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Process Development of Contact and Functional Layers for Novel Semiconductors

Permanent Link: http://ufdc.ufl.edu/UFE0024336/00001

Material Information

Title: Process Development of Contact and Functional Layers for Novel Semiconductors
Physical Description: 1 online resource (128 p.)
Language: english
Creator: Wright, Jonathan
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2009

Subjects

Subjects / Keywords: contact, device, gan, hydrogen, semiconductor, sensor, zno
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: This study is focused on process development regarding Ohmic and Schottky contacts to n-ZnO and the application of functional layers to GaN and InN nanostructures for Hydrogen sensing. Goals of this work are three fold: first is to develop low resistance contacts for n-ZnO with higher thermal stability than typical metal stacks. Second, cryogenic temperatures have been used to deposit metal contacts to n-ZnO in order to increase barrier height at the interface for improved Schottky behavior. Finally, metal functionalization layers of Pt or Pd have been attempted on GaN nanowires and InN nanobelts for Hydrogen sensing. Ohmic contacts to n-ZnO were fabricated using a variety of robust, refractory materials including TiB2, CrB2, and Ir. Boride contacts were rectifying for lower anneal temperatures but transition to Ohmic behavior at higher temperatures (700masculine ordinalC) and exhibit minimum specific contact resistivity as low as ~5x10-4 ?. cm. Higher temperatures led to severe contact metallurgy intermixing and an increase in specific contact resistivity. Ir contacts exhibited high thermal stability and minimum specific contact resistance of 3.6x10-5 ?.cm2 after a 1000masculine ordinalC anneal. Next, the effect of cryogenic temperatures during deposition of Pd, Pt, Ti, Ni, and Au on n-ZnO was investigated. Deposition at both room and low temperature produced contacts with Ohmic characteristics for Ti and Ni metallizations. By sharp contrast, both Pd and Pt contacts showed rectifying characteristics after deposition with barrier heights between 0.37-0.69 eV. Pd contacts showed an increase in barrier height along with a decrease in ideality factor with increasing annealing temperature. Au deposited at room temperature produced contacts with Ohmic characteristics while cryogenic deposition produced rectifying characteristics. The differences in contact behavior were stable to anneal temperatures of ~300?C. Finally, Pd and Pt functionalization layers were deposited to GaN nanowires and InN nanobelts for Hydrogen sensing. Both uncoated nanomaterials show little or no current response upon exposure to hydrogen gas. The addition of a functional layer is shown to dramatically affect response to H2, allowing for hydrogen to be detected down to the hundreds of ppm level. Pd exhibits a greater response to hydrogen than Pt in both cases.
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Statement of Responsibility: by Jonathan Wright.
Thesis: Thesis (Ph.D.)--University of Florida, 2009.
Local: Adviser: Pearton, Stephen J.

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2009
System ID: UFE0024336:00001

Permanent Link: http://ufdc.ufl.edu/UFE0024336/00001

Material Information

Title: Process Development of Contact and Functional Layers for Novel Semiconductors
Physical Description: 1 online resource (128 p.)
Language: english
Creator: Wright, Jonathan
Publisher: University of Florida
Place of Publication: Gainesville, Fla.
Publication Date: 2009

Subjects

Subjects / Keywords: contact, device, gan, hydrogen, semiconductor, sensor, zno
Materials Science and Engineering -- Dissertations, Academic -- UF
Genre: Materials Science and Engineering thesis, Ph.D.
bibliography   ( marcgt )
theses   ( marcgt )
government publication (state, provincial, terriorial, dependent)   ( marcgt )
born-digital   ( sobekcm )
Electronic Thesis or Dissertation

Notes

Abstract: This study is focused on process development regarding Ohmic and Schottky contacts to n-ZnO and the application of functional layers to GaN and InN nanostructures for Hydrogen sensing. Goals of this work are three fold: first is to develop low resistance contacts for n-ZnO with higher thermal stability than typical metal stacks. Second, cryogenic temperatures have been used to deposit metal contacts to n-ZnO in order to increase barrier height at the interface for improved Schottky behavior. Finally, metal functionalization layers of Pt or Pd have been attempted on GaN nanowires and InN nanobelts for Hydrogen sensing. Ohmic contacts to n-ZnO were fabricated using a variety of robust, refractory materials including TiB2, CrB2, and Ir. Boride contacts were rectifying for lower anneal temperatures but transition to Ohmic behavior at higher temperatures (700masculine ordinalC) and exhibit minimum specific contact resistivity as low as ~5x10-4 ?. cm. Higher temperatures led to severe contact metallurgy intermixing and an increase in specific contact resistivity. Ir contacts exhibited high thermal stability and minimum specific contact resistance of 3.6x10-5 ?.cm2 after a 1000masculine ordinalC anneal. Next, the effect of cryogenic temperatures during deposition of Pd, Pt, Ti, Ni, and Au on n-ZnO was investigated. Deposition at both room and low temperature produced contacts with Ohmic characteristics for Ti and Ni metallizations. By sharp contrast, both Pd and Pt contacts showed rectifying characteristics after deposition with barrier heights between 0.37-0.69 eV. Pd contacts showed an increase in barrier height along with a decrease in ideality factor with increasing annealing temperature. Au deposited at room temperature produced contacts with Ohmic characteristics while cryogenic deposition produced rectifying characteristics. The differences in contact behavior were stable to anneal temperatures of ~300?C. Finally, Pd and Pt functionalization layers were deposited to GaN nanowires and InN nanobelts for Hydrogen sensing. Both uncoated nanomaterials show little or no current response upon exposure to hydrogen gas. The addition of a functional layer is shown to dramatically affect response to H2, allowing for hydrogen to be detected down to the hundreds of ppm level. Pd exhibits a greater response to hydrogen than Pt in both cases.
General Note: In the series University of Florida Digital Collections.
General Note: Includes vita.
Bibliography: Includes bibliographical references.
Source of Description: Description based on online resource; title from PDF title page.
Source of Description: This bibliographic record is available under the Creative Commons CC0 public domain dedication. The University of Florida Libraries, as creator of this bibliographic record, has waived all rights to it worldwide under copyright law, including all related and neighboring rights, to the extent allowed by law.
Statement of Responsibility: by Jonathan Wright.
Thesis: Thesis (Ph.D.)--University of Florida, 2009.
Local: Adviser: Pearton, Stephen J.

Record Information

Source Institution: UFRGP
Rights Management: Applicable rights reserved.
Classification: lcc - LD1780 2009
System ID: UFE0024336:00001


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1 PROCESS DEVELOPMENT OF CONTACT AND FUNCTIONAL LAYERS FOR NOVEL SEMICONDUCTORS By JONATHAN WRIGHT A DISSERTATION PRESENTED TO THE GRADUATE SCHOOL OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 2009

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2 2009 Jonathan Wright

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3 To my family

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4 ACKNOWLEDGMENTS The research in this dissertation comes largely from receiving great support. I am deeply indebted to my advisor, Prof. Stephen J. Pearton, more than anyone else for the opportunities afforded to me throughout my studies, many of which appear in this work. His guidance, motivation and technical skill have helped transcend my abilities as a scientist and engineer. I would also like to thank my other supervisory committee members, Prof. Cammy R. Abernathy, Prof. David P. Norton, Prof. Rajiv Singh, and Prof. Fan Ren. I especially want to thank Prof. Ren for his time and assistance on nume rous aspects of my research. I thank the members of the research groups of Prof. Pearton, Prof. Ren, and Prof. Abernathy for their aid and technical support, including Dr. Brent Gila, Dr. Luc Stafford, Dr. Lars Voss, Dr. Hung t a Wang, Dr. Sam Kang, Dr. John Chen, Dr. Travis Anderson, Dr. Soohwan Jang, Dr. Mark Hlad, Wantae Lim, Nimo Chen, B yunghw an Chu, Yu lin Wang, Andrew Gerger and so many others who have made such an enriching impact on my experience at Florida. I also woul d like to thank Ivan Kravchenko, Bill Lewis and Al Ogden for their support in my work at the UF Nanofabrication Facility. My gratitude goes to Prof. Ant Ural from the Electrical and Computer Engineering Department. I thank his research group, particularly Jason Johnson, who provided ma ny of the nanostructured samples for this research. Additional t hanks go to Prof. Tim Anderson for providing me with my first graduate internship at NASA Glenn Research Center in Cleveland. My utmost thanks to my mentors there, Gus Fralick and Joh n Wrbane k, and all the other many people I had the privilege of working with. Thanks also go to the numerous people at Sandia National Laboratories with whom I had the opportunity to work in Albuquerque with on my second graduate internship I particularly thank my mentor Dr. Randy J. Shul, and my manager, Dale Hetherington, as well as

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5 all of the people I shared the clean room with including Dr. Jeff Stevens Sarah Rich, Tom Plut, Pat Archer, Carlos Sanchez, Mark Overberg and many others. Finally, I wish to expr ess my deep appreciation for the support from my family, my friends and my wife, Amanda. It is with their help, I have reached the end of this journey.

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6 TABLE OF CONTENTS page ACKNOWLEDGMENTS .................................................................................................................... 4 LIST OF TABLES ................................................................................................................................ 8 LIST OF FIGURES .............................................................................................................................. 9 ABSTRACT ........................................................................................................................................ 12 CHAPTER 1 INTRODUCTION ....................................................................................................................... 14 2 BACKGROUND ......................................................................................................................... 18 2.1 ZnO Properties .................................................................................................................. 18 2.2 Nitride Semiconductor Properties .................................................................................... 20 2.2.1 GaN Properties ...................................................................................................... 20 2.2.2 InN Properties ....................................................................................................... 21 2.3 Electrical Contacts ............................................................................................................ 22 2.3.1 Ohmic Contacts ..................................................................................................... 23 2.3.2 Schottky Contacts ................................................................................................. 25 2.4 Characterization Techniques ............................................................................................ 28 2.4.1 Current -Voltage Measurements ........................................................................... 28 2.4.2 Auger Electron Spectroscopy............................................................................... 29 2.4.3 X Ray Diffraction ................................................................................................. 30 2.4.4 Photoluminescence ............................................................................................... 30 3 THERMALLY STABLE OHMIC CONTACTS TO n -ZnO ................................................... 32 3.1 Ohmic Co ntacts ................................................................................................................. 32 3.2 Surface Treatment Investigation ...................................................................................... 33 3.3 Fabrication of Ohmic Contacts ......................................................................................... 35 3.3.1 Boride based Contact Deposition ........................................................................ 35 3.3.2 Iridium -based Contact Deposition ....................................................................... 36 3.4 TiB2-based Contact Study ................................................................................................. 36 3.5 ZrB2-based Contact Study ................................................................................................ 38 3.6 Iridium -based Contact Study ............................................................................................ 40 3.6.1 Nitrogen Annea ling of Contacts .......................................................................... 40 3.6.2 Oxygen Annealing of Contacts ............................................................................ 42 3.6.3 Iridium -based Contact Study Summary .............................................................. 43 3.7 Conclusions ....................................................................................................................... 44

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7 4 EFFECT OF CRYOGENIC DEPOSITION ON SCHOTTKY METAL CONTACTS TO n ZnO .................................................................................................................................... 63 4.1 Introduction ....................................................................................................................... 63 4.2 Experimental Details ......................................................................................................... 64 4.3 Results and Discussion ..................................................................................................... 65 4.3.1 Results and Discus sion Ti, Ni, Pt, Pd Contacts ............................................... 65 4.3.2 Results and Discussion Au Contacts ................................................................ 68 4.4 Conclusions ....................................................................................................................... 70 5 FUNCTIONALIZATION OF NANOMATERIAL DEVICES FOR H2 SENSING .............. 89 5.1 Introduction ....................................................................................................................... 89 5.2 InN Nanobelts .................................................................................................................... 90 5.2.1 Growth Process ..................................................................................................... 90 5.2.2 Pt -functionalization .............................................................................................. 90 5.2.3 Pd -functionalization.............................................................................................. 92 5.3 GaN Nanowires ................................................................................................................. 94 5.3.1 Growth Process ..................................................................................................... 94 5.3.2 Pd -functionalization .............................................................................................. 95 5.3.3 Pt -functionalization .............................................................................................. 96 5.4 Conclusions ....................................................................................................................... 99 6 CONCLUSION ......................................................................................................................... 115 LIST OF REFERENCES ................................................................................................................. 118 BIOGRAPHICAL SKETCH ........................................................................................................... 128

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8 LIST OF TABLES Table page 2 1 Bulk ZnO Material Properties. .............................................................................................. 18 2 2 Electrical Properties of bulk GaN and ZnO. ......................................................................... 21 3 1 Concentration of elements detected on the as received surface of TiB2/Pt/Au contacts to n ZnO (in atom%) .............................................................................................................. 45 3 2 Concentration of elements detected on the as received surface of ZrB2/Pt/Au contacts to n -ZnO (in atom%) ............................................................................................... 46 4 1 Summary of barrier height and ideality factors for contacts on ZnO. ................................ 72

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9 LIST OF FIGURES Figure page 2 1 Wurtzite crystal structure of ZnO .......................................................................................... 20 2 2 Typical schematic for linear TLM ........................................................................................ 28 2 3 Definition of resistances for typical linear TLM .................................................................. 29 3 1 Room temperature PL spectra for ZnO before and after various surface treatments. ....... 47 3 2 Specific contact resistivity of TiB2/Pt/Au contacts and ZnO sheet resistance as a function of annealing temperature. ....................................................................................... 48 3 3 Optical microscopy images of TiB2/Pt/Au contacts on ZnO. ............................................. 49 3 4 AES depth profiles from TiB2/Pt/Au contacts on ZnO. ...................................................... 50 3 5 Specific contact resistivity of ZrB2/Pt/Au contacts and ZnO sheet resistance as a function of annealing temperature. ....................................................................................... 51 3 6 SEM images of ZrB2/Pt/Au contacts at varying annealing temperatures ......................... 52 3 7 AES surface scans of ZrB2/Pt/Au contacts at varying annealing temperatures ............... 53 3 8 AES depth profiles of ZrB2/Pt/Au c ontacts at varying annealing temperatures ............... 54 3 9 Specific contact resistivity of Ir/Au contacts and ZnO sheet resista nce as a function of annealing temperature. ...................................................................................................... 55 3 10 SEM images of Ir/Au contacts at varying annealing temperatures .................................... 56 3 11 AES depth profiles of Ir/Au contacts at varying annealing temperatures ........................ 57 3 12 Specific contact resistivity and sheet resistance of the contacts annealed at 700C (N2 ambient) as a function of aging time at 350C. .................................................................... 58 3 13 Optical micrographs of the 700C annealed (N2 ambient) Ir/Au contacts after aging .... 59 3 14 Resistivity comparison of Ir/Au contacts between O2 and N2 anneal. ............................... 60 3 15 Optical micrographs of Ir/Au contacts after 700C anneal in O2 and N2 ambience ........ 61 3 16 AES depth profiles of Ir/Au contacts after 700C anneal in O2 and N2 ambience .......... 62 4 1 Optical microgr aph of Ti contacts deposited at either 300K. ............................................. 73

