cc Lat const. () 4.2117 a = 3.25 c = 5.206 Mol. Iight (g/mol) 40.30 81.39 Heat of fusion J/g 1910.5 642.6 Melting point 2826 1977 Refractive index 1.735 2.013/2.02 9 Band gap (eV) 7.8 3.2 Cohesive energy (eV) 10.41 8.9 Lattice energy (kJ/mol) 3791 3941 Ionic data Mg2+ Zn2+ O2Effect. radii for CN 6 72.0 pm 74.0 pm 140 pm Electronegativity 1.2 1.7 3.5

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61 CHAPTER 4 ZINC OXIDE DOPED WITH ARESENIC FROM ARSENIC TRI-OXIDE Results and Discussion To understand the effects of doping the film w ith As, I investigated the effects of several parameters, such as dopant concentration, back ground gas, substrate temperature, dopant source and the stability of the doped films. Because of the large si ze of As, it has been predicted theoretically that p-type doping is due to the formation of the AsZn-2VZn complex. It is also suggested that such a complex is more stab le at moderate temper ature and high growth pressure.94 Since As is very volatile, it is very diffi cult to control it incorporation into pulsed laser deposited films. PLD targets are usually sint ered at elevated temperature in air. Such treatment favors As5+ over As3+ which is thought to be respon sible for p-type conductivity in ZnO. Although growing As doped ZnO film at high growth pressure favors the formation of VZn and the As-related complex responsible for p-type doping, these conditions adversely affect film quality. As background gas pressure increases, collisions between the gas molecules and the plume generated by the laser/target interaction increases. As a result of the increase collision frequency, the plume broadens and the species ar rive at the surface with much lower kinetic energy. At moderate substrate temperature, lo w plume kinetic energy translates into low surface mobility and therefore islands form at the substr ates or growing film interface. Both defects density and non-uniformity increases with increasing growth pressure. I report on the transport and photolumines cence properties of ZnO epitaxial films doped with 0.2 atomic % of arsenic, in which the dopant source was arsenic trioxide and ozone was used as oxidant. The ozone/oxygen mixtur e was generated by flowing high purity oxygen through high intensity electrical discharge inside an ozone generator.

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62 The arsenic-doped ZnO epitaxial films were grown via pulsed laser deposition (PLD) on cplane sapphire substrates. The target was fabr icated using high-purity ZnO (99.9995%) with As2O3 (99.998%) as the doping agent. The arseni c doping level in the target was 0.2 atomic percent. Film growth was achieved using a KrF (248 nm) excimer laser. The background gas was an ozone/oxygen mixture produced by an oz one discharge generator. The concentration of ozone in the mixture was approximately 3%. The ablation targets were pr essed and sintered at 1000 C for 12 hrs in air. A laser repetition rate of 1 Hz was used, with a target to substrate distance of 6 cm and a laser pulse energy dens ity of 1-3 J/cm2. The ZnO growth chamber exhibits a base pressure of 10 -8 Torr. The total film thickness ranged from 500 to 700 nm. The films crystallinity, transport pr operties, optical properties a nd surface morphology were analyzed by X-ray diffraction, variable field Hall m easurements in a Van der Pauw geometry, photoluminescence and atomic force microscopy, respectively. Powder X-ray diffraction of the films grown at 400 C and 600C, shown in figures 4-1 and 4-2, reveal that the films are epitaxial and oriented in the ZnO [002] with no impurity phase present for all the O3/O2 growth pressures used. For the f ilms deposited at 400C, the c-axis lengths of the films are smaller than the c-axis of bulk ZnO, cbulk = 5.206, and the length of the c-axis decreases with in creasing growth pressure from 5.203 to 5.171 as indicated in table 4-1 and figure 4-3. For the film grown at 600C, the c-axis sp acing decreases from 5.223 for the film grown in 15 mTorr of ozone/oxygen to 5.171 for th e films grown in 30 mTorr. The decrease in clattice spacing with increasing pressure is most li kely associated with As segregating to the grain boundaries as grain density increases with increasing growth oxygen pressure. Surface morphology is also affected by increasing pr essure. For the films gr own at 400C, surface

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63 roughness increases from 8.88 to 12.52 nm as grow th pressure increases from 0.3 to 30 mTorr as shown in figure 4-3. The surface of the films grown at 600C were smoother compared to those grown at 600C. However, surface roughness incr eases with increasing growth pressure from 6.63 nm for the film grown at 15 mTorr of O3/O2 to 8.98 nm for the film grown at 30 mTorr. The photoluminescence properties of the films are shown in figures 4-4 and 4-5. At room temperature, the photoluminescence for the films grown at 600C shows a near band edge peak at 3.3 eV and shoulders at 3.09eV. For f ilms grown at 400 C, the room temperature luminescence peaks at 3.3 eV are broader, less intense and intensity decreases with increasing growth ozone/oxygen pressure The film grown at the 400C and 0.3 mTorr of oxygen/ozone pressure shows significant visible lumines cence due to defect states in the gap. The low temperature photoluminescence spectr um of the films deposited at 400C and 0.3 mTorr of oxygen/ozone growth pressure, figure 4-5, shows 2 broad peaks centered at 3.361 eV and 3.236 eV. The first peak, 3.361 eV, is also seen in the low temperature spectra of the undoped ZnO film and ZnO crystal substrate as we will see in Appendix A and is due to the donor-bound exciton emission. The DX has a should er at 3.329 eV which is associated to the acceptor-bound exciton, AX. Taking into account the intensities of the DX and AX, their ratio and the energy separation of the peaks at 3.236, 3.141 and 3.064 eV I believe that those peaks are longitudinal optical phonon replicas of the DX and DAP combined as indicated in the graph. For this film, the conductivity is strongly n-type as will be discussed later. In contrast, the low temperature photoluminescence of the film grow n at 600C in 15 mTorr of oxygen/ozone is 10 time more intense and shows the broad DX at 3. 361 eV. The AX is the most dominant feature of the spectrum at 3.316 eV. The Donor-Accepto r-Pair, DAP emission is at 3.242 eV and the

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64 DAP longitudinal optical phonon replicas, 70 meV apart from the DAP peak, at 3.179 and 3.113 eV. The activation energy of the dono r binding energy can be estimated by137: eV E E X Db D g o00534 0 4 1 ) ( (4-1) where ) ( X Dois the energy of the DX (3.361), b DE is the donor binding energy and Eg is ZnOs band gap at 16K. From this, the donor ener gy is estimated to be 58 meV. With the donor binding energy and the carrier concentra tion I can estimate the acceptor binding energy138, EA: ] ) ( [3 / 1 D b D g AN DAP E E E E (4-2) Here the term 3 / 1 DNis the Coulomb energy assuming th e average distance between donors is 3 / 1 DN ND is the electron concentration deri ved from the Hall measurements and is materials dependent constant and is 2.7 x 10-8 eVcm for (ZnO) = 8.6. From equation 4-2, I derived an estimated acceptor biding energy of 155 meV. This value is in line with binding energy for As related acceptor in As doped ZnO published in the literature.10, 26, 90, 92, 94, 139-141 The dependencies of Hall voltage on applied magnetic field of the films are shown in figure 4-7. The slopes of the linear fits to the data are the Hall coefficients (RH) and are used to determine the type, carrier concentration and mobility. For most of the films, including the film grown at 400C and 0.3 mTorr of ozone/oxygen, the Hall Voltages varies linearly with magnetic field. This confirms the peaks assignment in the low temperature spectrum of the film grown at 400C and 0.3 mTorr. For the film deposited at 600C and 15 mTorr of oz one/oxygen, there is a scatter in the Hall voltage when plotted against a pplied magnetic field. The large scatter in the data and the significant difference between the Ha ll coefficient of the f ilm grown at 600C and 15 mTorr and the others reflects mixed conduction, by electrons and holes, is occurring since 2 2 2) (n p n p Hn p e n p R (4-3)

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65 The changes in Hall voltage for the film at 600 C and 15 mTorr of ozone/oxygen is reflective of its low temperature luminescence spectra when DAP was observed. The room temperature Hall data for these films are shown in table 4-1. For films grown at 400C, the conductivity is heavily n-type and resistivity, on the order of 0.01 to 0.03 -cm, increases with increasing growth ozone/oxygen pressure as seen in figure 4-7 while the opposite happens for carrier density (on the order of 8.2 8.7 x1019cm-3). Hall mobilities of the films are low and decrease with increasing growth pressure reflecting the low crystalline quality of the films. The films grown at 600C also show n-t ype conductivity. However, the film grown at 15 mTorr shows a resistivity of 96 .cm and carrier concentration in the order of 1017/cm3. Hall mobility of the films are also relatively low, 0.24 cm2/(V.s) for the film grown at 15 mTorr and 9.6 cm2/(V.s) for the film grown at 30 mTorr. The increase in resistivity and carrier density indicates a reduction in poi nt defects such as Zni and Vo with increasing O3/O2 pressure while the decrease of carrier mobility is caused by the worsening of crystal quality of the film due increasing growth pressure. Other studies have also examined the photolum inescence properties of As-doped ZnO. In As-doped ZnO nanowires, the room-temperature near band emission was observed at 3.273 eV.141 At low temperature, the As-doped ZnO nanowir es did not exhibit a free exciton peak. Two predominant peaks were observed at 3.360 and 3.313 eV. Based on the temperature dependence and peak location, these peaks were associated with the donor-acceptor pair (DAP) transitions. There is also a broad peak accompanying the DAP at an energy 72 meV lower in energy. This was assigned as the phonon repl ica of DAP (DAP-1LO). In As-doped ZnO films, low temperature photoluminescence yielded peaks at 3.359, 3.322, and 3.273 eV for lightly doped films, 3.219 and 3.172 eV for heavily doped films.10 The peaks located at 3.359 eV were

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66 attributed to acceptor-bound exciton emissions (A0X), located at 3.322 and 3.273 eV as recombination emissions between free electrons and acceptor holes (FA), and peaks located at 3.219 and 3.172 eV as donor-to-acceptor recombinat ion emissions. For ZnO films converted to p-type through moderate temperat ure annealing, the low temperature PL showed a peak at 3.354 eV, which was assigned to the neutral-acceptor bound exciton (A0X) emission. A broad peak centered at 3.273 eV was assigned as a 1-LO phonon replica of the A0X emission. Note that a photoluminescence peak at energies of ~3.1 eV has also been observed in phosphorus-doped ZnO films.142 In the cases cited above, th e location of the DAP and A0X emission differ slightly in energy from the results reporte d in this work. This may reflect differences in dopant concentrations for the materials considered. Th e luminescent properties of nitrogen doped ZnO has been examined as a function of nitrogen content.143 The donor-acceptor pair luminescence was seen to be dependent on dopa nt concentration, shifting down ward in energy with increasing nitrogen concentration. The observation of strongly n-type films with low carrier mobility for as-deposited Asdoped ZnO films is similar to that observed for P-doped ZnO films deposited using similar conditions.89 Carrier doping in As-doped ZnO film s is strongly dependent on the growth conditions. While no previous study has consider ed the effect of ozone, the effect of oxygen partial pressure on ZnO p-type conduction has b een investigated previously by both experiment and theory. Xiong et al.144 reported evidence for p-type cond uction in undoped ZnO films grown at high oxygen partial pressure by reactive sputtering. The increased oxyge n chemical potential by electronic excitation raises the forma tion enthalpy of the intrinsic donor VO. Theoretical doping rules proposed by A. Zunger145 suggests that limitations of p-type doping can be

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67 overcome by manipulating the growth conditions, e. g. the use of the host anion-rich growth conditions to inhibit the formation of so-called h ole killer defects. Calc ulations from chemical potentials suggest that the enthalpy of forming anion vacancies decreases under cation-rich (zincrich) conditions. For all of the samples considered in this study, the Hall-effect results show ntype conductivity. Summary The transport and photoluminescence prope rties of arsenic-dope d ZnO films grown via pulsed laser deposition were examined. Arsenic trioxide was used as the source of As. The surface morphology and c-axis length the films are both dependant on growth conditions. The low temperature photoluminescence of the films show DX at 3.36 eV, AX at 3.31 eV and DAP at 3.24 eV with longitudinal optical phonon re plicas 70 meV apart. Acceptor binding energy, estimated from the DAP, is 155 meV. For the film grow the film grown at 600C and 15 mTorr of O3/O2, the AX intensity was higher than the DX and the presence of high concentration of As related acceptor was confirmed by confirme d by the scatter in the measured Hall voltages collected by mean of variable field Ha ll measurements and the relatively high |RH| of the film when compared to the other films. In additi on, the film grown at 600C and 15 mTorr of O3/O2 has the highest resistivity, lowest Hall mob ility and carrier concentration. The transport properties of the films are also affected by the deposition conditions. Resistivity and carrier concentration of the films deposited at 400C decrease with increasing O3/O2 growth pressure indicating reduction in point defects.

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68 Table 4-1. Room temperature Hall data calculat ed from average Hall Coefficients and surface roughness from 1 m2 area at 1 Hz AFM. Temp. C Pres. mTorr Carrier Density cm-3 Mobility cm2/V.s Resistivity /cm Rough. nm c-axis 0.3 8.67 x 1019 8.89 8.09 x 10-3 8.88 5.203 3 8.27 x 1019 9.37 8.06 x 10-3 12.32 5.201 400 30 8.20 x 1019 2.55 2.98 x 10-2 12.52 5.171 15 2.7 x 1017 0.24 96.10 6.63 5.223 600 30 1.5 x 1019 9.55 4.4 x 10-2 8.98 5.171

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69 2030405060 102103104105 Intensity (counts/s) 2Theta( o ) 400oC 0.3 mTorr 30 mTorr (Black line) 3 mTorr ZnO (002) Sapphire (006) Figure 4-1. Powder XRD patte rn for films grown at 400 C in 0.3, 3 and 30 mTorr of O3/O2. 2030405060 102103104105 10 I n t ens it y ( coun t s / s ) 2theta( o ) 600oC 30 mTorr 15 mTorr Sapphire (006) ZnO (002) Figure 4-2. Powder XRD patte rn for films grown at 600 C in 15 and 30 mTorr of O3/O2.

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70 051015202530 8.5 9.0 9.5 10.0 10.5 11.0 11.5 12.0 12.5 5.16 5.18 5.20 Growth Pressure (mTorr)c-axis length (A) S ur f ace roughness ( nm ) Surface roughness c-axis Figure 4-3. Changes in surface roughness and c-ax is length with growth ozone/oxygen pressure for the films grown at 400C. The filled s quares and solid lines indicate the changes in surface roughness; the empty triangles and dashed lines indicate the changes in caxis length with growth pressure and the f illed stars and dotted lines indicate the % strain in the films.

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71 3603804004204404604805005200.00 0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.00 0.05 0.10 0.15 0.20 0 25 IntensityWavelength (nm)400oC 3 mTorr 30 mTorr 0.3 mTorr 3.02eV 2.84 eV 3.3 eV 600oC 15 mTorr 30 mTorr 3.3 eV 3.09 eV Figure 4-4. Room temperature photoluminescence spectra. 3.003.053.103.153.203.253.303.353.400.01 0.1 3.003.053.103.153.203.253.303.353.400.1 1 10 3.362 eV D0XIntensity (a. u.)Energy (eV) 400oC; 0.3 mTorr 3.141 eV 3.236 eV DAP +1-LODoX 2-LODAP + 3-LODoX 1-LODAP + 2-LOD0X A0X 70 meV Intensity (a. u.) 600oC; 15 mTorr 3.362 eV D0X 3.113 eV 2-LODAP 3.179 eV 1-LODAP 3.242 eV DAP 70 meV 3.316 eV A0X Figure 4-5. Low temperature photo luminescence spectra at 16K.

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72 -1.0x104-5.0x1030.0 5.0x1031.0x104-4.0x103-2.0x1030.0 2.0x1034.0x103 400oC, 30 mTorr 400oC, 3 mTorr 400oC, 0.3 mTorr 600oC, 30 mTorrApplied Magnetic Field (G)1.1x1081.1x1081.1x1081.1x1081.1x1081.1x1081.1x108 VHRH= -23.047 600oC, 15 mTorr Figure 4-6. Room temperature variable magnetic fiel d Hall coefficient. The slope of the linear fit for the Hall voltage Vs the applied magnetic field is the average Hall coefficient.

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73 051015202530 0.01 0.015 0.02 0.025 0.03 8.2x10198.3x10198.4x10198.5x10198.6x10198.7x1019 2 4 6 8 10 Resisitivity (Ohm.cm)Growth Pressure (mTorr) Resistivity Carrier density Carrier MobilityCarrier Density (cm-3) Mobility (cm2/(V.s)) Figure 4-7. Changes of transport properties w ith growth ozone/oxygen pr essure for the films grown at 400C. The filled s quares and solid lines indicate the changes in resistivity; the red triangles and dashed lines indicate the changes in carrier density and the filled blue stars indicates the cha nges in carrier mobility.

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74 CHAPTER 5 ZINC OXIDE DOPED WITH ARESENIC FROM Zn2As3 AS THE SOURCE OF As Results and Discussion In this chapter, I analyze the transport, structural and photolum inescence properties of ZnO and (Mg, Zn)O epitaxial films doped with arsenic from Zn3As2. The films were pulsed laser deposited at various temperat ures using zinc arse nide as the source of arsenic source. The films doped with 0.02 and 0.2 at % As were grow n on c-sapphire. For the films doped with 2 at % As, a thin buffer layers of either ZnO (2050 nm) or MgO (5-20 nm) were grown at 450C on the sapphire substrates. The PLD systems used a KrF (248 nm) excimer pulsed laser at 4 Hz and energy density of 1-3 J/cm2. The films were grown at a 1 Hz laser repletion rate. The total film thickness ranged from 600 to 900 nm. O3/O2 gas mixture or ultra high purity oxygen was used as oxidant/background gas. I have also explored the effects of ozone as background/oxidant in low pressure on the As doped film. The ozone/oxygen mixture was generated by flowing high purity oxygen through high intensity electri cal discharge inside an ozone generator. Since ozone is more reactive than oxygen, lower pressure is needed for the oxidation. The electrical, structural and photoluminescence properties measurement methods were described in chapter 2. ZnO Films Doped with 0.02 at % As on c-Sapphire I first look at the effects of growth temperat ure on the electrical, optical and structural properties of ZnO films doped with 0.02 atomic pe rcent of arsenic using zinc arsenide as the arsenic source. The films were grown on bare c-sa pphire substrates with growth pressure of 3 mTorr ozone/oxygen mixture or 30 mTo rr of ultra-high purity oxygen. The powder X-ray diffraction patterns for the fi lms grown at the following conditions: A600C in 3 mTorr of O3/O2; B300C and 3 mTorr of O3/O2; C600C in 30 mTorr of O2; D500C and 30 mTorr of O2 are shown in Figure 5-1. The fi lms are single phase, epitaxial,

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75 textured and highly oriented in th e ZnO [002]. As pictured in fi gure 5-14, c-axis spacing of the films deposited in 3 mTorr of O3/O2 increases slowly from 5.195 for the film grown at 300C to 2.202 for the film deposited at 600C. Fo r the films grown in 30 mTorr of oxygen, c-axis ranged from 5.176 to 5.198 with the film grown at 600C having the longe st c-axis. The length of the c-axis of the films, although smalle r than bulk of the c-axis of the films (cbulk = 5.206 ), is affected by the presence of stress/strain tensile biaxial strain in the growth plane perpendicular to the c-axis due to the larg e lattice mismatch between sapphire and ZnO146 and cell expansion due to inco rporation of the large As ions in the lattice. Figures 5-2 and 3 show the room temperat ure photoluminescence spectra of the films deposited in O3/O2 and O2, respectively. The films grown in the O3/O2 gas mixture show near band edge (NBE) emission 375 nm or 3.31 eV. Peak intensity decreases with increasing growth temperature. The peaks for the films grown at 300 and 600C are sharp and possess shoulders around 400 nm. Emission from 450-700 nm which is usually another dominant feature in ZnO PL spectrum is present in our film. However, it decreases as growth temperature increases. For the films grown in 30 mTorr of oxygen, the NB E emission is dominates the spectra and its intensity increases with increasing temperature. This usually indicate s improvements in the lattice as growth temperature increases. Emissi on in the visible, from 450 to 700 nm, was also present in the film and its intensity also d ecreases as deposition temperature increases. Low temperature photoluminescence spectra we re collected at 16K and are represented in figures 5-4 to 5-9. Figure 5-4 shows the spectr um for an undoped ZnO film on sapphire grown at 700C and 3 mTorr of ozone. As I will discussed in the appendix, the spectrum of the undoped film is dominated by the donor-boun d exciton recombination at 3.361 eV with shoulders at 3.381 eV and 3.316 eV which are associated with reco mbination of free electrons from the conduction

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76 band and holes from the valence band recombination (Xa) as Ill as the longitudinal optical phonon replica of the DX (LO-DX), respectively. The DAP due to residual acceptor can be observed at 3.256 eV. I then compare this spect rum to those of the 0.02 at % As doped ZnO films. For the film grown at 300C and 3 mTorr of O3/O2, its low temperature PL spectrum, figure 5-5, is comprised of the DX emission at 3.361 eV, the AX at 3.309 eV and the DAP at 3.23 eV and the LO-DAP at 3.16 eV. Figure 5-6 shows the spectrum for the f ilm grown at 500C and 3 mTorr of O3/O2. In this spectrum I observe the DX at 3.361 eV with an Xa shoulder at 3.381 eV, the acceptor-bound excitonic emission (AX) peak at 3.311 eV. Due to the energetic separation between the peaks following the AX, the emission at 3.231, 3.194 and 3.162 eV can be associated to the donoracceptor pair (DAP) transition and its longitu dinal optical phonon replicas 40 meV apart. However, the presence of the Xa at 3.383 eV and the relative in tensity of the AX indicate that there is a possibility that those peaks could be combination of emission of recombination of free electrons from the conduction band with acceptors (FA) and DAP transitions. The peaks at 3.194 and 3.162 eV could also be a combination of lo ngitudinal optical phonon replicas of the FA and DAP. These possibilities should be taken into account because measurement limitations and low crystal quality of the films may cause broadening of those peaks and not allow us to distinguish them. In this case, I can approximate the neutral acceptor ground state energy (EA) from the FA by26, 90: e B o g AT k FA E E E 2 1 ) ( (5-1) Where Eg is ZnO band gap at 16K, Te is the equilibrium temperature of the free electrons in the lattice taken to be 16K and kB is the Boltzmann constant. The acceptor ground state energy,

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77 taken to be the average of the two transitions energy. From there I get an estimated binding energy of 200 meV. In the case of a DAP transition, I can estimate the acceptor binding energy EA from138: ] ) ( [3 / 1D b D g AN DAP E E E E (5-2) I have estimated the binding energy the donor res ponsible for the emission at 3.361 eV to be 58 meV in the previous chapter. ND is the carrier concentration derived from the Hall measurements, averaging 5 x 1018 cm-3 for the films and is materials dependent constant and is 2.7 x 10-8 eV.cm for (ZnO) = 8.6. The term3 / 1DNis the Coulomb energy assuming an average the distance between donors is3 / 1DN From equation 2, I derive an estimated acceptor biding energy of 194 meV. The low quali ty of the film and the measurement temperature may cause broadening and overlapping of the DAP and FA. Therefore, I estima te the binding energy of the acceptor to be the average of the two, 200 meV. This value is slightly higher but consistent with the 150 meV ionization energy of the AsZn-2VZn complex calculated by Limpijumnong and data published in the literature.10, 26, 90, 92, 94, 140, 141, 147 Figure 5-7 shows the low temperature photol uminescence spectrum for the film deposited at 600C and 3 mTorr of ozone/ oxygen. Two features dominate the figure, the DX emission at 3.361 eV and the AX at 3.331 eV. For this film, the DAP and LO-DAP are barely noticeable at 3.23 and 3.19 eV, respectively. Figure 5-8 shows the low temperature photoluminescence for the sample grown at 500C and 30 mTorr of oxyge n. The spectrum is dominated by the Xa emission at 3.383 eV, the DX at 3.361 eV and the A X at 3.311 eV and the DAP at 3.232 eV. The features of the low temperature spectrum of the film deposited at 600 C and 30 mTorr of oxygen are not as distinguishable as for the other ones, as shown in figure 5-9. The vertical lines in the graph represents the energy at which I have obser ved the emission in the previous graphs. I can

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78 see humps in the plot correspond ing to these lines at 3.361 eV (DX), 3.310 eV (AX), 3.23 eV (DAP) and 3.21 (LO-DAP). Transport properties of the films were measured through variable magnetic field Hall measurements at room temperature. For most of the films grown in the O3/O2 mixture they were n-type conductive with low resistivity except fo r the film grown at 500C. The film grown at 500C had a resistivity of 16.08 .cm. Figure 5-11 shows the changes in the measured Hall voltage as a function of the applied magnetic field. The data is scattered due to the multiple band conduction as I have seen in earlier chapters and is consistent with the film low temperature photoluminescence spectra where I observed both A X and DX. Since I was not able to estimate RH, I cannot accurately calculate th e carrier concentration and mob ility and therefore labeled the film indeterminate in terms of carrier type. The transport characteristics of th e films grown in 3 mTorr of O3/O2 are shown in figure 510 and Table 5-1. Resistivities of the films di d not follow any specific pattern with increasing substrate temperature and were less than 1 .cm. Carrier concentration first increases from 2. 56 x 1018 cm-3 for the film grown at 300C to 3.5 x 1019 cm-3 for the film grown at 400C and then decrease to 3.96 x 1018 cm-3 for the film deposited at 600C. Carrier mobility followed a similar pattern as carrier concentration; it first increases from 4.25 cm2/(V.s) for the film grown at 300C to 9.04 cm2/(V.s) for the film grown at 400 C and then decrease to 8.16 cm2/(V.s) for the film deposited at 600C. Those changes indicate decr ease of residual donor de fects with increasing growth pressure and possible activ ation of the acceptor state and optimal growth temperature is around 500C. The transport properties of the films doped w ith 0.02 at % As and grown in 30 mTorr of pure oxygen are shown in figures 5-12 and table 52. Overall, the films grown in 30 mTorr of

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79 oxygen have lower resistiv ity, on the order of 0.040 .cm; higher carrier mobility, on the order of 13 cm2/(V.s) than for the films grown in 3 mTo rr of ozone/oxygen while carrier concentration remains in the 1018 1019 cm-3 ranges and the data is less scatte red. For the film grown at 500C and 30 mTorr of oxygen; resistivity is 1.155 .cm, carrier concentration is 2.42 x 1018 cm-3 and carrier mobility is 2.23 cm2/(V s). Similar to the film grown at 500C and 3 O3/O2, scatter in the measured Hall voltage with applied magnetic fiel d is pronounced, reflective of the acceptor state present in the low temperature photoluminescence, indicating dual band conduction. The slope of the linear fit to the data, line in figure 5-13, is the average Hall coefficient and is used to derive the Hall mobility and carrier density of the film. ZnO Films Doped With 0.2 at % As on c-Sapphire In this section I focus on 0.2 at % of Arse nic doped zinc oxide thin films on bare csapphire grown as a function of substrate temper ature, from 300-600C in 100 increments, using in 3 mTorr of ozone/oxygen mixture and 30 mTorr of ultra high purity oxygen. The powder diffraction X-ray patterns for the films grown in 3 mTorr of ozone/oxygen gas mixture and 30 mTorr of oxygen are shown in fi gures 5-15 and 5-16, respectively. The films are single phase, epitaxial, textured and hi ghly oriented in the ZnO (002) for the growth temperature range. C-axis spacing of the film grown at 300C and 3 mTorr of O3/O2 is estimated at 5.222 As substrate temperature is raised up to 500C, c-axis length decreases to 5.183 For the film deposited at 600C and 3 mTorr of O3/O2, the c-axis increased to 5.205 as seen in figure 5-17. C-axis of the f ilms deposited in 30 mTorr of oxygen also depended on growth temperature. From 5.235 for the film grown at 300C, c-axis spacing decreases continually to 5.201 as substrate temp erature reached 600C. Room temperature photoluminescence spec tra for the films grown in ozone/oxygen mixture are represented in figure 5-18. The films show near band edge emission is at 376 nm or

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80 3.3 eV. Its intensity increases w ith increasing substrate temperatur e. Emission in the visible (VE) range, 400-700 nm, which are usually present in ZnO spectra (see the photoluminescence spectrum of a single crysta l substrate discussed in the appendix ) is absent in the spectra except for the film grown at 700C. For the films grown in 30 mTorr of oxygen, as shown in figure 518, NBE still occurs at 376 nm or 3.3 eV and its intensity increases with temperature. The NBE emission peak is the dominant feature of th e spectra of the films grown at 500-700C. As opposed to the films grown in the 3 mTorr of oz one/oxygen, emission in the visible for the films grown in 30 mTorr of oxygen is more significant and its intensity decreases with increasing temperature. Ultra violet emission strongly depe nds on the crystalline quality of the films and grain size. Ablated species in the plume arrive at the substrates surface with higher kinetic energy in lower background pressure causing the films grown in 3 mTorr of ozone/oxygen mixture to have higher near ul traviolet emission that for the films grown in 30 mTorr of oxygen at the same temperatures. The origin of visible emissions from 400-700 nm is still controve rsial in ZnO. Point defects such as interstitials and vacancies are t hought to be responsible for luminescence in this range. Photoluminescence in this range is very low for the films grown in ozone/oxygen mixture while in pure oxygen: it is dominant but decreas es with increasing growth temperature. These changes suggest either a decrease in point defects due to ozone reactivity with metal interstitials or lower grain boundary related defe cts as point defects tend to se gregate to the grain boundaries. Low temperature spectra for some the films were measure at 16K and are represented in figures 5-20 to 5-22. The spectrum of th e film grown at 500C and 3 mTorr of O3/O2, figure 520, is dominated by the AX, at 3.330 eV. Adj acent to the AX peak, is the DX peak at 3.361 eV. DAP transition is observed at 3.257 eV. Figure 5-21 shows the sp ectrum of the film grown at

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81 500C and 30 mTorr of oxygen. The spectrum is similar to the spect rum of the film deposited at the same temperature and 3 mTorr of O3/O2. The AX at 3.330 eV remains the peak with the highest intensity followed by the DX at 3.361 eV and the DAP is observed at 3.252 eV. As growth temperature increased to 600C for the films in 3mTorr of O3/O2, the AX red shifted to 3.304 eV and is no longer the dominant peak. The D X situated at 3.361 eV is the peak with the highest intensity as seen in fi gure 5-22. The energy of the AX and the c-axis length of the film grown at 600C and 30 mTorr of oxygen are similar to those of the 0.02 at % of As doped ZnO films grown at 600C as I show ed in the earlier section. However, the intensity of the NBE emission of the film doped at 0.02 at % of As and grown at 500C, 30 mTorr is higher than for the film doped at 0.2 at % of As deposited at similar conditions. Th is suggests a decrease in the arsenic concentration in bulk of the films due to segregation to the grain boundaries and, to a lesser extent, arsenic vaporization. The acceptor binding energy EA can be estimated from the DAP transitions of the films grown at 500C in both envi ronment using equation (5-2) 138: ] ) ( [3 / 1D b D g AN DAP E E E E (5-2) I estimated the binding energy the donor responsible for the emission at 3.361 eV to be 58 meV in the previous chapters. ND is the carrier concentration deri ved from the Hall measurements and is materials dependent constant and is 2.7 x 10-8 eV.cm for (ZnO) = 8.6. The term3 / 1DNis the Coulomb energy assuming an average the distance between donors is 3 / 1DN From equation 2, the binding energy is approximately 191 meV. Hall measurements data at 300K for the f ilms grown in the ozone/oxygen mixture is compiled in table 5-3. The resistivity, carrier concentration and mobility are plotted against growth temperature in figures 5-23 for the film s grown in 3 mTorr of ozone/oxygen. The films