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10 4 2 I-V characteristics of room temperature and cryogenically deposited diodes on linear and log plots. ......................................................................................................................... 74 4 3 I-V characteristics of Ohmic contacts deposited at either 300K or 77K. ........................... 75 4 4 I-V characteristics of Schottky diodes deposited at either 300K o r 77K on linear and log plots. ................................................................................................................................. 76 4 5 Schottky barri er height as a function of annealing temperature for diodes with Pd contacts deposited at either 77 or 300K. ............................................................................... 77 4 6 Reverse leakage current (@ 0.5V) as a function of annealing temperature for diodes with Pd contacts deposited at either 77 or 300K. ................................................................. 78 4 7 AES depth profiles from Ni on ZnO deposited at either 300K or 77K ............................ 79 4 8 AES depth profiles of Pt contacts on ZnO deposited at either 300K or 77K after annealing. ................................................................................................................................ 80 4 9 Optical microscope images of Au contacts on ZnO de posited at either 300K or 77K ..... 81 4 10 I-V characteristics of Au/n GaN Schottky diodes deposited at either 77K or 300K. ....... 82 4 11 I-V characteristics from the 300K deposited samples as a function of post -dep osition annealing temperature. ........................................................................................................... 83 4 12 I-V characteristics from the 77K deposited samples as a function of post -deposition annealing temp erature. .......................................................................................................... 84 4 13 Comparison of the I -V characteristics from both the 300K and 77K deposited sam ples after annealing at 300 ......................................................................................... 85 4 14 Surface characterization of Au contacts on ZnO deposited at 300K .................................. 86 4 15 Surface characterization of Au contacts on ZnO deposited at 7 7K and subsequently annealed. ................................................................................................................................ 87 4 16 AES depth profiles form Au deposited on ZnO at either 300K or 77K and annealed. .... 88 5 1 X ray diffraction spectrum of MOCVD grown InN nanobelts (the inset shows SEM images of the nanobelts). ..................................................................................................... 100 5 2 Current resp onses of Pt -coated InN nanobelt sensors. ....................................................... 101 5 3 Relat ive response of Pt -coated InN nanobelts exposed to a series of H2 concentrations (20300 ppm) in N2 ambient for 10 min at room temperature. ................ 102 5 4 Arrhenius plot of rate of resistance change (the inset shows the temperature dependence of resistance from Pt -coated InN nanobelts exposed to 200 ppm H2). ........ 103

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11 5 5 Current responses of Pd -coated InN nanobelt sensors. ...................................................... 104 5 6 Relative response of Pd-co ated InN nanobelts exposed to 300 ppm H2 after 10 minutes at different temperature. ........................................................................................ 105 5 7 Arrhenius plot of rate of resistance change (the inset shows the temperature dependence of resistanc e from Pd -coated InN nanobelts exposed to 300 ppm H2). ....... 106 5 8 SEM images of as -grown GaN nanowires. ........................................................................ 107 5 9 Measured resistance at 0.5 V bias as a function of time from multiple GaN nanowires exposed to varying H2 concentrations in N2 ambient at room temperature. ..................... 108 5 10 Response of Pd-coated GaN nanowires to varying H2 concentratio n. .............................. 109 5 11 Temperature dependence of resistance from Pd-coated multiple GaN nanowires exposed to 3000 ppm H2 in N2. ........................................................................................... 110 5 12 SEM images of GaN nanowires after Pt coating ................................................................ 111 5 13 Res ponse of Pt -coated GaN nanowires to varying H2 concentration. .............................. 112 5 14 Relative response of Pt -coated GaN nanowires exposed to 2000 ppm H2 after 10 minutes at different temperature. ........................................................................................ 113 5 15. Arrhenius plot of rate of resistance change (the inset shows the temperature dependence of resistance from Pt -coated GaN nanowires exposed to 1000 ppm H2). .... 114

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12 Abstract of Dissertation Presented to the Graduate School of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy PROCESS DEVELOPMENT OF CONTACT AND FUNCTIONAL LAYERS FOR NOVEL SEMICONDUCTORS By Jonathan Wright M ay 2009 Chair: Stephen J. Pearton Major: Materials Science and Engineering This study is focused on process development regarding Ohm ic and Schottky contacts to n ZnO and the application of functional layers to GaN and InN nanostructures for Hydrogen sensing. Goals of this work are three fold: first is to develop low resistance contacts for n -ZnO with higher thermal stability than typi cal metal stacks. Second, cryogenic temperatures have been used to deposit metal contacts to n -ZnO in order to increase barrier height at the interface for improved Schottky behavior. Finally, metal functionalization layers of Pt or Pd have been attempte d on GaN nanowires and InN nanobelts for Hydrogen sensing. Ohmic contacts to n -ZnO were fabricated using a variety of robust, refractory materials including TiB2, CrB2, and Ir. Boride contacts were rectifying for lower anneal temperatures but transition t o Ohmic behavior at higher temperatures (700C) and exhibit minimum specific contact resistivity as low as ~5x104 cm. Higher temperatures le d to severe contact metallurgy intermixing and an increase in specific contact resistivity. Ir contacts exhibit ed high thermal stability and minimum specific contact resistance of 3.6x105 2 after a 1000C anneal. Next, the effect of cryogenic temperatures during deposition of Pd, Pt, Ti, Ni, and Au on n ZnO was investigated. Deposition at both room and low tem perature produced contacts with Ohmic characteristics for Ti and Ni metallizations. By sharp contrast, both Pd and Pt contacts

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13 showed rectifying characteristics after deposition with barrier heights between 0.37 0.69 eV. Pd contacts showed an increase in b arrier height along with a decrease in ideality factor with increasing annealing temperature. Au deposited at room temperature produced contacts with Ohmic characteristics while cryogenic deposition produced rectifying characteristics. The differences in c ontact behavior were stable to anneal temperatures of ~300 Finally, Pd and Pt functionalization layers were deposited to GaN nanowires and InN nanobe lts for Hydrogen sensing. Both uncoated nanomaterials show little or no current response upon exposure to hydrogen gas. The addition of a functional layer is shown to dramatically affect response to H2, allowing for hydrogen to be detected down to the hundreds of ppm level. Pd exhibits a greater response to hydrogen than Pt in both cases.

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14 CHAPTER 1 INT RODUCTION The global sensor market has grown 4.5% annually for the last 8 years. This business sector should reach a value of US $61.4 billion by 2010 [1]. Chemical sensing represents the largest demographic accounting for at least US $11.5 billion each year. Th is specific market corresponds to species -specific elemental/ molecular detection in gas and liquid, flue gas and fire detection, and biosensors. While wet sensing is still a major component of total chemical sensors produced, semiconductor -based sensors are becoming increasingly desirable due to their highly controllable electronic character and increasingly inexpensive cost. The enormous knowledge base and relative ease of production gives Silicon -based devices in particular the largest market share of semiconductor chemical sensors. Silicon however, is not viable in harsh environments, does not stand up to high temperature or pressure, and is susceptible to corrosion. This greatly limits the usability in many new applications including gas se nsors in autos and spacecraft and chemical sensing in aggressive industrial areas. Wide bandgap semiconductors overcome the obstacles present with Silicon technology. These materials, such as Zinc Oxide (ZnO) and Gallium Nitride (GaN), have a larger bandg ap, higher electron mobility and higher breakdown field strength [2, 3]. They are chemically resistant and easily withstand high temperature use, making them great candidates for high power, high -temperature electronic devices and optoelectronics. Interes t in ZnO is derived from both its direct, wide bandgap (3.2 eV) and its large free exciton binding energy (~60 meV) [4, 5]. The lower exciton binding energy of other wide bandgap semiconductors such as GaN (~24 meV) leads to greater potential for excitons to dissociate due to heat or exciton scattering, resulting in the relative inability of these materials to be used in high temperature applications [6 9]. This is in contrast to ZnO where efficient

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15 excitonic emission processes can exist at room tempera ture and above, allowing ZnO -based LEDs to operate at temperatures much higher than GaN emitters. In addition, commercial -scale production of ZnO is far cheaper than for GaN. Thin film ZnO can be easily grown on most substrates, including glass, and a c ommercial ZnO substrate is readily available [10 11]. The use of ZnO has already been exhibited in transparent electrodes, piezoelectric transducers and some chemical sensing applications. There is also promise ZnO LEDs could be combined with phosphors to produce solid-state white lighting [7,8, 10 15]. This variety of new and emerging uses of ZnO is intended as a competitor for semiconductor applications that are complementary to currently used materials such as GaN. However in order to fully reach the potential of ZnO -based electronics, considerable improvements to both device processing and material quality are needed. Increasing demand for semiconductors in high temperature applications demands the development of reliable, thermally stable Ohmic cont acts [16 29]. These contacts must withstand extensive cyclic use at high temperature without device degradation or dramatic changes to contact structure. As high -purity ZnO becomes less expensive superior device characteristics will be limited less by m aterial quality and more by contact failure. Also critical to the improvement of ZnO based devices is the need for effective Schottky contacts. Typical metals used for Schottky contacts to ZnO often show extremely poor thermal stability and inferior Schot tky character [30 38]. Increasing the Schottky barrier height at the interface may alleviate these undesirable characteristics, although the mechanisms of increasing Schottky barrier height are still not fully understood. One potential method of engineer ing Schottky barrier heights is by cryogenic metal deposition at 77K. This technique has already been demonstrated for GaAs, InP, and InGaAs [39 41].

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16 The largest of roadblocks toward the full realization of ZnO devices is the lack of effective p type dopi ng. Efforts to make p -ZnO have involved several deposition methods including ion implantation and involved a number of different dopants; nevertheless these techniques have been met with limited success [42 45]. With both nand p -type material available however, semiconductor nitride (such as GaN/AlGaN ) based electronics are already in commercial production. Currently fabricated devices include high performance HEMTs, heterojunction bipolar transistors (HBTs) and metal oxide field effect transistors (MO SFETs). Especially of interest for potential sensing applications is the increasing ease of growth of GaN and InN 1 D nanostructures (nanowires, nanorods, nanobelts). The high surface to volume ratio, single crystalline structure and quantum effects of 1 D nanostructures all work to enhance the sensitivity of these materials in regards to environmental change [46 60]. Additionally, nanostructured sensors have very low power demands, minimal weight, and may often work at temperatures far below room tempera ture. Typically however, the 1 D nanostructure alone is not enough for sensor work. A catalytically active functional layer is necessary for most chemically sensitive nanodevices. The use of metal functionalization layers have been shown to be effective in detection of hydrogen at room temperature for carbon nanotubes (CNTs), ZnO, GaN SnO2, and In2O3 [61 64]. While nitride semiconductors are promising candidates for sensing, reports for functionalization of GaN or InN for use as chemical sensors are lim ited. The purpose of this work is two fold. First is to develop improved contact schemes, both Ohmic and Schottky, for ZnO electronic and optoelectronic devices. Ohmic contacts must be thermally stable and exhibit low specific contact resistance. Schottky contacts must have improved barrier heights over previous studies. Chapter 2 presents background regarding basic properties of ZnO, GaN and InN as well as an overview of the characterization techniques used

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17 throughout this work. The rest of this thesi s concerns specific experiments and analysis in improved contacts intended for sensor applications First, thermally stable Ohmic contacts to n ZnO using novel high temperature materials are presented. Next, the effect of cryogenic deposition on metal co ntacts to n ZnO is explored. Finally, GaN nanowires or InN nanobelts were used to fabricate Hydrogen gas sensors comparing Pd and Pt functional layers for each.

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18 CHAPTER 2 BACKGROUND 2.1 ZnO Properties As production of ZnO thin films and substrates hav e become increasingly inexpensive, research on ZnO semiconductor devices has greatly increased. Interest in ZnO is derived from both its direct, wide bandgap (3.2 eV) [64, 65] and its large free exciton binding energy (~60 meV) [67], making ZnO a potentia l material for blue/UV optoelectronics and light -emitting diodes (LEDs). ZnO LEDs can also be combined with phosphors to produce solid -state white lighting. ZnO has also been used in transparent electrodes, piezoelectric transducers and sensor applicatio ns [15, 68 73]. Basic material parameters of ZnO are given in Table 2 1. Table 2 1 Bulk ZnO m aterial p roperties. Property Value Lattice parameters at 300 K (nm) a 0 : 0.32495 c 0 : 0.52069 Density (g cm 3 ) 5.606 Stable phase at 300 K Wurtzite Melting point (C) 1975 Thermal conductivity 0.6, 1 1.2 Linear thermal expansion coefficient a 0 : 6.5 10 6 c 0 : 3.0 10 6 Static dielectric constant 8.656 Energy bandgap (eV) Direct, 3.37 Intrinsic carrier concentration (cm3) <10 6 max n -type doping: n ~ 1020 max p type doping: p ~ 10 17 Exciton binding energy (meV) 60 Electron effective mass 0.24 Electron Hall mobility, n -type at 300 K (cm 2 V 1 s 1 ) 200 Hole effective mass 0.59 Hole Hall mobility, p type at 300 K (cm2V1s1) 5 50

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19 The emerging use of ZnO is intended for use in semiconductor applications that are complementary to currently used materials such as GaN. The lower exciton binding energy of GaN (~24 meV) leads to greater potential for excitons to dissociate due to heat or exciton scattering resulting in the inability of GaN light emitters to be used in high -temperature applications [74]. This is in contrast to ZnO where efficient excitonic emission processes can exist at room temperature and above, allowing ZnO based LEDs to operate at tem peratures much higher than GaN emitters. The direct bandgap energy of ZnO is comparable to GaN and is transparent to visible light, with an operation range in the UV to blue part of the spectrum Room temperature Hall measurements give ZnO a Hall mobilit y of ~200 cm2/V -s, which is lower than that for GaN [75]. The saturation velocity of ZnO however is higher. In addition, commercial -scale production of ZnO is far cheaper than for GaN. Thin film ZnO can be easily grown on most substrates, including gla ss, and a commercial ZnO substrate is readily available. Although research for ZnO UV optoelectronic and LED applications has dramatically grown in popularity, there remain a number of roadblocks to improve device performance. The largest of these obstacl es is the lack of effective p type doping. Efforts to make p -ZnO have involved several deposition methods and a number of different dopants; however these techniques have been met with limited success. P -type carrier surface doping by ion implantation us ing N, P, and As dopants has sparked interest toward the realization of p ZnO. Hole concentrations have been achieved in the range 10151017 cm3. Crystal Structure and Conductivity of ZnO : At ambient conditions, thermodynamically stable ZnO has a hexagonal, wurtzite structure with lattice parameters a = 3.25 and c = 5.12 [76]. Zn atoms are tetrahedrally coordinated to four O atoms with alternating Zn -only and O only layers. The wurtzite structure of ZnO is shown in Figure 2 1.

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20 ZnO also exhibits bot h zinc -blende and rocksalt (NaCl) structures, however both alternate structures are uncommon [76]. Zinc -blende ZnO is stable only when grown on cubic substrates while rocksalt -structured ZnO is created at extremely high pressures. Figure 2 1. Wurtzite c rystal s tructure of ZnO Undoped ZnO is usually inherently n-type; intrinsic donors have been associated to Zn interstitials, O vacancies and hydrogen impurities. Asymmetric doping limitations of ZnO often result in the achievement of n -ZnO even after p typ e doping [7782]. While ntype conductivity for ZnO has been highly developed through the use of excess Zn or by Al, Ga or In doping, the goal of reproducible p -type ZnO remains elusive. 2.2 Nitride Semiconductor Properties Both IV -oxide and III nitride s emiconductors are excellent alternatives to traditional silicon for electronic devices due to numerous advantages such as improved chemical resistance, bandgap, and potential for high -temperature use. However, III nitride based semiconductor devices have many properties unique and preferred to ZnO. A summary of important electronic parameters for GaN as compared to ZnO are shown in Table 2 2 [66, 83, 84]. 2.2.1 GaN Properties Due to the relative ease of growth of both nand p type material, GaN technolog ies are much more developed than for ZnO. The wide bandgap of 3.475 eV is very close to that for

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21 ZnO (3.2 eV), making GaN a viable material for ultraviolet optoelectronics. Additionally, the use of In and Al in InGaN and AlGaN allows for tunable emission over a wide spectral range (0.7 eV to 6.2 eV) [8598]. Although the excitonic binding energy is lower than for ZnO, the wide bandgap of GaN produces a low intrinsic carrier concentration. While intrinsic concentrations of 1015 cm3 are reached at 300 for Si and 500 t his concentration until 1000 -stability gives GaN potential as a possible material for high temp applications when alloyed with materials having higher excitonic binding energy Table 2 2 Elec trical p roperties of bulk GaN and ZnO. Property GaN ZnO Direct bandgap energy (eV) 3.4 3.4 Electron mobility (cm2/Vs) 220 200 Hole mobility (cm 2 /Vs) 10 5.50 Electron effective mass 0.27 m0 0.24 m0 Hole effective mass 0.80 m 0 0.59 m 0 Exciton binding e nergy (meV) 28 60 Typically, the structural differences of GaN to its commonly used growth substrates result in significant lattice mismatching and can be detrimental to the electronic properties of the GaN. Sapphire and SiC for example, result in 13 and 3 % lattice mismatch, respectively. In contrast to GaN thin films, freestanding 1 D nanostructures are easier to grow in single -crystal forms without defects. These nanostructures are becoming promising replacement candidates for many potential applications involving thin -film GaN. 2.2.2 InN Properties Out of all the III -nitride semiconductors, InN is the least technologically developed. Recently, InN has attracted attention from the revision of its fundamental bandgap from the visible (1.8 2.1 eV) to i nfrared (0.7 0.8 eV) spectral range [99103]. This bandgap value is lower than for any other major nitride -based semiconductor and allows for a wider possibility of

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22 optoelectronic applications over large spectral ranges by alloying together GaN, InN, and AlN. Much of the reason for the change in property value has stemmed from the developing ease of producing high quality, single crystal material [104 109]. Provided the parameters of effective growth can be controlled, InN is an excellent potential materi al for low -cost, low -power, highly sensitive detectors due to its intrinsic surface charge accumulation. 2.3 Electrical Contacts Electrical contacts connect devices to an outside electrical input. They are the junction between the semiconductor and metal/ contact material and also include any specific metal/material layers above the interface itself (the stack). Electrical propert ies of the contact are material -controlled both by the interface and from the contents of the stack. Desired properties are gen erally achieved by annealing at elevated temperature allowing for the removal of structural defects and charge compensators like Hydrogen. Annealing also may create intermetallurgical phases of lower resistance. The choice of material within the stack may also be used to manipulate performance of the contact. A Gold overlayer for example, may be used for current spreading over the top of the contact to allow for even conduction over the whole area at the interface. Gold will also prevent room temperature oxidation of the contact. In -soluble materials such as Platinum are o ften put into the middl e of a contact stack to act as diffusion barriers, preventing undesirable metallurgical intermixing throughout the contact stack. Diffusion barriers also prevent outdiffusion of interfacial materials and semiconductor from reaching the surface of the contact. Most often, the choice of interface and stack materials comes from the need to achieve the lowest specific contact resistance during extended use, however specific applications may alter the desired contact properties. Light emitting diodes need transparent contacts or materials

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23 with high reflectivity in the range of their emission wavelength for maximum light emission. High temperature applications require thermally stable, minimally reactive contact materials. There are two types of contacts, both of which are critical toward the fabrication of a semiconductor device. It is necessary to understand the material properties of both. 2.3.1 Ohmic Contacts The d evelopment of connections with negligible contact resistance relative to the bulk or spreading resistance of a semiconductor for a device is essential. Ohmic contacts exhibit a relatively small voltage drop with current application as compared to drop ove r the active region of the device. These contacts are named Ohmic because of their current -voltage response that obeys Ohms Law (V=IR). In theory, these contacts should exhibit a linear relationship between current and applied voltage without signal dis tortion and consume nothing in power because of complete current transference. Most important to the viability of Ohmic contacts with any device is a low specific contact resistance. Specific contact resistance is an amalgam term which describes the resis tance at directly above, below and at (or through) the interface. The specific contact resistance c) for Ohmic contacts is given as: 1 0 V CV J (1 .1) where Low specific contact resistance reduces the required power of the contact and di minishes internal heating improving contact and device lifetime.

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24 Conduction across the metal -semiconductor interface for semiconductors with low to moderate doping occurs through thermionic -emission over the potential barrier. The specific contact resista nce for this conduction follows: kT q T qA kb C exp* (1 .2) where q is the electronic charge, A** is the Richardsons constant, T is t b is the barrier height, and k is the Boltzmanns constant. Apparent from this equation, metals with a low barrier height are desired for contacts with small specific contact resistance. In the case of p type conduction, metals with high er work functions than that of the semiconductor are preferred. There are few metals however, with higher work functions than ZnO. This complicates the creation of a p type conducting Ohmic contact. Usual approaches to reduce barrier height over the ZnO include surface cleaning or through the creation of a highly doped region near the surface. Increasing the carrier concentration at the surface can be done by ion implantation or through alloying the contacts. By highly doping the substrate underneath the contact, the Schottky barrier at the surface is reduced allowing enhancement of current transport by tunneling. The equation for contact resistance taking into account conduction by electron tunneling is : D b S CN m *2 exp (1 .3) where ND r efers to the doping concentration of the (n-type) semiconductor

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25 Ohmic contacts to n ZnO : A variety of metallization schemes for Ohmic contacts to n ZnO have been demonstrated. Pt Ga contacts were reported to have a specific contact resistance of 3 x 104 cm2 [15, 27]. Ti/Au Ohmic contacts on Al -doped epitaxial layers were found to c values of ~ 104 105 value of 9.0 x 107 cm2 was reported for n+ ZnO (n~1.7 x 1018 cm 3) with Ti/ Al contacts after 300C anneal [111]. The most common metal schemes for Ohmic contacts on n-type ZnO involve Ti/Au, Zn/Au, Al/Pt ,Re/Ti/Au, Ru and Pt/Ga, with typical specific contact resistances in the range 10 3 103) n type ZnO after annealing at < 500C. Minimum contact resistance is usually obtained after post depositing annealing at temperatures between 200 C and 300 C due to the increase in near -surface carrier concentration reached by annealing. Metals with high melting temperature and low reactivity are becoming promising candidates for n -ZnO devices due to the increased need for high, long term thermal stability of contacts. 2.3.2 Schottky Contacts By contrast to the ease of conduction from metal to semiconduc tor and vice versa over all voltages found in Ohmic contacts, Schottky contacts act as a switch, allowing no current to flow before reaching a critical voltage both for the positive and negative direction. A large enough forward bias results in conducti on while a large negative bias results in contact breakdown. During breakdown, a very large negative bias will permanently alter the conduction path and the Schottky contact can be destroyed. Schottky contacts are formed by lining up the Fermi levels of t he metal and semiconductor. For Fermi level align ment, a space charge or depletion region is formed at the interface and a current flow barrier is made. In order to assure a barrier is created in the absence

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26 of surface states, which would act as pathways for easy current flow. The barrier height of an n type semiconductor q bn follows as: m bnq q (1 .4) where q m is the metal work function and q is the electron affinity of the semiconductor. In practice, some amount of surface st ates will be present on the semiconductor surface. By the addition of surface states, the density of semiconductor at the surface is large enough to compensate for additional surface charges without having to move the Fermi energy level. In turn by not a ltering the Fermi level, the space charge region at the interface will not be affected, making conduction easier through the interface with lower forward bias. Thus, the barrier height for all semiconductors may be described as: 0 m bS q q (1 .5) where m is the metal electronegativity and 0 represents the contribution of surface states of bm it is dependent on the electronegativity difference between cation and anion of a compound semiconductor. Schottky contacts conduct via the majority carr ier. Similar to Ohmic contacts, there are two predominant transport mechanisms for Schottky type conduction at room temperature: by thermionic emission (TE), or by thermionic field emission (TFE) or tunneling. The current density for the TE model follows as: 1 exp exp2 *nkT qV kT q T A Jb (1 .6) where n is the ideality factor, a term describing the character of the contact to ideal Schottky behavior. [The theoretical value of the Richardson constant A** for ZnO is 32 Acm2K2.] Current density for the T FE model follows as:

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27 0 0exp E qV J JF (1 .7) where the saturation current density J0 is given by T k q T k q T k E k V q E T A JB B B B B B exp / cosh00 5 0 00 0 (1 .8) (EF-EV)/q equals the difference between the valence band maximum and the Fermi level and E0 = E00cot h( E00/kBT ) describes the characteristic energy as related to tunneling probability. Schottky Contacts to n-ZnO : Metallization schemes involving nonreactive metals such as Au, Ag, and Pd have been shown to form Schottky contacts to n-ZnO [112119]. Schott ky barrier heights for these contacts range between 0.6 0.8 eV, but do not match to the trend in metal work function, leading to the belief that intrinsic surface states or surface contamination have an influence on electrical conduction at the interface. Pt contacts were found to have a Schottky barrier height of ~0.7 eV upon UV treatment of bulk ZnO surface. Borides have also been investigated for rectifying behavior. W2B, W2B5, and CrB2 metals deposited by sputtering on ZnO were shown to exhibit nonr ectifying behavior in the as deposited state, but converted to Schottky behavior after annealing at 500 600C. Schottky barrier heights for Boride -based contacts were only in the range of ~0.4 0.5 eV, comparable to that expected from the electron affinity Numerous surface cleaning methods prior to contact deposition have been reported, including organic solvent rinsing and etching with concentrated phosphoric or nitric acid [37, 115121]. All methods intend to raise the Schottky barrier height at the int erface and limit the influence of deep recombination centers. Most literature reports high ideality factors, suggesting

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28 multiple transport mechanisms including tunneling and the large role of interface states on conduction. 2.4 Characterization Techniques 2.4.1 Current Voltage Measurements Specific contact resistance and Schottky diode parameters are determined using IV measurements by an Agilent 4156 Semiconductor Parameter Analyzer connected to a probe station. Two point probe measurements were generall y used When the resistance of the probes was too high or contacts exhibit prohibitive sheet resistance, a four point technique will be used. In the four point method, current is sent along the outer pads while voltage drop is measured between the two in ner probes. Both linear TLM and CTLM (circular transmission line method) pads were used depending on circumstance. A typical linear TLM pattern is shown in Figure 2 2. CTLM pads are generally used in cases where the substrate is difficult to etch Metal Pads Semiconductor film L1L2L3W Metal Pads Semiconductor film L1L2L3W Fi gure 2 2. Typical s chematic for l inear TLM The s pecific contact resistance is calculated by plotting resistance R as a function of TLM spacing distance L. The y intercept of the plot = 2RC (contact resistance). The slope = RS/W where RS is the sheet res istance and W is the width of the TLM pad. An example of this plot is shown in Figure 2 3. c is calculated via the sheet resistance and follows as: W L R R RS C T2 (1.9 )

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29 S C CR W R2 2 (1 .10 ) Distance LResistance R L1L2L3Slope=RS/W 2Rc Distance LResistance R L1L2L3Slope=RS/W 2Rc Figure 2 3 Definition of r esistances for t ypical l inear TLM In the case of Schottky behavior, the barrier height for samples, b, and diode ideality factor, n may be extracted from the relation for the thermionic emission over a barrier. 2.4. 2 A uger Electron Spectroscopy Auger Electron Spectroscopy (AES) determines the elemental composition of the surface of materials. In combination with an ion gun sputter source, AES can give compositional depth profiling from relative intensity thru material stacks. Both techniques will be utilized in this work in order to understand the role of diffusion in contact behavior before and after annealing. AES works by bombarding the surface of a material with a focused beam of electrons with energy between 3 keV to 30 keV. As electrons from the beam collide with atoms at the material surface, core level electrons are ejected creating energy level vacancies. From here, an electron from the outer shell relaxes to fill the empty lower energy state and releases energy causing the ejection of another electron, this time from an outer shell. The kinetic energy of the ejected Auger electron is characteristic of the element from which it was ejected (with exception of hydrogen and helium, which can not be detected) [122].