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82 were n-type and their resistiv ities were on the order of 0.02 .cm except for the film grown at 400C for which resistivity is 0.008 .cm. Resistivity follows a decreasing trend as growth temperature increases. Carrier concentra tion first increases from 2.3 to 3.2 x 1019 as growth temperature increased from 300 to 400C. Beyond 400C, carrier concentration decreases as temperature increases. Carrier mobility da ta for the films grown in 3 mTorr of O3/O2 is scattered, therefore, no relationship with grow th temperature can be derived. The transport properties the films grown in 30 mTorr of oxygen are shown in figure 5-24 and the data is compiled in Table 5-4. Resis tivity follows an increasing trend as growth temperature increases from 0.009 to 0.175 .cm for the films grow n at 300C and 600C, respectively. Carrier concentrati on steadily decreases from 1.14 x 1020 to 1.80 x 1018 cm-3 as substrate temperature is raised from 300 to 600C. Hall mobility of the films increases from 6.2 to 20 cm2/(V.s) as growth temperature increases excep t at 600C where it dipped from its value at 500C. The lengths of the c-axis of the films doped w ith 0.2 at % of As and grown in the 3 mTorr of O3/O2 are lower for the films grown in oxygen excep t for the films grown at 600C. The c-axis length of the film grown in both environments shows similar decreasing trend with increasing deposition temperature. Similarity can also be noti ced in the optical properties of the films. At 300K, the photoluminescence spectra show increasi ng NBE luminescence. At low temperature, the acceptor-bound exciton and donor-a cceptor transition emission have the same energy and the ratio of AX/DX peaks, which is a good measure of activated donor is closed for the films doped with 0.2 at % of As and grown at 500C. The 0.2 at % As doped ZnO films deposited in the ozone rich environment have lower resistivities. The resis tivity of the films follow opposite tren d with increasing growth temperature

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83 as compared to the films deposited in pure oxygen. Carrier concentration of the films grown in the two environment decreases with increasing deposition temperature. Ca rrier density of the films deposited in 30 mTorr of oxygen dropped about two orders of magnitude from 1020 to 1018 cm-3. However, the carrier density of the films gr own in 3 mTorr of ozone/oxygen remain in the 1019 cm-3 range. I believe those changes are due to the oxidizing nature of ozone. Ozone is known to be a strong oxidizer. Besi de reducing the Zni concentration, it may further oxidizes As(III) to As (V)148 near the surface as I cooled the films dow n causing a more conducting near surface layer. Although oxygen is less reactive than ozone, the oxygen pressure I used is 10 times higher than the ozone/oxygen pressure used. The increas ed interactions between the ablated species present in the plume and the oxygen molecules a nd the increase of reactivity of oxygen with increasing temperature may cause a greater decrease in the Zni. Even though As (III) may be more stable in oxygen than in the ozone/oxygen e nvironment, the concentration of As may be decreasing with increasing substrate temperatur e either by vaporization or out diffusion. The overall result is an increase in resistivity a nd Hall coefficient that becomes more negative. In terms of the dual-bands conduction model, conductivity ( ), Hall coefficient (RH) and Hall mobility ( H) are defined as19: ) (n pn p e (5-3), 2 2 2) (n p n p Hn p e n p R (5-4) and H HR (5-5)

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84 Where e is the charge of an electron; p are holes; n are electrons; n and p are electrons and holes mobility, respectively. Hall mobility does not change significantly with increasing temperature as the defects density of the films is high because of the large lattice mismatch between the films and the substrates. The effects of increasing arsenic concentration in the films can be seen by comparing the properties of the films doped a 0.02 and 0.2 at % in both ozone and oxygen. C-axis lengths of the films were closer to cbulk of ZnO (5.206 A) and were less was less affected by increasing temperature. For the 0.02 at % As concentration, the length of the c-axis of the films deposited in the ozone/oxygen mixture were long er than the c-axis length of the films deposited in oxygen. Caxis of the films deposited in both environments slowly increases with growth temperature to reach a value closer than to cbulk. At 0.2 at % of As, c-axis of the films grown in 30 mTorr of oxygen are higher than the caxis of the films grown in the ozone rich environment. The length of the c-axis of the films doped with 0.2 at % As is signifi cantly higher at low substrate temperature than cbulk than the c-axis spacing of the 0.02 at % As doped films. Moreover, c-axis length of the films doped with 0.2 at % As d ecreases with increasing deposition temperature approaching cbulk. The increase of the c-axis spacing of the films with increasing As content indicates an increase in cell as arsenic content increases, where as, the decrease of the length of the c-axis with increasing grow th temperature suggest a decreas e of the cell volume due to decreasing of As concentration in the lattice. As segregation to the grain boundaries, lost of arsenic due to vaporization and reducti on of stress/strain by formation of AsZn-2VZn complex or other thermodynamic processes stre ss/strain relief could be res ponsible of the decrease of c lattice constant with increas ing substrate temperature.

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85 Room temperature photoluminescence of th e films grown in the ozone/oxygen shows decreasing NBE emission with increasing substr ate temperature for the 0.02 % As doped films where as NBE emission for the 0.2 % As doped films increases with increasing deposition temperature. Deposited in 30 mTorr of oxygen, NBE emission intensity of the films doped 0.02 at % As increases with increasing growth temp erature where as the NBE emission of the films doped with 0.2 at % slightly de creases with increasing temper ature except for the film doped with 0.2 at% and grown at 300 C and 30 mTorr of oxygen. The AX emission and the DAP energies were not affected by the background gas. However, they blue-shifted from 3.311 eV to 3.331 eV as of the As concentration increased. Consequently, the estimated binding energy from the DAP decreases from 190 meV for the 0.02 at % As doped film to 180 meV for the 0.2 at % doped films. Table 5-5 summarizes the structur al, optical and transport proper ties of the films grown at 500C. When As concentration of the grown in 3 mTorr of O3/O2 increases from 0.02 to 0.2 at % As, the c-axis length decreases from 5.196 to 5. 183 ; AX and DAP emissions blue-shifts from 3.31 to 3.33 eV and from 3.23 to 3.25 eV, respec tively; the binding energy estimated from the DAP transition decreases slightly from 190 to 180 meV; resistivity decreases from 16.08 to 0.022 .cm. Since the film was highly compensate d I are not able to compare the carrier concentrations and mobilities. Fo r the films deposited at 500C is grown in 30 mTorr of oxygen, the c-axis length increases from 5.176 to 5.205 Similar changes occur in the AX and DAP emission energy and the binding energy estimated fr om the DAP transition as for the film grown in 3 mTorr of O3/O2. Resistivity decreases from 116 to 0.031 .cm, carrier density decreases from 2.4 to 1.8 x 1018 cm-3 and Hall mobility increases from 2.23 to20 cm2/(V.s) as the As concentration increases from 0.02 at % to 0.2 at % As.

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86 I believe that our Hall measurements of the ZnO films doped with 0.02 and 0.2 at % As are affected by the presence of a thin, highly disl ocated region generated from the sapphire/ZnO interface due to strain relief b ecause of the large lattice mism atch between ZnO and sapphire. This can cause a conducting layer near the interface. Di slocations and other structural defects may be responsible for the low mobility and the anomalous carrier density dependency of mobility. ZnO films Doped with 2 at % As grown on ZnO and MgO buffer The powder diffraction, not shown, reveals that the 2.0 at % As doped ZnO films are epitaxial and textured in the ZnO (002) for the growth temperatures (400-600C) and oxygen partial pressures (1-200 mTorr) considered. The length of the c-axis was 5.225 A for the film grown at 5 mTorr of oxygen. The film grown at 20 mTorr of oxygen has the longest c-axis spacing, 5.248 while the film grown 100 mTorr ha s the lowest c-axis le ngth of 5.180 as seen in figure 5-25. The graph suggests cell volume firs t increases with increasing growth pressure in the low pressure regime. With increasing growth over 20 mTorr, pressure, the grains becomes smaller facilitating As to migrate to the grain boundaries causing lattice re laxation and a decrease in cell volume. Figure 5-26 shows the high reso lution rocking curves of th e films grown at different pressure. The full width half maximum (FWHM) of the films, shown in figure 5-27, increase with increasing growth pressure. Figure 5-28 sh ows the effects of substrate temperature on the rocking curves of the 2 at % As doped f ilms grown in 150 mTorr of oxygen. The FWHM decreased from 2.323 for the film grown at 500C to 0.363 for the film deposited at 600C. The FWHMs also improved when a thin MgO buffer la yer was used instead of ZnO, from 0.975 to 0.795, as it can be seen in figure 5-29 for th e films grown at 500C and 5 mTorr of oxygen. The decrease in film quality can be attributed to th e formation of ZnO nanocrystals due to a decrease

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87 of plumes kinetic energy as the plume inter acts with the background ga s. Increasing substrate temperature increases the surface mo bility of the particles favoring 2D growth and the use of thin MgO buffer layer reduces the lattice mismatch between ZnO and sapphire to about half and reduces the tendency of domain misalignment. Figures 5-30 and 5-31 are the room temper ature photoluminescence spectra for the films grown at 500C and different oxygen pressures on ZnO and MgO, respectively. For the films grown on ZnO buffer layer, near band edge emissi on intensity increases w ith increasing growth PO2. When grown on thin MgO buffer layer, the films deposited at 0.1 and 1 mTorr of oxygen show weak luminescence. The film grown in 5 mTorr of oxygen on MgO buffer shows the highest NBE emission intensity. NBE emission in tensity of the films grown in 20 and 50 mTorr of oxygen is weak and decreases with increa sing growth pressure. Figure 5-32 shows the augmentation of near band edge luminescence with increasing growth temperature for the films grown at 150 mTorr and on ZnO buffer layer. Si nce the room temperature NBE emission in ZnO is attributed to the longitudi nal optical phonon replica of the shallow donor-bound exciton, the intensity of the observed NBE depends on crystallin e quality of the films, concentration of the residual donors and presence and concentration of acceptors. 149, 150, 104, 151. Low temperature photoluminescence was pe rformed on selected films at 20K, as represented in figures 5-33 a-f. At 20K, the photoluminescence spectrum of an undoped ZnO film grown at 700C and 3 mTorr of ozone/oxygen mixtur e, shown in figure 5-33(a), is dominated by the donor-bound exciton, DX, emissi on at 3.361 eV and its longitudinal optical phonon replica, 1-LO, around 3.316 eV and DAP at 3.256 eV. For the 2.0 at % As-doped film grown at 400C and 150 mTorr of oxygen (figur e 5-33(b)) the DX at 3.361 eV is broad indicating perhaps contribution from non-observed acceptor-bound emission centered in the

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88 range between the bars indicated in the plot. Ad ditional features are present in Figure 5-33(c) which shows the spectra for the film grown at 500C and 50 mTorr of oxygen. The DX the peak is located 3.361 eV and the AX is located at 3. 31 eV. I associates the peaks at 3.279, 3.252 and 3.321 eV to the Donor Acceptor Pair (DAP) tran sition and the Free electron to Acceptor (FA) recombination and the 1st and 2nd longitudinal optical phonon re plicas to the DAP and FA transitions, respectively. The spectrum of the film grown at 500C a nd 100 mTorr is shown in figure 5-33(d), a broad DX is seen at 3.361 with possible AX co ntribution as indicated by the bar in the figure. In figure 5-33(e), the low temperature spectrum of the film grown at 500C and 150 mTorr of oxygen is shown. One can observe a broad peak with shoulders at 3.361 eV (DX), 3.316 eV (AX), 3.261 eV (FA + DAP), 3.179 eV (1 -LO DAP and FA). The 30K photoluminescence spectrum of the film grown at 500C and 200 mTo rr of oxygen, figure 5-33(f), reveals the DX at 3.361 eV, the AX at 3.30 eV, the DAP at 3. 226 and the LO DAP at 3.047 eV. Figure 5-33(g) shows the spectrum for the film deposited at 500C and 5 mTorr of oxygen on MgO buffer layer. This spectrum shows the DX at 3.361 eV (DX ), the AX at 3.342 eV, the FA + DAP at 3.31 eV, and the humps at 3.235 and 3.135 eV caused by the LO DAPs. The DX emission is present in all the spectra, figures 5-33 (b)-(f), where as the acceptorbound exciton emission intensity and position, cons equently the DAP and FA as well, depends on the growth pressure and temperature. Fi gure 5-34 shows the dependency of the Acceptor bound exciton and DAP emission energies on grow th pressure. Both AX and DAP decreases with increasing growth oxygen pressure. The AX, as seen in figure 5-35 s hows an almost linear relationship with oxygen pressure which can be described by o oP X A410 02 2 343 3 (5-6)

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89 Using equations 5-126, 90 and 5-2138 : e B o g AT k FA E E E 2 1 ) ( (5-1) And ] ) ( [3 / 1D b D g AN DAP E E E E (5-2) I have estimated the acceptor binding energies Eb A of the films. I found that the Eb A decreases according to the following linear relationship with the acceptor-bound emission energy as shown in figure 5-36: b A oE eV X A 55 0 399 3 (5-7) The red-shifts of the AX energy with increasing growth oxygen pressure cause Eb A to increase with increasing depos ition pressure, as shown in figure 5-37 (a), following the 2364 0 102O b AP meV E (5-8) In the earlier section I showed that the acceptor binding energy decreased from 190 meV for the ZnO films grown from the target doped with 0.02 at % As to 180 meV for the films deposited from the target doped with 0.2 at % As. This suggests a decrease in activated As related acceptors w ith increasing growth oxygen pressure. When I plotted the acceptor binding energy agai nst hole concentration of the p-type films (figure 5-37 (b)), the accept or binding energy decreases linear with increasing p1/3 similar to a pattern first observed by Debye and Conwell152: where E0 is the binding energy at infi nite dilution. For the p-type ZnO films doped with As, E0 is estimated to be 163 15 meV. This value agrees well with the calculate d ionization energy of the AsZn-2VZn complex value of 160 meV predicted by Limpijumnong. 94 ) (3 / 1 0p E Eb A

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90 For the films showing only or predominant DX peaks, room temperature Hall measurements reveal that those films have lo w resistivity and are h eavily n-type. The room temperature Hall data for various films are show n in table 5-6 and 5-7 and figures 5-38 and 39 are the plots of resistivity, for the films grow n on ZnO buffer layer, plotted against growth pressure and temperature, respectively. Resist ivity did not follow any trend with increasing oxygen pressure. However, the films showing p-type behavior are more resistive. The dotted line in the figure reveals that the resistivity of th e films showing p-type c onductivity increases with increasing growth pressure. Increasing substrate temperature increases the resistivity of the doped films deposited on ZnO buffer layers in 150 mTorr of oxygen as s hown in figure 5-39 as we would expect. Those changes can be explained in terms of the dual-bands conduction model. In this model, conductivity ( ) ) (n pn p e (5-3), Where e is the charge of an electron; p are holes; n are electrons; n and p are electron and hole mobility, respectively. Increasing deposition oxygen pressure decreases the crystalline quality of the films, causes a decrease of the activated acceptor concentration and an increase of the acceptor binding energy as the omega rocking curv es and photoluminescence of the films show earlier. As I showed in chapter 3, carrier mob ility is inversely depende nt on the crystalline quality of the films and carrier concentration, dopant size and concentration among other factors. Residual donor defects such zinc interstitials also decreases with increasing growth pressure. For the films grown on MgO buffer layers, resi stivity increases with increasing growth oxygen pressure, as seen in figure 5-40.These resu lts are consistent with the activation of a relatively deep acceptor stat e that limits electron carrier concentrations.

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91 Carrier mobilities of the films are low and decreases with increasi ng pressure or temperature when comparing carrie rs of the same type. This is due primarily to the changes in crystal quality that result from increasing pressure or temperature. Carrier concentration and type were also affected by changing the deposition para meters. Hole concentration for the films with p-type conductivity decreases w ith increasing oxygen pressure where as the already low hole mobility of the films is not as affected the growth pressure. Further Hall characterization of the films grown at 500C in 5 and 150 mTorr of oxygen were carried out in our PPMS. The dependency of the Hall voltages against applied magnetic field of the films is shown in figure 5-41. The scatter in the data is a good indication of the degree of compensation in the films. The decrea se in Hall coefficient is due to the changes caused by the increase in pressure 2 2 2) (n p n p Hn p e n p R (5-4) As I have seen earlier, an increase in growth pressure causes n, p, n and p to decrease. Since As concentration (p) and holes mobility ( p), native donor defects (n) and electron mobility by and n decreases with increasing growth pressure, consequently, RH decreases as well. The samples exhibited weak p-type conductivity with carrier concentrat ion in the range of 2 1016 cm-3 and a mobility of about 7.4 cm2/(V s) for the film grown at 5 mTorr and 6 1015 cm-3 and a mobility of about 0.4 cm2/(V s) for the film grown at 150 mTorr. These values differs from the values reported in table 5-6 and are attri buted to difference in films resistivity when the measurements were performed in air (T able 5-6) and in vacuum (PPMS). The semi-logarithmic plots of resistivity a nd carrier concentration as a function of temperature for the films doped with 2 at % As deposited at 500C at 5 and 150 mTorr of oxygen are shown in figure 5-42. The slope of linear fit to the plots of log of carrier concentration

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92 against reciprocal of temperature suggest Accep tor carrier thermal activation energy of 87 14 meV and 140 8 meV for the films grow n at 5 and 150 mTorr, respectively. Using the charge-balance equation with all the assumptions it entails7, 15, 19, 153, I attempted to estimate the acceptor activation energy by anal yzing the temperature dependence of carrier concentration on temperature. The least-squa res fits to the Arrhenius plot of p/T3/2, shown in Figure 5-43, give an estimate acceptor thermal activation energy, EA, of 70 14 meV and 150 12 meV for the samples grown at 5 and 150 mTo rr, respectively. These values are in good agreement with the ones I calcula ted in the paragraph and the es timated binding energy using the DAP and FA energies, 99 and 150 meV for the film s grown at 5 and 150 mTorr, respectively. Moreover, interpolation using the acceptor ther mal activation energies to estimate the acceptor thermal activation energy at infinite dilution gives 158 meV. This value is in good agreement with the estimated value for E0 I earlier found, with data repor ted in the literature, 120 to 180 meV8, 90, 147 and most importantly very close to th e estimated 160 meV ionization energy of the AsZn-2VZn complex predicted by Limpijumnong. 94 The effects of pressure on surface morphology are shown in Figure 5-44 for selected samples. Table 5-8 is a compilation of the m ean roughness collected by AFM for some of the samples. The smoothest, with a mean roughness of 2.1 nm, films obtained we re at 100 mTorr of oxygen while the roughest were at 150 mTorr with a roughness of 10.7 nm. Figure 5-45 shows how oxygen growth pressure af fects resistivity, Hall coefficient, Hall mobility, carrier density and acceptor activa tion energy for the films showing p-type conductivity. Resistivity and Hall coefficient a nd acceptor activation energy increase with increasing growth pressure while mobility, alrea dy low, is not as affected by growth pressure. The dependency of hole concentration on binding energy and growth pressure is better seen in

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93 figure 5-46. Hole density decrea ses with increasing carrier ac tivation energy and growth oxygen pressure. I believe these changes are due a decrease in the change in activated As concentration in the films either due to decrease of As content due to increased interaction between the plume and the background oxygen pressure or segregati on of As to the grain boundaries as grain density increases with increasing growth pressure. This is suggest ed by decrease of c-axis length of the films from 5.243 to 5.193, which is close to the c-axis length of the films doped with 0.2 at % As as I have seen in the pr eceding chapter and in the plot of the acceptor binding energy with As concentration shown in figure 5-48. From figure 5-37, I have estimated the accepto r binding energy of a film grown at 500C and 30 mTorr of oxygen and plot it along with th e estimated binding energy for the films doped with 0.02 and 0.2 % As. The figure reveals that binding energy changes with Arsenic concentration according to the following: As b AC E 229 0 246 5 ln (5-9) Solving for the acceptor binding energy: ) exp( 3 1 190As b AC meV E This is a similar patter as I have seen earlier for the dependence of the acceptor optical binding energy on hole concentration. The change in the ionization en ergy with carrier density follow a similar pattern recently observed by O. Lopatiuk-Tirpak et al.96 in Sb doped ZnO, in Mg doped GaN154, 155 and in Si.156, 157 Its nature remain a controversy, however Pearson157attributed this effect to a residual potential energy of attraction between the holes and the im purity ions. Others attr ibutes the decrease of activation energy to screening of the trapping center by the free carri ers, formation of a band-tail state that extend into the forbidden gap, th e broadening of the acceptor band in the gap. .96, 154-157

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94 Effects of Target Concentration I further investigate the effects of target concentration on the prope rties of the films by growing two samples on thin MgO buffer layer from the targets doped with 0.2 and 2.0 at % As from Zn3As2 as the source of As. The transport properties of the films are repr esented in table 5-9 and figure 5-50. Both films are p-type. Increasing the As content from 0.2 to 2 at % As caused the resistivity of the films to decrease from 71.15 to 60.57 .cm, carrier concentration to increase from 0.43 to 3.25 x 1017 cm-3 and, as I would expect, carrier mob ility to decreases from 2.04 to 0.89 cm2/(V.s). Variable magnetic field Hall measurements of the films shows increase scattering in the plots of the Hall voltages with applied magnetic field as concentration increased from 0.2 to 2.0 at % As, as shown in figure 5-49. This suggests e ither increasing compensation, increasing Asi with increasing As concentration. Room temperatur e photoluminescence of th e films is shown in figure 5-51. The figure shows that both spectra contain similar features except that NBE emission for the films doped with 0.2 at % is much broader than for the film doped with 2 at % As. Increasing the growth pressure from 5 to 150 mTorr of the films doped with 0.2 at % As and deposited at 500C causes: a carrier type conversion from p-type conductivity, for the film grown at 5 mTorr of oxygen with a resistivity of 71.15 .cm, carrier concentration and mobility of 4.3 x 1016 cm-3 and 2.04 cm2/(V.s), to n-type conduction fo r the film grown at 150 mTorr of oxygen with a resi stivity of 22.09 .cm a carrier density of 3.25 x 1018 cm-3 and carrier mobility of 0.09 cm2/(V.s), as seen in table 5-10. The resis tivity of the films doped with 0.2 at % As deposited at 500C and 150 mTorr is much lower that the resistivity of the p-type films doped with 2.0 % As and grown at sim ilar conditions. Analysis of the va riable field Hall measurements

PAGE 95

95 data, shown in figures 5-49 and 5-52, shows incr eased scattering in the data with increasing pressure. Effects of Persistent Photoconductivity The effects of pers istent photoconductivity153, 158 on the resistivity of the film doped with 2 at % As deposited at 500C and 5 mTorr on MgO buffer layer after being placed in the dark and in vacuum can be observed in Figures 5-52 to 5-56 The plot of resistivity with time, figure 5-52, can be divided into 3 sections. When the film is first place in the dark, its resistivity increases from around 14 to about 27 .cm after 21 hours. I believe th at the increase is due to photorelaxation/surface reconstr uction and oxygen adsorption. Ot hers have observe a similar phenomenon in ZnO and they attributed th is prolonged photoconductivity effect to oxygen adsorption in the absence of UV illumination eith er on the surface or on crystalline interfaces which reduces its conductivity. Si nce adsorption rate is limite d by oxygen concentration in the vacuum, the decrease in c onductivity is extremely slow.20, 53, 159-161 A region of saturation is observe d from 21 to 31 hours where resistivity of the film is more or less stable. Beyond 31 hours, the resistivity of the films begi ns to slowly decrease and eventually to level off after approximately 100 hours. I believe the decrease in resistivity in this region is cause either by H2 and reaction with oxygen adsorbed at the surfaces.43, 44, 46, 162-165 Such reaction would cause an increase in compensation in the films. The plots of Hall voltages with applied magnetic field for the film measured af ter 24, 37 and 42 hours being placed in the dark, figures 5-53 to 5-56, show progressi ve increases in the scatter in th e data. The slope of the linear fit to the data, the Hall coefficient, decreases with time, indicating an increase of compensation with time. When the film doped with 2 at % As grown at 500C and 100 mTorr is placed in the dark and Hall measurements were performed in air, th e change in resistivity after 24 hrs wasnt as

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96 dramatic. Resistivity only increased from 3.6 to 3.9 .cm, as shown in figure 5-57. However the plots of the Hall voltage with appl ied magnetic field performed at 0. 5, 12 and 24 hrs, as seen in figures 5-58 to 5-60, shows process that I descri bed earlier occurs much faster when the samples are measure in air and 1 atm. Effects of MgO Alloying on ZnO:As0.002 In this section I report on the electrical, optical and structural propert ies 0.2 at % As doped MgxZn(1-x)O films grown on sapphire using oxygen as background gas and oxidant. The targets were fabricated from mixing high pu rity powders of MgO, ZnO and Zn3As2 with the desired concentration and then sintered in Air at 1000C for 12 hrs. The films were deposited at 500C and 600C in oxygen in pressure ranging from 30-120 mTorr. Figures 5-61and 5-62 show X-ray diffraction patterns of the MgxZn(1-x)O:As films. For both As-doped and undoped (not shown) films, only the wurtzite phase is present in the films. The films are epitaxial and well aligned in the MgxZn(1-x)O [002] with a 2 between 34.48 and 34.53 degrees depending on growth pressure as opposed to 34.57 and 34.61 degrees for ZnO films. The changes in c lattice length as a functio n of growth pressure is shown in figures 5-6.3 and 5-64. c-lattice constant remains constant for the films grown at 500C from 30 to 90 mTorr of oxygen while there is a decrease in c-latt ice spacing for the films grown at 500C and 120 mTorr of oxygen and 600C and 60mTorr. Chemi cal analysis by WDS shows that composition of the films grown at 500C a nd 30-mTorr of oxygen have Mg c ontent of 5.93 at % while the films grown at 500C an d 120 mTorr has slightly higher Mg concentration, 6.83 at %. The increase of Mg content due to the increase of s ubstrate temperature is to the difference in the sticking coefficient between Mg and Zn. The c-axis of the films is affected by the Mg and As content of the films.

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97 Although the films remain wurtzite at all th e pressures studied, the crystalline quality decreases as background pressure increases. Figure 5-65 shows the broadening of the high resolution -rocking curves of the films grown at 500C and 30, 60 and 120 mTorr of oxygen. FWHM, as seen in figure 5-66, increases linear ly with increasing growth oxygen pressure. The room temperature photoluminescence of th e films are shown in figure 5-67. The films exhibits low luminescence at room temperature a nd near band edge emission appears at 3.48 eV. Photoluminescence measurements at 25K for an undoped Mg0.05Zn0.95O film deposited at 500C and 60 mTorr of oxygen is shown in figure 5-68. The film shows a peak at 3.471 eV. Since the film is n-type and its Mg concentration is low, I assume the films exhibits strong ZnO character and attribute the peak to the DX. The peak wi dth is significantly broader than the DX peaks that I usually observe in ZnO films. The broade ning of the peak is cause d by scatter in the band gap energy due to the statistic al distribution of the two ki nds of cations (Mg and Zn).166 The spectrum for the 0.2 at % As doped f ilm grown at 500C and 60 mTorr, shown in figure 5-69, shows the DX at 3.47 eV with a shoulde r at 3.42 eV and appears to be two peaks, at 3.34 and 3.25 eV that I believe are the result s acceptor-bound exciton (AX) and Donor-Acceptor Pair transition (DAP) emission. The low temperat ure spectrum for the 0. 2 at % As doped film grown at 90 mTorr of oxygen, shown in figure 570, is dominated by a broad peak centered at 3.47 eV and with a shoulder at 3.49 (Xa). Figure 5-71 shows the low temperature photoluminescence spectrum of the 0.2 at % As doped films grown at 120 mTorr of oxygen and it shows the DX at 3.48 eV. Since donor binding energy and coulomb interaction energy for (Mg, Zn)O are unknown, I was not able to estim ate the exciton binding energy of the donor nor the acceptor using the equati ons I earlier described.

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98 The presence of compensating defects in the 0.2 at % As doped Mg0.05Zn0.95O is confirmed by the plots of the Hall voltages against applied magnetic field in figures 5-73 and 5.74. For the films grown at 30 mTorr of oxygen, Hall voltage is low and the da ta is scatted preventing us from estimating the Hall coefficient. Compen sation was less pronounced for the films grown at 60, 90 and 120 mTorr of oxygen. The Hall voltage of the film grown at 600C and 60mTorr changed linearly with the applied magnetic fi eld indicating little to no compensation. The changes in resistivity of the films with growth pressure are seen in figure 5-72. Resistivity increases with increasing growth pressure up to 90 mTorr of oxygen where resistivity is 67.2 x 103 .cm. For the films grown at 500C in120 mTo rr and 600C and 60 mTorr of oxygen, their resistivities are 56 and 0.43 .cm, respectively. The transport properties of the films are shown in Tables 5.11 and 5.12. The undoped film gr own at 500C and 60 mTorr of oxygen has a resistivity of 0.3 .cm, a carrier concen tration of 2.44 x 1018 cm-3 and carrier mobility of 8.53 cm2/(V.s). I was not able to estimate the Hall coef ficient of the As doped (Mg, Zn)O film grown at 500C and 30 mTorr of oxygen due to the scatter in the plot of the Hall voltage against applied magnetic field, I classify the tran sport properties of this film as indeterminate. The film grown at 500C and 60 mTorr of oxygen is weakly ptype with carrier density of 4.83 x 1015 cm-3 and a mobility of 0.24 cm2/(V.s). The film grown at 500C a nd 90 mTorr was too resistive to accurately measure its transport properties. Both films grown at 500C and 120 mTorr and were n-type and had resistivities, carrier concentrations and mobilities of 56.23 and 0.43 .cm, 4.61 x 1019 and 1.46 x 1018 cm-3, 0.13 and 9.93 cm2/(V.s), respectively. The results indicate a p-type transition window at 500C between 30 mTorr and 90 mTorr. The photoluminescence properties of the 0.2 at % As doped MgxZn(1-x)O films were also affected by growth pressure and temperature. The room temperature photoluminescence for these

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99 films were low and the shows an NBE emissi on around 3.48 eV. Mg compositional variations accounts for the shifts and width of the D X in the films. The 0.2 at % As doped MgxZn(1-x)O film deposited at 500C and 60 mTorr showed an AX at 3.34 eV and a DAP 3.25 eV. I was not able to estimate the binding energy for the As related Acceptor in the films since the donor binding energy and the effects of Mg allo ying on coulomb interactions are unknown. Mg.10Zn.90O Doped with 0.02 at % As The structural and electronic properties of 0.02 % doped MgxZn(1-x)O films deposited on bare sapphire at 0.1 mTorr of O2 are described below. The target was made of high purity powder materials in the following ratio: 90 at % of ZnO, 0.10 at % MgO and 0.02 at % As from Zn2As3 as the arsenic source. The powder diffraction patterns of the films are shown in figure 5-75. The diffraction pattern of films show n multiple phases at 400 and 500C and those secondary phases disappear and only one phase of the MgxZn(1-x)O is present at 600 and 700 C. In the high resolution rocking curve of the MgxZn(1-x)O:As deposited at 400 C and 0.1 mTorr of O2, as seen in figure 576, on can observe the at least 3 peak in the figure. Backscattered electron images of the MgxZn(1-x)O:As deposited at 400 and 700C are shown in figure 5-77 and 5-78, respectively. At 400C, one can observed segregated ZnO rich particle in the shape of hexagon, bright particles. At 700C, the particles are no longer present and the surface of the film appears to be uniform. The magnesium content of the films was es timated by WDS. Mg concentration of the MgxZn(1-x)O films increases with increasing temperature from 15.9 at % for the film deposited at 400C to over 29 at % for the films deposited at 700C, figure 5-79. From the diffraction patterns, I estimated the c-axis length of the films. The length of the films decreased with increasing Mg composition from 5.245 to 5.190 for the MgxZn(1-x)O films having 16 to 29 at % of magnesium. The band gap of the films wa s estimated from the work of A. Ohtomo.108 The

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100 band gap of the films deposited at 400C, with an Mg content of 16 at % was estimated to be 3.65 eV. With increasing Mg conten t to 29 at % for the film grown at 700C, the band gap of the films increases to 3.90 eV, as shown in figure 5-80. Since c-axis decreases with increasing growth temperature, the bang gap of the films, therefore, decreases with increasing c-axis length, as seen in figure 5-81. The transport properties of the films are show n in figure 5-82. The f ilms were n-type. The resistivities of the films grown at 400 (16 at % Mg) and 500C (18 at % Mg) were on the order of 0.023 and 0.13 .cm, respectively. Resistivity increased for the film deposited at 600C (24 at % Mg) to 0.03 .cm. Resistivity rapidly increases when s ubstrate temperature was raised to 700C (29 % Mg) to about 100 .cm. Carrier density of the grow n at 400C was on the order of 5 x 1019 cm-3. Carrier concentration first increases for the films grown at 500C to 9 x 1019 cm-3, however, it decreases when substrate temp erature was further raised to 1 x 1018 cm-3 for the films deposited at 700C. Hall mobility of th e films was low, on the order of 5 cm2/(V.s) and it decreases slightly with incr easing growth temperature. Summary The transport properties of the films depe nd on the film composition, buffer layer, substrate temperature and oxygen gr owthpressure. Increasing As composition from 0.02 to 0.2 at % As decreases the resistivity and Hall mobility a nd increases carrier concentration of the films in 3 mTorr O3/O2 and 30 mTorr of oxygen. For the films doped with 2.0 at % As, there is large scatter in the data of resistivity of the films and carrier type varied sporadically. For the films that were p-type conductive, resis tivity increases whil e carrier concentra tion decreases with increasing growth pressure and decreasing As co ntent. The resistivity of the As doped (Mg, Zn)O films also increases with increasing growth pressure.

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101 The effects of persistent photoconductiv ity were analyzed by performing Hall measurements in the dark, in vacuum and air. Re sistivity of the films increases with time. Three regions were depicted in the resistivity: a re gion where resistivity incr eases exponentially with time controlled by oxygen adsorption at the su rface and crystalline in terfaces, photorelaxation and surface reconstruction; a region of saturation were resistivity is more or less stable and a third region where resistivity decreases with ti me which I believe is due to hydrogen reaction with adsorbed oxygen at the surface. The effects of increasing arsenic concentration in the films can be seen by comparing the properties of the films doped at 0.02 and 0.2 at % doped in both ozone and oxygen. As I have seen in this chapter, c-axis lengths of the ZnO films are affected by pressure of the background gas, substrate temperature and composition of the alloy. The length of the c-axis increases with increasing As composition and decreases with increasing growth temperature and background gas and those changes are attributed to the increase in cell volume due to the large As ions, concentration fluctuations due to evaporation and difference in sticking coefficient of the different ions, segregation to grain boundaries among others. Room temperature photoluminescence of th e films grown in the ozone/oxygen shows decreasing NBE emission with increasing substr ate temperature for the 0.02 % As doped films where as NBE emission for the 0.2 % As doped films increases with increasing deposition temperature. Deposited in 30 mTorr of oxygen, NBE emission intensity of the films doped 0.02 at % As increases with increasing growth temp erature where as the NBE emission of the films doped with 0.2 at % slightly de creases with increasing temper ature except for the film doped with 0.2 at% and grown at 300 C and 30 mTorr of oxygen. When doped with 2.0 at % As, NBE

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102 of the ZnO films deposited on ZnO and MgO buffe r layer at 500C increases with increasing temperature and growth pressure. The AX emission and consequently the DAP and binding energies were affected by the composition of the target, substrate temperatur e and growth pressure. AX blueshiftes with increasing As concentration and decreasing gr owth oxygen pressure. AX follows the following relationship with growth pressure: o oP X A410 02 2 343 3 (5-6) The estimated binding energy of the dopant also ch anged because shifts of the AX energies and the following relationship was depicted: b A oE eV X A 55 0 399 3 (5-7) The changes in binding energy with growth oxygen pressure and As composition can be described by the following equation: 2364 0 102O b AP meV E (5-8) and As b AC E 229 0 246 5 ln (5-9) The changes in binding energy with increasing growth oxygen were confirmed by temperature dependent Hall measurements. The acceptor thermal activation energy were in the order of 70 14 meV and 150 12 meV for the films doped with 2. 0 at % As ZnO films deposited at 500C in 5 and 150 mTorr of oxygen, respectively. The plot of acceptor binding energy against hol e concentration and the interpolation from the acceptor thermal activation energy allowed me to estimate the As-acceptor inonization energy at infinite dilution to be about 160 meV/

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103 Table 5-1. Hall data for the films doped with 0.02 at % of As and grown at 3 mTorr of ozone/oxygen. Table 5-2. Hall data for the films doped with 0. 02 at % of As and grown at 30 mTorr of oxygen. Table 5-3. Hall data for the films doped with 0.2 at % of As and grown at 3 mTorr of ozone/oxygen. Table 5-4. Room temperature Hall data for the fi lms doped with 0.2 at % of As and grown in 30 mTorr of oxygen. Growth Temperature C Resistivity .cm Carrier Concentration (x 1018 cm-3 Mobility cm2/(V s) Type 300 0.573 2.56 4.25 n 400 0.0197 35.1 9.04 n 500 16.076 Indeterminate 600 0.193 3.96 8.16 n Growth Temperature C Resistivity .cm Carrier Concentration x 1018 cm-3 Mobility cm2/(V s) Type 300 0.049 9.27 13.73 n 400 0.044 9.14 15.52 n 500 1.155 2.42 2.23 n 600 0.041 9.00 16.91 n Growth Temperature C Resistivity .cm Carrier Concentration x 1019 cm-3 Mobility cm2/(V s) Type 300 0.027 2.29 10.05 n 400 0.008 3.21 22.8 n 500 0.022 2.19 12.8 n 600 0.018 1.43 24.6 n Growth Temperature C Resistivity .cm Carrier Concentration x 1019 cm-3 Mobility cm2/(V s) Type 300 0.009 11.40 6.2 n 400 0.095 1.00 6.6 n 500 0.031 .61 33 n 600 0.175 .18 20 n

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104 Table 5-5. Summary of the structural, optical and transport properties of the films grown at 500C grown in 3 mTorr of O3/O2 and 30 mTorr of O2. Gas Comp. % As c-axis AX eV DAP eV EA meV Resis. .cm Carr. Conc. cm-3 Hall Mob. cm2/(V s) 0.02 5.196 3.31 3.23 194 16.08 Indeterminate O3/O2 0.2 5.183 3.33 3.25 180 0.022 2.2 x 1019 12.8 0.02 5.176 3.31 3.23 194 1.16 2.4 x 1018 2.23 O2 0.2 5.205 3.33 3.25 180 0.031 1.8 x 1018 20 Table 5-6. Summary of the average Hall measurem ent data complied for samples doped with 2 at % As and grown on ZnO buffer layer. Table 5-7. Hall measurement data for samples doped with 2 at % As and grown at 500C on MgO buffer layer. Growth Temp. C Growth Pressure mTorr Avg. Hall Coefficient cm-3/C Resistivity .cm Carrier Concentration cm-3 Mobility cm2/(V s) Type 400 150 -0.91 12 7 x 1018 .1 n 500 5 41.3 53 1.5 x 1017 0.8 p 20 -15.7 62.6 4.0 x 1017 0.25 n 50 476 677 1.3 x 1016 .7 p 100 -4.3 9.24 1.5 x 1018 .5 n 150 7977 7629 7.8 x 1014 .9 p 200 -763 3019 8.2 x 1015 0.3 n 600 150 -91.8 308 6.7 x 1016 0.3 n Growth Pressure mTorr Avg. Hall Coefficient cm-3/C Resistivity .cm Carrier Concentration cm-3 Mobility cm2/(V s) Type 0.1 -1.81 .113 3.43 x 1018 16.1 n 1 -1.82 .169 3.42 x 1018 10.8 n 5 223 89.6 2.8 x 1016 2.5 p 20 -0.072 0.0081 8.7 x 1019 8.9 n 50 39037 25659 1.6 x 1014 1.52 p

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105 Table 5-8. Surface roughness, in nm, of several films doped with 2 at % As. Table 5-9. Transport properties of the ZnO film s doped with 0.2 and 2.0 at % As and grown at 500C and 5 mTorr of oxygen on MgO buffer layer. Table 5-10. Transport properties of the ZnO fi lms doped with 0.2 at % As and grown 500C and on MgO buffer layer. The Target was doped with Zn3As2 and sintered at 1000C for 10 hrs. Table 5-11. Transport properties of the undoped Mg0.05 Zn0.95O film grown at 500C and 60 mTorr of oxygen on sapphire. Growth Temperature C Growth Pressure mTorr Buffer Layer Roughness nm 400 150 ZnO 7.868 none 5.438 500 5 MgO 3.125 ZnO 5.005 20 6.110 50 4.778 100 2.114 150 10.675 200 5.435 600 150 8.214 Nominal Comp, at % Resistivity .cm Hall Coefficient C/cm3 Carrier Density cm-3 Mobility cm2/(V s) Type 0.2 71.15 145.26 4.30 x 1016 2.04 p 2.0 60.57 54.24 1.15 x 1017 0.89 p Growth Pressure mTorr Resistivity .cm Hall Coefficient C/cm3 Carrier Density cm-3 Mobility cm2/(V s) Type 5 71.15 145.26 4.30 x 1016 2.04 p 150 22.09 -1.92 3.25 x 1018 0.09 n Growth Temp. Growth Pres. mTorr of O2 Resistivity .cm Carrier D ensity c m-3 Mobility Cm2/(V.s) Type 500C 60 0.3 2 .44 x 1018 8.53 n

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106 Table 5-12. Transport properties of the Mg0.05Zn0.95O:As0.002 films grown on sapphire. Growth Temp. Growth Pres. mTorr of O2 Resistivity .cm Carrier Density cm-3 Mobility Cm2/(V.s) Type 30 932.14 Indeterminate 60 5.33 x 103 4.83 x 1015 0.24 n 90 67.2 x 103 Too resistive 500C 120 56.23 4.61 x 1019 0.13 n 600C 60 0.43 1.46 x 1018 9.93 n

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107 3040506070 102103104105 2Intensity (counts/s)C o n d i t i o n s A B C D ZnO (004) ZnO (002) Al2O3 (006) Figure 5-1. Powder X-ray diffractio n of: A600C in 3 mTorr of O3/O2; B300C and 3 mTorr of O3/O2; C600C in 30 mTorr of O2; D300C and 30 mTorr of O2. 400500600700 300o C 400oC 500o C 600o C Intensity (a. u.)Wavelength (nm) 3.31 eV Figure 5-2. Room temperature photoluminescence of the 0.02 at % As ZnO films grown on sapphire and in 3 mTorr of O3/O2.

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108 400500600700 0 1 2 3 4 5 6 7 400o C 500o C 600oCIntensity (a. u.)Wavelength (nm) 300oC 3.3 eV Figure 5-3. Room temperature photoluminescence of the 0.02 at % As doped ZnO films grown on sapphire and 30 mTorr of O2. 3.123.183.243.303.363.42 0.01 0.1 Intensity (a. u.)Energy (eV) 3.316 eV 1-LO 3.361 eV DoX Xa 3.256 eV DAP Figure 5-4. Photoluminescence spectrum at 16K for an undoped film grown at 700C and 3 mTorr of (3 mol % 03 and 97 mol % 02) gas mixture.

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109 3.103.153.203.253.303.353.403.4 5 1 I n t ens it y ( a. u. ) Ener gy ( eV ) 3.309 eV AoX 3.361 eV DoX LO DAP DAP Figure 5-5. Low temperature pho toluminescence spectrum for the 0.02 at % As doped ZnO film grown at 300C and 3 mTorr of O3/O2. 3.153.203.253.303.353.4 0 0.05 0.1 0.15 0.2 0.25 40 meV Intensity (a.u.)Energy (eV) 3.361 eV DoX 3.311 eV AoX 3.194 eV LO-DAP 3.231 eV DAP or FA 3.162 eV LO-DAP 40 meV3.383 eV Xa Figure 5-6. Low temperature pho toluminescence spectrum for the 0.02 at % As doped ZnO film grown at 500C and 3 mTorr of O3/O2.

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110 3.153.203.253.303.353.403.45 0.01 0.02 0.03 0.04 0.05 Intensity (a. u.)Energy (eV) 3.363 eVDoX 3.331 eV AoX LO DAP DAP Figure 5-7. Low temperature pho toluminescence spectrum for the 0.02 at % As doped ZnO film grown at 600C and 3 mTorr of O3/O2. 3.153.203.253.303.353.403.45 0.01 0.1 Intensity (a.u.)Energy (eV) 3.383 eV Xa 3.362 eV DoX 3.232 eVDAP 3.301 eV AoX Figure 5-8. Low temperature pho toluminescence spectrum for the 0.02 at % As doped ZnO film grown at 500C and 30 mTorr of oxygen.

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111 3 .003.053.103.153.203.253.303.353.403.4 5 AoX3.324 eV 3.361 eVDoX LODAP DAP Energy (eV) As-grown 3.259 eV 3.190 eV Figure 5-9. Low temperature pho toluminescence spectrum for the 0.02 at % ZnO:As grown at 600C and 30 mTorr of oxygen. 3003504004505005506000.01 0.1 1 10 101710181019 0 2 4 6 8 10 ResistivityMobility (cm2/(V.s))Resistivity (Ohm.cm)Growth Temperature (oC)Carrier Density (cm-3) Carrier Density Mobility Figure 5-10. Resistivity, carrier concentration and mobility of the films doped with 0.02 at % As and deposited at 3 mTorr of O3/O2 as a function of substrate temperature.

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112 -6.0x103-4.0x103-2.0x1030.0 2.0x1034.0x1036.0x103-2.8x106-2.6x106-2.4x106-2.2x106 Applied Magnetic Field (G)Hall Voltage (VH) Figure 5-11. Variable magnetic field Hall measurem ent of 0.02 at % As doped ZnO film grown at 500C and 3 mTorr of O3/O2. 300350400450500550600 0.1 1 2x10184x10186x10188x10181019 2 4 6 8 10 12 14 16 18 ResistivityMobility (cm2/(V.s))Resistivity (Ohm.cm)Growth Temperature (oC)Carrier Density (cm-3) Carrier Density Mobility Figure 5-12. Changes in resistivity, carrier conc entration and mobility fo r the films doped with 0.02 at % As and deposited at 30 mTorr of O2 as a function of s ubstrate temperature.

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113 -4.0x1030.0 4.0x103-2x104-1x1040 1x104 Magnetic Field (G) x Hall Voltage RH = -2.581 Figure 5-13. Variable magnetic field Hall measur ement for the film doped with 0.02 at % As doped ZnO film grown at 500C and 30 mTorr of oxygen. 300350400450500550600 5.175 5.180 5.185 5.190 5.195 5.200 5 205 c-ax i s L eng th (A) Substrate Temperature (oC) O3/O2c-axis O2 c-axis Figure 5-14. Effects of growth temperature on the c-axis length for the 0.02 at % As doped ZnO film grown in 3 mTorr of O3/O2 and 30 mTorr of O2. The filled squares with solid lines represent the c-axis length of the films in grown in 3 mTorr of O3/O2 where as the filled circles and solid lines represent th e c-axis length of the films in grown in 30 mTorr of O2.

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114 203040506070102103104105 500oC 600oC 400oCZnO (004)T e m p e r a t u r e2 Theta ZnO (002)Al203 (006) 300oC Figure 5-15. Powder X-ray diffraction of the films doped with 0.2 at % As and grown on sapphire and at 3 mTorr of (3 mol % ozone and 97 mol % oxygen). 203040506070102103104105 ZnO (004) 400oCT e m p e r a t u r e2 600oC ZnO (002) Al2O3 (006) Figure 5-16. Powder X-ray diffraction of the Zn O films doped with 0.2 at % As and grown on sapphire and at 30 mTorr of oxygen.

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115 300350400450500550600 5.180 5.185 5.190 5.195 5.200 5.205 5.210 5.215 5.220 5.225 5.230 5.235 5.240 c-axis Length (A)Substrate Temperature (oC) O3/O2 c-axis O2 c-axis Figure 5-17. C-axis length for the ZnO films doped with 0.2 at % As as a function of growth temperature. The filled squares with solid li nes represent the c-axis length of the films in grown in 3 mTorr of O3/O2 where as the filled circles and solid lines represent the c-axis length of the films in grown in 30 mTorr of O2. 400500600700 0.0 0.5 1.0 1.5 2.0 2.5 600oC 500oC 400oCIntensity (a.u.)Wavelength (nm) 300oC 3.3 eV Figure 5-18. Room temperature photoluminescence spectra for the ZnO films doped with 0.2 at % As and grown at 3 mTorr of O3/O2.

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116 40050060070 0 600oC 500oC 400oC 300oCIntensity ( a. u. ) Wavelength (nm) 3.3 eV Figure 5-19. Room temperature photoluminescence spectra for the ZnO films doped with 0.2 at % As and grown at 30 mTorr of oxygen. 3.13.23.33.4 0.01 0.1 Intensity (a. u.)Energy (eV) 3.331 eV AoX 3.254 eV DAP 3.365 eV DoX Figure 5-20. 16K photoluminescence spectrum for the ZnO film doped with 0.2 at % As and grown at 500C and 3 mTorr of O3/O2 mixture.

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117 3.103.153.203.253.303.353.403.450.01 0.1 1 Intensity (a. u.)Energy (eV) 3.330 eV AoX 3.361 eV DoX 3.252 eV DAP Figure 5-21. PL at 16K for the ZnO film doped with 0.2 at % As and grown at 500C and 30 mTorr of Oxygen. 3.103.153.203.253.303.353.403.4 5 0.01 0.1 Intensity (a. u.)Ener g e y ( eV ) 3.307 eV AoX 3.361eV DoX Figure 5-22. PL at 16K for the ZnO film dope d with 0.2 at % As and grown at 600C and 3mTorr of O3/O2.

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118 300350400450500550600 0.01 0.015 0.02 0.025 0.03 2x10194x1019 8 10 12 14 16 18 20 22 24 26 ResistivityMobility (cm2/(V.s))Resistivity (Ohm.cm)Growth Temperature (oC)Carrier Density (cm-3) Carrier Density Mobility Figure 5-23. Transport properties of the ZnO fi lms doped with 0.2 at % As and deposited at 3 mTorr of O3/O2 as a function of substrate temperature. 300350400450500550600 0.01 0.1 101810191020 4 8 12 16 20 24 28 32 36 ResistivityMobility (cm2/(V.s))Resistivity (Ohm.cm)Growth Temperature (oC)Carrier Density (cm-3) Carrier Density Mobility Figure 5-24. Transport properties of the ZnO fi lms doped with 0.2 at % As and grown at 30 mTorr of O2 as a function of substrate temperature.

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119 0204060801001201401605.180 5.190 5.200 5.210 5.220 5.230 5.240 5.250 c-axis Length (A)Growth Pressure (mTorr) Figure 5-25. Effects of growth oxygen pressure on the c-axis spacing of the ZnO films doped with 2 at % As and deposited on thin ZnO buffer layer. 051015202530 102103 Intensity A ngle (o) 5 20 50 150 200 Figure 5-26. High resolution rocki ng curves for ZnO films doped w ith 2 at % As and grown at 500C and different oxygen partial pressures.

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120 050100150200 0 2 4 6 8 10 Full Width at Half Maximum (o)Growth Pressure (nm) Figure 5-27. Effects of growth pressure on the crystalline qualities (FWHM) of the ZnO films doped with 2 at % As and grow n at 500C on ZnO buffer. 131415161718192021102103104 IntensityAngle (o) 600oC 500oC Figure 5-28. High resolution rocki ng curves for ZnO films doped w ith 2 at % As and grown at 500 and 600C and 150 mTorr of oxygen. FWHM are 2.323 and 0.363, respectively.

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121 151617181920 102103 I ntens i tyAngle (o) MgO Buffer ZnO Buffer Figure 5-29. High resolution rocki ng curves for ZnO films doped w ith 2 at % As and grown at 500C and 5 mTorr of oxygen on thin ZnO and MgO buffer layers. FWHM are 0.975 and 0.795, respectively. 340360380400420440460480500 I n t ens it y ( a.u ) Wavelength (nm) 50 100 5 20 200 150 Figure 5-30. 300K PL spectra for ZnO films doped with 2 at % As and grown on ZnO buffer at 500C.

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122 340360380400420440460480500 50 mTorrIntensity (a. u.)Wavelength (nm) 5 mTorr 20 mTorr 0.1 mTorr 10 mTorr Figure 5-31. PL at 300K spectra for ZnO films doped with 2 at % As and grown on MgO buffer layer at 500C. Figure 5-32. Room temperature PL for ZnO film s doped with 2 at % As and grown in 150 mTorr PO2 and various temperatures (C) on thin ZnO buffer layer. 350375400425450475500 Wavelength (nm) 600 500 400

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123 3.123.183.243.303.363.42 0.01 0.1 Intensity (a. u.)Energy (eV) 3.316 eV 1-LO 3.361 eV DoX Xa 3.256 eV DAPa: undoped films AoX 3.103.153.203.253.303.353.403.453.5 0.06 0.08 0.1 0.12 0.14 0.16 0.18 0.2 0.22 0.24 Intensity (a. u.)Energy (eV)b: 400oC, 150 mTorr 3.386 eV DoX AoX

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124 3.103.153.203.253.303.353.403.45 0.02 0.04 0.06 0.08 0.1 Intensity (a. u.)Energy (eV) c: 500oC, 50 mTorr 3.252 eV LODAP + LOFA 3.211 eV 2LODAP + 2 LO FA 3.331 eV AoX 3.279 eV DAP + FA 3.361 eV DoX 3.103.153.203.253.303.353.403.45 1 10 100 Intensity (a. u)Energy (eV) 3.362 eV DoXd: 500oC, 100 mTorr AoX

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125 2.93.03.13.23.33.4 0.1 Intensity (a. u)Energy (eV) 3.362 eV DoX 3.316 eV AoX 3.179 eV LO(DAP+ FA) 3.261 eV DAP + FA e: 500oC, 150 mTorr 2.93.03.13.23.33.4 0.01 0.1 1 Intensity (a. u.)Energy (eV) 3.361 eV DOX 3.30 eV AoX 3.226 eV DAP 3.047 eV LO DAP f: 500oC, 200 mTorr

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126 3.103.153.203.253.303.353.40 1 10 2LO-DAP + 2LO-FAIntensity (a. u.)Energy (eV) 3.342 eV (AoX) 3.314 eV DAP+ FAg: 500oC, 5 mTorr on MgO 3.361 eV DoX LO-DAP + LO-FA Figure 5-33. Low temperature Ph oto Luminescence spectra of se lected 2 at % As doped ZnO films collected at 30K. Conditions are indicated on the each graph. 050100150200 0.10 0.12 0.14 0.16 0.18 3.20 3.25 3.30 3.35 Energy (eV)Growth Pressure ( mTorr ) AoX DAP Eb A Figure 5-34. Dependence of Acceptor-bound excit on (AX), Donor-Acceptor Pair transition (DAP) emission energies and Acceptor binding energy (EbA) on oxygen growth pressure.

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127 050100150200 3.30 3.31 3.32 3.33 3.34 A0X (AoX) = 3.34272.02 x 10-4PoAoX (eV)Growth Pressure (mTorr) Figure 5-35. Dependence of Acceptor-bound exci ton (AX) emission energy on growth oxygen pressure. 3.303.313.323.333.34 0.10 0.12 0.14 0.16 0.18 Eb A (AoX) = 3.399 0.55Eb A Binding Energy (eV)Acceptor-bound exciton (eV) Figure 5-36. Dependency of Acceptor binding energy (Eb A) on acceptor-bound exciton (AX) emission energy.

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128 050100150200 100 120 140 160 180 Eb A Eb A = 102 meV + 0.364PO2Acceptor Binding Energy (meV)Growth Oxygen Pressure (mTorr)a 1x1052x1053x1054x1055x10590 100 110 120 130 140 150 Eb a= E0 (p1/3) Eo= 163 +/15 meVBinding Energy (meV)p1/3 Figure 5-37. Dependency of acceptor binding energy (Eb A) on growth pressure (a) and hole concentration (b). b

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129 020406080100120140160180200220101102103 nIndep pResistivity (ohm.cm)Growth Pressure ( mTorr ) p n Figure 5-38. Resistivity as a func tion of growth Pressure for the ZnO films doped with 2.0 % As and grown on ZnO buffer layer. 39042045048051054057060010 100 1000 Resistivity (Ohm*cm)Growth Temperature ( C)p n n Figure 5-39. Resistivity as a func tion of growth temperature for the ZnO films doped with 2 at % As and grown on ZnO buffer layer at 150 mT orr of oxygen. See Table 5.1 for details.

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130 01020304050 10-1100101102103104 Resistivity (ohm.cm)Growth Pressure (mTorr) Figure 5-40. Resistivities of the ZnO:As0.02 films grown on MgO buffer layer at 500C. -4x107-4x107-4x107-4x107-4x107-4x107 500C; 150 mTorr O Slope = Rh = 41.182 -6.0x104-4.0x104-2.0x1040.0 2.0x1044.0x1046.0x104-1.4x108-1.2x108 500C; 5 mTorr O Slope = Rh = 222.962Magnetic Field (Oe)VH Figure 5-41. Hall voltages the ZnO:As0.02 films grown at 5 and 150 mTorr of oxygen pressure.

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131 345678910 10121013101410151016 102103 Carrier Density Ea A= 87 +/14 meVCarrier Concentration(cm-3)1000 T-1(1000K-1)Resistivity (Ohm.cm) Resistivity 5 mTorr of O2a 3.03.54.04.55.0 101410151016 103104 Carrier Density Ea A= 140 +/8 meV Resistivity (Ohm.cm)Carrier Concentration (cm3 )1 000 T-1 ( K-1 ) Linear Fit 150 mTorr of O2 Resistivity b Figure 5-42. Carrier density and resistivity Vs reciprocal temp erature for the ZnO films doped with 2 at % As and grown at 500C and 150 (b) and 5 (a) mtorr of oxygen.

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132 34567 8 1E11 1E12 a: 5000C and 5 mTorr O2 Ea A= 70 +/-14 meVP(T)/T3 (cm-3. K-3 )1000/T(K) 3.03.54.04.55.0 1E11 1E12 1E13 b: 500oC and 150 mTorr O2 Ea A= 154 meV +/12 meVp(T) T-3/21000/T (1000K-1) Figure 5-43. Charge balance equation analys is of the hole concentration for the ZnO:As0.02 samples grown at (a): 500C and 5 mTo rr of oxygen on MgO Buffer layer and (b): 150 mTorr of oxygen on ZnO buffer layer.

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133 0501001502004 6 8 10 Growth Pressure (mTorr) Surface Roughness Figure 5-44. Surface roughness as a function of gr owth temperature for the ZnO film doped with 2 at % As and grown at 500C and on ZnO buffer layer. 050100150 10-1100101102103104 101510161017 Growth Oxygen Pressure (mTorr) Carrier Concentration (cm-3)Carrier Mobility Hall Coefficient Resistivity Eb A Figure 5-45. Changes in car rier concentration, resi stivity, Hall coefficient and Hall mobility as a function of growth pressure for the ptype ZnO films doped with 2 at % As.

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134 100110120130140150 1E15 1E16 1E17 Hole Concentration (cm-3)Binding Energy (meV) 5 mTorr 50 mTorr 150 mTorr Figure 5-46. Dependency of hole con centration on Acceptor binding energy (Eb A). 020406080100120140160 5.19 5.20 5.21 5.22 5.23 5.24 5.25 c-axis Length(A)Growth Pressure (mTorr) Figure 5-47. C-axis length of the ZnO films doped with 2 at % As and grown at 500C at 5, 50 and 150 mTorr of oxygen.

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135 0.00.51.01.52.0 Eb A lnEb A = 5.246 0.229CAsAcceptor Binding Energy, Eb A, (meV)Arsenic Concentration (at. %) Figure 5-48. Dependency of Acceptor binding energy (Eb A) on Arsenic concentration. -10.0k-5.0k0.05.0k10.0k 9.4x1079.5x1079.6x1079.7x1079.8x107 Applied Magnetic field (G) V H 0.2 at. % As 1.2x1061.6x1062.0x1062.4x106 2.0 at % As Figure 5-49. Hall voltage as a function of applie d magnetic field for the ZnO films doped with 0.2 and 2.0 at % As grown at 500C and 5 mTorr of O2 on MgO buffer.

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136 0.00.51.01.52.0 1.0 1.5 2.0 60 80 100 120 140 160 5x101610171.5x101 As Concentration (nominal target As at %) Eb A(meV) Hall Coefficient Carrier Density Mobility Resistivity Figure 5-50. Transport properties of the ZnO film s doped with 0.2 and 2 at % As grown at 500C and 5 mTorr of oxygen on MgO buffer. 350400450500550600650700 0.1 Intensity (a. u.)Wavelength (nm) 2.0 at. % As 0.2 at % As Figure 5-51. PL at 300K spectra of the ZnO film s doped with 0.2 and 2.0 at % As and grown at 500C and 5 mTorr of oxygen on MgO buffer layer.