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30 2.4.3 X -Ray Diffraction X Ray Diffraction (XRD) is an analytical, non intrusive characterization technique used to determine qualitative and quantitative data regarding chemical composition, crystal structure, and crystallographic orientation. In this work, XRD will be used to confirm material identities of nanostructures after growth and determine defects present if any. XRD works by subjecting a powdered or polycrystalline substrate to monochromatic xradiation. The individual orientation of each cr ystal diffracts (reflects) the incoming x ray beam as the angular position of the beam is rotated about the center of the sub produced is characteristic of the material sample. 2.4.4 Photoluminescence Photoluminescence (PL) is an analytical technique used to gain information about the optical properties of a substrate. From this data, the structural quality, including the influence of possible surface states and deep level traps, of the substrate can be inferred. PL emitted light is generated by impinging a light source with energy larger than the bandgap energy of the semiconductor being studied onto the semiconductor surface. This light source, such as a He Cd, Ar or Kr laser, creates free electron -hole pairs within the semiconductor. These excess carr iers recombine both via radiative and nonradiative recombination. Radiative recombination results in emission of specific wavelengths, characteristic of the recombination mechanism that created it [122]. Excitonic (electrons and holes bound to one another) luminescence is observed only at low temperatures in highly pure materials. As temperature increases, excitons break into free carriers from thermal energy and recombine via the band to -band process. An increase in doping concentration may also cause excitons to become dissociated. Because some electrons may not

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31 lie at the bottom of the conduction band, their recombination may yield a high-energy tail in the luminescence spectrum. Strict band -to -band recombination will result in a sharp cutoff at the wavelength corresponding to the material band gap. 2.4. 5 Scanning Electron Microscopy Scanning electron microscopy (SEM) produces surface images via a rastered (scanning) electron beam at high magnification (10,000X to 1,000,000X) and with high resolution (up to ~2 nm). A beam of electrons is accelerated from an electron gun and passed through a series of condenser and objective lenses to focus the beam. This electron beam hits and penetrates the top layers of the surface of the substrate and electrons are emitted from the surface interaction volume. These emitted electrons include backscattered, secondary, x ray, and Auger types. Signals collected are converted to an image that can give information regarding composition, topography and morphology [122].

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32 CHAPTER 3 THERMALLY STABLE OHM IC CONTACTS TO N ZNO 3.1 Ohmic Contacts For n -type ZnO, the most common approaches to minimizing contact resistance have been increasing the doping concentration in the near -surface region or preparing the surface to redu ce the metal semiconductor barrier height. Common metal schemes for Ohmic contacts on n type ZnO have included Ti/Au, Zn/Au, Al/Pt ,Re/Ti/Au and Pt/Ga, with typical specific contact resistances in the range 102107 2 on unintentionally (~1017 cm3) n -type ZnO after annealing at < 500C. However these typical contact stacks suffer from poor reliability at high temperature ranges needed in aggressive applications such as automobile sensors and LEDs. To improve th e reliability of contact metals at high temperature there is interest in more thermodynamically stable contact stacks. One promising materials group for thermally stable contacts on n -ZnO are stoichiometric diborides which have high melting temperatures (e .g. 3200C for ZrB2) and thermodynamic stabilities at least as good as comparable nitrides or silicides [123 125]. They also exhibit good corrosion resistance but are susceptible to oxidation during thermal processing. This oxidation may be easily counter ed however by depositing an overlayer of a metal such as Au in the same deposition chamber. Another possible material for improving contact stability is Iridium, which has not attracted much attention for use with ZnO. Ir is expected to have a relatively low barrier height on ZnO from the electron affinity of ZnO ( ZnO =4.35 eV) and work function of Ir ( Ir = 5 5.76 eV) [125]. The use of Iridium for thermally stable, low resistance Ohmic contacts has been shown for p -type GaN. In this case, minimum speci fic contact resistance for Ir contacts was ~ 9x102 2 [126]. These contacts remained stable up to annealing temperatures of 600 C. Ir -

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33 based contacts have also been shown to extend the available range of operating temperatures for AlGaN/GaN High Elect ron Mobility Transistors (HEMTs) up to 550 C [127]. 3.2 Surface Treatment Investigation A variety of typical pre -deposition surface treatments to bulk ZnO were considered for this study. Since contact behavior depends heavily on surface quality of the substrate, it may be desirable to add preparatory surface cleaning steps before contact fabrication if Ohmic character can be improved. Most surface treatments focus primarily on the removal of hydro-carbons and large free particulates, usually involving wa shing the surface with volatile solvents (IPA, acetone, and methanol). There are some reports however in which species -preferential etchants (both dry and wet) have been used in order to induce electrically desirable damage to the surface For example, t he removal of surface oxygen from ZnO has been attempted by dry etching in BCl3 plasma via the formation of volatile B -O compounds [165] This removal results with a higher concentration of charge carrying O vacancies Ar -based plasma has also been shown to increase the number of O vacancies through energetic ion bombardment of the ZnO surface [166] There is a great degree of surface roughening created by this method however, and may not result in improved Ohmic conduction of surface contacts Our sampl es of ZnO included surface treatments by exposure to Ar, Oxygen, or BCl3 plasma (via ICP dry etching ), wet etching using 1% H3PO4, treatment to Ozone (O3), or pre de position annealing at 300 Treated samples were inspected using photoluminescence spectroscopy at room temperature using a He -Cd laser excitation source A simple Ti (200)/Au (800) metallization scheme was sputter deposited for each substrate and sub sequent I -V character measured. The PL spectrum observed for all samples including untr eated ZnO is shown in Figure 3 1 Only Ozone exposure and H3PO4 wet etching resulted in an increase in photoluminescence

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34 while pre -deposition annealing led to little co nsiderable difference from the untreated substrate. Ozone is expected to remove the surface of C contamination through conversion to volatile CO and CO2 products, which may have resulted in the increase in photoluminescence observed. The reason for the i ncrease in PL by wet etching is unclear, h owever could also be due to desorption of C from the surface. The use of H3PO4 for the removal of surface contaminates on ZnO has been previously reported [38 ]. I -V measurements exhibited similar current for cont acts deposited on ozone -treated ZnO and for wet etched samples as for untreated substrates, helping confirm the existence of a clean interfacial layer by the removal of surface hydrocarbons. Large current through contacts on ZnO exposed to H3PO4 could als o be due to incorporation of hydrogen at the surface in addition to preferential removal of oxygen. Pre -deposition annealing of ZnO had little effect on I -V response of Ti/Au contacts. All dry etching trials resulted in a decrease in PL both at the excit onic -near -band -edge energy level (~3.2 eV) and at the deep -trap level in the middle of the band gap (~2.25 eV) suggesting an increase in the number of nonradiative recombination centers by pre deposition ICP etching. This is further supported by a sharp drop in current for contacts deposited to substrates receiving these treatments. In the case of contacts on ZnO exposed to oxygen plasma, I-V response changed from Ohmic to Schottky type conduction, probably due to a drastic reduction in quality of the s ubstrate surface due to dry etching and from oxygen re -incorporation, lowering conductivity at the interface. As dry etching had deleterious effects on conductivity and photoluminescent character of bulk, high -quality ZnO and other surface treatments resul ted with little noticeable effect on I -V behavior of typical contacts, no surface treatment method was employed for use in the fabrication of Ohmic contacts for this study other than degreasing in order to remove gross, loose

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35 hydrocarbons from the substrat e surface. A 3 min Ozone treatment was used for cryogenically deposited Schottky contacts (Chapter 4) to achieve superior surface quality before deposition. 3.3 Fabrication of Ohmic Contacts All n ZnO samples used were (0001) undoped grade I quality bulk, single -crystal ZnO crystals from Cermet. They were epiready with one -side Zn -face -polished by the manufacturer. The room temperature electron concentration and mobility established by van der Pauw measurements were 1017 cm3 and 190 cm2/V -s, respectively No special surface treatment was done except degrease in acetone and methanol prior to the metal deposition. 3.3 .1 Boride -based Contact Deposition For both Boride -based contact studies, a circular transmission line method (C TLM) pattern was created by liftoff of the deposited metals. CTLM patterning helped to prevent possible current spreading over the surface of the substrate. The nominal contact pad spacing varied from 5 to 45 mm and the precise length was determined by SEM measurements after patter ning. A metallization scheme of XB2 (500) / Pt (200) /Au (800) was used in all experiments. Gold was added to lower contact sheet resistance by outsourcing incoming current to all areas of the contact. Platinum was added as a diffusion barrier to pre vent metallurgical intermixing upon annealing. All of the metals or compounds were deposited by Ar plasma assisted rf sputtering at pressures of 15 40 mTorr and rf (13.56 MHz) powers of 200250 W. The sputter rates were held constant at 1.4.sec1 for al l of the metals or compounds. The contacts were patterned by standard photolithography and liftoff. Samples were annealed at temperatures up to 950 C for 1 min in a flowing N2 ambient in a RTA furnace. AES depth profiling was used to determine layer int ermixing and interfacial chemistry. As-deposited contacts showed sharp interfaces between the various metals. SEM was also used

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36 to examine contact morphology as a function of annealing temperature. The specific contact resistivity was derived from the C TLM based on current voltage measurements. 3.3 .2 Iridium based Contact Deposition An Ir(250)/Au(800) metallization scheme was deposited on all samples by Ar plasma assisted rf sputtering at pressures of 15 40 mTorr and rf (13.56 MHz) powers of 200250 W. Contacts were patterned by liftoff and annealed at 200 1100 C for 1 min within a flowing N2 or O2 ambient in a RTA furnace. N2 annealed contacts exhibiting low specific contact resistance were aged for a period of 30 days at 350C on a heater plate in air. Current -voltage response was measured every 1 3 days. I-V characteristics of the resulting diodes were measured on an Agilent 4145B parameter analyzer. The contact properties were obtained from TLM measurements on 300300 m pads with spacing of 2.5, 5.0, 7.5, 10, and 17.5 m. Depth profiling of the as -deposited and annealed contacts was done by AES using a Physical Electrons 660 Scanning Auger Microprobe. Optical micrographs were used to examine contact morphology. 3. 4 TiB2-based Contact Study TiB2-based Ohmic contacts were fabricated in order to improve thermal stability at high temperature. All contacts were rectifying to annealing temperatures of 700C. At higher temperatures there was a transition to Ohmic behavior and over a narrow temperatur e window both low specific contact resistivity and smooth morp hology were achieved. Figure 3 2 shows the variation of the specific contact resistivity and sheet resistance of the ZnO under the TiB2/Pt/Au structured contact as a function of anneal temperature. A minimum specific contact resistivity of 5x104 C anneals. This is a typical value for Ohmic contacts on undoped, lightly n type ZnO but is achieved at much higher temperatures and suggests the TiB2/ZnO interface is much more stable than for the more typical metallizations

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37 m entioned earlier. Higher temperatures led to a severe increase in sheet resistance as the ZnO dissociated and therefore led to the degradation of contact resistivity. Figure 3 3 shows optical micrograph images of the metal contacts before and after anneal ing at either 800C or 900C. The as -deposited samples highlight the small scratches and other morphological features typical of currently available ZnO substrates and are a result of imperfect polishing of the surface. The TiB2/Pt/Au shows some roughening after 800C anneals and has an almost completely reacted appearance after 900C anneals. This thermal stability is far superior to the conventional Ti or Al contacts with Au overlayers, where the contacts are not stable above 550600C. The change in a ppearance of the contacts after high temperature annealing is due to the outdiffusion of Ti, Pt, Zn and B as is clear from the AES surface composition data in Table 3 1. The increase in surface oxygen content is probably due to the oxidation of out -diffus ed Ti. C noted at the surface is adventitious and expected. Any change in surface color of the contacts is mostly from outdiffusion of oxidizing Ti. The AES depth profiles from these same three samples are shown in Figure 34 The as deposited contact s tack shows sharp interfaces between the metals. After 800C anneals, the Ti, B and Pt show diffusion through the Au while after the 900C the contact metallurgy is completely intermixed. The mechanism for Ohmic contact formation may involve formation of a more heavily doped near -surface region as result of the dissociation of the ZnO. Both TiB2 and Ti are expected to have very low barrier heights on ZnO in any case from the electron affinity of ZnO ( ZnO =4.35 eV) and work function of Ti ( Ti =4.33 eV) and B (4.3 eV) [17]. In addition, thin titanium oxide layers can be formed at the interface when Ti is contacted to ZnO since the titanium has a higher affinity with oxygen than Zn [125]. As a result, the oxygen vacancies, which are effective electron donors, increase the carrier concentration near the ZnO surface

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38 promoting the tunneling phenomena through the extremely thin oxide barrier. Our AES depth profiles do not have the resolution to provide clear evidence of the formation of a TiOx layer, but others ha ve observed this in Ti/Au contacts on ZnO [17]. TiB2/Pt/Au contacts on bulk ntype ZnO provide improved thermal stability relative to the more commonly used metal schemes published previously. The minimum contact resistance is obtained after annealing at 8 00C and suggests that these boride based contacts may be promising for high temperature ZnO device applications. 3.5 ZrB2-based Contact Study ZrB2-based contacts were also fabricated in order to achieve thermal stability at high temperature and to compare with previous Boride based contacts to n -ZnO. The contacts were rectifying to anneal temperatures of 600C but displayed Ohmic behavior between annealing temper atures of 600800C. Figure 35 shows the variation of the specific contact resistivity and t he sheet resistance of the ZnO under the contact, as a function of annealing temperature. A minimum specific contact resistivity of 9x103 C anne als. This specific contact resistivity is within the range of typical values for Oh mic contacts on undoped, lightly n type ZnO but is achieved at much higher temperatures and suggests the ZrB2/ZnO interface is relatively stable. Annealing temperatures above 800 to rectifying behavior as sheet resistance o f the ZnO increased. F igure 3 6 shows SEM images of the metal contacts before and after annealing at either 300C, 600C or 800C. The ZrB2/Pt/Au does not show any roughening until 800C anneals and is more stable than similarly prepared TiB2/Pt/Au Ohmic c ontacts on ZnO in which the contact morphology was completely destroyed by 800C [128]. This thermal stability is far superior to the conventional Ti or Al contacts with Au overlayers, where the contacts are not stable above 550600C. The change in appea rance and darkening of the contacts after high temperature

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39 annealing is due to the outdiffusion of Pt and Zn. The scratches on the metal surface are from the probes used for obtaining the I -V measurements. AES surface scans from the as -deposited and anneal ed samples are shown in Figure 3 7 A summary of this AES surface composition data is shown in Table 32. The surface scans show little change until 800C where Zn is detected. Once again this data shows the boride -based contacts are much more thermally s table than conventional metal stacks. For example, AES depth profiles of Ti/Al/Pt/Au contacts performed after annealing even at 250C showed the initially sharp interfaces between the different metals were degraded by reactions occurring, especially betwe en the Ti and the ZnO to form Ti O phases and between the Pt and Al [21]. In that case, the O appears to diffuse outward while the Pt diffuses inward. Anneals at 600 C almost completely intermixed the contact metallurgy. The AES depth profiles from these same three samples are shown in Figure 3 8 The as deposited contact stack shows sharp interfaces between the metals, although both O and C are present in the boride layer. This data points out the susceptibility of the borides to gettering of water vapor during sputter deposition [123, 124] The Pt starts to diffuse through the Au overlayer after 600 the diffusion barrier layer show some diffusion through the Au. As with TiB2-based contacts, the mechanism for Ohmic c ontact formation may involve formation of a more heavily doped near -surface region as result of the dissociation of the ZnO. ZiB2, like TiB2, is expected to have very low barrier heights on ZnO in any case from the electron affinity of ZnO ( ZnO =4.35 eV) and work function of ZrB2 ( Ti = 3.94 eV) [21, 125]. Unlike TiB2based contacts however, Zr has a low tendency to readily oxidize. As such, there was no evidence of formation of thin zirconium oxide layers formed at the interface or at the surface of the contacts even after high