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137 -1x1040 1x104-4x104-3x104-2x104-1x1040 1x1042x104 VHApplied Magnetic Field (G) 0.2 at. % As, 150 mTorr O2 RH = -1.92 Figure 5-52. Hall coefficient as a function of applied magnetic field for the ZnO films doped with 0.2 at % As and grown at 500C a nd 150 mTorr of oxygen on MgO buffer layer. 02040608010012 0 15 20 25 Resistivity (hm.cm)Time (hr) O2 Surface adsoption Hydrogen adsortion Saturation Figure 5-53. Effects of Persis tent PhotoConductivity on the resistivity when the ZnO films grown at 500C and 5 mTorr of oxygen on MgO buffer is measured in vacuum (0.1 mTorr).

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138 -6.0x104-4.0x104-2.0x1040.0 2.0x1044.0x1046.0x104-1.3x10-3-1.2x10-3-1.1x10-3-1.0x10-3-9.0x10-4-8.0x10-4-7.0x10-4-6.0x10-4-5.0x10-4-4.0x10-4 24 hrs Applied Magnetic Filed (G)VH Figure 5-54. Variation of Hall voltage with a pplied magnetic field after the ZnO sample doped with 2 at As and grown at 500C in 5 mTorr oxygen pressure on MgO buffer layer was placed in the dark and in vacuum for 24 hrs. -6.0x104-4.0x104-2.0x1040.0 2.0x1044.0x1046.0x104-3.8x10-3-3.8x10-3-3.7x10-3-3.6x10-3-3.6x10-3-3.5x10-3-3.5x10-3-3.4x10-3-3.4x10-3 V HApplied Magnetic Field (G) 37 hrs Figure 5-55. Variation of Hall voltage with a pplied magnetic field after the ZnO sample doped with 2 at As and grown at 500C in 5 mTorr oxygen pressure on MgO buffer layer was placed in the dark and vacuum for 37 hrs.

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139 -6.0x104-4.0x104-2.0x1040.0 2.0x1044.0x1046.0x104-5.0x10-3-4.9x10-3-4.9x10-3-4.8x10-3-4.8x10-3-4.7x10-3 VH App lied Ma g netic Field ( G ) 42 hrs Figure 5-56. Variations of Hall voltage with applied magnetic field after the ZnO sample doped with 2 at As and grown at 500C in 5 mTorr oxygen pressure on MgO buffer layer was placed in the dark and in vacuum for 42 hrs. 05101520253.65 3.70 3.75 3.80 3.85 3.90 3.95 Resistivity (Ohm.cm)Time (hrs) Figure 5-57. Resistivity as a function of time in the dark and measured in AIR for the ZnO sample doped with 2 at % As grown at 500C and 100 mTorr of oxygen on ZnO buffer layer.

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140 -1.0x104-5.0x1030.0 5.0x1031.0x104-6.5x105-6.0x105-5.5x105-5.0x105-4.5x105-4.0x105 30 minHall voltage VH App lied Ma g netic Field ( G ) Figure 5-58. Variation of Hall voltage with a pplied magnetic field for the ZnO sample doped with 2 at % As grown at 500C and 100 mTorr of oxygen on ZnO buffer layer after being placed in the dark for 30 minutes in AIR. -10.0k-5.0k0.05.0k10.0k -7x105-6x105-5x105-4x105-3x105 12 hrsApplied Magnetic Filed (G)VH Figure 5-59. Variation of Hall voltage with a pplied magnetic field for the ZnO sample doped with 2 at % As grown at 500C and 100 mTorr of oxygen on ZnO buffer layer after being placed in the dark for 12 hours in AIR.

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141 -10.0k-5.0k0.05.0k10.0k-3.0x105-2.5x105-2.0x105-1.5x105-1.0x105-5.0x1040.0 Hall Voltage VHApplied Magnetic Field (G) 24 hrs Figure 5-60. Variation of Hall voltage with a pplied magnetic field for the ZnO sample doped with 2 at % As grown at 500C and 100 mTorr of oxygen on ZnO buffer layer after being placed in the dark for 24 hours in AIR.

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142 34.483 41.683 34.483 41.683 34.483 41.683 34.533 41.683 20304050607080102103104105 MgZnO (004)Intensity (counts/s)MgZnO (002) Indium contactsAl2O3(002) a: 500oC, 30 mTorr of O2 20304050607080 b: 500oC, 60 mTorr of O2 MgZnO (004)Al2O3(006)MgZnO (002)2 20304050607080102103104105 c: 500oC, 90 mTorr of O2 MgZnO (004)MgZnO (002)Al2O3(006) 20304050607080 d: 500oC, 120 mTorr of O2 MgZnO (004)Al2O3(006) MgZnO (002) Figure 5-61. Powder XRD of the Mg0.05Zn0.95O:As0.002 films grown in O2 at 500C [a, b, c, d: 30, 60, 90 and 120 mTorr of oxygen, respectively]. 20253035404550556065707580102103104105 34.533 41.683 MgZnO (004) MgZnO (002) Al2O3(006)Intensity (counts/s)2 600oC, 60 mTorr Figure 5-62. Powder XRD of the Mg0.05Zn0.95O:As0.002 films grown at 600C and 60 mTorr of O2.

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143 204060801001205.191 5.192 5.193 5.194 5.195 5.196 5.197 5.198 5.199 c-axis length (A)Growth Pressure (mTorr) Mg0.5Zn.95O at 500oC 5.93 at % Mg 6.83 at % Mg Figure 5-63. c-axis spacing as a function of growth O2 pressure for the Mg0.05Zn0.95O:As0.002 films grown at 500C on sapphire. 500600 5.191 5.192 5.193 5.194 5.195 5.196 5.197 5.198 5.199 5.200 c axis length (A)Growth temperature (oC) 60 mTorr Figure 5-64. c-axis spacing as a func tion of growth temperature for the Mg0.05Zn0.95O:As0.002 films grown in 60 mTorr of O2.

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144 121416182022 101102103104 Intensity (counts/s)120 mTorr of O2 30 mTorr of O2 60 mTorr of O2 Figure 5-65. -rocking curves of the Mg0.05Zn0.95O:As0.002 films grown at 500C on in 30, 60 and 120 mTorr of oxygen and on sapphire. 20406080100120 1.4 1.6 1.8 2.0 2.2 2.4 2.6 FWHMGrowth Pressure (mTorr of O2) Figure 5-66. FWHMs of the -rocking curves of the Mg0.05Zn0.95O:As0.002 films grown at 500C on in 30, 60 and 120 mTorr of oxygen and on sapphire.

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145 400500600700 3.48 eV 3.48 eV 3.48 eV 600oC and 60 mTorr O2500oC and 120 mTorr O2500oC and 90 mTorr O2IntensityWavelength (nm) 500oC and 60 mTorr O2 3.48 eV Figure 5-67. Room temperature luminescence of the Mg0.05Zn0.95O:As0.002 films grown in oxygen and on sapphire. 3.13.23.33.43.53.6 0.02 0.04 0.06 0.08 0.1 0.12 0.14 Intensity (a. u.)Energy (eV) Undoped, 500oC, 60 mTorr of O2 3.471 eV DoX Figure 5-68. PL at 25K of an undoped Mg0.05Zn0.95O films grown in 60 mTorr of O2.

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146 3.153.203.253.303.353.403.453.503.55 Energy (eV) 3.49 eV Xa 3.469 eV DoX LO DAP 3.424 eV AoX 500oC and 60 mTorr DAP Figure 5-69. Photoluminescence at 25K of the Mg0.05Zn0.95O:As0.002 films grown 500C and 60 mTorr of oxygen on sapphire. 3.203.253.303.353.403.453.503.55 0.02 0.03 Intensity (a. u.)Ener gy ( eV ) 3.502 eV Xa 3.468 eV DoX? 500oC, 90 mTorr of O2 Figure 5-70. Photoluminescence at 25K of the Mg0.05Zn0.95O:As0.002 films grown 500C and 90 mTorr of oxygen on sapphire.

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147 3.13.23.33.43.53.6 0.016 0.018 0.02 0.022 0.024 0.026 0.028 0.03 Intensity (a. u.)Energy (eV) 120 mTorr of O2 3 483 e V DoX Figure 5-71. PL of the Mg0.05Zn0.95O:As0.002 films grown at 500 C and 120 mTorr of O2. 20406080100120101102103104 Resistivity (Ohm.cm)Growth Pressure (mTorr of O2) Figure 5-72. Resistivity against O2 growth pressure for the Mg0.05Zn0.95O:As0.002 films grown at 500C.

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148 -10.0k-5.0k0.05.0k10.0k-2.7x107-1.8x107 x VH a: 30 mTorr of O2 RH is indeterminate-10.0k-5.0k0.05.0k10.0k0.0 5.0x107 b: 60 mTorr of O2 RH= 1293.07-10.0k-5.0k0.05.0k10.0k-1.0x1013 c: 90 mTorr of O2-10.0k-5.0k0.05.0k10.0k-2.0x105 d: 120 mTorr of O2 RH = -7.382Applied Magnetic Field (G) Figure 5-73. Dependence of Hall voltage on applied magnetic field for the Mg0.05Zn0.95O:As0.002 films grown at 500C [a, b, c, d: 30, 60, 90 and 120 mTorr of oxygen, respectively]. -10k010k -3x104-2x104-1x1040 1x1042x1043x1044x104 Applied Magnetic Field (G) 600C, 60 mTorr of O2 RH= -4.268Hall Voltage Figure 5-74. Hall voltage Vs app lied magnetic filed for the Mg0.05Zn0.95O:As0.002 films grown at 600C and 60 mTorr of O2.

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149 32364044 102103104105 MgxZn1-xO (002)Silver Painta: 400oCIntensity (counts/s) Al2O3(006) MgO(111)32364044 MgxZn1-xO (002) MgO(111)b: 500oCAl2O3(006)Silver Paint2 32364044 102103104105 Silver PaintSilver Paint Silver Paintc: 600oCMgxZn1-xO (002)Al2O3(006) 32364044 102103104105106 MgxZn1-xO (002)Silver Paintd: 700oCAl2O3(006) Figure 5-75. Powder XRD patterns of the Mg0.1Zn0.9O:As0.002 films grown in 0.1 mTorr of O2 [a: 400, b: 500, c: 600 and d: 700C]. Figure 5-76. High resolution w-rocking curve of the Mg0.1Zn0.9O:As0.002 deposited at 400C and 1E-4 mTorr of O2.

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150 Figure 5-77. Backscattered electron image of the Mg0.1Zn0.9O:As0.002 deposited at 400C and 1E4 mTorr of O2. Figure 5-78. Backscattered electron image of the Mg0.1Zn0.9O:As0.002 deposited at 700C and 1E4 mTorr of O2.

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151 400450500550600650700 14 16 18 20 22 24 26 28 30 Mg Composition (at. %)Temperature (oC) Figure 5-79. Variations of Mg cont ent with substrate temperature. 1618202224262830 5.19 5.20 5.21 5.22 5.23 5.24 5.25 3.60 3.65 3.70 3.75 3.80 3.85 3.90 3.95 c-axis lengthc-axis (A)Composition (At. %)Estimated Band Ga p ( eV ) Estimated band gap Figure 5-80. Variations of c-axis of the f ilms and estimated band gap with Mg content.

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152 5.195.205.215.225.235.245.25 3.65 3.70 3.75 3.80 3.85 3.90 Estimated Band Gap (eV)c-axis (A) Figure 5-81. Variations the band ga p with c-axis of the films. 1416182022242628300.01 0.1 1 10 100 101810191020 3 4 5 6 7 400oC 500oC600oCResistivity (hm.cm)Mg content (at. %) Resistivity Carrier Density Mobility700oC Carrier Density (cm-3) Mobility (cm2 V-1s-1) Figure 5-82. Transport properties of the film s as a function of Mg content and growth temperature.

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153 CHAPTER 6 STABILITY OF THE ARSENIC DOPED ZnO FILMS Results and Discussion In this chapter I discuss the st ability of the arseni c doped films and the various factors that affects the films properties. The films were depo sited from targets doped with different arsenic sources, concentration and proce ssed differently. Rapid thermally annealed were performed in nitrogen and oxygen. Tube furnace anneals we re done in one atm of flowing Ar and O2 to understand the thermal stability of the films. I expl ored the effects of agin g and coating the films with photoresist. The electrical structural and optical char acterizations were performed according to the procedures outlined earlier. Low Temperature RTA study RTA Temperature Study For this experiment, several samples doped with 2 at % As from zinc arsenide were grown at 500C and 50 mTorr of oxygen pa rtial pressure on a thin Zn O buffer layer and were post growth rapid thermal annealed for 200 to 500C at 100C increment for 60 seconds in nitrogen. The samples were then placed in the dark befo re Hall measurements were performed according to the procedure described in chapter 2. The high-resolution omega-rocking curves, figure 6-1, of the as-grown film and the film RTA at 500C for 60 seconds in nitrogen show that crystalline quality of the film was not affected by the anneal. The FWHM of the as-grown film is 2.495 as compared to the 2.322 for the FWHM of the film subjected to an RTA at 500C. The large FWHMs are attributed to low growth temperature (500C) and relatively high growth pressure (50 mTorr of oxygen). At high background pressure, collisions between the pa rticles in the plume and the gas increase dramatically causing the ablated species to arrive at the substrates su rface with lower kinetic

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154 energy. It is not clear whether the slight difference in FWHM is due to sample inhomogeneity or annealing effects. The 60 seconds RTAs did not alter the room temperature photoluminescence spectra of the films as shown in figure 6-2. Both near band edge at 3.3 eV, and visible emissions of the films remain the same. At 30K, photolum inescence measurements of the f ilms, as shown in figure 6-3, reveal broad peaks comprised of the DX at 3.361 eV and the near by AX at 3.328eV. The presence of the AX becomes more apparent with increasing RTA temperature as the shoulder at 3.328 eV becomes more pronounced from 300 to 500C. The transport properties of the films measur ed by Hall measurements of the films doped with 2 at % As from zinc arsenide and gr own at 500C and 50 mTorr of oxygen and RTA for 60 seconds in nitrogen are summarized in table 6-1. Resistivity of th e films decreases with increasing N2 RTA temperature, as show n in figure 6-3, from 510 .cm to 1.6 .cm for the film annealed at 400C. The Hall coe fficient decreases with increasi ng RTA temperature and there is a carrier type conversion from p to n for annea ling temperature greater than 200C as shown in table 6-1. The plots of the variation of the Hall voltage of the films with applied magnetic field, figure 6-5, show carrier type c onversion from the compensated as-g rown film to the n-type film when subjected to an RTA at 400 C. The carrier concentration of the as-grown film could not be calculated due to the scatter in the data. Carrier density and mobility of the film RTA at 200C were 1.22 x 1017 cm-3 and 0.25 cm2/(V.s), respectively. For RTA temperature above 200C, there is a carrier type convers ion from pto n-type and the magnitude of both carrier density and Hall mobility increases with increasing RTA temper ature as shown in figure 6-6 and 6-7. The transport properties of the films RTA at 400 a nd 500C did not change si gnificantly. Resistivity, carrier concentration and mobility were es sentially identical, on the order of 1.6 .cm, 18.91 x

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155 1017 cm-3 and 4.3 cm2/(V.s), correspondingly. In the case of a compensated semiconductor, the Hall coefficient is defined as: 2 2 2) (n p n p Hn p e n p R (6-1) where n, p, n and p are electrons, holes and electron and hole mobility, respectively. The results indicate increasing electron in the film causing the RH to change sign and |RH| to increase since carrier concentration is inve rsely proportional to the Hall voltage. RTA Time Study For this experiment, several samples, which were grown at 500C and 20 mTorr of oxygen partial pressure on a thin ZnO buffer layer from a target doped with 2 at % As from zinc arsenide, were rapid thermal annealed at 200 C for 30-300 seconds in nitrogen. The samples were then placed in the dark for Hall measuremen ts to be performed according to the procedure described in chapter 2. Photoluminescence measurements collected at room temperature, as represented in figure 6-8, reveal that the spectra of the films subj ected to RTA at 200C from 30 to 300 seconds are almost identical with low NBE em ission at 3.3 eV and emission in the visible present as well. Photoluminescence at 30K reveal broad DXs, DAP and LO-DAP at 3.36, 3.324 and 3.169 eV, respectively. I estimated the acceptor binding energy EA from the DAP transitions of the films grown at 500C in both envi ronment using equation (6-2) 138: ] ) ( [3 / 1D b D g AN DAP E E E E (6-2) I have estimated the binding energy the donor re sponsible for the emission at 3.361 to be 58 meV in the previous chapters. ND is the carrier concentration derived from the Hall measurements and is materials dependent constant and is 2.7 x 10-8 eV.cm for (ZnO) = 8.6.

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156 The term3 / 1DNis the Coulomb energy assuming an av erage the distance between dopants is 3 / 1DN From equation 2, I find an estimated an acceptor binding energy of 152 meV. The activation energy is closed to the value predicted for E0 I found in the earlier chapter. This suggests very low acceptor concentration. Howeve r, the accuracy of the current estimate is limited by the high resistivity of the films. The temperature dependence of photolumin escence of the film RTA at 200C for 60 second is shown in figure 6-10. The DX broadens and red-shifts with increasing temperature from 3.36 eV at 30K to 3.30 eV at 300K. The DAP peak red-shift was less severe and its intensity increases with increas ing temperature, suggesting in creasing acceptor concentration with increasing temperature. At 50K, the FA p eak emerges. Its intensity first increases with increasing measurement temperature and then is no longer observable as sample temperature approaches room temperature. The FA at 3.25 eV indicates an acceptor activation energy around 153 meV. Table 6-2 lists the resistivities of the films as a function of annealing time in seconds. Since most of the films e xhibited resistivity above 104 .cm, I did not calculate carrier concentrations and mobilities as they are pron e to errors. Figure 611 shows the plots of resistivity as a function of RTA time. The data is scattered. Howeve r one can observe an increasing trend of resistivity with increasing RTA temperature. It is not clear whether the trend is due to samples inhomogeneity. Table 6-3 summarizes the transport properties fo r several other films annealed at 200C for 60 seconds. Growth conditions are indicated in th e table. For film A, annealing yielded little change the film properties. For film B, both resistivity and the Hall coefficient decreased reflecting a decrease in both carrier concentratio n and mobility. Both films remained n-type. For

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157 the remaining films (C, D, E), which were p-type resistivities decreased dramatically with the 200C RTA. Carrier concentration increased up to 3 orders of ma gnitude (for E) while carrier mobility decreased. Above, I have presented the effects of low temp erature rapid thermal anneal in nitrogen on various ZnO films doped with arsenic and gr own under different conditions. High resolution rocking curves show minimal change in the microstructure after 500C RTA and photoluminescence, measured at 300K and 30K, of th e films show identical feature in the spectra regardless of the temperature of the 60 sec RTAs or the length of the 200C RTAs. Hall measurements of the films reveal increasing compensation in the films and carrier type conversion when annealed above 200C. It is clear that RTA at low temperature did not alter the bulk properties of the films. I suspect that changes in resistivity and RH are surface related. Possible explanations include: absorption of hydrogen; migration of hydrogen to the surface, formation of an electron accumulation layer at the surface; desorption of o xygen and/or arsenic near the surface and/or surface relaxation or reconstruction. High Temperature Annealing Nitrogen RTA at 800C A set of p-type films doped with 2.0 at % As and deposit ed at 500C and 150 mTorr of oxygen on thin ZnO buffered sapphire were subjected to RTA in nitrogen at 800C. The films became n-type and conductive even for an RTA as short as 5 seconds. Figures 6-12 and 6.13 show the changes in resistivity, carrier density and mobility as a function of 800C N2 RTA duration. Prior to annealing, an as-grown film had a resistivity, carri er concentration, and mobility on the order of 8 x 103 .cm, 7.8 x 1014 cm-3 and 0.9 cm2/(V.s), respectively. The resistivity of the annealed films first rapidly d ecreases with RTA length for the first 60 seconds.

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158 Beyond 60 seconds, the changes in resistivity were less severe and it levels off at around 120 seconds. The carrier concentration and mobility of the f ilms increase with RTA duration. Carrier concentration rapidly increases wi th RTA time for the first 60 s econds, then slowly decreases as RTA time increases. The Hall mobility of the films, however, increases continuously with increasing RTA time from 0.9 to 24 cm2/(V.s) from the as-grown film to the film RTA for 180 seconds. The room temperature PL spectra are sh own in figures 6-14 and 6-15. Near band edge luminescence of the films remain s at 3.3 eV and its intensity increases with increasing RTA length. There is an apparent blueshift in the emission in the visible and its intensity drops significantly for the 120 and 180 seconds RTAs. Effects of Growth Temperature on Post-Annealing Properties When a series of the films doped with 0.02 at % of As, grown on ba re sapphire and in 3 mTorr of ozone/oxygen were annealed in flowin g high purity oxygen in a tube furnace at 800C for 10 min, resistivity increased several orders of magnitude as seen in figure 6-18. The resistivity of the films doped with 0.02 at % of As and grown at 300 and 400C was on the order of 0.5 .cm as-grown. This increases to over 100 k .cm when annealed. The changes in resistivity are severe for the f ilms grown at 500 and 600C. Resistivities went from 16 and 0.2 .cm of the 600C as-grown to 998 and 5.84 .cm when annealed, respectively. The data is tabulated in table 6-7. The room temperature photoluminescence of the as-grown and annealed films are shown in figure 6-19 and 6-20, respectively. For the as-gro wn films NBE emission is the dominant feature of the room temperature spectra while visibl e emission becomes the dominant feature in the spectra of the annealed films except for the anneal ed film grown at 300C. For the films grown at 300C and annealed at 800C in oxygen for 10 mi nutes, NBE remains present and significant and visible emission is more significant when compar ing the as-grown and annealed PL spectra.

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159 When two of the films doped with 0.02 at % As grown on bare sapphire and 30 mTorr of O2 were tube furnace annealed at 800C for 10 min in flowing oxygen, the transport properties also changed as seen in table 6-8. The film grown in 30 mTorr of O2 at 300C becomes semiinsulating once annealed while the film resist ivity grown at 600C remains semiconducting. Resistivity, carrier concentration and mobility of the film deposited at 600C and 30 mTorr of O2 changes from 0.041 to 23.13 .com, 9 to 1.8 x 1018 cm-3 and 16.91 to 21.02 cm2/(V.s). The room temperature photoluminescence spectra, figure 6-21, shows the NBE emission of the films is considerably reduced compared to the as-grown and visible emission blue-shifts and becomes the most dominant feature for the spec tra of the annealed films. Photoluminescence of the 0.02 at % As doped film and deposited at 500C and 30 mTorr of O2 collected at 20K, figure 6-22, reveals a more distinct AX peak at 3.324 eV when the film is annealed as opposed to the as-grown film. The Hall properties of a series of films dope d with 0.2 at % As on bare sapphire and 3 mTorr of ozone/oxygen were annealed in flowin g oxygen at 800C for 10 min are listed in table 6-9 and the changes of resistivit y due to annealing as a function of growth temperature is shown in figure 6-23. Upon annealing, the resistivity of the 0.2 at % As doped films increases several orders of magnitude from 0.027 and 0.018 .cm of the as-grown state of the films grown at 300 and 600C to 103.64 and 62.24 .cm, respectively, when anneal ed at 800C in flowing oxygen. The carrier concentration and mobility of the 0.2 at % As doped films changed from 2.3 and 1.4 x 1019 cm-3 and 10 and 15 cm2/(V s) for the as-grown state of films deposited at 300 and 600C to 1.3 and 0.3 x 1013 cm-3 and 0.45 and 3.21 cm2/(V s), respectively. Fo r the 0.2 at % As doped film grown at 600C an d 3 mTorr of ozone/oxygen, the film becomes semi-insulating when annealed.

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160 The room temperature photoluminescence spectr a of both the as-grown and annealed films are shown in figure 6-24. NBE emission intensity of the annealed films is considerably lower compared to the as-grown films and its intens ity appears to decrease with increasing growth temperature. Visible emission, from 400 to 700 nm, dominates the spectr a and increases with increasing growth temperature. Low temperature photoluminescence of the f ilms doped with 0.2 at % deposited at 600C and 30 mTorr of ozone/oxygen is shown in figure 6-25. The AX peak blue shifted after the anneal to 3.325 eV from 3.310 eV The AX beco mes the most intense p eak of the spectrum. The DX peak can be observed at 3.3 61 eV and a DAP shoulder at 3.250 eV. When a set of films doped at 0.2 at % As, grown on bare sapphire and in 30 mTorr of oxygen, were annealed in flowing oxygen at 800C for 10 min, transport properties were also altered as seen in figures 6-27 (a ) & (b) and table 6-10. Overall, the resistivity of the annealed films increased several orders of magnitude from the as-grown values. Resistivity of the annealed films decreases with increasing growth temperature. Although th e carrier concentration of the annealed films are lower than the as-grown, Hall mobility of the annealed films were lower as compared to the annealed films. Room temperature luminescence of the films, as seen in figure 6-26, shows NBE emission peaks of the films grown at 300 a nd 400C have merged with emissi on in the visible, while there is no NBE emission for the film grown at 500C. Effects of Growth Pressure on F ilm Properties After Ar Annealing Figure 6-28 shows the changes of resistivity as a function of growth pressure of a series of films doped with 2 at % As and grown at 500C and different growth pressure. The Hall characteristics of those films were explored in chapter 5. Upon a 600C, 10 min, flowing Ar anneal, the films becomes very n-type conductive The films grown at 5, 50 and 150 mTorr of

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161 oxygen which were resistive and p-type as-gro wn with carrier concentration in the 1014 1017 cm-3 changed to n-type with carrier concentration in 1018 cm-3. It is worth nothing that even though the resistivity of the annealed films (filled circle in the plot) dropped significantly from the as-grown values of the films (filed square s), the dependency of re sistivity in growth conditions remains the same despite the 10 min Ar annealed. The transport properties of the films are summarized in table 6-15. Effects of Target Processing The targets used in this experiment were doped with 2 at % As from zinc arsenide. Target 1 and 2 were doped with 2 at % As from zinc arse nide. Target 1 was sinter ed in air at 1000C for 10 hrs and Target 2 was sintered in Ar at 600C for 8 hrs and ta rget. Target 3 was doped 2 at % As from zinc di-arsenide and sintered in Ar at 600C for 8 hrs. The samples were deposited on thin MgO buffer layer. I labeled Sample A and Sample B the films deposited from Target 1 and 2, respectively. Several samples were grown at once to perform the RTA. Powder diffraction x-ray, not shown, reveals that the samples were epitaxial and highly textured in the ZnO [002]. Room temperatur e photoluminescence measured for the as-grown samples A and B are represented in figure 6-16. Th e spectra of the two films are very different. NBE emission occurs at 377 nm, 3.3 eV, for both films. Emission in the visible, is more pronounced for Sample A, while for sample B, vi sible emission intensity remains at the noise level. Since Target 2 was sintered at a lowe r temperature and in an inert atmosphere, I anticipated it to be less dense than Target 1, richer in As and Zn. Non-emission beyond 400 nm for sample B is associated with the formati on of non-radiative defect s centers during growth. Emission in the visible is usually associated to points defects su ch as Zn inters titials and oxygen vacancies as reported in Chapter 3.

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162 Hall measurements of a series of films, gr own at 500C on MgO buffer layer from a target doped with 2.0 at % As from zinc arsenide and sintered at 600 C for 8 hrs, indicate increasing resistivity and carrier mobility whereas the carrier concentrati on of the films decreases with increasing growth oxygen pressure. The films grown at 5 and 50 mTorr of oxygen are p type while the films grown at 150 mTorr of oxygen is n-type. The da ta is shown in table 6-11. Va riable field Hall measurements show increased scattering in the Hall voltage data with increasing oxygen pressure suggesting increasing compensation in the materials and carrier type conversion as seen in figures 6-28 to 30. Room temperature photoluminescence of the f ilms, shown in figure 6-33 shows a decrease in NBE emission with increasing growth oxygen. This is the opposite of was seen in the room temperature photoluminescence of the films deposite d from the target doped with 2 at% As but sintered at 1000C for 10 hrs, shown in figure 6-32. The effects of target processing can be seen by comparing the transport properties of the films deposited at similar growth conditions but fr om the two targets as seen in figure 6-34 and 6.35. The films deposited from the target sinter ed in flowing Ar at 600C at 8 hrs are less resistive and resistivity is more affected by the increase in pressu re than the films doped from the target sintered in air for 10 hrs. Carrier con centration of the film gr own at 500C and 5 mTorr from the target sintered in Ar at 600 C for 8 hrs is unusually high, 3.4 x 1019 cm-3. Repeated measurements at different excitation currents sh ow similar results. Carrier densities of films grown at 500C from the same target at 50 and 150 mTorr of oxygen growth pressure are lower, on the order of 1015 and 1016 cm-3 respectively. Carrier concentr ation of the films doped from the target sintered at 1000C fo r 10 hrs also decreases with incr easing growth pressure from 1017

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163 to 1014 cm-3. Mobility of the films doped from both targ ets is low and is not as affected by the growth pressure as other parameters. When a series of films grown at 500C and 5 mTorr of oxygen from the target doped with 2.0 at % As from zinc arsenide and sintered at 600C for 8 hrs, are rapid thermal annealed at 800C in oxygen, resistivity increases rapidly wi th RTA duration. Carrier concentration of the films decays rapidly also upon rapid thermal anneal while carrier mobility increases modestly, as it can be seen in figures 6-36 and 6-37. The film s remained p-type upon annealing. Resistivity of the films increases from 2.75 .cm for the as-grown film to 507.2 .cm for the films RTA in oxygen for 90 seconds. Carrier concen tration decreased from 3.40 x 1019 to 5.32 x 1015 cm-3. Hall mobility increased from 0.07 to 2.31 cm2/(V.s). Transport properties of the as-grown films ar e reported in table 6-5. Both samples were p-type when measured. Sample A is highly resistive, over 26 k .cm, with carrier concentration of 1.6 x 1014 /cm-3 and carrier mobility of 1.52 cm2/(V s). The resistivity of the as-grown Sample B is an order of magnitude lower than for Sample A, 1.65 K /cm and displays a carrier concentration of 4.1 x 1015 cm-3 and carrier mobility of 0.9 cm2/(V s). Upon rapid thermally annealing of both films at 200C for 60s, the electrical properties of both films changed dramatically. The resistivity of both films dropped to 1.5 and 0.74 k /cm, Carrier concentrations increased to 2.1 x 1016 and 1.8 x 1016 cm-3 and Hall mobilities decreased to 0.2 and 0.5 cm2/(V s) for Sample A and B, respectively. Rapid thermal annealing in oxygen of the samples at 800C for 5 seconds also altered their electrical prope rties. Upon the RTA, sample A becomes semiinsulating. Its resistivity increased from 25.6 k .cm before anneal to 84 k .cm when annealed, making it challenging to measure its transport prop erties. Resistivity for Sample B also increases when subjected to the RTA in oxygen at this temperature and time, from 1.6 k .cm for the as-

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164 grown film to 2.1 k .cm for the annealed film. There was a carrier type convers ion from pto nwith an electron density of 5.3 x 1016 cm-3 and mobility of 0.6 cm2/(V.s). The effects of annealing temperature of a 5 second oxygen RTA for several sample (A)s are shown in table 6-6. When subjected to the RTA in oxygen at 400C, carri er type converted to n and resistivity decreased to 8.8 k .cm. When subjected to RT A temperature higher than 400C, resistivity increases with RTA temperatur e as it can be seen in figure 6-17. Since the resistivity of the films approached the limits of our Hall measurements system, I was unable to accurately measure the transport pr operties, therefore, only the ch anges resistivity are shown in the figure. When a film grown at 500C and 50 mTorr of ox ygen from the same target is subjected to RTA at 800C in oxygen for 5 sec, carrier type convers ion from nto poccu rs and resistivity of the film almost doubled while carrier concentrati on slightly increases and mobility changed little. Effects of Arsenic Source Table 6-14 shows the effects of arsenic sour ce and target sintering conditions on the electrical properties of the f ilms. The films were grown at 500C and 5 mTorr of oxygen on an MgO buffer layer. Target 1 was doped with 2% As from Zn3As2 and sintered in air at 1000C for 10 hrs, Target 2 was doped with 2% As from Zn3As2 600C for 8 hrs and Target 3 was doped with 2% As from Zn3As and sintered at 600C for 8 hrs. The film grown from Target 2 has the lowest resistivity, 2.745 .cm, the highest carrier concentration, 5.7 x 1019cm-3 and lowest carrier mobility 0.040 cm2/(V s). The film grown from Target 3 wa s semi-insulating with resistivity, on the order of 49 x 103 .cm. Effects of Aging Nine months after growing the films described in chapter 5, the transport properties of the films grown at 5 (Sample C) and 150 mTorr of oxygen (Sample D) were re-measured. The

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165 results are tabulated in table 6-16. For the film grown at 150 mTorr of oxygen, resistivity decreased from 7.6 .cm to 1.6 .cm; carrier mobility increased 9 x 1014 to 2 x 1016 cm-3 and Hall mobility decreased from 0.9 to 0.18 cm2/(V.s). Photoluminescence spectra of the as-grown and aged film reveal the spectra are identical The opposite happens when the film grown at 5 mTorr of oxygen is aged for the same period of time. Resistivity of the aged film is 180 times higher than the as-grown film from 53 .cm to 9.7 k .cm; carrier mobility is 10 times higher from 0.78 to 7.4 cm2/(V.s) and carrier concentration decrea sed 3 orders of magnitude from 1.5 x 1017 to 8.7 x 1013 cm-3. Similar changes in the transport properties occurs when I re-measured one month old film which was-grown at 5 mT orr of oxygen, although the target used was processed differently. Resistivity and carrier mob ility of the aged film were higher than for the fresh sample while carrier concentration decrease d almost 5 fold from its original value. Nineteen days after Sample D, a 0.2 at % As doped film grown at 500C and 5 mTorr of oxygen, was grown and first measured, I re-measured its Hall characteristics. Its resistivity dropped considerably while both carrier concentration and mobility increased. The plots of the Hall voltages as a function of the applied magnetic fiel d for samples C, D and F are shown in figures 6-38 (a), 6.39 and 6.40. The graphs indicate incr eased scattering in the data for sample C (150 mTorr) with time but the opposite for samples D and F (5 mTorr, 2 and 0.2 at % As). Effects of Photoresist Coat ing on Transport Properties I further investigate the eff ects of surface conduction on the bulk electrical properties by looking at the effects of photoresis t on the transport properties of th e film. A series of films were coated with polymer-based photoresist (PR) a nd soft baked at 120C for 10 min. Before the coating, indium contacts were at tached to films surface to ensure ohmic contact. Once the sample reached room temperature after cu ring the PR, I performed I-V test to ensure the contacts were ohmic and then measured the electrical characte ristics of the film by Hall measurements. Table

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166 6-17 sums up the as-grown, right after PR cu re and 48 hrs after curing the PR electrical properties of films. Sample G is 0.2 at % As doped film grown at 500C and150 mTorr of oxygen. Sample H is a 2 at % of As doped film grow n at 500C and 150 mTorr of oxygen. Sample I is a film dope 2 at % As and gr own at 500C in 5 mTorr of oxygen. J is a film doped with 2% As grown at 500C, 5 mTorr and RTA in oxygen at 800C for 2 s before being cover with PR. Sample K was an undoped ZnO film grown at 500C, 20 mTorr on MgO buffer for reference. The electrical properties of all the As doped f ilms were greatly affect ed by the PR coating. Resistivity changed regardless of As concentratio n, or partial pressure while mobility changed little for those films. For the undope d film, sample K, resistivity sl ightly increased when the PR was applied. Mobility for the undoped film also increased with the PR application. The variations of Hall coefficient for the film, as-g rown, 30 and 60 min after the PR bake and 48 hrs later, are shown in figures 6-41 to 6-54. The measurements were very noisy for the As doped films regardless of As content and growth pressure. Hall coefficient for the undoped film remained linear with the application of PR. Summary I presented in this chapter several factors aff ecting both the transport and optical properties of arsenic doped zinc oxide films. Since arsenic is very volatile, the as both the grown and post anneal properties of the films depends on ta rget composition, target sintering conditions, substrate temperature and backgr ound/oxidizer gas and pressure. A nnealing the films in N or Ar turn the films n-type, conductiv e and carrier mobility genera lly improves with increasing annealing time or temperature, independent of annealing methods. Often the resistivity of the films annealed in an inert environment show s similar dependence on growth pressure or temperature. When the films are annealed in oxygen, resistivity increa ses several orders of

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167 magnitude from the as-grown value while the carri er mobility does not always improve. RTA in oxygen as short as 5 second turns so me of the films semi-insulati ng. The changes in resistivity were even greater for the films deposited at lowe r temperature, growth pressure and from targets with higher As content. X-ray diffraction and low temp erature photoluminescence of th e films at RTA in Nitrogen from 200 to 500C suggest modest to no change in the bulk properties of the films annealed while Hall measurements showed that the transp ort properties of the f ilms were significantly altered. The films turned n-type for a nitrogen RTA turned the films n-type. Those changes point out to the formation or activation of an n-type conductive layer when the films are annealed in oxygen deficient environments. The rapid changes in resistivity and carri er concentration point to oxygen acting as electronic traps at the films su rface or grain boundaries surfaces167. Since oxygen has a high affinity for electron, when the films are annealed in oxygen, oxygen is quickly adsorped at the surface to form O-, O2 and O3 which acts as electronic traps. It is possible that oxygen adsorption process is amplified by higher concentration of As at the surface and crystallite interfaces. Upon heating As up in oxygen environment, As will further oxidizer from its As3+ to As5+. In inert environments, such as N or Ar, oxyge n is desorpted from the surface releasing the electron in the conduction band. There are othe r that may cause the changes include: hydrogen diffusion to the surface or hydrogen adsortion in oxygen deficient environment, formation of (Zn, O, C, H) complexes at the surface or grain boundaries20 due to atmospheric contamination, hydroxylation167, 168 of the films surface and changes in su rface composition since As sublimes at temperature as low as 300C.169

PAGE 168

168 Analysis of the variations of the Hall volta ge with applied magnetic field reveals increase scatter in the data with increasing As concentr ation, growth pressure and decreasing target density. The noise in the data shows how difficult is it to accu rately evaluate the transport properties. Since the Hall coefficient is taken to be the slope of linear fit to the plot of measured field Hall voltage x applied magnetic field, meas urement uncertainty is greatly affected by the deposition parameters. The scatter in the data of Hall voltage is also affected by compensation and structural defects since multiple bands conduction depends on the concentration of electron, hole and their mobility. Hall mobility is inversely proportional to resistivity, our suspicions of surface conduction dominating the bulk electrical properties of the doped films is fuel by the significant changes in resistivity of the film while hall mobility remain more subtle. I have also seen the electrical properties of th at film changing with time as short as 19 days after growth. It seems that films grown at low growth pressure their conductivity increases with time while the opposite is happens in the film s grown at high pressure. Low temperature photoluminescence of a sample, measured soon after the sample was grown and after aging for 9 months, reveals that the spectra were identical I believe upon abortion of UV light from day light oxygen the surface or grain boundaries20-22, 161 is photodesorbted releas ing an electron to the conduction band when the films are grown at high er deposition pressure where as for films grown at lower pressure the films adsorpt oxygen over time. A. B. Anderson168 hydrogen adsorption on O2 terminated surface releasing an electron to the conduction band by e OH O H2 On Zn terminated surfaces, H reacts with Zn2+ to form (ZnH)+ H Zn ZnH Zn H2 .) (

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169 I have seen in chapter 3 when the Fermi le vel approaches the conduction band, for p-type sample, formation energy of the H related defects decreases spontaneously43. This would also explain the larger resistivity increase for the films grown lower deposition pressure. Surface effects on the transport properties of the doped films became more apparent when I coated some of the films with photoresist. Va riable Hall measurements of the As-doped and coated films showed great increased in the scatter in the data with the application of the PR. As for the undoped films, resistivity remained almost unaffected while mobility slightly increased. Hence, I strongly believe that su rface conduction plays a major ro le in the measurement of As doped sample and may be preventing us from meas uring the true bulk tran sport properties of the films as O. S. Schmidt170 speculated.

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170 Table 6-1. Transport properties of the films doped with 2 at % as from zi nc arsenide and grown at 500C and 50 mTorr of oxygen and RTA for 60 seconds in Nitrogen. RTA Temperature (C) Hall Coefficient (C/cm-3) Resistivity ( .cm) Carrier Concentration (cm-3) Hall Mobility (cm2/(V.s)) Type 0 Indeterminate 510 200 51 201 1.22 x 1017 0.25 p 300 -15.4 66 4.05 x 1017 0.23 n 400 -7.01 1.61 8.91 x 1017 4.37 n 500 -6.95 1.62 8.98 x 1017 4.28 n Table 6-2. Resistivity of the films doped with 2 at % as from zinc arse nide and grown at 500C and 50 mTorr of oxygen and RTA at 200C in Nitrogen. RTA time (s) Resistivity ( .cm) 0 8.29E+3 30 3.74E+5 60 1.04E+5 120 1.48E+4 180 6.57E+5 240 8.23E+5 300 7.69E+4 Table 6-3. Before and after nitrogen rapid ther mal anneal at 200C for 60s for several films grown at different conditions. Targets 1 and 2 were both doped with 2.0 atomic percent As from Zn3As2, however, Target 1 was sinter ed in air at 1000C for 10 hrs and Target 2 sintered in Ar at 600C for 8 hrs. Target 3 was doped with 0.2 atomic percent As from Zn3As2 sintered in air at 1000C for 10 hrs. 60s N RTA at 200C Resistivity ( .cm) Hall Coefficient (cm3/C) Carrier Density cm3 Hall Mobility (cm2/(V-s)) A: 500C, 1 x 10-6 Torr Po on MgO buffer layer, Target 1 as-grown 0.113 -1.81 (n) 3.43E+18 16.10 After 0.115 -1.72 (n) 3.63E+18 14.97 B: 500C, 1 mTorr Po on MgO buffer layer, Target 1 as-grown .169 -1.82 (n) 3.42E+18 10.81 After .083 -7.04 (n) 8.9E+17 8.41 C: 500C, 50 mTorr Po on MgO buffer layer, Target 1 as-grown 25659 39037 (p) 1.6E+14 1.52 After 1543 296.26 (p) 2.1E+16 0.19 D: 500C, 50 mTorr Po on MgO buffer layer, Target 2 as-grown 1658.00 1513.80 (p) 4.12E+15 0.91 After 738.90 353.63 (p) 1.76E+16 0.48 E: 500C, 5 mTorr Po on MgO buffer layer, Target 3 as-grown 907.80 839.03 (p) 7.44E+15 0.92 After 1.09 0.99 (p) 6.32E+18 0.91

PAGE 171

171 Table 6-4. Hall data for the as-grown and RTA in nitrogen at 800C Films. The films were doped with 2 at % As doped and grown at 5 00C and 150 mTorr of oxygen on ZnO buffer layer. RTA Length s Resistivity .cm Carrier Density cm-3 Mobility cm2/(V.s) Type 0 7977 7.8 x 1014 0.9 p 5 57 1.7 x 1017 0.6 n 30 5.14 3.0 x 1017 4.2 n 60 0.21 4.2 x 1018 7.5 n 120 0.12 2.5 x 1018 19.5 n 180 0.12 2.0 x 1018 24 n Table 6-5. Transport properties of films grown from Targets 1 a nd 2. Both films were grown at 500C and 50 mTorr of oxygen for 2 hrs on thin MgO buffer layer. RTA Conditions Resistivity .cm Hall Coefficient C/cm3 Carrier Density cm-3 Mobility cm2/(V s) Type as-grown 25659 39037 1.6 x 1014 1.5 p 200C, N, 60s 1543 296.26 2.1 x 1016 0.2 p Target 1 Sample A 800C, O2, 5s 84028 as-grown 1658.0 1513.8 4.1 x 1015 0.9 p 200C, N, 60s 738.9 353.6 1.8 x 1016 0.5 p Target 2 Sample B 800C, O2, 5s 2108.0 -1174.6 5.3 x 1015 0.6 n Table 6-6. Hall measurements for the samples A were RTA for 5 s in oxygen. Negative sign indicates n-type conductivity. RTA Temp C Resistivity .cm Hall Coeff. C/cm3 Car. Conc. x 1014 cm-3 Mobility cm2/(V s) Type as-grown 25659 39037 1.6 1.52 p 400 8817 -13629 4.6 1.54 n 600 19772 -58464 1.1 2.97 n 700 19010 800 84028

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172 Table 6-7. Hall data for the film doped with 0.02 at % As on bare sapphire and 3 mTorr of ozone/oxygen. The films were annealed in flowing oxygen at 800C for 10 min. Conditions Growth Temp C Resistivity .cm Carrier Density cm-3 Mobility cm2/(V s) Type 300 0.573 2.56 x 1018 4.25 n 400 0.0197 3.51 x 1019 9.04 n 500 16.076 8.05 x 1018 0.05 N As Grown 600 0.193 3.96 x 1018 8.16 n 300 100 x 103 400 157 x 103 500 9.98 x 103 4.56x 1018 1.37 n 800C 10 min O2 Anneal 600 5.84 1.01 x 1017 9.67 n Table 6-8. Hall data for the film doped with 0.02 at % As on bare sapphire and 30 mTorr of oxygen. The films were annealed in flowing oxygen at 800C for 10 min. Conditions Growth Temp C Resistivity .cm Carrier Density cm-3 Mobility cm2/(V s) Type 300 0.049 9.27 x 1018 13.73 n As Grown 600 0.041 9.00 x 1018 16.91 n 300 insulating 800C, 10 min O2 Anneal 600 23.17 1.28 x 1018 21.02 n Table 6-9. Hall data for the film doped with 0.2 at % As on bare sapphire and 3 mTorr of ozone/oxygen. The films Ire annealed in flowing oxygen at 800C for 10 min. Conditions Growth Temp C Resistivity .cm Carrier Density cm-3 Mobility cm2/(V s) Type 300 0.027 2.29 x 1019 10.05 n 500 0.022 2.19 x1019 12.8 n As Grown 600 0.018 1.43 x 1019 24.6 n 300 103.64 1.33 x 1017 0.45 n 500 235 x 103 800C, 10 min O2 Anneal 600 62.24 3.12 x 1016 3.21 n

PAGE 173

173 Table 6-10. Hall data for the film doped with 0.2 at % As on bare sapphire and 30 mTorr of oxygen. The films Ire annealed in flowing oxygen at 800C for 10 min. C onditions Growth Temp C R esistivity .c m C arrie r D ensity c m-3 M obility cm 2/(V s) T ype 400 0.009 1.14 x 1020 6.2 n 500 0.095 1.00 x 1019 6.2 n As Grown 600 0.031 6.10 x 1016 33 n 400 209.6 3.68 x 1014 0.083 n 500 197 6.96 x 1015 3.57 n 800C, 10 min O2 Anneal 600 128 4.30 x 1015 7.24 n Table 6-11. Transport properties of the films doped with 2 at % As and grown 500C and on MgO buffer layer. The Target was doped with Zn3As2 and sintered at 600C for 8 hrs. Growth Pressure mTorr Resistivity .cm Hall Coefficient C/cm3 Carrier Density cm-3 Mobility cm2/(V s) Type 5 2.745 0.11 5.67 x 1019 0.040 p 50 1658.00 1513.80 4.12 x 1015 0.91 p 150 176.74 -90.55 6.89 x 1016 0.51 n Table 6-12. 800C oxygen RTA of the film 500C and 5 mTorr of oxygen for 2 hrs on MgO buffer layer from the ZnO target doped with 2% As from Zn3As2 and sintered in Ar at 600C for 8 hrs. 800C O2 RTA Time s Resistivity .cm Hall Coefficient C/cm3 Carrier Density cm-3 Mobility cm2/(V.s) Type As grown 2.75 0.1835 3.40 x 1019 0.07 p 2 29.7 0.73 8.55 x 1018 0.02 p 30 7807 3564.6 1.75 x 1015 0.46 p 90 507.2 1173.6 5.32 x 1015 2.31 p Table 6-13. Effects of an 800C, O2, 5 seconds RTA on the transport properties of a film grown at 500C and 50 mTorr of O2 from the ZnO target doped with 2% As from Zn3As2 and sintered in Ar at 600C for 8 hrs. RTA Conditions Resistivity .cm Hall Coeff. C/cm3 Car. Conc. cm-3 Mobility Cm2/(V s) Type as-grown 1658.00 1513.80 4.12 x 1015 0.91 p 800C, O2, 5 s 2108.00 -1174.60 5.31 x 1015 0.56 n

PAGE 174

174 Table 6-14. Transport properties of the films gr own from different targ ets at 500C and on MgO buffer layer. Target 1 was doped with 2% As from Zn3As2 and sintered in air at 1000C for 10 hrs, Target 2 was doped with 2% As from Zn3As2 600C for 8 hrs and Target 3 was doped with 2% As from Zn3As2 and sintered at 600C for 8 hrs. Target Resistivity .cm Hall Coefficient C/cm3 Carrier Density cm-3 Mobility cm2/(V s) Type 1 52.99 41.35 1.51 x 1017 0.78 p 2 2.745 0.11 5.67 x 1019 0.040 p 3 49336 p Table 6-15. Transport properties of 10 min Ar ann ealed films at 600C. The type of the as-grown films is indicated in the parentheses. The Ha ll characteristics of th e films Ire analyzed in chapter 5. Growth Pressure mTorr Resistivity .cm Hall Coefficient C/cm3 Carrier Density cm-3 Mobility cm2/(V.s) Type 5 (p) 0.179 -2.26 2.76 x 1018 12.63 n 20 (n) 0.219 -4.15 1.5 x 1018 18.95 n 50 (p) 0.028 -0.75 8.32 x 1018 26.79 n 150 (p) 10.3 -5.83 1.07 x 1018 0.57 n Table 6-16. Transport properties of as-grown an d aged 2 at % As doped samples grown at 500C: C 150 mTorr of oxygen, D 5 mTorr from the target doped using Zn3As2 and sintered at 1000C in Ar for 10 hrs, E 5 mTorr from the target doped using Zn3As2 and sintered at 600C in Ar for 8 hrs, F 0.2 at % As doped and 5 mTorr of oxygen. Sample Age months Resistivity .cm Carrier Density cm-3 Mobility cm2/(V.s) Type As Grown 7626 9 x 1014 0.92 p C 9 1573 2.25 x 1016 0.18 p As Grown 53 1.5 x 1017 0.78 p D 9 9670 8.7 x 1013 7.38 p As Grown 7.13 7.7 x 1019 0.01 p E 1 29.91 2.1 x 1014 0.1 p As Grown 907.80 7.46 x 1013 0.92 p F .633 53.36 5.77 x 1016 2.03 p

PAGE 175

175 Table 6-17. Transport properties of as-g rown and PR covered films: Sample G is 0.2 at % As doped grown at 500C and150 mTorr of oxygen. Sample H was grown from the target doped with 2 at % of As at 500C and 150 mTorr of oxygen. I 2% As and grown at 500C in 5 mTorr. J 500C, 5 mTorr from and RTA in oxygen at 800C for 2 s before being cover with PR. Kwas an Undoped ZnO Film grown at 500C, 20 mTorr on MgO buffer. Mobility is too lo to be measured. Samples Age months Resistivity .cm Carrier Density cm-3 Mobility cm2/(V s) Type As Grown 53.363 5.70 x 1016 2.05 p After PR cured 22.847 3.74 x 1018 0.073* n G 48hrs PR 22.231 8.79 x 1017 0.32 n As Grown 898 4.76E+15 1.46 n After PR cured 1919 3.70E+16 0.088 n H 48hrs PR 1835 Indeterminate As Grown 9969 6.09E+14 1.03 p After PR 35749 7.41E+13 2.36 p I 48hrs PR 37957 Indeterminate As Grown 23.284 2.39E+18 0.11 p After PR cured 15.13 Indeterminate J 48hrs PR 15.899 Indeterminate As Grown 0.108 2.55E+18 22.68 n After PR cured 0.157 1.45E+18 27.48 n K 48hrs PR 0.142 1.56E+18 28.18 n

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176 161820 0 200 400 600 800 1000 Intensity (counts/s)Omega (Degree) a: FWHM= 2.495 b: FWHM= 2.322 Figure 6-1. High-resolution omega rocking curves fo r the films doped with 2 at % As : A) is the as-grown film (no RTA) with FWHM of 2.495 and B) is for the film annealed at 500C and has a FWHM of 2.322. 400450500550600650700 200C 300C 500C 400C 3.3 eV Figure 6-2. Photoluminescence spectra of the film s doped with 2 at % As from zinc arsenide were grown at 500C and 50 mTorr of oxygen partial pressure on a thin ZnO buffer layer and RTA in nitrogen at low temperature.

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177 3.203.253.303.353.403.45 A0Xa: 200C c: 400C D0X 3.328 eV3.203.253.303.353.403.45 3.328 eVA0X D0X 3.203.253.303.353.403.45 b: 300C 3.328 eV D0X A0X 3.203.253.303.353.403.45 3.328 eVA0Xd: 500C Energy (eV) Figure 6-3. PL at 30K of the films doped with 2 at % As from zinc arse nide and grown at 500C and 50 mTorr of oxygen partial pressure on a thin ZnO buffer and RTA in nitrogen at low temperature [a, b, c, d are for RT A at 200, 300, 400 and 500C, respectively]. 0100200300400500 0 100 200 300 400 500 Resistivity (Ohm.cm)RTA Temp (C) Figure 6-4. Resistivity of the films doped with 2 at % As and grown at 500C in 50 mTorr of O2 on a thin ZnO buffer layer.

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178 -1.0x104-5.0x1030.0 5.0x1031.0x104-6.0x107-4.0x107-2.0x1070.0 2.0x1074.0x1076.0x107 Applied Magnetic Field (G)VHa: as grown-1.0x104-5.0x1030.0 5.0x1031.0x104-3.0x106-2.5x106-2.0x106-1.5x106-1.0x106-5.0x1050.0 5.0x105 b: 200C-1.0x104-5.0x1030.0 5.0x1031.0x1043.4x1063.6x1063.8x1064.0x1064.2x1064.4x1064.6x1064.8x106 c: 300C-1.0x104-5.0x1030.0 5.0x1031.0x104-8.0x104-6.0x104-4.0x104-2.0x1040.0 2.0x1044.0x1046.0x1048.0x1041.0x105 d: 400C Figure 6-5. Field dependent Hall voltage with ann ealing temperature of the films doped with 2 at % As RTA in N for 60 sec [a: as grown; b, c, d are RTA at 200, 300 and 400C respectively].

PAGE 179

179 200300400500-1.0x1018-8.0x1017-6.0x1017-4.0x1017-2.0x10170.0 2.0x1017 Carrier Density (cm-3)RTA Temperature (oC) p-type n-type Figure 6-6. Effects of annealing temperature on carrier density of the films doped with 2 at % As RTA in N for 60 sec. 200300400500 -5 -4 -3 -2 -1 0 1 n-typeHall Mobility (cm2/(V.s))RTA Temperature (C) p-type Figure 6-7. Hall mobility as a function of the 60 seconds N2 RTA temperature of the films doped with 2 at % As.

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180 400450500550600650700 30s 300s 240s 60s 120s 180s 3.3 eV Wavelength (nm) Figure 6-8. Photoluminescence spectra of the films doped with 2 at % As and RTA in N2 at 200C RTA at 200C for 30-300s. 3.153.203.253.303.353.403.45 3.103.153.203.253.303.353.403.45 LODAP a: 60 sec 3.361 eV DoX 3.234 eV 3.169 eV DAP 3.328 eV AoX3.328 eV AoX 3.169 eV 3.234 eV LODAP DAP b: 120 sec Energy (eV)

PAGE 181

181 3.103.153.203.253.303.353.403.45 3.328 eV AoX c: 180 sec 3.169 3.234 eV DAP LODAP energy (eV) 3.103.153.203.253.303.353.403.45 3.103.153.203.253.303.353.403.45 3.328 eV AoX d: 240 sec 3.169 eV 3.234 eV DAP LODAP Energy (eV) 3.328 eV AoX 300 sec 3.169 eV 3.234 eV DAP LODAP Figure 6-9. PL at 30K of the films doped with 2 at % As and RTA in N2 at 200C [a: 60s; b: 120s; c: 180s; d: 240s; e: 300s].

PAGE 182

182 2.82.93.03.13.23.33.4 1 2 LODAP FA 250K 300K 200K 150K 120K 90K 50KIntensity (a. u.)Energy (eV) 30K D0X DAP AoX (3.3 eV) Figure 6-10. Temperature dependence of the photoluminescence from 30-350K for the film doped with 2 at % As and RTA in N2 annealed for 60s at 200C. 050100150200250300 104105106 Resistivity (ohm.cm)RTA Time (s) Figure 6-11. Plot of resistivity as a function of annealing time fo r the films doped with 2 at % As and RTA in N2 at 200C.

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183 -20020406080100120140160180200 10-1100101102103 1015101610171018 Carrier Density (cm-3)Resistivity (Ohm.cm)RTA time (s) Figure 6-12. Changes in resistivity and carrier concentration as a function of RTA time for the 2 at % As doped films grown at 500C and 150 mTorr of oxygen on ZnO buffer layer. 0408012016020 0 -25 -20 -15 -10 -5 0 5 Carrier Mobility (cm 2 /(V s))RTA Time (s) n-type p-type Figure 6-13. Changes in carrier mobility as a f unction of RTA time for the 2 at % As doped films grown at 500C and 150 mTorr of oxygen on ZnO buffer layer.

PAGE 184

184 4005006007000.01 0.1 Intensity (a. u.)Wavelength (nm) Figure 6-14. As grown room temperature photolumi nescence of the film 2 at % As doped film grown at 500C and 150 mTorr of oxygen on ZnO buffer layer. 350400450500550600650700 1 Intensity (a. u.)Wavelength (nm) 5s 30s 60&120s 180s Figure 6-15. Room temperature photoluminescence spectra of the annealed 2 at % As doped films grown at 500C and 150 mTorr of oxygen on ZnO buffer layer.

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185 350400450500550 IntensityWavelength (nm) Sample A Sample B x Figure 6-16. 300K photoluminescence spectra for as-grown samples A and B. 0200400600800104105 Resistivity (Ohm.cm)RTA Tem p erature ( oC ) Figure 6-17. Resistivity as a function of O2 RTA temperature for sample A.

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186 400450500550600650700 1E-3 0.01 0.1 1 10 100 1000 10000 100000 Resistivity (Ohm.cm) As grown 800C, 10 min Anneal Growth Temperature (C) Figure 6-18. Changes in resistivity as a function of growth temp erature for the film doped with 0.02 at % As on bare sapphire and in 3 mTorr of O3/O2. The films were annealed in flowing oxygen at 800C for 10 min. 400500600700 400oC 500oC 600oC 700oCIntensity (a. u.)Wavelen g th ( nm ) Figure 6-19. Room temperature photoluminescence of the as-grown films doped with 0.02 at % ZnO:As films on sapphire and 3 mTorr of O3/O2.

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187 4005006007001E-3 0.01 0.1 1 700C 400CIntensityWavelength (nm) 400C 500C 700C 600C 500 & 600C Figure 6-20. Room temperature photoluminescence of the 0.02 at % ZnO:As films grown on sapphire and 3 mTorr of O3/O2 and annealed in flowing oxygen at 800C for 10 min. 40050060070010-210-1100 Wavelen g th ( nm ) b: 800oC, 10 min oxygen anneal 10-210-1100 700 C 700C400CIntensity (a. u.)400Ca: As grown Figure 6-21. Room temperature photoluminescence sp ectra of the 0.02 at % ZnO:As films grown on sapphire and 30 mTorr of O2 (a) and annealed in flowing oxygen at 800C for 10 min (b).

PAGE 188

188 3.003.053.103.153.203.253.303.353.403.45 AoX3.324 eV3.361 eVDoX LODAP DAP3.361 eVDoX Energy (eV) a: As-grown 3.259 eV 3.190 eV3.003.053.103.153.203.253.303.353.403.45 AoX3.324 eVb: Annealed Figure 6-22. Photoluminescence sp ectra at 20 K of the as-grown (a) and annealed (b) films doped with 0.02 at % ZnO:As grown at 700C on sapphire and 30 mTorr of O2.

PAGE 189

189 300350400450500550600 2x10-22.5x10-2102103104105 5x101610171.5x10171.2x10191.4x10191.6x10191.8x10192x10192.2x10192.4x1019 0 4 8 12 16 20 24 Growth Temperature (oC)Resis. Mob. Annealed As gorwn Car. Con. Annealed Car. Con. Mob. As gorwn Resis. Figure 6-23. Transport properties for the as-grown and the 10 min 800C O2 annealed films doped with 0.2 at % As grown in 3 mTorr of O3/O2. 3504004505005506006507000.01 0.1 1 Intensity (a. u.)Wavelength (nm) a: As deposited 700C 400C 600C0.01 0.1 1 400 & 600C 700Cb: 800oC, 10 min oxygen anneal 600C Figure 6-24. PL at 300K spectra of the 0.2 at % ZnO:As films grown in 3 mTorr of O3/O2 (a) and annealed in flowing O2 at 800C for 10 min (b).