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40 temperature annealing. As a result, this probably rules out the formation of oxygen vacancies, which are effective electron donors, and which would increase effective n type doping at the ZnO surface. ZrB2/Pt/Au contacts on bulk ntype ZnO provide improved thermal stability relative to the more commonly used metal schemes published previously. The minimum contact resistance is obtained after annealing at 700C and suggests that these boride based contacts may be promising for high temperature ZnO device applications. The ZrB2 produces slightly higher contact resistances than comparable CrB2-based contacts but the ZrB2 itself is more thermodynamically stable on ZnO. 3.6 Iridium based Contact Study Ir-based contacts were also attempt ed on n -ZnO in order to improve the thermal stability at high temperature. As previous literature has noted a difference between Ir contacts annealed in N2 versus O2 ambient [126] both annealing conditions were investigated for separate samples. 3.6 .1 Ni trogen Annealing of Contacts All contacts were Ohmic for all annealing tempera tures up to 1000 C. Figure 3 9 shows the variation of specific contact resistivity and the sheet resistance of the ZnO under the contact, as a function of annealing temperature in N2 ambient. A specific contact resistivity of ~5x105 2 was achieved for all contacts from as -deposited up to 900 C anneals. This specific contact resistivity is within the range of typical values for Ohmic contacts on undoped, lightly n-type ZnO but is achieved up to much higher temperatures, suggesti ng the Ir/ZnO interface is highly thermally stable. After annealing at 1000 C, the minimum specific contact resistivity drops to 3.6x106 2. After annealing at 1100 C the contact morphology degraded and electrical measurements were no longer consis tent as agglomerates (likely Iridium) collected at the surface. The sheet resistance remained at ~5

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41 Figure 3 10 shows SEM image s of the metal contacts before and after annealing (under N2 ambient) at either 800C, or 1000C. An optical micrograph of the metal cont acts after annealing at 1000C is also provided for clarity. The Ir/Au does not show any roughening until 1000C anneals and is more stable than similarly prepared boride Ohmic contacts on ZnO in which contact morph ology was completely destroyed by 900C or less [128]. This thermal stability is far superior to the conventional Ti or Al contacts with Au overlayers, where the contacts are not stable above 550 600C. The change in appearance of the contacts after high temperature annealing is not due to the outdiffusion of Ir, as confirmed through AES surface scans, but there is some dissociation of the ZnO with the appearance of Zn on the surface. Any scratches on the metal surface are from probes used for obtaining t he I -V measurements. A comparison of the depth profiles for as -deposited versus annealed at 500C, 800C (bottom left), or 1000C (bottom right) Ir/Au contacts on ZnO is given in Figure 3 11. The depth profile from AES of the as -deposited contacts shows a sharp interface between the deposited Ir and the ZnO. Ir appears to diffuse outward toward the surface while Au appears to diffuse inward toward the interface beginning after 500C anneals. Once again this data shows the Ir contacts are much more therma lly stable than conventional metal stacks. For example, AES depth profiles of Ti/Al/Pt/Au contacts performed after annealing even at 250C showed the initially sharp interfaces between the different metals were degraded by reactions occurring, especially b etween the Ti and the ZnO to form Ti O phases and between the Pt and Al [21]. In that case, the O appears to diffuse outward while the Pt diffuses inward. Anneals at 600 C almost completely intermixed that contact metallurgy. The depth profiles between con tacts annealed at 500C and contacts annealed at 800C show little difference in the diffusion of Ir toward the surface. As suggested by their similar

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42 specific contact res istances in Figure 3 8 these similar profiles are not surprising. After the 1000C anneal, both the Ir from the interface and the Au from the surface layer show heavy interdiffusion. This corresponds to a dramatic decrease in specific contact resistanc e at 1000C. In contrast to previous boride contacts, there is no evidence of a thin Iridium oxide layer or dissociation of ZnO at the interface, likely ruling out Ohmic conduction by the formation of oxygen vacancies or by a heavily doped near -surface region, respectively. Figure 3 12 shows the room temperature specific contact resistivi ty and sheet resistance of contacts annealed at 700C, as a function of time spent at 350C. This time trial was completed in order to approximate the thermal stability of a working Au Ir -ZnO device under aggressive conditions. These contacts exhibit exc ellent thermal stability over the entire 30 days. Specific contact resistivity shows a small initial drop after day 1 and then fluctuates between 2.5 and 3 x 104 2. Sheet resistance shows almost no change with annealing time, remaining at ~60 for the entire trial. In addition to being thermally stable over 30 days, the contacts also showed little change in morphology with aging time, as seen in Figure 3 1 3 Optical micrographs are shown before annealing after 15 days, and after 30 days annealing. There is a slight pitting of the surface in the annealed contacts, similar in appearance to the N2 annealed contacts shown in Fi gure 3 9 3.6.2 Oxygen Annealin g of Contacts Specific contact resistivity and sheet resistance are compared for O2 and N2 anneals in Figure 3 14. Previous reports demonstrated a dramatic difference in contact response for Ir contacts to p -type GaN between annealing in O2 and N2 ambient attributing the change in resistance to the creation of a thin IrO2 layer [126]. In our case, both specific contact resistivity and sheet resistance increase with annealing temperature under the O2 anneal by orders of magnitude while the N2 annealed con tacts remain thermally stable up to 1000C. Jang et al.

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43 explained a decrease in contact resistivity for p -type GaN after O2 annealing due to the reduction of the barrier height for injection of holes by an IrO2 phase. In our case, creation of a conducting oxide phase would not explain the data and thus an increase of barrier height on the ntype material might be a likely cause of the change in contact resistance. AES scans confirmed a higher percentage of Oxygen on the surface of contacts annealed in O2 a s compared to those annealed in an N2 ambient, supporting the existence of an IrO2 layer after annealing in O2. Surface morphology of the contact appears roughly the same between O2 or N2 anneals however, as shown in the optical micrographs taken after annealing at 700C in O2 and N2 in Figure 3 15. A comparison of the AES dep th profiles shown in Figure 3 16 also gives little as to differences in contact stability between the O2 versus N2 annealed samples. Both profiles show slight intermixing of the Au/I r and Ir/ZnO interfaces and little, if any, dissociation of ZnO. There appears to be a slightly greater diffusion of Au into the Ir layer in the case of the N2 anneal. 3.6.3 Iridium based Contact Study Summary Ir/Au contacts on bulk ntype ZnO provide hig hly improved thermal stability relative to the more commonly used metal schemes published previously. The minimum contact resistance is obtained after annealing at 1000 C (N2 ambient) and suggests that these Ir -based contacts may be promising for high te mperature ZnO device applications. The Ir produces comparable contact resistances to previous boride -based contacts, yet the Ir is much more thermally stable than any of the previously tested Borides. These Ir/Au contacts showed very little change in res istance under aggressive aging (350C) for 30 days. Resistance is increased by orders of magnitude after annealing in O2 ambient, which may be related to formation of an interfacial IrO2 layer.

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44 3.7 Conclusions Boride based and Ir/Au Ohmic contacts have be en fabricated to n ZnO. All materials tested were superior to contact stacks currently used commercially for n ZnO due to their extreme thermal stability shown via high temperature annealing. Particularly, Ir/Au contacts out performed Boride -based contac ts both in terms of low contact resistance and thermal stability at high temperatures. These results show temperature stability and effective resistance is a function of contact material, therefore it is critical to choose materials for contacts which are less chemically reactive and exhibit low thermal solubility.

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45 Table 3 1 Concentration of elements detected on the as -received surface of TiB2/Pt/Au contacts to n ZnO (in atom%) Surface element C O B Ti Zn Pt Au As deposited 41 4 nd nd nd nd 55 8 00 48 11 7 7 1 6 20 900 50 16 7 13 1 7 6

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46 Tabl e 3 2. Concentration of elements detected on the as -received surface of ZrB2/Pt/Au contacts to n ZnO (in atom%) Sample ID C O Zn Au As deposited 53 9 nd 38 300C annealed 55 7 n d 38 600C annealed 62 2 n d 36 800C annealed 68 4 4 24

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47 1.5 2.0 2.5 3.0 3.5 0.01 0.1 1 Intensity (a.u.)Energy (eV) As-deposited Annealed 300C O3 exposure H3PO4 wet etch O2 plasma BCl3 plasma Ar plasma Figure 3 1. Room temperature PL spectra for ZnO before and after various surface treatment s

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48 Figure 3 2 Specific contact resistivity of TiB 2 /Pt/Au contacts and ZnO sheet resistance as a function of annealing temperature.

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49 Figure 3 3 Optical microscopy images of TiB2/Pt/Au contacts on ZnO A) As -deposited. B) A fter annealing at 800C. C) After annealing at 9 00C

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50 Figure 3 4 AES depth profiles from TiB2/Pt/Au contacts on ZnO A) As -deposited. B) After annealing at 800C. C) After annealing at 900C.

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51 Figure 3 5 Specific contact resistivity of ZrB2/Pt/Au contacts and ZnO sheet resistance as a fu nct ion of annealing temperature

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52 Figure 3 6 SEM images of ZrB2/Pt/Au contacts A) A s deposited. B ) A fter annealing at 300C. C) After annealing at 600C. D) After annealing at 800C.

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53 Figure 3 7 AES surface scans of ZrB2/Pt/Au contacts A) As -deposited. B) After annealing at 300C. C) After annealing at 600C. D) After annealing at 800C.

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54 Figure 3 8 AES depth profiles of ZrB2/Pt/Au contacts A) As -deposited. B) After annealing at 300C. C) After annealing at 600C. D) A fter annealing at 800C.

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55 Figure 3 9 Specific contact resistivity of Ir/Au contacts and ZnO sheet resistance as a function of annealing temperature.

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56 Figure 3 10. SEM images of Ir/Au contacts A) A s -deposited B) A fter annealing at 800C. C) After annealing at 1000C D) Optical micrograph after annealing at 1000C.

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57 Figure 3 1 1 AES depth profiles of Ir/Au contacts A) A s -deposited B) A fter annealing at 500C. C) After annealing at 800C. D) After annealing at 1000C.

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58 0 3 6 9 12 15 18 21 24 27 30 2x10-42.5x10-43x10-43.5x10-44x10-44.5x10-45x10-4 Days Annealed @ 350 CSpecific Contact Resistance ( cm2) 20 40 60 80 100 Sheet Resistance ( /square) Figu re 3 12. Specific contact resistivity and sheet resistance of the contacts annealed at 700 C (N2 ambient) as a function of aging time at 350C.

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59 Figure 3 13. Optical micrographs of the 700C annealed (N2 ambient) Ir/Au contacts A) Before aging. B) A fter 15 days aging C) After 30 days aging.

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60 Figure 3 1 4 Resistivity comparison of Ir/Au contacts between O2 and N2 anneal. A) Specific contact resistivity. B) S heet resistance

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61 Figure 3 15. Optical micrographs of Ir/Au contacts after 700C anneal A) After O2 anneal. B) After N2 anneal

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62 Figure 3 1 6. AES depth profiles of Ir/Au contacts after 700C anneal A) After O2 anneal. B) After N2 anneal 0 100 200 300 400 500 600 700 800 0 10 20 30 40 50 60 70 80 90 100 N 2 Anneal Sputter Depth ( ) Atomic Concentration (%) C O Ir Au Zn Sputter Depth ( ) 0 100 200 300 400 500 600 700 800 900 1000 0 10 20 30 40 50 60 70 80 90 100 O 2 Anneal Atomic Concentration (%) C O Au Ir Zn

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63 CHAPTER 4 EFFECT OF CRYOGENIC DEPOSITION ON SCHOTTKY METAL CONTACTS TO N ZNO 4.1 Introduction The development of reliable and thermally stable Ohmic and Schottky contacts to ZnO is one of the critical issues related to the fabrication of ZnO based UV light emitters/detectors and field effect transistors. To date, a number of differe nt metallization schemes have been examined for Ohmic and rectifying contacts on n -ZnO [13, 3039, 129133]. Metals such as Au, Ag and Pd form rectifying contacts with low Schottky barrier heights in the 0.6 0.8 eV range [115121]. In addition, the therm al stability of these contacts is usually extremely poor, with degradation occurring even at 60C for Au/n -ZnO. The trend in barrier heights often does not correlate with the metal work functions, indicating that surface states or surface contamination is playing an important role in determining the electrical transport properties of the contacts. One approach to achieving increased Schottky barrier heights that has proven successful for other semiconductors such as GaAs and InP is the use of cryogenic dep osition temperatures [39, 40, 134136]. The mechanism for the barrier height enhancement is still not firmly established. In the case of Au contacts on InP [136], room temperature deposition produced an ideality factor of 1.02 which was nearly independent of temperature and the current transport was controlled by thermionic emission (TE). In the case of 77K deposition of the Au, the ideality factor was increased and the current transport was controlled by thermionic field emission (TFE). The barrier heig ht enhancement and the difference in transport mechanism were both attributed to the formation of an amorphous layer at the metal/semiconductor interface. This amorphous layer was suggested to act as an insulator to create a metal -insulator semiconductor (MIS) like structure. Alternative mechanisms for the enhanced barrier properties

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64 have been suggested, including the presence of an inhomogeneous Schottky barrier height in the diodes due to a dependence of the local interface dipole on local interfacial s tructure [137]. In this study, we report on the effect of cryogenic temperature during metal deposition on the contact properties of Au, Pd, Pt, Ti, and Ni on bulk single -crystal n -type ZnO. Some differences in the electrical behavior are noted as well a s problems due to peeling/cracking in some of the contacts deposited at low temperature. The effect of pos t -deposition annealing is to improve the electrical performance of the Pd contacts deposited at cryogenic temperatures. 4.2 Experimental Details The samples used were (0001) undoped bulk ntype ZnO crystals from Cermet Inc. The samples were epiready, one -side -Zn -face -polished by the manufacturer. As determined via van der Pauw measurements, room temperature electron concentration and mobility were ~1 017 cm3 and 190 cm2/V.s, respectively. Full backside Ohmic contacts with a metallization scheme of Ti (200) / Au (2000) were deposited on all samples by Ar plasma assisted rf sputtering at pressures of 15 40 mTorr and rf (13.56 MHz) powers of 200250 W These contacts were annealed at 450C for 1 minute in O2 ambient. The typical specific contact resistivity is <104 2. Schottky contacts were deposited on all samples using an evaporator system with a load lock. The loadlock maintained a backgroun d pressure in the deposition chamber of ~1010 Torr. Using these very low background pressures, the rate of gas molecule impingement onto the sample surface is significantly reduced and monolayer formation time is extended by ~12 hours. Front -side contact s were deposited at 77K or 300K using metallization schemes of (1000) Au, Ti, Ni, Pt, and Pd. Prior to insertion in the evaporator, a ll samples were exposed to a 3 min UV ozone t reatment. Contacts ranged in diameter from 200shadow -mask.