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190 3.003.053.103.153.203.253.303.353.403.45 as grown, 600oC Energy (eV) D0X 3.36 eV 3.33 AoX 3.31 AoX 800oC, 10 min O2 Anneal DAP or FA Figure 6-25. PL at 16K spectrum of the 0.2 at % ZnO:As films grown on sapphire and at 600C and 3 mTorr of O3/O2 and annealed in flowing O2 at 800C for 10 min compared to the as grown (dashed lines). 4005006007000.01 0.1 4005006007000.01 0.1 600CWavelength (nm)b: O2 annealed at 800C for 10 min. 600C 500C 400C 400C 500C 600CIntensitya: As grown 390 nm Figure 6-26.PL at 300K spectrum of the 0.2 at % ZnO:As films grown on sapphire in 30 mTorr of O2 (a) and annealed (b) in flowing O2 at 800C for 10 min.

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191 400450500550600 0.01 0.1 1017101810191020 400450500550600 4 6 8 10 12 14 16 18 20 22 24 26 28 30 32 34 36 Carrier densityMobilityGrowth temperature (C) resistivity Carrier density MobilityResistivity (Ohm.cm) a: as grown400450500550600150 200 1015 0 2 4 6 8 MobilityGrowth temperature (C) Resistivity Carrier densityCarrier densityResistivity Mobility b: Annealed Figure 6-27. Transport properties of: (a) as-grown and (b) ann ealed in flowing oxygen at 800C for 10 min as a function of growth temper ature for films doped with 0.2 at % As grown on sapphire and 30 mTorr of O2. The filled circles, squares, triangles and circles represent resistivity, carrier density and mobility, respectively.

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192 020406080100120140160 10-310-210-1100101102103 p pResistivity (Ohm.cm)Growth Pressure (mTorr of O2) As grown 600C, 10 min, Ar anneal all n p Figure 6-28. Resistivity as a func tion of growth Pressure for the films doped with 2 at % As and grown on ZnO buffer layer. The Hall charact eristics of the films were analyzed in chapter 5. -1.0x104-5.0x1030.0 5.0x1031.0x104-2x104-2x104-2x104-2x104 VHApplied Magnetic Field (G) 2.0 at % As, 5 mTorr RH = 0.11 Figure 6-29. Hall coefficient as a function of appl ied magnetic field for the films doped with 2 at % As and grown at 500C and 5 mTorr of oxygen on MgO buffer layer.

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193 -1.0x104-5.0x1030.0 5.0x1031.0x104-4x107-3x107-2x107-1x1070 1x1072x1073x107 VHApplied Magnetic Field (G) 50 mTorr O2, 2 % As RH = 1514.37 Figure 6-30. Variable Hall voltage measurement of the films doped with 2.0 at % of As and grown at 500C and 50 mTorr of oxygen on MgO buffer layer. The target was doped using zinc arsenide and sintered at 600C in Ar for 8 hrs. -1x1040 1x104-4.0x106-2.0x1060.0 2.0x106 Applied Magnetic Field (G)VH 150 mTorr O2, 2 % As RH= -69.88 Figure 6-31. Hall coefficient as a function of a pplied magnetic field for the films doped with 2.0 at % of As and grown at 500C and 150 mTorr of oxygen on MgO buffer layer. The target was doped using zinc arsenide and sintered at 600C in Ar for 8 hrs.

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194 360400440480520560600640680 Wavelength (nm) 50 mTorr 5 mTorr 150 mTorr 3.3 eV Figure 6-32. PL at 300K spectra of films doped w ith 2 at % As and grown on ZnO buffer layer at 500C from the Target sintered in 1000C for 10 hrs. 350400450500550600650700 150 mTorr 50 mTorrWavelength (nm)5 mTorr 3.3 eV Figure 6-33. PL at 300K spectra of films deposite d at 500C from the ZnO target doped with 2.0 at % of As and sintered at 600C in Ar for 8 hrs.

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195 2040608010012014016010 100 1k 1016101710181019 0.0 0.2 0.4 0.6 0.8 1.0 MobilityC.C. ResistivityResistivityGrowth Pressure ( mTorr ) Carrier Density Mobility Figure 6-34. Transport properties of the films gr own at 500C on MgO buffer layer from a target doped with 2 at % As and sintered in flowing Ar at 600C at 8 hrs. 020406080100120140160100101102103 101510161017 0.70 0.75 0.80 0.85 0.90 ResistivityResistivityOxygen Pressure (mTorr) Carrier Density Mobility Carrrier Density Mobility Figure 6-35. Transport properties of the films gr own from a target doped with 2 at % As and sintered in air at 1000C for 10 hrs.

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196 0306090 101102103104 Resistivity (Ohm.cm)RTA time(s) Figure 6-36. Resistivity of films O2 RTA at 800C. The films was grown at 500C and 5 mTorr of oxygen on MgO buffer layer from the Zn O target doped with 2% As from Zn3As2 sintered in Ar at 600C for 8 hrs. 02040608010010151016101710181019 0.0 0.5 1.0 1.5 2.0 2.5 Mobility (cm 2 /(V s))Carrier Concentration (cm-3) RTA Time (s) Figure 6-37. Carrier concentration and mobility of the films rta at 800C in oxygen. The films 500C and 5 mTorr of oxygen for 2 hrs on MgO buffer layer from the ZnO target doped with 2% As from Zn3As2 sintered in Ar at 600C for 8 hrs.

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197 -1.0x104-5.0x1030.0 5.0x1031.0x104-1x1070 3x1084x1085x108 VHApplied Magnetic Field (G) As Grown RH = 7045.1 Aged 6 months, RH = 60.04 2.82.93.03.13.23.33.43.5 0.1 I n t ens it y ( a. u. ) Energy (eV) As grown done at 25 K Aged done at 30K 3.360 eV DoX 3.314 eV AoX 3.258eV DAP LO DAP Figure 6-38. Effects of aging on: (a) Variable Hall measurements and (b) low temperature luminescence of sample C in Table 6-16. The film was doped with 2.0 at % of As and grown at 500C and 150 mTorr of oxygen on ZnO buffer layer.

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198 -1.5x104-1.0x104-5.0x1030.0 5.0x1031.0x1041.5x104-7x1010-7x10101.0x1061.5x1062.0x106 VH as grown RH = 41.35 After 9 moths RH = 71403.51Applied Magnetic Field (G) Figure 6-39. Variable Hall measurements of an as -grown and aged (D in Table 6-16) film doped with 2.0 at % of As and grown at 500 C and 5 mTorr of oxygen on MgO buffer layer. -1.0x104-5.0x1030.0 5.0x1031.0x104-3.5x107-3.0x107-2.5x107-2.0x107-1.5x107-1.0x107 As grown RH = 836.31VHApplied Magnetic Field (G) 7.05x1077.20x1077.35x10 7 19 days old RH = 108.14 Figure 6-40. Variable Hall measurements of an as -grown and aged (F in Table 6-16) film doped with 0.2 at % of As and grown at 500 C and 5 mTorr of oxygen on MgO buffer layer.

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199 -10000-500005000100007.0x1077.1x1077.2x1077.2x1077.3x1077.3x1077.3x107 Applied Magnetic Field (G)VH As grown Rh = 108.14394 Figure 6-41. Variable Hall measurements of the as-grown Sample G. The Hall data are shown in Table 6-17. The film was-grown from the ta rget doped with 0.2 at % of As at 500C and 150 mTorr of oxygen on MgO buffer layer. 5 000-10000-500005000100001500 0 5.0x1041.0x1051.5x1052.0x1052.5x1053.0x105 VHApplied Magnetic Field (G) after curing PR RH = -7.09948 Figure 6-42. Variable Hall measurements of Samp le G right after curing the PR. The Hall data are shown in Table 6-17. The film was grow n from the target doped with 0.2 at % of As at 500C and 150 mTorr of oxygen on MgO buffer layer.

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200 -10000-500005000100000.0 5.0x1041.0x105 V HApplied Magnetic Field (G) 0.2% As, PR, 48 hrs Rh = -1.66906 Figure 6-43. Hall Variable Hall measurements for Sample G 48 hrs after curing the PR. The Hall data are shown in Table 6-17. The film was grown from the target doped with 0.2 at % of As at 500C and 150 mTorr of oxygen on MgO buffer layer. -10000-50000500010000-8.0x107-7.5x107-7.0x107-6.5x107-6.0x107-5.5x107-5.0x107-4.5x107-4.0x107 VH 2s, 150 mTorr, as grown Rh = -1310.12 App lied Ma g netic Field ( G ) Figure 6-44. Variable Hall measurements for Samp le H right after curing the PR. The Hall data are shown in Table 6-17. The film was grown from the target doped with 2 at % of As at 500C and 50 mTorr of oxygen on MgO buffer layer.

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201 -1.0x104-5.0x1030.0 5.0x1031.0x104-2x107-1x1070 1x1072x1073x1074x1075x107 VH 2% as, 150 mTorr, affter PR cure RH = -1717.89Magnetic Field (G) Figure 6-45. Variable Hall measurements for Sample H 24 hrs after curing the PR. The Hall data are shown in Table 6-17. The film was grown from the target doped with 2 at % of As at 500C and 150 mTorr of oxygen on MgO buffer layer. -15000-10000-5000050001000015000-1.0x107-5.0x1060.0 VH 2% As, 150 mTorr, PR, 48 hrs Rh = -168.66 Magnetic Field (G) Figure 6-46. Variable Hall measurements of Samp le H 48 hrs after curing the PR. The Hall data are shown in Table 6-17. The film was grown from the target doped with 2 at % of As at 500C and 150 mTorr of oxygen on MgO buffer layer.

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202 1.5x104-1.0x104-5.0x1030.0 5.0x1031.0x1041.5x104-2.0x1090.0 2.0x1094.0x1096.0x10 9 VH 2% As, 5 mTorr, PR aft. Cure Rh = 83727.27 Magnetic Field (G) Figure 6-47. Variable Hall measurements of as Sa mple I after PR cure. The Hall data are shown in Table 6-17. The film was grown from the target doped with 2 at % of As at 500C and 5 mTorr of oxygen on MgO buffer layer. 15000-10000-500005000100001500 0 -2.0x1090.0 2.0x10 9 V H 2% As, 5 mTorr, PR, 48 hrs Ma g netic Field ( G ) Figure 6-48. Variable Hall measurements of Sample I 48 hrs after PR cured. The Hall data are shown in Table 6-17. The film was grown from the target doped with 2 at % of As at 500C and 5 mTorr of oxygen on MgO buffer layer.

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203 -15000-10000-5000050001000015000 -1x105-1x105-1x105-9x104-8x104-7x104-6x104-5x104-4x104-3x104 VH 2% As, 5 mTorr after 800C O RTA Rh = 2.62022Magnetic Field (G) Figure 6-49. Variable Hall measurements of as-g rown Sample J. The Hall data are shown in Table 6-17. The film was grown from the targ et doped with 2 at % of As at 500C and 5 mTorr of oxygen on MgO buffer layer. Th e Sample was RTA in oxygen at 800C for 2 s before being cover with PR. -15000-10000-5000050001000015000 -5.0x104-4.5x104-4.0x104-3.5x104-3.0x104-2.5x104-2.0x104 VH after PR cured Applied Magnetic Field (G) Figure 6-50. Variable Hall measurements of Samp le J right after curing the PR. The film was grown from the target dope d with 2 at % of As at 500C and 5 mTorr of oxygen on MgO buffer layer. The Sample was RTA in oxygen at 800C for 2 s before being cover with PR.

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204 -10000-50000500010000 -6.0x104-5.5x104-5.0x104-4.5x104-4.0x104-3.5x104-3.0x104-2.5x104-2.0x104-1.5x104-1.0x104 VH 48 hrs after PR Magnetic Field (G) Figure 6-51. Variable Hall measurements of Samp le J 48 hrs after curing the PR. The film was grown from the target dope d with 2 at % of As at 500C and 50 mTorr of oxygen on MgO buffer layer. The Sample was RTA in oxygen at 800C for 2 s before being cover with PR. -15000-10000-5000050001000015000-3.0x104-2.5x104-2.0x104-1.5x104-1.0x104-5.0x1030.0 5.0x1031.0x1041.5x1042.0x1042.5x1043.0x104 VH Undoped ZnO as Grown RH = -2.44995 Magnetic Field (G) Figure 6-52. Variable Hall measurements of as grown Sample K. The Hall data are shown in table 6.17. The film was undoped and grown 500C and 20 mTorr of oxygen on MgO buffer layer.

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205 15000-10000-500005000100001500 0 -6.0x104-4.0x104-2.0x1040.0 2.0x1044.0x1046.0x104 V H Sample E After PR cur e Rh = -4.31439 Magnetic Field (G) Figure 6-53. Variable Hall measurements of Samp le K right after curing the PR. The Hall data are shown in table 6.17. The film wa s undoped and grown 500C and 20 mTorr of oxygen on MgO buffer layer. -1.0x104-5.0x1030.0 5.0x1031.0x104-6.0x104-4.0x104-2.0x1040.0 2.0x1044.0x1046.0x104 VH Sample E, PR, 48 hrs RH = -4.00158 Magnetic Field (G) Figure 6-54. Variable Hall measurements of Samp le K 48 hrs after curing the PR. The Hall data are shown in table 6.17. The film wa s undoped and grown 500C and 20 mTorr of oxygen on MgO buffer layer.

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206 CHAPTER 7 OPTICAL CHARACTERIZATI ON OF OXIDIZED ZNNXO(1-X) FILMS Results and Discussion In this study, the properties of N-doped Zn O films grown from zinc nitride doped ZnO ablation targets were examined. Several ZnNxO(1-x) targets were fabricated from 99.9995 % pure ZnO and 99 % pure Zn3N2 powders. After mixing the powder w ith the desired concentration to homogenous solid solutions, the targ ets were isostatically pressed. The targets with 0.2 and 2.0 at % N were sintered at 1000C for 10 hrs while th e one doped with 10 at % N was sintered in 1 atm flowing nitrogen gas at 800C for 16 hrs in an attempt to avoid complete zinc nitride decomposition/oxidation. ZnO Films Doped with 0.2 and 2 at % N on c-Sapphire A series of films were deposited on c-sapphi re at temperature ra nging from 500 to 700C, oxygen pressure from 30 to 60 mTorr of oxygen. wen addition, growth was also examined at 500C using 120 mTorr of nitrogen as background gas. Powder X-ray diffraction of the ZnO films deposited from both targets shows no evidence of secondary phases (Zn3N2 peaks). The films are epitaxial and oriented ZnO [0001] even for the grown in 120 mTorr of N2. The X-ray diffraction pattern of the film deposited at 500C and 120 mTorr of nitrogen is shown in figure 7.1. Figure 7.2 shows the room temperatur e photoluminescence spectrum of an undoped single crystal ZnO substrate. The spectrum is dom inated by the near band edge emission peak at 376 nm (3.3 eV), which is associated to th e longitudinal optical phonon replica of the native donor-bound exciton recombination, and a broad peak emission peak in the visible, between 425 nm and 700 nm centered around 500-510 nm. The nature of emission in the visible is still controversial; however, emission in this range is usually associated to Zni or VO. The room

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207 temperature photoluminescence spectra for the film s are represented in figures 7.3and 7-4. For the films grown in oxygen, near band edge emissi on at 376 nm or 3.29 eV dominates the spectra. Its intensity decreases and br oadens with increasing target nitrogen content, substrate temperature and growth pressure. There is a blue shift in the NBE peak of the film grown at 700C and 30 mTorr of oxygen to 374 nm or 3.3 eV. This blue shift is most likely caused by improvement in the lattice and a decrease in N content in the film due to oxidation. For the film grown at 500C and 120 mTorr of nitrogen, NBE intensity is low and is no longer the dominating feature; the broad peak at 465 nm or 2.67 eV with a FWHM of 25 nm, associated with defects located deep into the band gap, is the most intense peak in the spectrum. Low temperature PL spectra were collected at 20K for the films. The results are show in figures 7.5 to 7-10. All of the films show ne utral donor bound exciton recombination, DX, at 3.36 eV. AX emission peak occurs at 3.31 eV and its relative intensity increases with increasing N content, decreasing substrate temperature and growth oxygen pressu re. For the film doped with 2 % N and grown at 500C and 120 mTorr of nitrogen, the AX intensity surpasses the DX intensity. These are indications of greater nitr ogen retention in the films deposited at lower growth temperatures and oxygen pressure. Also th e nitrogen content is higher in films grown with targets possessing high nitroge n content and/or grown in N2. The films also show neutral donor acceptor pair transition at 3.24 3.27 eV and longitudinal optical phonon replicas, 1LODAP, around 3.17 eV. Hall measurements for the films were carried out at room temperature. The results are shown in table 7-1. The films were conductive an d n-type. Resistivity of the films grown in oxygen is less than 0.12 .cm. The film grown 500C and 60 mTorr of oxygen from the target with 2 % N has the lowest resistivity, 0.032 .cm, a carrier concentration of 7.98 x 1018 cm-3 and

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208 the highest carrier mobility of 22.89 cm2/(Vs). The film grown at 500C and 120 mTorr of N2 has a resistivity of 1.08 .cm, carrier density 2.75 x 1018 cm-3 and carrier mobility of 2.05 cm2/(Vs). From the equation171: ] ) ( [3 / 1D D gap AN DAP E E E E (7-1) The N-acceptor binding energy can be esti mated. Assuming a donor biding energy, ED of 52 meV172 and taking ND, to be numerical average of the carri er density of the films, 7.03 x 1018 cm3, = 2.7 x 10-8 eV.cm and an average E(DAP) = 3.240 eV, EA is estimated to be 167meV. This value is in great agreement with th e reported value in the literature.172-175 Thermal Oxidation of ZnNxO(1-x) Films Upon annealing the above films in flowi ng oxygen at 600C for 15 minutes, the film grown in nitrogen turned insula ting while the resistivity of th e other films increases to 100 K .cm range. Upon further anneal, the resistivity of the films dropped cons iderably and the film turned n-type conductive. Since Zn3N2 decomposes in temperature as low as 724C, I therefore suspected the changes is in resistivity for annea ling past 30 min is due to complete oxidation of the Zn3N2 and reduction Zn interstitials left by a the decomposition of zinc nitride: 2 2 2 33 2 3 N ZnO O N Zn (7-2) To ensure enough nitrogen in the films, I deposited two series of films from the target doped with 10 at % N at 500C in 1 and 120 mTorr of N2 and subsequently performed thermal oxidation annealing. Thermally oxidized zi nc nitride films deposited by DC magnetron sputtering and PCVD were reported to show p-type ZnO formation with carrier concentration in the 1017 cm-3 range176-179. Powder diffraction X-ray of the films, show n in figures 7.12 and 7-13, confirms the presence of zinc oxide and zi nc nitride phases, ZnO (002), Zn3N2 (332) and (661), present in the

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209 as-grown films for both growth pressure. U pon annealing at 700C fo r 35 minutes, the zinc nitride phases disappear. However, I see the em ergence of ZnO (101) and (200). Upon further annealing, the ZnO (101) and (200) peaks di sappear and only ZnO (002) is present. The as-deposited films were translucent a nd their photoluminescen ce spectra, shown in figures 7.14 and 15, are flat U pon annealing at 700C for 35 min, the films become transparent and NBE emission is present in the photolumines cence spectra. The spectrum of the film grown in 1 mTorr of N2 shows the typical characteristics of ZnO films while only NBE emission is present in the spectra of the annealed film grown in 120 mTorr of nitrogen. Low temperature photoluminescence of both films annealed for 35 minutes and one for 45 minutes were measured at 20 K and are shown in figures 7.16 and 7.17. The spectrum of the film grown in 1 mTorr N2 and annealed for 35 min shows th e DX exciton emission at 3.36 eV, the AX emission at 3.31 eV. The donor-accepto r pair transition and it longitudinal optical phonon emission take place at 3.24 and 3.15 eV, resp ectively. The spectrum for the film grown at 120 mTorr and annealed for 35 min is shown in figure 7.17. Since the spectrum was taken at a much higher resolution, we are able to observe a broad emission peak, from 3.35 to 3.36 eV, for the DX. The AX at 3.31 eV, is the dominati ng feature of the spectra The DAP and LO-DAP peaks are observed at 3.27 and 3.15 eV, respectiv ely. There is a broad peak with its center around 3.27 eV, judging from the size of the AX, th is peak is due to th e recombination of free electron from the conduction band with an Acceptor FA. There is a shoulder at 3.381 eV which is attributed to the recombinat ion of free electron in the conduc tion band and holes in the valence band, Xa. From The FA, the acceptor energy can be estimated using equations 7-326, 90: e B o g AT k FA E E E 2 1 ) ( (7-3)

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210 Where Eg is the band gap at the measured temperature, kb and Te are the Boltzmann constant and the temperature of the electron in the lattice, respectively. Te is taken to be the measurement temperature. From equation 7-3, the binding energy of the N-related acceptor is estimated to be 167 meV. This value is very clos ed the value that I found earlier, hence confirming that the peak at 3.269 is due to FA. The low temperature photoluminescence of th e film grown at 500C in 120 mTorr of N2 and annealed for 45 minutes, figure 7.18, reveals the DX at 3.362 eV as the dominant peak of the spectrum. The intensity of the AX at 3.31 eV was greatly reduced. Both the DAP and LODAP peaks at 3.24 and 3.17 eV, respectively, are still present but at much lower intensities. The higher ratio of AX/DX intensity for the fi lm grown at 120 mTorr and annealed at 700C for 35 min indicate higher N acceptors present in the film as to compared to the other samples. The presence of high compensating defect conc entrations in film grown in 120 mTorr of N2 and annealed for 35 min is confirmed by the high degree of scattering in the plot of the Hall voltages as a function of applied magnetic filed. It is worth pointing out that compensation is not the only contributor to the scatter in the data. Th e low Hall voltage due to the high resistivity of the film makes the results vulnerable errors. Ta ble 7-2 shows the results of Hall measurements for the films. Upon annealing for 15 and 35 minute s, the films grown at both nitrogen pressures become insulating. Further annealing to 45 mi n turns the films n-type and reduces their resistivities from thei r insulating states when annealed at shorter time to 0.642 and 44.35 .cm for the films grown in 1 mTo rr and 120 mTorr, respectively. The changes in the X-ray diffraction patte rns, photoluminescence spectra confirm oxidation of zinc nitride upon ann ealing at 700C. However it is not completely clear why the films turns semi-insulating while low temperature PL shows evidence of compensation. The

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211 wide DX peak with a flat top in figure 7.17 su ggests the presence of multiple donor bands. Since the purity of zinc nitrid e powder that I used in the target is only 99 %, I suspect one of the impurities acting as a deep impurity or the forma tion of deep N-impurity complex that turns the films insulating. The biding energy of donors in ZnO can be calculated by137: ) 00534 0 4 0 (eV E E Eb D b D g (7-4) An estimated biding energy of 31 me V is derived from equation 3. This value is closed to the reported hydrogen and Zni binding energy in ZnO.43, 44, 163, 165 The Zni are due to the decomposition of Zn3N2. Hydrogen is known to passivate N-doped ZnO. Both theoretical and experimental studies of annealed ZnO doped w ith nitrogen points to No-H complex to be responsible for the passivation.44, 45, 180 It is possible that hydrogen incorpor ated in the films during the deposition process or diffusion during the thermal oxidation/a nnealing is causing the f ilms to turn insulating upon annealing. Summary Although I failed to achieve p-type conduc tivity by thermal oxidation/annealing of ZnNxO1-x, I have observed acceptor bound emission due to oxygen substitution by nitrogen at 3.31 eV, DAP emission at 3.24eV and FA at 3. 27 eV. From the DAP and FA emission, I was able to estimate the binding energy of the N -related acceptor to be 167 meV. Unlike the As doped ZnO films, the acceptor binding energy did not change with nitrogen concentration. The high resistivity of the films when annealed is associated to hydrogen passivation of the Nacceptor.

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212 This experiment shows that it maybe possible to attain p-type c onductivity by doping of ZnO with nitrogen and post thermal/oxidation of th e films. The high resistivity of the films when annealed may also be due some deep impurity from the low purity of the Zn3N2 powder. Our experiment maybe hampered by difficultly to control the thermal decomposition of Zn3N2.

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213 Table 7-1. Room temperature Hall data for the 0.2 and 2 % doped films. Substrate Temp. Conditions N content, Pressure Resistivity .cm Carrier Density x 1018 cm-3 Mobility Cm2/(Vs) Type 0.2 % N, 60 mTorr O2 0.062 12.65 8.01 n 2 % N, 30 mTorr O2 0.11 5.32 11.35 n 2% N, 60 mTorr O2 .034 7.98 22.89 n 500C 2 % N, 120 mTorr N2 1.08 2.75 2.05 n 700C 2 % N, 30 mTorr O2 .115 6.45 8.41 n Table 7-2. Room temperature Hall data for the as-grown and thermally oxidized films grown from the target doped with 10 at % N. growth Conditions Annealing Duration (min) Resistivity .cm Carrier Density cm-3 Mobility Cm2/(Vs) Type As grown 0.81 7.77E+179 n 15 7.54E+05 35* 1.65E+05 500C 1 mTorr N2 45 0.642 8.34E+1711.66 n As grown 63.14 3.05E+170.32 n 15 3.69E+06 35 1.45E+05 500C 120 mTorr N2 45 44.354 2.29E+166.14 n

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214 20304050607080102103 Intensity (counts/s)2ZnO (002) ZnO (004) Silver paint Figure 7.1. Powder X-ray diffraction pattern fo r the film grow at 500C and 120 mTorr of N2. 350400450500550600650700 0.0 0.2 0.4 0.6 0.8 Intensity (a. u.)Wavelen g th ( nm ) 3.3 eV

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215 Figure 7.2. Photoluminescence spectrum at 300K of an undoped single crystal ZnO substrate. 400500600700 3.29 eV3 2 9 e V3.29 eV2 at. % N, 700oC, 30 mTorr O22 at. % N, 500oC, 30 mTorr O22 at. % N, 500oC, 60 mTorr O2Wavelength (nm)Intensity (a. u.) 0.2 at. % N, 500oC, 60 mTorr O2 3.30 eV Figure 7.3. Photoluminescence spectra at 300K for the films grown in oxygen. 40050060070 0 Intensity (a. u.)Wavelength (nm) 2 % N at 500oC and 120 mTorr N2 3.26 eV 2.67 eV Figure 7.4. Photoluminescence spectra at 300K fo r the film doped with 2 % nitrogen and grown at 500C in N2.

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216 3.103.153.203.253.303.353.40 0.1 1 10 Intensity (a. u.)Energy (eV) 3.314 eV AoX 3.357 eV DoX 3.270 eV DAP Figure 7.5. Photoluminescence spectrum at 16 K fo r the film doped with 0.2 % N and grown at 500C and 60 mTorr of oxygen. 3.103.153.203.253.303.353.40 0.05 0.1 0.15 Intensity (a. u.)Energy (eV) 3.36 eV DoX 3.237 eV 3.381 eV Xa 3.31 eV AoX DAP-1LO DAP Figure 7.6. Photoluminescence spectrum at 16 K for the film doped 2% N and grown at 500C and 30 mTorr of oxygen

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217 3.103.153.203.253.303.353.40 0.1 1 3.2452 eV DAP 3.364 eV DoXIntensity (a. u.)Energy (eV) 3.317 eV AoX Figure 7.7. Photoluminescence spectrum at 16 K for the film doped with 2 % N and grown at 500C and 60mTorr of oxygen. 3.103.153.203.253.303.353.403.45 0.1 1 Intensity (a. u.)Energy (eV) 3.361 eV DoX 3.237 eV DAP AoX LODAP Figure 7.8. Photoluminescence spectrum at 16 K for the film doped with 2 % N and grown at 700C in 30 mTorr of O2.

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218 3.103.153.203.253.303.353.403.45 0.01 0.02 0.03 0.04 0.05 Intensity (a. u.)Energy (eV) 3.306 eV DoX 3.312 eV AoX 3.236 eV DAP DAP-1LO Figure 7.9. Photoluminescence spectrum at 16K fo r the film grown at 50 0C in 120 mTorr of N2. 3.123.183.243.303.363.42 0.01 0.1 Intensity (a. u.)Energy (eV) 3.316 eV 1-LO 3.361 eV DoX Xa 3.256 eV DAP Figure 7.10. Photoluminescence spectrum at 16K for an undoped film grown at 700C and 3 mTorr of (3 mol % 03 and 97 mol % 02) gas mixture.

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219 203040506070 102103104105 ZnO (101) ZnO (200)ZnO (004)Intensity (counts/s) ZnO (002) Indium contacts Al2O3(006) b: 500oC and 1 mTorr N2Annealed at 700oC for 35 min 203040506070 ZnO (101)ZnO (200) d: 500oC and 120 mTorr N2Annealed at 700oC for 35 minZnO (004)Al2O3(006)Indium contactsZnO (002) 2 203040506070 102103104 ZnO (004) ZnO (002) Al2O3(006) a: As grown in 1 mTorr N2 Zn3N2 (332) 203040506070 102103 Zn3N2c: As grown in 120 mTorr N2ZnO (004) Al2O3(006)ZnO (002) (332) (661) Figure 7.11. As grown and for the film grown from the target doped with 10 at % N at 500C 1(a) and 120 (c) mTorr of N2 and annealed in oxygen at 700C for 35 min (b and d).

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220 2530354045505560657075 102103104 ZnO (004)Intensity (counts/s)2ZnO (002) Al2O3(006) 500oC, 1 mTorr, Annealed at 700oCfor 45 min Figure 7.12. X-ray diffraction pattern of the film grown from the ta rget doped with 10 at % N at 500C and 1 mTorr of N2 and annealed in oxygen at 700C for 45 min. 202530354045505560657075 101102103104 500oC, 120 mTorr, Annealed at 700oCfor 45 min ZnO (004) ZnO (002)Intensity (counts/s) 2 Figure 7.13. X-ray diffraction patte rn of the film grow n from the target do ped with 10 at % N2 at 500C and 120 mTorr of N2 and annealed in oxygen at 700C for 45 min.

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221 350400450500550600650700 0.05 0.10 0.15 10 at. % N,500oC, 1 mTorr N2As grown Annealed at 700oC for 35 minIntensity (a. u.)Wavelength (nm) 3.3 eV Figure 7.14. Photoluminescence at 300K of the before and after 35 minutes annealed at 700C for the film grown at 50 0C and 120 mTorr of N2 from the 10 % N2 doped target and annealed in oxygen at 700C for 35 min. 35040045050055060065070 0 0.010 0.015 0.020 0.025 0.030 3.27 eV 10 at. % N,500oC, 120 mTorr N2As grown Annealed at 700oC for 35 minIntensity (a. u.)Wavelength (nm) Figure 7.15. Photoluminescence at 300K of the before and after 35 minutes annealed at 700C for the film grown at 500C and 1 mTorr of N2 from the 10 % N2 doped target and annealed in oxygen at 700C for 35 min.