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65 The current voltage (I -V) characteristics of the resulting diodes were measured on an Agilent 4145B parameter analyzer. The barrier height for the n type sample b, and diode ideality factor, n were extracted from the relation for the thermionic emission over a barrier (see section 2.3.2). I -V curves were obtained as a function of post -deposition annealing temperature (up to 300C, 30 min anneals) for Pd. Depth profiling of the as -deposited and annealed contacts was done by Auger Electron Spectroscopy (AES) using a Physical Electronics 660 Scanning Auger Microprobe. Optical micrographs were used to examine contact morphology 4.3 Results and Discussion 4.3.1 R esults and Discussion Ti, Ni, Pt, Pd Contacts The effect of cryogenic temperature during metal deposition on the contact properties of Pd, Pt, Ti, and Ni on bulk single crystal n type ZnO was investigated. Differences in the electrical behavior were noted as well as physical problems due to peeling/cracking in some of the contacts deposited at low temperature. Figure 4 1 shows an optical micrograph of Ti contacts deposited at 300K. The contacts deposited at 77K showed similar appearance. This appearance w as typical of almost all contacts deposited; however both Pt and Pd exhibited peeling/cracking of some of the contacts after deposition at 77K. We carried out our electrical measurements only on the contacts that did not show these problems. Figure 4 2 dis plays the I -V characteristics for all metal contacts deposited at room temperature or low temperature in both linear and log form. These I -V curves show Ohmic behavior for Ni and Ti after both room temperature and low temperature deposition, while Schottk y like behavior is exhibited for Pt and Pd contacts. This difference in electrical behavior is in reasonable agreement with the Schottky -Mott model by comparison of the work function of the metals with the electron affinity of ZnO (~ 4.5 eV), with only Ni diverging from the expected behavior. Ti has a work function of 4.3 eV, suggesting Ohmic behavior by the

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66 Schottky Mott model. The work function for both Pd and Ni is 5.1 eV and for Pt is 5.7 eV [125], suggesting rectifying behavior [138]. The I -V respon se of Pd contacts (along with other factors discussed later) suggests more of a MIS behavior than a true Schottky contact [41, 125, 135, 139]. Figure 4 3 shows the I -V respon se of the Ohmic contacts of Ni or Ti after deposition at 300K or 77K. There is a clear increase in resistance for low temperature deposited contacts. In Ni, resistance increases by 10 20% for low deposited contacts. In Ti, the resistance increases by 30% or more with low temperature deposition. The resistance of the contacts decrease d with increased contact area, as expected. Figure 4 4 shows the I -V response of the Schottky like contacts of Pt or Pd after deposition at 300K or 77K. Ideality factor and barrier height values for Pt and Pd contacts from room temperature and low tempe rature deposition are given in Table 4 1. In nearly all cases, the ideality factor for the low temperature deposited contacts was lower than those deposited at room temperature. Barrier heights were greater in the low temperature deposited samples for Pt and Pd contacts. It has been suggested the more abrupt metal -semiconductor interface for low temperature deposited contacts causes the difference in barrier height [41]. The ideality factors of Pd contacts decreased with subsequent anneals for all contact sizes over all temperatures, typically by 0.20.3. The ideality factor of Pd contacts after annealing at 300C was, as expected, in the range of 1 2. The barrier height increased with subsequent anneals for both the room and low temperature deposited sa mples, as shown in Figure 4 5. After annealing at 300C, the increase in barrier height was ~0.6 eV in comparison to as -deposited contacts for all samples. This increase in barrier height is consistent with the formation of a growing interfacial oxide la yer [135137]. This is sharp contrast with the results for Au (a

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67 difficult metal to oxidize), in which case barrier height changed little as a function of annealing temperature [140]. Capacitance -voltage and X -Ray Photoelectron Spectroscopy (XPS) measure me nts on these annealed diodes may mor e firmly establish the barrier heights and conduction mechanisms. Figure 4 6 shows the change in leakage current (@ 0.5V) with annealing temperature. Leakage current decreased by an order of magnitude after subsequent annealing for contacts deposited at room temperature. However for low temperature deposited contacts, leakage current remained low, irrespective of annealing. This is supportive of the more abrupt metal semiconductor interface for the low temperature dep osited contact. As annealing temperature increased, the effect of surface area on leakage current diminished. A comparison of the typical depth profile for the metals deposited on ZnO is given in Figure 4 7. AES surface scans of Ni as deposited after room or low temperature deposition had only three components: Ni, oxygen from a native oxide on the Ni, and carbon from atmospheric exposure. The depth profile from AES of the as -deposited contacts shows a sharp interface between the deposited metal and the ZnO. There is no major difference in the depth profile between contacts deposited at 300K versus those deposited at 77K. A comparison of the depth profile for Pd contacts on ZnO after annealing is shown in Figure 4 8. After annealing, contacts deposite d at low temperature show no detectable difference in interface from those deposited at room temperature. More unexpected is the apparent lack of difference between annealed and as deposited contacts, however our AES depth profiles do not have the resoluti on to provide clear evidence of the formation of a thin oxide layer. There are a number of noteworthy similarities and differences in the electrical characteristics of these various metals on ZnO dependant on the temperature of contact

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68 deposition. Ni and Ti samples deposited at both 300 and 77K showed Ohmic behavior, while Pt and Pd samples exhibited rectifying behavior for both deposition temperatures. Contacts deposited at lower temperature showed higher resistance for all samples. Barrier heights for P t and Pd deposited at room temperature were ~0.4 eV and above 0.5 eV for contacts deposited at low temperature. In the case of Pd, barrier height increased by ~0.6 eV after annealing up to 300C. Leakage current (at 0.5V) in the Pd contacts decreased as a function of annealing temperature for room temperature deposited contacts, however remained low for low temperature deposited contacts regardless of anneals. Possible mechanisms for the differences in behavior as a function of contact deposition tempera ture include the formation of an interfacial oxide layer or a sharper interface between metal and ZnO. Our measurements do not allow us to differentiate between the two mechanisms; more detailed XPS measurements with samples annealed at varying temperatur es would be helpful in deciding the appropriate mechanism. 4.3 .2 Results and Discussion Au Contacts Au is known to have great potential as an effective Schottky barrier to nZnO at room temperature; however because of the extremely poor thermal stability of Au, Au contacts begin thermal degradation at temperatures as low as 60 contacts to n -ZnO was investigated to improve thermal stability and Schottky character. Figure 4 9 shows optical microscope images of Au c ontacts deposited at 300K or 77K, with both exhibiting excellent morphology. We did not observe any cracking or peeling of the contacts deposited at 77K. Figure 4 10 shows I -V characteristics of Au/n GaN Schottky diodes deposited at either 77K or 300K in b oth linear and log form. There is a clear decrease in both reverse and forward bias current for the diodes with Au deposited at 77K, consistent with an increase in the effective Schottky barrier height. Indeed, the contacts deposited at room

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69 temperature ar e essentially Ohmic. The reverse current was more than an order of magnitude lower in the 77K deposited diodes, but in both types of diodes was proportional to both the perimeter and area of the rectifying contact, suggesting that both surface and bulk contributions are present in this voltage range. Therefore, low temperature deposition does not seem to reduce Fermi level pinning by surface states in ZnO. For the 77K deposited samples, we extracted a barrier height of 0.37 eV, with an ideality factor of > 2 which indicates there are multiple current transport mechanisms present. The barrier height is lower than reported previously, but it should be noted that ZnO may exhibit an electrically conducting layer at the surface depending on the measurement ambien t [141]. It appears that there can be an electrically conducting surface channel present in vacuum that disappears upon exposure to air and this may affect the Schottky barrier properties. Figures 4 11 and 4 12 show the I -V characteristics from the 300K a nd 77K deposited samples as a function of post -deposition annealing temperature. The characteristics become more similar in terms of magnitude of current in both directions as the annealing temperature increases. Figure 4 13 shows a direct comparison of the samples after annealing at 300 temperature, the difference in reverse current is down to approximately a factor of 2, which we consider to be insignificant because this is typical of sample to -sample variations. The sample deposited at 300K dis played a patchy appearance, as shown in the SEM micrograph at the top of Figure 4 14. The AES surface scan of this sample showed only Au, oxygen from a native oxide on the Au and adventitious carbon from atmospheric exposure. The sample deposited at 77K and annealed at 300 at the top of Figure 4 15, but the AES surface scan was similar to that from the 300K sample and the near -surface stoichiometry was also similar (the average composition of the top 100 from

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70 the surface was 36% C, 3% O and 61% Au). The barrier height of the cryogenically deposited sample showed little change with annealing and was in the range 0.39+/ 0.02 eV for all annealing temperatures in the range 25 300 d improve with annealing, with a value of 1.3 after 300 The AES depth profiles from the 300K and 77K samples were similar, as shown in Figure 4 16. Previous x ray reflectivity data on our samples has shown that the main difference between Au deposited at 77K and room temperature is a decreased interfacial roughness between the Au and the semiconductor. As the diodes are annealed to 300 height and interfacial roughness is lost. Previous electron microscopy and c hemical bonding studies are also supportive of interfacial layer differences being the cause of the differences in barrier height, with a more abrupt metal -semiconductor interface for low temperature deposited contacts [41]. The most commonly suggested mec hanisms for such data include the formation of an amorphous layer at the cryogenic diode interface to create a metal insulator semiconductor (MIS) structure [135, 136] and inhomogeneous Schottky barrier heights due to a dependence of the local interface dipole on the local interface structure [40, 137]. It has been reported that the crystal structure and grain size of the low temperature deposited metal are different from metals deposited at room temperature. 4.4 Conclusions There are a number of notewor thy similarities and differences in the electrical characteristics of various metals on ZnO dependant on the temperature of contact deposition. Ni and Ti samples deposited at both 300 and 77K showed Ohmic behavior, while Pt and Pd samples exhibited rectif ying behavior for both deposition temperatures. Au samples deposited at 77K show rectifying behavior with barrier heights around 0.4 eV, while room temperature deposition produces Ohmic behavior. Contacts deposited at lower temperature showed higher resi stance for

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71 all samples. Barrier heights for Pt and Pd deposited at room temperature were ~0.4 eV and above 0.5 eV for contacts deposited at low temperature. In the case of Pd, barrier height increased by ~0.6 eV after annealing up to 300C. Meanwhile, dif ferences in electrical properties of Au contacts are reduced by annealing and are stable to ~300 0.5V) in the Pd contacts decreased as a function of annealing temperature for room temperature deposited contacts, however remained low for low temperature deposited contacts regardless of anneals. Possible mechanisms for the differences in behavior as a function of contact deposition temperature include the formation of an interfacial oxide layer or a sharper interface between metal and ZnO.

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72 Table 4 1 Summary of barrier height and ideality factors for contacts on ZnO. Deposition n b (eV) Pt Room Temp ~2 0.37 Low Temp 1.6 0.5 0 Pd Room Temp 1.9 0.44 Low Temp 1.7 0.69

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73 Figure 4 1. Optical micrograph of Ti contacts deposited at either 300K. The smallest contact

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74 -1.0 -0.5 0.0 0.5 1.0 -80 -60 -40 -20 0 20 40 60 80 Current (A/cm2)Voltage (V)ZnO 400 m device room temp deposition Ti Ni Pt Pd -1.0 -0.5 0.0 0.5 1.0 10-410-310-210-1100101102 Current (A/cm2)Voltage (V)ZnO 400 m device room temp deposition Ti Ni Pt Pd -1.0 -0.5 0.0 0.5 1.0 -100 -75 -50 -25 0 25 50 75 100 Current (A/cm2)Voltage (V)ZnO 400 m device low temp deposition Ti Ni Pt Pd -1.0 -0.5 0.0 0.5 1.0 10-610-510-410-310-210-1100101102 Current (A/cm2)Voltage (V)ZnO 400 m device low temp deposition Ti Ni Pt Pd Figure 4 2. I-V charact eristics of room temperature and cryogenically deposited diodes. A) 300K deposition linear plot. B) 300K deposition log plot. C) 77K deposition linear plot. D) 77K deposition log plot.

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75 -0.2 -0.1 0.0 0.1 0.2 -40 -30 -20 -10 0 10 20 30 40 Current (A/cm2)Voltage (V)Ni // ZnO as-deposited RT200 m LT200 m RT400 m LT400 m RT800 m LT800 m -0.6 -0.4 -0.2 0.0 0.2 0.4 0.6 -40 -20 0 20 40 Current (A/cm2)Voltage (V)Ti // ZnO as-deposited RT200 m LT200 m RT400 m LT400 m RT800 m LT800 m Figure 4 3. I-V characteristics of Ohmic contacts depos ited at either 300K or 77K. A) Ni contacts. B) Ti contacts.