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222 3.103.153.203. 253.303.353.40 0.1 1 Intensity (a. u.)Energy (eV) 3.314 eV AoX 3.239 eV DAP 3.365 eV DoX DAP-1LO Figure 7.16. Photoluminescence at 20K of th e film grown at 500C and 1 mTorr of N2 from the 10 % N doped target and annealed in oxygen at 700C for 35 min. 3.103.153.203.253.303.353.40 0.02 0.03 0.04 0.05 0.06 0.07 0.08 Intensity (a. u.)Energy (eV) 3.311 AoX 3.35-3.36 DoX 3.240 eV DAP 3.269 eV FA 3.150 eV DAP-LO 3.381 e V Xa Figure 7.17. Photoluminescence at 20K of th e film grown at 500C and 120 mTorr of N2 from the 10 % N doped target and annealed in oxygen at 700C for 35 min.

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223 3.153.203.253.303.353.403.450.1 1 Intensity (a. u.)Energy (eV) 3.362 eV DoX 3.312 eV AoX3.242 eV DAP3.170 eV DAP Figure 7.18. Photoluminescence at 20K of th e film grown at 500C and 120 mTorr of N2 from the 10 % N2 doped target and annealed in oxygen at 700C for 45 min. -8.0x104-6.0x104-4.0x104-2.0x1040.0 2.0x1044.0x1046.0x1048.0x104-8.5x1011-8.0x1011-7.5x1011-7.0x1011-6.5x1011 RH = [9.5 +/1.9] x 105 cm-3/C Applied Magnetic Filed (Oe)VH Figure 7.19. Changes in Hall voltage as a function of applied magnetic field for the film grown from the target doped with 10 at % N at 500C and 120 mTorr of N2 and annealed in oxygen at 700C for 35 min. is a films thickness dependent conversion factor.

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224 CHAPTER 8 GROWTH AND CHARACTERIZATION OF Nb AND Ta DOPED ZnO Results and Discussion I reported in the previous chapters on the gr owth of p-type ZnO films doped with As. Ptype doping of ZnO has been report ed in the literature using various elements such as N, P, As, Sb. Understanding the mechanism of p-type doping remains a challenge. It was widely assumed that oxygen substitution by the group V dopant lead s to p-type because of the difference in valence between the two ions. Th eoretical calculations, however, s uggest that substitution on the oxygen site by P, As or Sb requi res high energy and is predicted to yield a deep acceptor of energy around 950 meV. S. Limpijumnong122, 181 recently proposed AsZn-2 VZn or SbZn-2 VZn defect complex (equation 1), gui ded by strain relief and Coulomb interaction, as a possible answer to the experimental obs ervations of p-type conductivity of Asor Sb-doped ZnO. ) 2 (2 2 3 Zn Zn Zn Zn ZnV As V V As (8-1) Where AsZn is an arsenic sitting on a Zn site and VZn is a Zn vacancy. Their calculated 150 meV activation energy for the As related complex is consistent with published experimental data. They also predict that the use of moderate temperature and high O2 pressure during deposition would favor the formation of such a complex ove r other isolated As re lated defects such as: donor defects (AsZn), deep acceptor (AsO), or amphoteric interstitial (Asi).64, 87, 147, 182 It is widely known when As is heated in the presence of O2, oxidation to As5+ is favored over As3+.183 During the preparation of our targets, I si ntered the As-doped ZnO targets in air at elevated temperature, 1000C, for 10 to 12hrs to obtain dense targets. P-type ZnO could be realized under narrow proce ssing conditions leaving as an open question the role of As5+ in arsenic doped ZnO films. In this experiment, I use Ta and Nb as dopants in an attempt to determine if Ta5+ or Nb5+ substitution Zn site yields evid ence for acceptor state formation. If

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225 As5+ substitution on Zn site plays a role in the form ation of the acceptor state, it would also be of interests to understand the effects of ot her +5 valence state cations on Zn site. In addition having similar valence to As5+, Nb and Ta have similar ionic radius and slightly lower electronegativities. Nb has an atomic ra dius of 146 pm, an elec tronegativity of 1.6 and effective ionic radius of 64 pm in the 5+ stat e in the coordination number 6 while Ta has an atomic radius of 146 pm, an electronegativity of 1.5 and effective radius of 64 pm in the coordination number 6. These numbers are cl ose to As atomic radius of 124.8 pm, electronegativity of 2.0 and an effective ionic radius of 45 pm in the coordination number 6. Since these ions are also similar to Zn which ha s an 134 pm atomic radius an electronegativity of 1.7 and an effective ionic radius of 74 pm in the coordination number 6, I expect high solid solubility of Nb and Ta in ZnO.184, 185 The phase diagram of ni obium oxide and zinc oxide confirms high solid solubility of Nb in ZnO with solubility up to 25 mol. % in the wurtzite phase.186 In particular, I examine the transport, structur al and optical properties of 0.1 at % Nb, 0.1 and 1.0 at % Ta doped ZnO films. The ta rgets, pressed and sintered at 1000 C for 12 hrs in air, were fabricated using high-pur ity ZnO (99.9995%) doped the desire d dopant concentration using high purity (99.999%) Nb2O5 or Ta2O5. For some of the films, a thin buffer layer of MgO (5-20 nm) was pre-deposited at 450C and 1 mTorr of O2 on the c-sapphire substrates. Total film thicknesses range from 600-700nm. ZnO films Doped with 0.1 at % Nb and Ta Powder diffraction x-ray of the 0.1 at % doped films, not shown, reveals that the films are single phase, epitaxial, textured a nd highly oriented in th e ZnO [002]. The length of the c-axis of the 0.1 % Nb doped films increases linearly with increasing substr ate temperature when the films are deposited in 30 mTorr of O2 and on bare sapphire from 5.193 for the film grown at 400C

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226 to 5.204 for the film deposited at 800C, as seen in figure 8.1. C-axis of the films deposited at 500C also changes with increasing growth O2 pressure. The c-axis leng th of the films deposited at 30 and 60 mtorr of O2 were 5.193 while the film deposited at 90 and 120 mTorr of O2 had a c-axis of 5.190 and 5.205 respectively. The room the room temperature photolumin escence spectra for the films grown at 30 mTorr of O2 and doped with 0.1 at % Nb on bare sapphire substrates are show n in figure 8.2. All the films shows near band edge recombination emission peak at 375 nm, 3.31 eV, except for the film grown at 800C which show an NBE peak sh ift to 378 nm, 3.28 eV a nd the intensity of the NBE emission increases with increasing growth te mperature. The increase in intensity of the NBE recombination peaks and reduction of their FWHMs with increasing deposition temperature is an indication of improvements in the lattice with increasing growth temperature. The intensity of emission in the visible, howev er, decreases with increasing temperature. The effects of growth pressure on the photoluminescence propertie s of the 0.1 at % doped films grown at 500C are shown in figure 8.3. NB E recombination is at 375 nm, 3.31 eV. NBE recombination decreases while emission in the visible increases with increasing growth O2 pressure except for the films grown at 90 mTorr. Low temperature photoluminescence was perfor med at 20K and the spectra of the 0.1 % Nb doped films are shown in figure 8.4. The 0.1 Nb doped films grown at 500C and 30 mTorr of O2 shows a donor-bound exciton D1X emission, due to residual donor D1, at 3.353 eV followed by a second DNbX caused by the incorpor ation of Nb in the lattice. Around 3.25 eV is a shoulder in the plot which is attributed to the tran sition between a pair of D1 and residual acceptor (D1AP). Transition between Nb related donor and residual acceptor is observed at DNbAP at 3.153 eV. The same feat ures are seen in spectrum of the 0.1 at % Nb doped films

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227 grown at 500C and 90 mTorr of O2, figure 8.4 (b), however there is a decrease in the intensity of the DNbX. Increasing the growth pr essure to 120 mTorr of O2 causes most of the features of the low temperature photoluminescence spectrum of th e film doped with 0.1 at % Nb at 500C, figure 8.4 (c), to disappear leav ing a broad DX centered at 3.352 eV. The spectrum of the 0.1 at % Nb doped films grown at 700C and 30 mTorr of O2 shows a DX at 3.345 and the LODAP at 3.159 eV. The decrease in Nb concentration in th e lattice and lattice wors ening due to increasing growth pressure are responsible for the decreasing DNbX emission and peak broadening with increasing growth pressure. Where as at 700C and 30 mTorr of O2, an increase in Nb concentration due to the differen ce in sticking coefficient between Zn and Ta is responsible for the peak shift. Those changes are reflected in the changes in c-axis length as seen in figure 8.1. The binding energy of the donor can be estimated by137: eV E E X Db D g o00534 0 4 1 ) ( (8-1) Where ) ( X Dois the energy of the DNbX (3.314), b DE is the donor binding energy and Eg is ZnOs band gap at 16K. I get an estimat ed Nb related donor binding energy, Eb Nb of 86 meV. With the donor binding energy and the carrier con centration I estimated the residual Acceptor binding energy138, EA: ] ) ( [3 / 1 D b D g AN DAP E E E E (8-2) Here the term 3 / 1DNis the Coulombic Energy assuming an average the distance between donors is 3 / 1DN ND is the electron concentration deri ved from the Hall measurements and is materials dependent constant and is 2.7 x 10-8 eVcm for (ZnO) = 8.6. Using both DAPs and donor binding energies I derive an estimated acceptor biding energy of 210 meV. This value is good

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228 agreement with reported value of residual acceptor binding energy reported in the literature.163, 187, 188 The room temperature photoluminescence for th e films doped with 0.1 at % Ta grown at 400 and 500C at 30 mTorr on bare sapphire substr ates are shown in figure 8-5 and the low temperature PL spectrum for the film at 500C and 30 mTorr of O2 is shown in figure 8-6. At room temperature, NBE emission at 375 nm, 3.3 eV, dominates the spectra for both film. NBE intensity decreases with increasing growth temperature. At 20K, donor-bound exciton D1X emission, same observe in the 0.1 at % Ta doped films, is seen at 3.353 eV followed by a second DTaX caused by the incorporation of Ta in the lattice at 3.321 eV. The D1AP remains around 3.24 eV while the DTaAP is observed at at 3.163 eV. Th e Ta related donor binding energy, Eb Ta, is estimated to be 78 meV. Hall measurements of the doped films were carri ed out at room temperature. The results are tabulated in tables 8-1 and 8-2 and are shown in figures 8-7. and 8-8. The films are n-type and conductive. Resistivity of the 0.1 at % Nb doped films deposite d in 30 mTorr of O2, on the order of 0.15 .cm, increases with growth temperature from 0.138 .cm for the grown at 400C to 0.176 .cm for the deposited at 800C. The carrie r concentration of th e 0.1 at % Nb doped films increases with increasing substrat e temperature up to 700C from 4.5 x 1018 cm-3 for the film grown at 400C to 1.07 x 1019 cm-3 for the film grown at 700C. For the Nb doped film deposited at 800C, there is a decreas e in carrier density to 2.51 x 1018 cm-3. The Hall mobility decreases with increasing substrate temperature from 10.01 for the film deposited at 400C to 4.01 cm2/(V.s) for the films deposited at 700C. At 800C, Hall mobility increases to 14.12 cm2/(V.s).

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229 The effects of deposition pressure on the electri cal properties of the films doped with 0.1 at % Nb and grown at 500C are shown in figures 8-8 (a) and (b) and the data is arranged in table 8-2. The resistivity of the 0.1 at % Nb dope d films increases linear ly with increasing O2 pressure from 0.15 .cm for the film grown in 30 mTorr of O2 to 1.2 .cm for the film deposited at 120 mTorr. Both Carrier concentration and mobility decrease from 5.33 to 1.23 x 1018 cm-3 and 10 to 4.1 cm2/(V.s), respectively, when growth pressure is raised from 30 to 120 mTorr of oxygen. Hall measurements of the 0.1 at % ZnO:Ta films grown at 400 and 500C in 30 mTorr of oxygen reveal that the film are highly conductive and n-type. The re sults are presented in table 83. Resistivity and carrier density increase fr om 0.023 to 0.026 Ohm/cm and 1.7 to 3.81 x 1018 cm-3 while mobility decreases from 13.6 to 6.3 cm2/(V.s). As I have seen earlier, those changes are probably due to a reduction of intrinsi c defects with increasi ng temperature and the difference in vapor pressure between ZnO and Ta2O5 or an reduction in Tai and increase TaZn as temperature is raised. The transport properties of se veral undoped films ar e shown in figures A-5 and A-6 in the appendix. For the undoped films grown at 500C in 30 and 60 mTorr of oxygen, figure A-5, resistivity incr eased from 0.108 .cm for the film grown in 30 mTorr to 0.508 .cm for the film grown at 60 mTorr of O2. Carrier density and Hall mobility of the films decreased from 2.55 x 1018 cm-3 and 22.68 cm2/(V.s) to 9.78 x 1017 cm-3 and 19.1 cm2/(V.s) when growth pressure is raised from 30 to 60 mTorr of oxygen. The surface became rougher from around 6 nm to 10 nm as growth pressure was raised For the film deposited in 30 mTorr of oxygen at 500 and 600C, figure A-6, resistivity and Hall mobility of the films increased from 0.108 .cm and 22.68 cm2/(V.s) to 0.364 .cm and 33.41 cm2/(V.s) for the film grown at 600C. Carrier density of the

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230 films decreased from 2.55 x 1018 cm-3 to 5.13 x 1017 cm-3 as deposition temperature was raised from 500C to 600C. I have seen that increasing the substrate temp erature of the 0.1 at % Nb doped films causes the length of the c-axis of the films to increase towards the c-axis length of bulk ZnO, 5.206 The improvement in crystallinity could be s een in the room temperature photoluminescence where NBE emission intensity increases and emissi on in the visible, believed to be caused by points defects, decreases with increasing growth temperature. The changes in the transport properties reflect the multiple process occurring in the films. As substrate temperature increases, the concentration of residual donor s decreases and the mi crostructure improves. However, the concentration of the dopant increases because of the difference in sticking coefficient between Zn and Ta. As temperature increases, the Nb or Ta related donor becomes the dominant donor in the films. This is seen in the redshift in the DX in the low temperatur e spectra of the Nb doped films with increasing temperature. The D1X at 3.352 eV seen for th e film grown at 500C and 30 mTorr is no longer then dominant feature for the film deposited at 700C. The changes in the transport properties of the films were small. Re sistivity, carrier concentration and Hall mobility of the films remains on the order of 0.1-0.2 .cm, 5 x 1018 cm-3 and 10 cm2/(V.s) for the films doped with Nb and 0.02 .cm, 2 x 1019 cm-3 and 10 cm2/(V.s) for the films doped with Ta, respectively. Although, increasing the oxygen growth pressure of the films doped with Nb has adverse effects on the crystalline quality of the films. However, it reduces the concentration of the residual donor in th e films. Room temperature photol uminescence of the films show decreasing band edge luminescence and emission in the visible with incr easing oxygen pressure. The decrease in point defects cau ses carrier concentration to d ecrease with increasing growth pressure. The increase in grain de nsity and other microstructural de fects with increasing growth

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231 oxygen pressure causes carrier mobility to decr ease. Resistivity increases because of the reduction of residual donors and the decrease of carrier mobility. Figure 8-9 shows the low temperature PL spect ra of a 0.1 % Nb doped ZnO film (solid line), 0.1 % Ta doped film (dashed lines) a nd an undoped ZnO film (dotted line). Doping the film with Nb and Ta causes a redshift of the donor-bound exciton (D1X) from 3.362 eV for the undoped film to 3.353 eV for the doped films. As it can be seen, The DNbX is deeper than the DTaX at 3.311 and 3.327 eV, respectively. The shift in D1X causes by the incorporation of Ta and Nb cause D1AP to shift accordingly while the DxAP shift according to the position of the Nb or Ta doping energy level. DNbAP has emission energy than DTaAP since DNb is deeper in the bandgap than DTa. The difference in binding energy between the Nb and Ta related donors cause the films doped with Ta to be more conductive w ith carrier concentrati on about an order of magnitude higher for the same doping level and growth conditions. ZnO films Doped 1 at % Ta At 1.0 at % Ta doping, the films remain epitax ial, single phase and hi ghly textured in the ZnO [002] as it can be seen in fi gure 9-10. The high resolution x-ray -rocking curves for the films grown at 1 and 90 mTorr of oxygen, shown in figure 8-18, show that the peaks are broad with FWHMs of 1.347 and 2.105 for the films depos ited at 1 and 90 mTorr correspondingly. The surface morphology of the films also worsens with increasing growth pressure. Surface roughness steadily increases with in increasing pressure, as shown in figure 8-11 and table 8-13, from 1.218 to 13.072 nm as pressure increa sed from 1 mTorr to 90 mTorr of oxygen. The decrease of crystalline and surf ace quality with increasing pressure is a result the decrease of the kinetic energy of the ablated part icles present in the plume due to increase interaction between background gas and the plume.

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232 The degradation of the crystalline qualit y can be observed in the photoluminescence spectra of the films shown in figure 8-12. NBE luminescence, at 376 nm or 3.30 eV, decreases with increasing growth oxygen pressure. Emission in the visible, due to Zn interstials and O vacancies, decreases with increasing growth pres sure. For the film grown in at 500C and 60 mTorr, defects deep into the band ga p can be seen at 2.97 and 2.84 eV. The transport properties of the films are shown in figure 8-13 a nd table 8-2. Increasing the growth oxygen pressure from 1 mTorr to 90 mTorr cau ses the resistivity of the films 1.0 at % Ta doped ZnO films grown on thin MgO buffer layer, figure 8-13 (a) to increase from 4.85 to 102 .cm and Hall coefficient of the fi lms, Figure 8-13 (b), to decrease s as pressure increases. As a result of the changes in Hall coefficient and resi stivity, carrier concentr ation decreases from 4.6 to 1.54 x 1018 cm-3 and Hall mobilities for the films, on the order of 0.3 cm2/(V.s), slightly increases with increasing growth pressure, as se en in figure 8-13 (c) a nd (d). The electrical characteristic of the 1.0 at % Ta dop ed films grown at 500C and 30 mTorr O2 did not fall inline with the other films As I have seen have seen in the earlier sect ion, Ta introduces a re latively deep donor lever compare to the residual donor le vel. With increasing growth pre ssure, the concentration of point defects decreases. The conductivity ( ), Hall coefficient (RH) and Hall mobility ( H) are defined as: n Ta n Ta Ren n n e ) (, (8-3) Ta Ta R Hen n n e R 1 ) ( 1 (8-4) and H HR (8-5)

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233 Where e is the charge of an electron; nR and nTa are electrons contribution from residual donors and Ta, respectively and n is the electron mobility. As oxygen pressure is raised, nR decreases and nTa becomes the majority carri er causing resistivity and |RH| to increase. Mobility increases slightly because of the effects of the decrease in carrier concen tration is negated by the increase in structural scattering. Further Hall analysis of the film grown at 30 mTorr was performed in our Quantum Design physical properties measurements system. The plot of resistivity as a function of reciprocal temperature, shown in figure 8-14, confirms th e semiconducting nature of the film. Resistivity increases exponential as temperature is lowered. The plot Hall coeffi cient as a function of temperature, figure 8-15, reveals a singl e band conduction region from 350-100K where electrons are the dominant carrier and a region of mixed conduction as temperature is further lowered below 100K. Beyond this temperature th e data become noisy and Hall coefficient changes sign irregularly indicati ng a region of mixed conduction as 2 2n pn p (8-6) Carrier density, N is experimentally calculated by HeR 1 where e is the electronic charge. Figure 8-16(a) is the plot of ca rrier concentration against reci procal temperature for the [100350]K temperature range. The plot shows carri er concentration increases with decreasing temperature. This is because the residual donors have lower activation en ergy than the activation energy of the Ta related donor. The data is fair ly linear and activation energy of the residual donor can be estimated from the slope to the linear fit (Ea/2kb), where Ea is the residual donor thermal activation energy and kb is the Boltzmann constant. Th e derived activation energy, Ea, is 19 meV +/5 meV, which is in good agreement with the activation energy found in Figure 8-16 (b). In the case of dual conduction bands, the charge balance equation [

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234 ) / exp(2 3T k E T const nb a ] is more appropriate to derive the carrier activation energy. Ea is derived from the slope of linear f it, solid line of figure 8-16 (b), of the Arrhenius plot of N x T3/2 against reciprocal of temperature. The slope, Ea/kb, is 0.2245 indicating Ea is 19 meV +/0.3 meV. The plot of Hall mobility as a function of temperature is shown in figure 8-17. Carrier mobility decreases as temperature is decrease down to 100K as pictured in the inset of in the figure. Beyond this temperature, the data become noisy as mobility is proportional to the Hall coefficient. This is an indication of multiple ba nds conduction where electrons and holes coexist. Thermal Stability of the 1 at % Ta Doped ZnO films I have performed several anneals in air, oxygen and Ar at different temperatures and lengths of time on several sets of samples of 1.0 at % Ta doped of ZnO grown at 500C and MgO buffered c-sapphire. The microstuctural, electr ical, optical properties of the annealed films are presented below. Improvements in the microstructure by anneal ing can be seen by AFM and the changes in the FWHMs of the high resolution -rocking curves. Figures 8-33 to 8-35 shows grain growth and evolution with annealing time for 1000C ann eals of the film grown at 500C and 90 mTorr of oxygen. Surface morphology of the films was analyzed by means of AFM for a 1 m2 and the results are presented in figure 8-29 and tabl e 8-13. Surface roughness decreases continuously with increasing annealing temper ature from 13.072 for the as-gro wn film to 0.053 nm for the film annealed at 1200C. Annealing for extend ed period of time at 1000C also changed the surface morphology of the annealed films significan tly from 13.073 nm of the as-grown film to 3.497 nm for the film annealed at 1000C for 60 min. One can also s ee the effects of ZnO evaporation of the surf ace in picture 8-36.

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235 The high resolution -rocking curves of the as-grown (90 mTorr of oxygen) and some of the annealed films are shown in figure 8-18. Th e as-grown FWHM of the 1.0 at % Ta doped ZnO film deposited at 500C and 90 mTorr of oxyge n was 2.104. When annealed at 1000C and 1200C for 5minutes, FWHM decreased to 0.598 and 0.464, respectively. Annealing the films at 1000C for 30 minutes causes the FWHM to br oadens to 0.802. Further annealing at 1000C for 60 minutes causes the emergence of a second peak at = 16.5 with a 0.576 FWHM. The FWHM of the original peak decreases to 0.756 from 0.802 for the film annealed at 1000C for 30 min. Powder X-ray Diffractions for some of the films are shown in figures 8-19 and 8-20 show only a single phase are present in the film s aligned in the ZnO [002] even for the film annealed at 1000C for 60 min. For this reason I believe the peak observed at = 16.5 in figure 8-18 is from a MgxZn1-xO layer resulting from the inter-di ffusion of the MgO buffer layer and the ZnO film at the interface. MgxZn1-xO can remain in the wurtzite structure with a Mg content up to 33%. Figure 8-21 is the room temperature photolumin escence spectra for the 1.0 at % Ta doped films grown at 500C and 1 mTo rr of oxygen annealed in air for 5 min from 900C to 1200C. The photoluminescence of the films change comp letely compared to the as-grown. When annealed in air, NBE emission completely disappears while emission between 400 and 700 nm dominates the spectra. The emission in this range is clearly different from the as-grown film. Intensity declines with increasing annealing temp erature. When the films grown at 90 mTorr of oxygen are annealed at the same conditions, NBE emission is also absent but this time the emission peak is narrower and less intense. For the 1200C anneal, there were almost no emission present in the spectra as seen in figur e 8.22. Those changes indicate increasing of nonradiative defects centers as annea ling temperature increases.

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236 Annealing the films in air resulted in semi-insulating behavior. Hall voltages were too low to be accurately measured. The graph of resistivity of the films as a func tion of temperature for a 5 min anneals is shown in Figure 8.23. Resistivity for both films increased at the same rate from as grown films; from 4.85 to 3.22 x 106 .cm for the film grown at 1 mTorr and from 101.88 to 5.13 x 106 .cm for the film grown at 90 mTorr. Re sistivity for both films decreases from thereon as annealing temperature increases. Wh en some of the films grown at 90 mTorr of oxygen were annealed in air at 1000C for a longe r period of time, resistivity decreases from 5.31 x 105 .cm for the 5 min anneal to 1.02 x 103 .cm after a 30 min anneal and 103 .cm for the 60 min anneal as shown in figure 8-24. The dr op in resistivity for temperatures higher than 900C and for 1000C annealing duration are probably due to changes in films composition near surface region of the films. ZnO has a tendency to evaporate at high temp erature and segregation of defects to the grain boundaries. When a series of 1.0 at % Ta doped ZnO fi lms, deposited on MgO buffered sapphire at 500C and 30 mTorr of oxygen, were tube furnace a nnealed at different te mperatures and 1 atm of flowing oxygen, resistivity increases rapidly with increa sing annealing temperature. Annealing at 450C in flowing oxygen cause resis tivity to increase dramatically from the as grown state from 1.58 to 329 .cm, with carrier concentr ation changing from 33.73 x 1019 to 8.74 x 1015 cm-3and mobility from 0.1 to 2.17 cm-2/(V.s). As annealing temperatures increases, resistivity, as pictured in figur e 8-25, continue to increase st eeply and the films become semiinsulating preventing us from meas uring the Hall characteristics of the annealed films. The room temperature photoluminescence spectra for the fi lms annealed at 450 and 900C are shown in figure 8-26. NBE luminescence of the film annealed at 450C is to a large extent lower than for

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237 the as-grown film. For the 900C oxygen anneal NBE is further reduced and lower energy emission is almost nonexistent. Annealing in flowing Ar was performed on a se ries of films grown at 500C and 30 mTorr of oxygen at temperature ranging from 700 to 1000C for 5 min. Resistivity of the films drops rapidly from 2.43 .cm for the as-grown film to 0.16 .cm, then slowly increases as annealing temperature increases, as seen in figure 8-27. Carrier mobility incr eases continuously with increasing annealing te mperature from 2.18 cm2/(V.s) for the as-grown film to 29.74 cm2/(V.s). Carrier density increased from 1.2 to 8.1 x 1018 cm-3. Room temperature photoluminescence spectra of the 5 min Ar annealed films, figure 8-28, reveals increase in the NBE emission intensity for the annealed films as compared to the asgrown. Luminescence in the visible also increased with annealing. For the film annealed at 900C, emission in the visible becomes the domin ating feature of the spectrum. The FWHM and the intensity of NBE emission of n-type films ar e good indications of the crystalline quality of the films. The nature of emission in the visible is still disputed in ZnO, however, it is thought to be cause by Zn interstials and oxygen vacancies. I performed air/Ar and Ar/air back-to-back furn ace annealing of a set of 1.0 at % Ta doped film at 1000C for 5 min in each environment. Th e results are presented in table 8-12. After the Air/Ar anneal, the films resistivity, car rier concentration and mobility were 0.34 .cm, 1.10 x 1018 cm-3 and 18.28V cm2/( s) as opposed to 90.51 .cm, 3.05 x 1015 cm-3 and 22.58 cm2/(V s) after Ar/air. The final Hall mobilities of the two films are closed. However, the carrier concentrations and resistivities of the films are 3 to 4 orders of magnitude apart. The changes in electrical properties of the Ar/air differ from the transport properties of the films after annealing

PAGE 238

238 in air Ar alone as opposed to the air/Ar. The trans port properties of the films annealed in air/Ar is similar to the properties of the film s annealed in fl owing Ar only. It was also seen that performing oxygen RTA on a series of films grown at 60 mTorr of oxygen for 5 seconds also turned the films semi -insulating. Figure 831 shows the graph of resistivity as a function of O2 RTA temperature. Resistivity of the films increased from 72 .cm for the as-grown film to 105 .cm for the annealed films. The resistivity follows a pattern similar to the resistivity of the films annealed in air for 5 minutes, as I have seen earlier and can be seen in figure 8-23. To further understand the films effects of O2 on the transport properties of the films and isolate any effects of surface Zn-C-Ta-O comp lexes or other complexes due to post-growth atmospheric contamination, I grew some sample s of 1.0 at % Ta doped ZnO film 500C and 30 mTorr of oxygen. After deposition, I quickly fi lled the growth chamber with 400 mTorr of O2 while maintaining the film at growth temper ature. Once the chamber reached 400 mTorr, the films were rapidly cooled the sample by turning the heating quartz lamp off. The resistivity, carrier concentration and mobility were 1 x 104 .cm, 5.8 x 1015 cm-3 and 0.76 cm-2/(V s) as opposed to 1.58 .cm, 5.8 x 1019 cm-3 and 0.11 cm-2/(V s) for the as-grown film. Figure 8-30 shows that NBE photoluminescence for the film is near the noise level. Upon rapid thermal annealing in N2 at 900C for 2 sec, resistivity, carrier concentration and mobility decrease to 22.52 .cm, 9.16 x 1014 cm-3 and an unreal 300 cm-2/(V.s), respectively. A rapid increase of resistivity of the films to the semi-insulating state is observed when annealed the 1.0 at % Ta doped films. This not typical of undoped ZnO films, as we have seen for the annealed undoped film. If the residual donors are cause by oxygen vacancies or zinc interstitials, annealing in oxygen should have comparable change s in the Ta doped films as in the

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239 undoped ZnO film. The rapid increase in resistivity when annealed at 450C in flowing oxygen, RTA in oxygen at 5 sec and when the film was c ooled in 400 mTorr of Oxygen as opposed to 30 mTorr of oxygen points towards el ectronic traps being introduced by oxygen. I believe the films becomes insulating and less luminescent when ann ealed in an oxygen rich environments because of the fast diffusion of oxygen through the grai n boundaries and its adsorption at the surface and the grain boundaries surfaces causing surface elec tron depletion or electron traps. The surface adsorption process is amplified by surface Ta and th e migration of Ta to the grain boundaries and Tas preferential TaO6 coordination.189 When an undoped ZnO film, deposited at 500C, 20 mTorr of oxygen and on MgO buffered sapphire, wa s annealed in flowing oxygen for 5 min at 800C, changes in transport parameters were le ss severe. The films resistivity and carrier mobility increased from 73.02 to 320 .cm and 0.26 to 10.35 cm-2/(V.s), in that order, while carrier concentration de creased from 3.29 x 1017 to 1.85 x 1015 cm-3. When the films are annealed in Ar, oxygen first desorbs from the surface and boundary surfaces causing surface conductivity to increase and oxygen deficient grains.20, 160, 161 Figure 8-32 compares the low temperature photoluminescence spectra of: an undoped ZnO film (dotted line); 0.1 at % Nb and Ta doped (das hed lines); some As doped films (solid lines). ZOA1 is an 0.2 at % As doped film doped from As2O3; ZOA2 is 0.2 at % As doped ZnO from Zn3As2 and ZOA3 is an 0.02 at % As doped ZnO film from Zn3As2. One can observe that the Nb related DX has similar energy has the AX of so me of the the films doped with As from zinc arsenide. Since there is a redshift in the DxX from switching from Nb to Ta and the main difference between the to ion is their electrone gativity (1.6 for Nb and 1.5 for Ta), I would expect further red-shiting for the As5+ related donor since the electrone gativity of As is 2.0. I see two shoulders in the spectra of the ZOA2 film (0.2 at % As doped film from Zn3As2) and the

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240 ZOA1 (0.2 at % As from As2O3) film at 3.287 eV. In the case this shoulder is due to As5+ I estimate the binding energy of donor to be 103 meV. Summary In the preceding chapter, I anal yzed the electrical, optical and structural properties of the 0.1 and 1 at % Nb and Ta doped ZnO films grow n on both bare c-sapphire substrate and thinly MgO buffed c-sapphire. The films were n-type c onductive regardless of dopant concentration. The films doped with 0.1 at % Nb were more re sistive than the 0.1 % Ta doped films grown at similar conditions. Low temperature shows the Nb related donor-bound exciton emission (3.311 eV) to be deeper in the band gap than ZnO than Ta (3.327 eV). Calculations from the PL peaks position allowed us to estimate the optical bi nding energy residual donor to be around 52 meV; residual acceptor optical binding energy to be around 200 meV; Nb and Ta optical binding energies to be around 80 and 90 meV, respectively. Increasing the Ta content of the film to 1. 0 at % decreased the carrier mobility due to higher impurity scattering despite the improvement of the crystalline quality of the films due to the use of a thin MgO buffer layer. The resistivity of the films increased and carrier concentration of the 1.0 at % Ta doped film s decreased almost two orders of magnitude increasing growth oxygen pressure while hall mob ility remains on the same order. Temperature dependant Hall dependant reveals that the 1. 0 at % Ta doped films display semiconducting behavior and the residual donors have an estima ted of thermal activati on energy 20 3 meV. Annealing the 1.0 at % Ta doped films in Ar increases the conductivity and carrier concentration of the films where as annealing in oxygen rich environments, furnace anneal or RTA, rapidly increase the resistivity of the f ilms to a semi-insulating state. I believe oxygen adsorption at the films surface a nd crystallite interfaces due Ta acts as electronic traps causing

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241 the steep increase in resistivit y of the films and the decrease of the photoluminescence of the films, even though they remain transparent. When comparing the low temp erature photoluminescence of th e films doped with Nb and Ta to some of the As doped film, I first observe that the Nb related DX has similar energy as the AX of the some of the films doped with arsenic fr om zinc arsenide. I also observe a shoulder in the spectra of some of the film at 3.287 eV assuming this shoulder is an As5+ related donor I estimated the binding energy of this donor to be 103 meV.

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242 Table 8-1. Room temperature Hall measurements data for the 0.1 at % Nb doped ZnO films grown on sapphire at 30 mTorr of oxygen partial pressure. Temperature C Resistivity /cm Carrier Density cm-3 Mobility Cm2/(V.s) Type 400 0.138 4.52 x 1018 10.01 n 500 0.117 5.34 x 1018 9.99 n 700 0.142 1.07 x 1019 4.11 n 800 0.176 2.51 x 1018 14.12 n Table 8-2. Room temperature Hall measurements data for the 0.1 at % Nb doped ZnO films grown on sapphire at 500C and diff erent oxygen partial pressures. Pressure (mTorr) Resistivity ( /cm) Carrier Density cm-3 Mobility Cm2/(Vs) Type 30 0.117 5.33E+18 9.99 n 60 .444 3.16E+18 4.44 n 90 0.698 4.32E+18 2.07 n 120 1.24 1.23E+18 4.1 n Table 8-3. Room temperature Hall measurements data for the 0.1 at % Ta doped ZnO films grown on sapphire at 30 mTorr of oxygen partial pressure. Temperature C Resistivity /cm Carrier Density cm-3 Mobility Cm2/(V.s) Type 400 0.023 1.95 x 1019 13.6 n 500 0.026 3.81 x 1019 6.3 n Table 8-4. Room temperature Hall measurements data for the 1.0 at % Ta doped ZnO films grown at 500C on sapphire at 30 mTorr of oxygen partial pressure. The exact type could not determine accurately. Pressure (mTorr) Resistivity ( /cm) Hall Coefficient Carrier Density cm-3 Mobility Cm2/(V.s) Type 1 4.85 -1.36 4.60 x 1018 0.28 n 30* 1.58* 0.17* 3.73 x 1019* 0.11* 60 72.14 -27.38 2.28 x 1017 0.38 n 90 101.88 -40.41 1.54 x 1017 0.40 n

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243 Table 8-5. Temperature dependent Hall data for the 1.0% Ta doped film grown at 500C and 30 mTorr of oxygen on MgO buffer layer. Temperature (C) Resistivity ( /cm) Hall Coefficient Carrier Density cm-3 Mobility Cm2/(Vs) Type 10 3.83 1.39 4.48 x 1018 0.36 n 20 3.40 -1.75 -3.57 x 1018 0.51 n 30 3.14 4.30 1.45 x 1018 1.37 n 40 2.95 -0.92 -6.82 x 1018 0.31 n 50 2.80 0.29 2.18 x 1019 0.10 p 60 2.67 0.04 1.70 x 1020 0.01 p 70 2.55 -0.78 -8.04 x 1018 0.30 n 80 2.45 0.11 5.47 x 1019 0.05 p 90 2.36 0.12 5.34 x 1019 0.05 p 100 2.109 -0.0831 -7.51 x 1019 0.0394 n 150 1.738 -0.295 -2.12 x 1019 0.170 n 200 1.451 -0.3107 -2.01 x 1019 0.214 n 250 1.235 -0.6081 -1.03 x 1019 0.492 n 300 1.073 -0.2.73 -2.29 x 1019 0.254 n 350 0.987 -0.806 -7.74 x 1018 0.817 n Table 8-6. FWHMs of the High resolution -rocking curves shown in Figure 8-18. Conditions FWHM A: as-grown 1 mTorr O 1.347 B: as-grown 90 mTorr O 2.105 C: Annealed at 1000C for 5 min 0.598 D: Annealed in air at 1000C for 30 min 0.802 E: Annealed in air at 1000C for 60 min E1: 0.756 E2: 0.576 F: Annealed in air at 1200C for 5 min 0.464 Table 8-7. Hall measurements for the 1.0 at % Ta doped ZnO films grown in 1 mTorr of oxygen and annealed in air for 5 min. Annealing Temperature (C) Resistivity ( /cm) Hall Coefficient Carrier Density cm-3 Mobility Cm2/(Vs) Type 900 3.22 x 106 1000 2.05 x 106 1100 4.13 x 106 1200 9.36 x 103 -16278 3.83 x 1014 1.74 n

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244 Table 8-8. Hall measurements for the 1.0 at % Ta doped ZnO films grown in 90 mTorr of oxygen and annealed in air for 5 min. Annealing Temperature (C) Resistivity ( /cm) Hall Coefficient Carrier Density cm-3 Mobility Cm2/(V*s) Type 900 5.13 x 107 1000 5.31 x 105 1100 3.31 x 106 1200* 4.27 x 104 -75704 8.24 x 1013 1.77 n Table 8-9. Hall data for the films grown at 90 mT orr and annealed in air at 1000C from 0 to 60 min. Temperature (C) Resistivity ( /cm) Hall Coefficient Carrier Density cm-3 Mobility Cm2/(V s) Type As grown 101.88 -40.41 1.54 x 1017 0.40 n 5 5.31 x 105 Too resistive 30 1.02 x 103 -2.27 x 104 2.75 x 1014 22.30 n 60 1.30 x 102 -2.79 x 103 2.23 x 1015 21.60 n Table 8-10. Resistivity data for the films anneal ed in 1 atm of flowing oxygen. The films were 1.0 at % Ta doped ZnO grown at 500C and 30 mTorr of oxygen. Temperature (C) Resistivity ( /cm) 0 2.43 450 329 700 1.70 x 106 900 7.90 x 106 Table 8-11. Hall data for the films annealed in 1 atm of flowing Ar. The films were 1.0 at % Ta doped ZnO grown at 500C and 30 mTorr of oxygen. Temperature (C) Resistivity ( /cm) Hall Coefficient Carrier Density cm-3 Mobility Cm2/(V*s) Type As grown 2.43 -5.30 1.18 x 1018 2.18 n 700 0.16 -1.66 3.76 x 1018 10.35 n 900 0.19 -4.74 2.50 x 1018 24.95 n 1000 0.26 -7.61 8.08 x 1018 29.74 n

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245 Table 8-12. Summary of the eff ects of annealing on the electrica l properties of the films (upper) and isochronal Ar/Air and Ai r/Ar anneals (5 min/5 min) at 1000C and 1 atm (lower). Anneals Resistivity ( /cm) Hall Coefficient Carrier Density cm-3 Mobility Cm2/(V*s) Type As grown 101.88 -40.41 1.54 x 1017 0.40 n Air 5.31 x 105 Ar 0.26 -7.61 8.08 x 1018 29.74 n Isochronal Anneals Ar/air 90.51 -2043.7 3.05 x 1015 22.58 n air/Ar 0.34 -6.20 1.01 x 1018 18.28 n Table 8-13. Surface roughness of the as-grown a nd some of the annealin g films from a 1 x 1 m. As grown films 5 min air anneals 1000C air anneals Pressure (mTorr O2) Roughness (nm) Temperature (C) Roughness (nm) Time (min) Roughness (nm) 1 1.22 900 12.53 5 12.53 30 3.22 1000 8.76 30 4.80 60 12.37 1100 1.59 60 3.50 90 13.07 1200 0.053 Table 8-14. Hall data for an undoped ZnO film grown on a thin MgO buffer layer at 500C and 20 mTorr of Oxygen. The anneals were pe rformed in flowing high purity oxygen or argon at 1 atm for 5 min. Conditions Resistivity ( /cm) Carrier Density cm-3 Mobility Cm2/(V*s) Type As grown 73.02 3.29 x 1017 0.26 n 800C, O2 320 1.85 x 1015 10.35 n 800C, Ar 0.11 2.55 x 1018 22.68 n Table 8-15. Effects of cooling pr essure on the electrical propert ies of 1.0 at % Ta doped films grown at 500C and 30 mTorr of oxygen. Cooling Pressure Resistivity ( /cm) Carrier Density cm-3 Mobility Cm2/(V*s) Type 30 mTorr O2 1.58 3.73 x 1019 0.11 n 400 mTorr O2 1.41 x 104 5.8 x 1015 0.76 n Table 8-16. Effects of 5 sec oxygen rapid thermal a nnealing on the electrical properties of the of 1.0 at % Ta doped films grown at 500C and 60 mTorr of oxygen. RTA Temperature (C) Resistivity ( /cm) 800 8.01E+5 900 1.51E+5 1000 1.2E+5

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246 4005006007008005.192 5.194 5.196 5.198 5.200 5.202 5.204 C-axis Length (A)Growth Temperature (oC) 0.1 at. % Nb, 30 mTorr of O2 b20406080100120 5.188 5.190 5.192 5.194 5.196 5.198 5.200 5.202 5.204 5.206 0.1 at. % Nb, 500oC Oxygen Pressure (mTorr) a Figure 8-1. C-axis lengths of th e ZnO films doped with 0.1 at % Nb deposited at 500C as a function of growth oxygen pressure (a); a nd as a function of growth temperature grown in 30 mTorr of oxygen (b).

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247 400500600700 800oC 700oC 500oCIntensity (a.u.)Wavelength (nm) 400oC 3.3 eV Figure 8-2. Effects of growth temperature on the room temperature photoluminescence of the 0.1 at % Nb doped ZnO films on bare sapphire substrates. 400500600700 60 mTorr 120 mTorr 90 mTorrIntensity (a.u)Wavelength (nm)30 mTorr 3.18 eV Figure 8-3. Room temperature photoluminescence spectra of the 0.1 at % Nb doped ZnO films grown at 500C and different growth pressure on bare sapphire substrates.

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248 3.13.23.33.4 0 .01 0.1 3.13.23.33.4 0.1 1 3.13.23.33.4 0.01 0.1 3.23.4 0.1 b: 500oC, 90 mTorr 3.353 eV DoX3.153 eV DNbAP 3.311 eV Do NbX 3.325 eV D1AP3.325 eV D1AP3.352 eV D1 0X +DoNbX c: 500oC, 120 mTorr 3.325 eV D1AP XaEnergy(eV)d: 700oC, 30 mTorr 3.345 eV D1 oX+ DoNbX 3.159 eV DNbAP 3.154 eV DNbAP a: 500oC 30 mTorr 3.311 eV DoNbX 3.3534 eV D1 oX Xa Figure 8-4. 20K photoluminescence spectra for sele cted films doped with 0.1% at Nb on bare sapphire substrate [a, b,c were grown at 500C in 30, 60 and 90 mTorr of O2, respectively and d was grown at 700C and 30 mTorr].

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249 400500600700 400 C 500 CIntensityWavelength (nm) Figure 8-5. PL at 300K spectra for sele cted films doped with 0.1% at Ta. 3.103.153.203.253.303.353.40 Intensity (a. u.)Energy (eV) 3.353 eV D1 oX 3.304 eV D0 TaX 3.225 eV D1AP 3.147 eV DTAAP Xa Figure 8-6. PL at 16K of the film doped with 0. 1% at Ta deposited at 500C and 30 mTorr of O2.

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250 400500600700800 0.11 0.12 0.13 0.14 0.15 0.16 0.17 0.18 Growth Temperature (oC)resistivity (Ohm.cm)(a) 400500600700800 10181019 3 4 5 6 7 8 9 10 11 12 13 14 15 Mobility (cm /(V*S)) Carrier Density (cm-3)Growth Temperature (oC) Figure 8-7. Changes in Resistivity [opaque squares and solid line], Carrier density [triangles and solid line] and Mobility [open stars with dashed line], (a) and (b), respectively, as a function of Growth Temperature for the films doped with 0.1 at % Nb on sapphire and 30 mTorr. Data are ta bulated in Table 8.1.

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251 20406080100120 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 resistivity (Ohm.cm)Growth Pressure (mTorr) (a) 20406080100120140 2x10184x10186x10188x10181019 2 4 6 8 10 Mobility (cm2/(V*S))Carrier Density (cm-3)Growth Pressure (mTorr) (b) Figure 8-8. Resistivity [opaque s quares and solid line in (a)], Carri er density [triangles and solid line in (b)] and Mobility [ope n stars with dashed line (b)] as a function of deposition pressure for the films doped with 0.1 at % Nb on sapphire at 500C. Data are tabulated in Table 8.2.

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252 3.003.053.103.153.203.253.303.353.403.45 0.01 0.1 D1 oX 3.327 eV D0 TaXIntensity (a. u.)Energy (eV) ZnO:Nb ZnO:Ta Undoped ZnO 3.147 eV DNbAP 3.174 eV DTaAP 3.311 eV D0 NbX D1AP Undoped ZnO ZnO:Ta Figure 8-9. 20K photoluminescence spectra of. undoped ZnO film (dotted line); 0.1 at % Ta doped ZnO film and 0.1 at % Nb doped ZnO film. 102103104105 3040506070 ZnO(004) 2 Thetha ZnO (002) Al2O3 (006) 90 mTorr 1 mTorr Figure 8-10. Powder X-ray diffraction of the 1. 0 at % Ta doped ZnO grown on thin MgO buffer layer on sapphire.

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253 020406080100 0 2 4 6 8 10 12 Surface Roughness (nm)Growth Pressure (mTorr) Figure 8-11. Surface roughness as a function of growth pressure for the 1.0 at % Ta doped ZnO grown on thin MgO buffer layer on sapphire. 350400450500550600650 Intensity (a. u.)Wavelength (nm) 1 mTorr 30 mTorr 60 mTorr 90 mTorr 2.97 eV 2.84 eV 3.30 eV Figure 8-12. PL at 300K of selected films doped with 1 at % Ta on MgO buffer layer.

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254 0204060800 20 40 60 80 100 10161017101810191020 0.10 0.15 0.20 0.25 0.30 0.35 0.40 Growth Pressure (mTorr) Resistivity Mobility Carrier Density Figure 8-13. Resistivity (a), Hall coefficient (b), carrier concentration (c) and Hall Mobility (d) as function of growth pressure for the f ilms deposited on thin MgO buffer layer at 500C and doped with 1 at % Ta.

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255 020406080100 1 2 3 4 Resistivity (Ohm-cm)1000/Temperature (1000/K) Figure 8-14. Logarithmic plots of resistivity versus reciprocal te mperature for the film grown at 500C and 30 mTorr of oxygen on thin MgO on sapphire. 050100150200250300350 -2 -1 0 1 2 3 4 Single Band ConductionHall coefficient (c.cm-2)Temperature (K) p-type n-typeMixed Conduction Figure 8-15. Hall coefficient vers us temperature for the film gr own at 500C and 30 mTorr of oxygen on thin MgO on sapphire.

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256 0.0020.0030.0040.0050.0060.0070.0080.0090.0100.011 1E19 Slope = Ea/(2k) = 112 +/31.03 Ea = 19.4 +/5 meVC.C (cm-3)T-1 (K-1) (a) 234567891011 1E15 1E16 Ea/k = 0.22425; Ea = 19.3 +/3 meVN x T-3/2 (cm-3*K-3/2)1000/T (1000/K) Figure 8-16. Arrhenius plot of carrier concentra tion versus reciprocal temperature (from 350 to 100K) of carrier concentra tion (a), C.C., and Charge Balance Equation (b) [ ) / exp(2 3T k E T const nb a ] for the film grown at 500C and 30 mTorr of oxygen on thin MgO buffer/c-sapphire.

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257 050100150200250300350 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 501001502002503003500.0 0.2 0.4 0.6 0.8 Mobility (cm2/(V*s))Temperature (K) [100-350]K Mobility Linear fit to mobility data Figure 8-17. Carrier mobility as a function of temperature for the film grown at 500C and 30 mTorr of oxygen on thin MgO buffer/c-sapphire. Insert is the plot of mobility versus temperature for the single carrier range. 15.516.016.517.017.518.018.519.019.520.0 E1E2DBACIntensity (counts/s)Omega (o) F Figure 8-18. High resolution rock ing curves: (A) As grown; 500 C and 1mTorr of oxygen; (B) As grown; 500C and 90 mTorr Oxygen; (C) Annealed in air at 1000C for 5 min; (D) Annealed in air at 1000C for 30 min; (E) Annealed in air at 1000C for 60 min. The films were doped wit 1.0 at % Ta and grown at 500C on MgO buffer layer.

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258 20304050607080 102103104 Intensity 2thetha(Degree) 1000C, 60 min, air Figure 8-19. Powder XRD pattern of the ZnO:Ta0.001 film annealed in air at 1000C for 60 min. 3040506070 1 2 3 4 5 6 7 A n n e a l i n g co n d i t i o n s2 ThetaIntensity (count/s) Figure 8-20. Powder XRD of: (1) annealed in ai r at 900C for 5 min, (3) annealed in air at 1200C for 5 min; (5) annealed in air at 1000C for 60 min; (7) annealed in Ar at 1000C for 5 min. The ZnO:Ta0.001 films were grown at 500C on MgO buffer layer.

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259 350400450500550600650700 0 1 2 3 4 5 6 7 8 9 10 11 12 1100C 1200C 900CIntensity ( a. u.)Wavelength (nm)1000C Figure 8-21. PL of the ZnO:Ta0.001 films grown at 500C and 1 mTorr of O2 and annealed in air. 300350400450500550600650700750 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1200C 1000 C 900CIntensity (a. u.)Wavelength (nm) 1100C Figure 8-22. PL for the ZnO:Ta0.001 films grown at 500C and 90 mTorr of O2.

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260 020040060080010001200 100102104106 10 Resistivity (ohm.cm)Annealing Temperature (C) 1 mTorr of O2 90 mTorr of O2 Figure 8-23. Resistivity as a func tion of air annealing temperatur e for the films grown at 1 and 90 mTorr of oxygen. 0102030405060 102103104105 Resistivity (Ohm.cm)Annealing Time (min) Figure 8-24. Resistivities as a function of air 1000C annealing time for the ZnO:Ta0.001 films grown at 90 mTorr of O2.

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261 0200400600800100010-1101103105 10 7 Resistivity (Ohm.cm)Annealing Temperature (C) Figure 8-25. Resistivities as a function of oxygen anneali ng temperature for the ZnO:Ta0.001 films grown on MgO buffered sapphire at 30 mTorr of oxygen. 350400450500550600650700 0.04 0.06 0.08 0.10 0.12 0.14 0.16 0.18 900CIntensity (a. u.)Wavelength (nm) 450C Figure 8-26. PL spectra of the films annealed in flowing oxygen (1 atm) for 5 min. The ZnO:Ta0.001 films were deposited on MgO buffere d sapphire at 30 mTorr of oxygen.

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262 02004006008001000 0.1 1 Resistivity (ohm.cm)Annealing Temperature (C) c0 4 8 12 16 20 24 28 Mobility (cm2/(VS))b1018 Carrier Density (cm-3)a Figure 8-27. Transprot properties of the 5 min Ar annealed ZnO:Ta films grown in 30 mtorr od oxygen and on MgO buffer [carrier density (a), carrier mobility (b) and resistivity (c)].

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263 350400450500550600650700 Intensity (a. u.)Wavelength (nm) 900C 700C as grown Figure 8-28. PL spectra of the f ilms annealed in flowing argon (1 atm) for 5 min. The films were 1.0 at % Ta doped and deposited on MgO buffered sapphire at 30 mTorr of oxygen. 01020304050602 4 6 8 10 12 14 Surface Roughness (nm)Air Annealing Time (min)b9001000110012000 2 4 6 8 10 12 14 Air Annealing Temp (C)a Figure 8-29. Effects of annea ling temperature (a) and duratio n (b) on the films surfaces.

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264 350400450500550600650 Intensity (a. u.)Wavelength (nm)30 mTor r 400 mTor r Figure 8-30. Effects of cooling pr essure on the room temperature PL of 1.0 at % Ta doped films grown at 500C and 30 mTorr of oxygen. 02004006008001000102103104105 10 6 Resistivity (Ohm.cm)RTA Temperature Figure 8-31. Effects of 5 sec oxyge n rapid thermal annealing on th e electrical properties of the 1.0 at % Ta doped films grown at 500C and 60 mTorr of oxygen.

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2653.003.053.103.153.203.253.303.353.40 ZOA3 ZOA23.287 eV As5+?Ta5+Nb5+Intensity (a. u.)Ener gy ( eV ) D1AP DxAP D0 XX Do 1X Xa ZOA1 und. ZnO A0X Figure 8-32. Comparison between the low te mperature photoluminescence spectra of: undoped ZnO film(dotted line); 0.1 at % Nb an d Ta doped (dashed lines); As doped films (solid lines). ZOA1 is an 0.2 at % As doped film doped from As2O3; ZOA2 is 0.2 at % doped ZnO from Zn3As2 and ZOA3 is an 0.02 at % As doped ZnO film from Zn3As2.

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266 Figure 8-33. As grown film AFM image. Figure 8-34. AFM image for the film annealed at 1000C for 5 min. Figure 8-35. AFM image of the film annealed at 1000C for 30 min.

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267 Figure 8-36. AFM image of the film annealed at 1000C for 60 min.

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268 CHAPTER 9 FUTURE WORK In the preceding chapters I have demonstrated p-type conductivity in ZnO using arsenic as a dopant. Optimal growth temperature seems to be around 500C. Most of the films low temperature photoluminescence spectra showed a cceptor related exciton emission. The binding energy of the acceptor-bound exciton is found to be dependent on the As content of the films. The use of a buffer layer improved the quality of the films and only films grown on thin ZnO or MgO buffer layers showed p-type conductivity. One major draw back in our experiments is the lack of control of th e concentration of As. Since As is volatile, I was not able to accurately determined the concentration of As in the films using XPS nor WDS. This have pr evented us from calculating the ratio of activated As. I have also showed that I need greater co ntrol of the surface of the films. It is possible that As in near the surface region of the films di ffused out of the film during the slow cooling of the film causing an arsenic depleted region. A depleted surface region would make it more challenging to measure the transport properties of the films. A possible solution to this problem would be to attach an As source to the system. The external As source would help us maintain and control the concentration of As in the films during the de position process and duri ng cooling and heating. The next logical step from this work is the fabrication of device stru ctures. I have showed that films deposited at 500C and 5 mTorr of oxygen have the optimal transport and optical properties. Although our attempt to grow several devices failed, I be lieve those failure are due to a lack of control and unders tanding of the surface prope rties of the films.

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269 APPENDIX: CHARACTERIAZATI ON ON UNDOPED ZNO FILM Photoluminescence of Single Crystal ZnO To identify the peaks present in our films, I have performed photoluminescence on a single crystal ZnO and an undoped film grown at 700C and 3 mTorr of O3/O2. Figure A-1 is the photoluminescence spectrum at 300K of an undoped single crystal ZnO substrate. The spectrum is dominated by the near band edge emission peak at 376 nm (3.3 eV) and a broad peak between 425 and 700 nm centered around 520 nm. Point x in th e figure is an inst rument artifact and therefore is not real. The NBE in undoped films is due to the free exciton of the optical phonon replicas emission while the nature of peak cen tered around 520 nm is st ill disputed but is generally attributed to Zn in terstitials and oxygen vacancies. The photoluminescence at 25 K is represented in figure A-2. The spectrum is dominated by the donor-boun d exciton emission, DX, at 3.36 eV. Residual acceptors in the subs trates cause a DAP and its longitudinal optical phonon replica to be observed at 3.290 and 3.216 eV, respectively. The binding energy of the residual acceptor in the single cr ystal from our supplier is esti mated at 200 meV. B. K. Meyer el al163, 187, F. Tuomisto et al.190 and K. Thonke et al.138, 188 have observed similar DAP caused by residual acceptor with similar binding energy in single crystal ZnO. The nature of the residual acceptor was found to be dependent on the suppl ier and the crystal growth technique and conditions. NO and VZn are the most common causes. Photoluminescence of Undoped ZnO film Figure A-3 shows the room temperature photoluminescence spectrum of an undoped ZnO film grown at 700C and 3 mTorr of (3 mol % 03 and 97 mol % 02) gas mixture. The spectrum is dominated with the NBE emission due to free ex citon of the optical phonon replicas at 376 nm (3.3 eV). The high intensity of the NBE is a good indication of the crystallin e quality of the film. Emission in the visible is low as to compar e to the NBE emission. The low temperature

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270 photoluminescence of the undoped film is shown in figure A-4. DX emission is centered at 3.36 eV with a shoulder at 3.316 eV due to the l ongitudinal optical phonon of the DX and a DAP is at 3.256 eV. Unlike in the single crystal I do not observe the LODAP due to low crystalline quality of ZnO films grown on sapphire. The esti mated binding energy of the residual acceptor is 150 meV. The binding energy of the residual acceptor in our films indicates that the chemical nature of the contaminant in our film is diffe rent from the one causing the residual acceptor in the ZnO single crystal from our supplier. Since I deposit As and P doped films in our chamber, I suspect As or P contamination from the chamber walls during growth. Judging from the width of the peak at 3.316 and the size of the DAP, the p eak at 3.316 maybe a combination of the LODX and the contaminant related acceptor-bound emission. Hall measurements of Undoped ZnO films Figure A-5 shows the transport properties of a two undoped films grown at 500C in 30 and 60 mTorr of oxygen, respectively. Resistivity increased from 0.108 .cm for the film grown in 30 mTorr to 0.508 .cm for the film grown at 60 mTorr of oxygen. Carrier density and Hall mobility of the films decreased from 2.55 x 1018 cm-3 and 22.68 cm2/(V.s) to 9.78 x 1017 cm-3 and 19.1 cm2/(V.s) when growth pressure is raised from 30 to 60 mTorr of oxygen. The surface became rougher from around 6 nm to 10 nm as gr owth pressure was raised. For the film deposited in 30 mTorr of oxygen at 500 and 600C, figure A-6, Resistivity and Hall mobility of the films increased from 0.108 .cm and 22.68 cm2/(V.s) to 0.364 .cm and 33.41 cm2/(V.s) for the film grown at 600C. Carrier density of the films decreased from 2.55 x 1018 cm-3 to 5.13 x 1017 cm-3 as deposition temperature was raised from 500C to 600C. Those changes occur because of the reduction in Zni and Vo with increasing growth pressure and temperature. The changes in mobility reflects the negative effects of raising pressure on the lattice quality of the films and lattice improvements by raising growth temperature.

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271 350400450500550600650700 0.0 0.2 0.4 0.6 0.8 Intensity (a. u.)Wavelength (nm) 3.3 eV X due to grating Figure A-1. Photoluminescence spectrum at 300K of an undoped single crystal ZnO substrate. 3.153.203.253.303.353.403.4 5 0.1 1 10 I n t ens it y ( a. u. ) Energy (eV) undoped single crystal 3.360 eV DoX 3.216 eV LODAP 3.290 eV DAP Figure A-2. Photoluminescence spectrum at 25K of an undoped single crystal ZnO substrate.

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272 400500600700 0 1 2 3 4 5 Intensity (a. u.)Wavelength (nm) 3.3 eV Figure A-3. Photoluminescence spectrum at 300K of an undoped film grown at 700C and 3 mTorr of (3 mol % 03 and 97 mol % 02) gas mixture. 3.123.183.243.303.363.42 0.01 0.1 Intensity (a. u.)Energy (eV) 3.316 eV 1-LO 3.361 eV DoX Xa 3.256 eV DAP Figure A-4. Photoluminescence spectrum at 16K for an undoped film grown at 700C and 3 mTorr of (3 mol % 03 and 97 mol % 02) gas mixture.

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273 30405060 0 .1 0 .2 0 .3 0 .4 0 .5 10181.5x10182x10182.5x10183x1018 6 8 10 12 14 16 18 20 22 Growth Pressure ( mTorr ) Resistivity 500oC Surf. rough Mobility Carrier Density Figure A-5. Room temperatur e Hall properties of undoped ZnO films grown at 500C.

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274 500600 0.1 0.15 0.2 0.25 0.3 0.35 0.4 1016 22 24 26 28 30 32 34 Substrate Temperature (oC) Resistivity Mobility Carrier Density Figure A-6. Room temperature Hall properties of undoped ZnO films grown in 30 mTorr of oxygen.

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285 BIOGRAPHICAL SKETCH Jean-Marie G. Erie was born in New York on June 18, 1978 and raised in Haiti. He returned to the U.S. in 1996 where he completed his B.S. in chemical engineering at Florida International University. He was awarded Natio nal Science Foundation / South East Alliance for Graduate Education and the Prof essoriate Fellowship at the Un iversity of Florida where he earned his M.S. and Ph.D. in mate rials science and engineering. His graduate research focused on understanding of arsenic, nitrogen, niobium, ta ntalum and gallium doping in zinc oxide and magnesium zinc oxide thin films. He has partic ipated in a number of conferences, published several papers and received awards for his scie ntific work. Outside his research, Jean-Marie enjoys cooking, exercising, r eading and serving his community.