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76 -1.0 -0.5 0.0 0.5 1.0 -25 0 25 50 75 100 125 150 Current (A/cm2)Voltage (V) Pt // ZnO as-deposited RT200 m LT200 m RT400 m LT400 m RT800 m LT800 m -1.0 -0.5 0.0 0.5 1.0 10-510-410-310-210-1100101102 Current (A/cm2)Voltage (V) Pt // ZnO as-deposited RT200 m LT200 m RT400 m LT400 m RT800 m LT800 m -1.0 -0.5 0.0 0.5 1.0 0 20 40 60 80 100 120 Current (A/cm2)Voltage (V) Pd // ZnO as-deposited RT200 m LT200 m RT400 m LT400 m RT800 m LT800 m -1.0 -0.5 0.0 0.5 1.0 10-710-610-510-410-310-210-1100101102 Current (A/cm2)Voltage (V) RT200 m RT400 m RT800 m LT200 m LT400 m LT800 m Pd // ZnO as-deposited Figure 4 4. I-V characteristics of Schottky diodes deposited at either 300K or 77K A) Pt contacts linear plot. B) Pt contacts log plot. C) Pd contacts linear plot. D) Pd cont acts log plot

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77 Figure 4 5. Schottky barrier height as a function of annealing temperature for diodes with Pd contacts deposited at either 77 or 300K. 50 100 150 200 250 300 0.4 0.5 0.6 0.7 0.8 0.9 1.0 1.1 1.2 1.3 1.4 1.5 Barrier Height (eV)Annealing Temperature ( C)Pd // ZnO RT200 m LT200 m RT400 m LT400 m RT800 m LT800 m

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78 0 50 100 150 200 250 300 -3.0 -2.5 -2.0 -1.5 -1.0 -0.5 0.0 Leakage Current (mA)Annealing Temperature ( C) Pd // ZnO @ 0.5V RT200 m LT200 m RT400 m LT400 m RT800 m LT800 m Figure 4 6. Revers e leakage current (@ 0.5V) as a function of anneali ng temperature for diodes with Pd contacts deposited at either 77 or 300K.

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79 Figure 4 7. AES depth profiles from Ni on ZnO. A) 300K deposition. B) 77K deposition 0 500 1000 1500 0 10 20 30 40 50 60 70 80 90 100 Atomic Concentration (%) C O N i Zn Sputter Depth ( ) 0 500 1000 1500 0 10 20 30 40 50 60 70 80 90 100 Sputter Depth ( ) Atom ic Concentration (%) C O Ni Zn

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80 Figure 4 8. AES depth profiles of Pt contacts on ZnO after annealing. A) 300K dep osition. B) 77K deposition 0 100 200 300 400 500 600 700 800 900 1000 0 10 20 30 40 50 60 70 80 90 100 Sputter Depth ( ) Atomic Concentration (%) Pd Zn O 0 100 200 300 400 500 600 700 800 900 1000 0 10 20 30 40 50 60 70 80 90 100 Sputter Depth ( ) Atomic Concentration (%) O Pd Zn

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81 Figure 4 9. Optical microscope images of Au cont acts on ZnO. A) 300K deposition. B) 77K deposition The largest contact diameter is 400m.

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82 Figure 4 10. I-V characteristics of Au/n GaN Schottky diodes deposit ed at ei ther 77K or 300K A) L inear plot. B) L og plot -0.4 -0.2 0.0 0.2 0.4 -0.10 -0.05 0.00 0.05 0.10 Current(A)Voltage(V)ZnO 400u device Ohmic(Ti/Au) room T 77K -0.4 -0.2 0.0 0.2 0.4 1x10-41x10-31x10-21x10-11x1001x1011x102 Current(A/cm2) Voltage(V)Au/ZnO as-grown devices 400um room T 77K

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83 Figure 4 11. I-V characteristics from the 300K deposited samples as a function of post deposition annealing temperature A) L inear plot. B) Log plot -0.4 -0.2 0.0 0.2 0.4 -80 -60 -40 -20 0 20 40 60 80 Current(A/cm2)room T device Au/ZnO 400um as-grown 50C 100C 150C 200C 250C 300C Voltage(V) -0.4 -0.2 0.0 0.2 0.4 1x10-41x10-31x10-21x10-11x1001x1011x102 Current(A/cm2)room T device Au/ZnO 400um as-grown 50C 100C 150C 200C 250C 300C Voltage(V)

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84 Figure 4 12. I-V characteristics from the 77K deposited samples as a function of post -deposition annealing temperature A) L inear plot. B) Log plot -0.4 -0.2 0.0 0.2 0.4 -80 -60 -40 -20 0 20 40 60 80 Current(A/cm2)Voltage(V)77K device Au/ZnO 400um as-grown 50C 100C 150C 200C 250C 300C -0.4 -0.2 0.0 0.2 0.4 1x10-41x10-31x10-21x10-11x1001x1011x102 77K device Au/ZnO 400um as-grown 50C 100C 150C 200C 250C 300CCurrent(A/cm2) Voltage(V)

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85 Figure 4 13. Comparison of the I -V characteristics from both the 300K and 77K deposited samples after annealing at 300 A) L inear plot. B) L o g plot. -0.4 -0.2 0.0 0.2 0.4 -80 -60 -40 -20 0 20 40 60 80 Au/ZnO 300C annealed devices 400um room T 77K Current(A)Voltage(V) -0.4 -0.2 0.0 0.2 0.4 1x10-41x10-31x10-21x10-11x1001x1011x102 Au/ZnO 300C annealed devices 400um room T 77K Current(A)Voltage(V)

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86 Figure 4 14. Surface characterization of Au contacts on ZnO deposited at 300K. A) SEM micrograph. B) AES surface scan

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87 Figure 4 15. Surface characterization of Au contacts on ZnO deposited at 77K and subsequently annealed. A) SEM micro graph. B) AES surface scan.

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88 Figure 4 16. AES depth profiles form Au deposited on ZnO. A) 300K deposition. B) 77K deposition and subsequently annealed at 300

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89 CHAPTER 5 FUNCTIONALIZATION OF NANOMATERIAL DEVICES FOR H2 SENSING 5.1 Introduction Continued research into the use of H2 as an energy source has necessitated the development of robust, low power hydrogen-selective gas se nsors [61, 142 -144]. Because H2 is a hazardous, odorless, and flammable gas, H2 gas sensors have many possible niche uses, particularly for combustion gas detection for fuel leak detection in spacecraft, autos and aircraft, fire detectors, and industrial process emiss ions [145148]. Central to realizing these next generational sensors is the ability to detect hydrogen with minimal power consumption at or much below room temperature. Nitride -based semiconductors such as GaN and InN are possible popular materials for H2 gas sensing in part to their sensitivity to surface charge and wide temperature stability [51, 52, 149157]. While numerous groups have already reported the use of H2 sensors based on CNTs, ZnO nanorods, and SnO2 nanowires with excellent response and re covery characteristics, there are few studies on H2 gas sensors based on GaN and InN nanostructures, which may offer excellent environmental stability. Critical to the func tionalization of nanomaterial H2 sensors is the addition of a thin metallic catalyti c coating. While the mechanism of this functional layer for H2 sensing is well understood, the choice of material still seems a matter of preference. Although both Pt and Pd have typically been accepted as choice metals for functional layers on oxide an d nitride -based sensors in numerous reports, there have been few studies to compare Pt and Pd metal functionalities. In this section, the hydrogen sensing properties of Pt and Pd -coated multiple GaN nanowires and multiple InN nanobelts at different hydro gen concentrations are evaluated and compared.

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90 5.2 InN Nanobelts 5.2.1 Growth Process The growth of InN nanobelts was performed by Metal Organic Chemical Vapor Deposition on Si substrates using trimethylindium (TMIn) and ammonia (NH3) as In and N sources, respectively, while nitrogen was used as the carrier gas. The nucleation sites for catalyst -driven growth of the nanobelts were formed by depositing 2 nm of Au film on SiNxcoated Si substrates by direct current sputtering. InN nanobelts were synthesized at 500C and 5 Torr pressure for 2 hours. The complete growth process and optical properties for InN nanobelts has been described in detail elsewhere [158, 159]. Figure 5 1 shows 2 X ray diffraction (XRD) spectrum of as -grown InN nanobelts. The diffrac tion peaks can be indexed to the hexagonal structure with lattice parameter of a=3.597 and c=5.703 calculated from Braggs law (JCPDS 021450, a=3.537 and c=5.704). The nanobelts have well h in the inset of Figure 51. 5.2.2 Pt -functionalization A 10 nm thick layer of Pt was deposited by sputtering onto the InN multiple nanobelts to verify the effect of catalyst on gas sensitivity. A shadow mask was used to pattern nonalloyed, electron beam evaporated Ti(20 nm)/Au(80 nm) electrodes on the InN nanobelts. The separation between both contacts was ~30 package. The InN nanobelts sensors were exposed to different H2 concentrations (20300 ppm H2 in N2 ambient) at 25 190C. The current -voltage ( I -V ) characteristics from the nanobelts both before a nd after Pt deposition were linear under our operating conditions. I-V characteristics of both uncoated and Pt -coated InN nanobelts in N2 ambient at room temperature were linear, as shown in Figure 5 2 A ). Pt coating of the nanobelts improved the

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91 current b y a factor of ~8 relative to the uncoated case. This is due to the presence of the metal (Pt), resulting in higher effective conductivity of the nanobelts. This result is in contrast to the I V response of Pt -coated InN nanorods as measured by Kryliouk e t al. who noticed a decrease in current with Pt coating attributing the decrease to sputter damage [160]. This difference might be caused either by the inability to precisely measure amount of coating deposition at low thicknesses and/or quality of the In N nanomaterials. Figure 5 2 B ) shows the change in current from Pt -coated InN nanobelts at a fixed bias of 0.1V when exposed to different H2 concentrations (20300 ppm) from air. The addition of Pt on InN nanobelts produced a strong response to H2 while t here was no detectable change in current for the uncoated nanobelts under the same gas concentrations. The most likely mechanism for the current increase is the adsorption of atomic hydrogen dissociated catalytically from hydrogen molecules (H2). Such atom ic hydrogen rapidly diffuses through Pt and is adsorbed at the Pt/InN interface, leading to changes in depletion depth by the exchange of charges between adsorbed gas species and the surface of the nanobelts [51, 52] While the exact chemical constructs o f the sensing mechanism(s) are unclear, previous literature suggests the adsorption of atomic hydrogen over thin oxidic layers between the metal and semiconductor surface resulting in the creation of a dipole layer [163]. This dipole layer is then balance d by a reduction in the depletion area underneath the semiconductor surface [164]. The current change was slow (~ 20s) in the beginning of the exposure to hydrogen and then sharply increased, which could be due to the Pt being covered with native oxide. A s the native oxide was removed by H2 exposure, the rate of current change increased with time. However, the rate began to decrease after 100s because the Pt surface was gradually saturated with H2 and limited the supply to the Pt/InN interface. Figure 5 3 shows the relative response of

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92 Pt -coated InN nanobelts exposed to a series of hydrogen concentrations in N2 ambient for 10 min at room temperature. Note that the InN sensor detected hydrogen down to tens of ppm concentration levels with relative responses of ~1.2% at 20 ppm and 4% at ~300 ppm. The resistance was recovered to approximately 90% of initial value within 2 min upon the removal of hydrogen from the ambient. This is faster than previous results using InN nanowires [160]. The maximum power consumpt ion during the measurement was ~ 0.5 mW, indicating that InN nanobelts are attractive for long term and low -cost hydrogen sensing applications. The rate of resistance change for Pt -coated InN nanobelts exposed to 200 ppm H2 in N2 at a fixed bias of 0.1V w as measured at different temperature range (25 190C), as shown in Figure 5 4. An adsorption activation energy was found to be 0.094 eV from the slop of Arrhenius plot, which is smaller than those obtained from PdGaN (0.097 eV) [55] and Pd-ZnO (0.12 eV) [ 52] nanowires. The low hydrogen adsorption activation energy of the sensor explains its high sensitivity and fast response to hydrogen gas. The inset shows the resistance change during the exposure to hydrogen of 200 ppm at different temperature. The res istance change was increased with the measurement temperature. Similarly, the relative responses were increased from 3.1% at room temperature to 4.5% and 5.6% at 90C and 150C, respectively, resulting from the increase in catalytic dissociation rate of mo lecular hydrogen (H2) or diffusion rate of atomic hydrogen into Pt/InN interface. 5.2.3 Pd -functionalization A ~7 nm thick layer of Pd was sputter deposited onto some of the nanobelts. A shadow mask was used to pattern non alloyed, rf -sputtered Ti(50 nm)/ Al (80 nm)/Pt (40 nm)/Au (300 bonded to the contact pads for device packaging. The InN nanobelts sensors were exposed to

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93 different H2 concentrations (1001000 ppm H2 in N2 ambient) at 25 130C. Sensor response to high concentration of CO2, C2H6, NH3, and O2 (all in N2 ambient) was also examined Linear current voltage ( IV ) characteristics were observed for both uncoated and Pdcoated InN nanobelts in air at room te mperature, as shown in Figure 5 5 A ). The addition of the Pd -coating increased the effective conductivity by a factor of ~5. Figure 5 5 B) shows the change in relative response from Pd-coated InN nanobelts at a fixed bias of 0. 1V upon 10 minute exposure t o different H2 concentrations (1001000 ppm) for 10 minutes at 130C from air. With Pd functionalization, the InN sensor detected hydrogen down to 100 ppm concentration levels (equipment limit) with a relative response of ~8% at 100 ppm and ~9.5% at 1000 ppm after 5 minute exposure. The response recovered to ~ 50% of initial value 10 minutes after removal of hydrogen from the ambient. This is slower than previous reports using Pt coated nanobelts [160, 161], likely due to Pd having a greater affinity for hydrogen. By sharp contrast, there was no detectable change in current for uncoated nanobelts under the same conditions. In addition, there was no current response of the nanobelts to CO2, C2H6, NH3 or O2 gas either before or after deposition of the Pd functional layer. The sharp increase in current change upon exposure to hydrogen was nearly instantaneous (< 2s). This is a marked difference from previous Pt -coated InN nanobelts which showed a much slower (~ 20s) response time.19 In both cases, response rate begins to decrease as the functional layer becomes saturated with H2 (~ 100s), limiting the supply to the (Pt/Pd)/InN interface. The relative response for Pd -coated InN nanobelts exposed to 300 ppm H2 in N2 at a fixed bias of 0.1V after 10 minutes wa s measured over the temperature range (25 130C), as shown in Figure 5 6. Relative response increased with measurement temperature from 4.3% at

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94 room temperature up to 9.4% at 130C, resulting from the increase in catalytic dissociation rate of molecular h ydrogen (H2) or diffusion rate of atomic hydrogen into the Pd/InN interface. This response is at least 30% larger for all temperatures in comparison to similar Pt -coated InN nanobelts. An extrapolation of the trend line suggests there would still be at l east a 1% relative change in current at temperatures as low as ~220 K. This result is supportive of the viability of Pd -coated InN nanobelt H2 sensors for low temperature applications. The activation energy for H detection on Pd -coated InN nanobelts of 0. 07 eV was calculated from the slop e of the Arrhenius plot, shown in Figure 5 7. This is slightly lower than energies obtained from Pd GaN (0.097 eV) [52] and Pd ZnO (0.12 eV) [60] nanowires and previous Pt InN (0.094 eV) nanobelts but still larger than v alues expected for a typical diffusion process. Thus the dominant sensing mechanism is likely chemisorption of the H to the Pd InN interface [52] The low hydrogen adsorption activation energy of this sensor is indicative of its high sensitivity and inst antaneous response to H2 in gas. 5.3 GaN Nanowires 5.3.1 Growth Process For GaN nanowire growth, a growth substrate was prepared by e beam evaporating 15 gold onto a clean piece of (100) Si with 100 nm thermally grown oxide. A Gallium metal source (99.999%) was poured into a quartz boat and placed into a tube furnace. The growth substrate was positioned within 3 cm downstream of the Ga metal source. The growth chamber was purged with Ar for 10 min at room temperature to remove any residual oxygen. The substrate was heated up to 850 Au catalyst nanoparticles on the sample surface. After annealing, high purity NH3 (99.999%) and H2 (99.999%) were introduced for growth. The GaN nanowires were grown for ~3 h at 850

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95 oxidation. Typical length of the resultant GaN nanowires was 2typical as -grown nanowires are shown in Figu re 5 8. X ray diffraction, high resolution transmission electron microscopy, and photoluminescence showed the nanowires to be single crystal wurtzite GaN. 5.3.2 Pd -functionalization A shadow mask was used to pattern Ti (20 nm)/Al (80 nm)/Pt (40 nm)/Au (80 nm) Ohmic contacts by e -beam evaporation. The contacts were annealed at 350 C for 60 seconds in flowing N 2 ambient in a rapid thermal annealing furnace. Au wires were bonded to the contact pads for device packaging. A 10 nm thick Pd functional layer was deposited by sputtering onto the nanowires to verify the effect of catalyst on gas sensitivity. The device was exposed to varying H 2 concentrations (200 3000 ppm H2 in N 2 ambient) at 25 150 C. Current -voltage characteristics from multiple nanowires were linear with a maximum power consumption of ~0.6 mW under our operating conditions. Figure 5 9 shows the measured resistance at a bias of 0.5 V as a function of time from Pd -coated and uncoated multiple GaN nanowires exposed to a series of H2 concentration s (200 1500 ppm) in N 2 for 10 min at room temperature. Pd coating of the nanowires improved the sensitivity to ppm level H2 by a factor of up to 11. The addition of Pd appears to be highly effective in catalytic dissociation of molecular hydrogen. The res istance change depended on the gas concentration but the variations were small at H2 concentrations above 1000 ppm. Initial resistance was restored to approximately 90% of initial level within 2 min of exposing the nanowires back to air. Pd -coated GaN nano wires exhibited relative responses of ~7.4 % at 200 ppm up to ~9.1% at 1500 ppm H2 in N2 as shown in Figure 5 10 A ). Relative responses of uncoated nanowires under the same conditions were ~0.48% and 1.2%, respectively. Figure 5 10 B) shows the time

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9 6 depen dent resistance change in Pd -coated multiple GaN nanowires as the gas ambient is switched from air to various concentrations of H2 in N2 (200 1500 ppm) and then back to air. The rate of resistance change was sharply increased and then reached maximum at th e exposure time of 1 min, which might be due to the Pd being covered with native oxide. As the native oxide was removed by H2 exposure, the rate increased with time. However, the rate decreased after 1 min as the Pd surface was gradually saturated with H2 and limited in supply of atomic hydrogen to the Pd/nanowire interface. The resistance of the nanowires was changed back to their initial values after switching from H2 to air. Similar to Pd -coated multiple InN nanobelt sensors, the recovery of resistance is most likely dominated by the removal of hydrogen atoms from the Pd/nanowire interface. The temperature dependence of resistance from Pd-coated multiple GaN nanowires exposed to 3000 ppm H2 in N2 is shown in Figure 5 11. The relative response saturated a bove a measurement temperature of 100 C, which could be due to the limitation in supply of atomic hydrogen. The rate of resistance change increased dramatically with time at 150C and was a factor of about 5 faster than that at room temperature. The inset figure shows the Arrhenius plot of rate of nanowire resistance change. Adsorption activation energy = 2.2 kcal mol was estimated from the standard Arrhenius equation. This value is smaller than that of multiple ZnO nanorods, confirming the higher sensit ivity and faster response of the GaN nanowire sensor. 5.3.3 Pt -functionalization A shadow mask was used to pattern non alloyed rf -sputtered Ti(50 nm)/Al (80 nm)/Pt (40 nm)/Au (300 nm) Ohmic contacts on the GaN nanowires with a contact separation of ~50 Au wires were bonded to the contact pads for device packaging. Finally, a ~7 nm thick functional layer of Pt was sputter -deposited onto the nanowire device. The GaN nanowire sensors were exposed to different H2 concentrations (2002000 ppm H2 in N2 ambi ent) at 25 -

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97 100 -coated GaN nanowire H2 sensors [162]. Figure 5 12 shows scanning electron micrographs (SEMs) of the nanowires with the added Pt -coating. The current -voltage characteristics from multiple nanowires were linear with a maximum power consumption of <0.3 mW under our operating conditions. Measured resistance at a bias of 0.1 V as a function of time is shown in Figure 5 13 A ) for Pt -coated multiple GaN nanowires exposed to varying H2 co ncentrations (2002000 ppm) in N2 for 10 min at room temperature. The measured resistance for uncoated multiple GaN nanowires exposed to 2000 ppm H2 concentration is shown for comparison. The addition of the functional Pt coating is critical to H2 detec tion for these sensors. Specifically, it is believed Pt, like Pd, works to catalytically dissociate H2 into molecular hydrogen. Diffusion of atomic hydrogen to the metal/GaN interface alters the surface charge by depletion and induces a current response at a fixed bias. Resistance change depends on the gas concentration and recovers to 80% of the initial level within 2 min. Resistance recovery is probably controlled by the removal of hydrogen atoms from the Pt/nanowire interface. The rate of resistance change of GaN nanowire sensors upon exposure to H2 is time dependent. Resistance decreases slowly in the first minute after H2 exposure, likely due to the Pt functional layer being covered with native oxide. As the native oxide was removed with continued H2 exposure, the rate increased with time. However, after ~2 minutes the rate decreased as the Pt surface became gradually saturated with H2 and limited the supply of atomic hydrogen to the Pt/nanowire interface. Previously, we reported that Pd -coated mu ltiple GaN nanowires showed high sensitivity to low H2 concentrations (~7.4% at 200 ppm). Uncoated nanowire sensors displayed little to no response at similar concentrations of H2. By comparison, Pt -coated GaN nanowires gave relative

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98 responses of ~1.7% a t 200 ppm, ~1.8% at 1200 ppm, and ~1.9% at 2000 ppm H2 in N2. As the maximum power consumption in either case is <0.6 mW, both Pt and Pd functionalized GaN nanowire sensors are suitable for long -term hydrogen sensing applications. Figure 5 13 B) shows th e time dependent relative response of Pt -coated multiple GaN nanowires at varying concentrations of H2 in N2 ambient. Pt -coated GaN nanowire sensors are less sensitive to hydrogen than the same nanowires with a Pd coating. Pd coatings have also shown to have higher sensitivity than similar Pt layers for multiple InN nanobelt H2 sensors. This is in contrast to results for ZnO nanowire H2 sensors. The relative response for Pt -coated GaN nanowires exposed to 1000 ppm H2 in N2 at a fixed bias of 0.1 V for 10 minutes measured over a temperature range (25100 Figure 5 14. Relative response increased with measurement temperature from 1.8% at room temperature up to 4% at 100 molecul ar hydrogen (H2) or diffusion rate of atomic hydrogen into the Pt/GaN interface. This response is at least 50% smaller for all temperatures in comparison to similar Pd -coated GaN nanowires. An extrapolation of the trend line suggests there would still be at least a 1.5% relative change in current at temperatures as low as 0 of GaN nanowire H2 sensors for low temperature applications. An Arrhenius plot of rate of nanowire resistance change was used to calculat e an adsorption activation energy of 7.3 kcal mol1, shown in Figure 5 15. This value is larger than that of previous Pd -coated multiple GaN nanowires (2.2 kcal mol1), indicative of the lower sensitivity and slower response of the Pt functional layer. T his difference between Pt and Pd coated nanomaterial sensor activation energy has also been noticed for multiple InN nanobelts while the opposite has been shown for multiple ZnO nanowires. This is suggestive of a material

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99 difference at the (Pd/Pt)/nanowir e interface between nitride and oxide based nanoscale sensors and will need more data to be understood. The inset figure displays the temperature dependence of resistance from 1000 ppm H2 exposure. The rate of resistance change is drastically increased with increasing temperature. Resistance change is a factor of about 10 faster at 100 supportive of an increase in atomic hydrogen reaching the Pt/GaN interface with increasing temperature. 5.4 Conclusions Both InN nanobelts and GaN nanowires benefitted from the addition of Pt or Pd functionalization layers for hydrogen sensing. Pd-coated nanostructures showed an improvement in sensitivity by up to a factor of ~10X larger than uncoated controls while Pt coatings on ly increased sensitivity by up to ~5X. The reason for this difference is unclear and may be indicative of nitride -based material issue. Calculated activation energies are all relatively similar in the hundredths of eV range and are all lower than previou s ZnO nanomaterial hydrogen sensors. Regardless, Pd or Pt -coated nanostructures are promising for hydrogen gas sensors with high sensitivity, low -power consumption and may be operable at considerably low temperatures.

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100 Figure 5 1 X ray diffraction spe ctrum of MOCVD grown InN nanobelts (the inset shows SEM images of the nanobelts).

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101 Figure 5 2 Current responses of InN nanobelt sensors. A) I -V characteristics of both uncoated and Pt -coated InN nanobelts in N2 ambient B) T he change in current from Pt coated InN nanobelts at different H2 concentration.

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102 Figure 5 3 Relative response of Pt -coated InN nanobelts exposed to a series of H2 concentrations (20300 ppm) in N2 ambient for 10 min at room temperature.

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103 Figure 5 4. Arrhenius plot of rate of resistance change (the inset shows the temperature dependence of resistance from Pt -coated multiple InN nanobelts exposed to 200 ppm H2).

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104 Figure 5 5. Current responses of InN nanobelt sensors. A) I -V characteristics of both uncoated and Pd -coated InN nanobelts in N2 ambient. B) The change in current from Pdcoated InN nanobelts at different H2 concentration measured at 130C. -1.0 -0.5 0.0 0.5 1.0 -40 -30 -20 -10 0 10 20 30 40 Current (mA)Voltage (V) InN nanobelts before coating after Pd-coating 0 200 400 600 800 1000 1200 0 2 4 6 8 10 Air H2 in N2 ambient |dR|/R (%)Time (sec) 100 PPM 300 PPM 500 PPM 1000 PPM Base line

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105 Figure 5 6. Relative response of Pd-coated InN nanobelts exposed to 300 ppm H2 after 10 minutes at different temperature. 20 40 60 80 100 120 140 4 5 6 7 8 9 10 |dR|/R (%)Temperature (C) 280 300 320 340 360 380 400 420 Temperature (K)

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106 Figure 5 7. Arrhenius plot of rate of resistance change (the inset shows the temperature dependence of resistance from Pd -coated multiple InN nanobelts exposed to 300 ppm H2).

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107 Figure 5 8. SEM images of as -grown GaN nanowires.

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108 Figure 5 9. Measu red resistance at 0.5 V bias as a function of time from Pd-coated and uncoated multiple GaN nanowires exposed to varying H2 concentrations in N2 ambient for 10 minutes at room temperature.

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109 Figure 5 10. Response of Pd-coated GaN nanowires to varying H2 concentration. A) Relative response with time. B) T ime dependent resistance change.

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110 Figure 5 11. Temperature dependence of resistance from Pd-coated multiple GaN nanowires exposed to 3000 ppm H2 in N2. Inset figure shows Arrhenius plot of resistanc e change for the highest change in resistance (i.e. from 1 20 seconds).

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111 Figure 5 12. SEM images of GaN nanowires after Pt coating

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112 Figure 5 13. Response of Pt -coated GaN nanowires to varying H2 concentration. A) Time dependant resistance measured at an applied bias of 0.1 V B) Relative response 0 5 10 15 20 25 30 40.8 41.0 41.2 41.4 41.6 41.8 52 54 No Pt-coating 2000 ppm H2 Resistance (Ohms)Time (min.) Pt-coated 200 ppm 400 ppm 1200 ppm 1600 ppm 2000 ppm 0 400 800 1200 1600 2000 0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4 1.6 1.8 2.0 |dR| / R (%)Time (sec.) 200 ppm 400 ppm 1200 ppm 1600 ppm 2000 ppm

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113 Figure 5 14. Relative response of P t -coated GaN nanowires exposed to 2000 ppm H2 after 10 minutes at different temperature. 20 30 40 50 60 70 80 90 100 110 1.5 2.0 2.5 3.0 3.5 4.0 Temperature (C)|dR| / R (%)

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114 Figure 5 15. Arrhenius plot of rate of resistance change (the inset shows the temperature dependence of resistance from P t -coated multiple GaN nanowires exposed to 1000 ppm H2).

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115 CHAPTER 6 CONCLUSION In this work, several processes were investigated to improve the base electrical characteristics of compound semiconductor de vices for n ZnO, GaN and InN. A series of refractory materials was deposited to improve thermal stability of Ohmic type contacts to bulk nZnO. Cryogenic temperature deposition of simple metal contacts was examined in order to improve Schottky electrical behavior for n -ZnO devices. Finally, the effect of Pt and Pd functionalization layers for chemical sensing was explored with GaN nanowires and InN nanobelts. The use of TiB2, ZrB2, and Ir as Ohmic contact materials for bulk n ZnO was studied in hopes of obtaining low specific contact resistance and high thermal stability. Both boride studies used an XB2/Pt/Au metallization scheme. In this case, Pt acted both as a diffusion barrier between the boride and Au and improved adhesion of Au to the lattice mism atched contact. Ir contacts did not use a Pt layer. Both TiB2and ZrB2-based contacts exhibited thermal stability well in excess of previous studies for typical contacts to bulk ZnO, although at a cost of higher specific contact resistance. This behavi or is suggestive of the influence of detrimental interfacial layers due to metallurgica l intermixing from annealing. TiB2 and ZrB2 both showed rectifying behavior before high temperature annealing. Minimum specific contact resistance for both borides was = ~104 2. Meanwhile, Ir -based contacts annealed in Nitrogen were stable up to 1000 resistance in the range of 106 2. Unlike the boride contacts, there was no evidence of a thin Iridium oxide layer or dissoc iation of ZnO at the interface, likely ruling out Ohmic conduction by the formation of oxygen vacancies or by a heavily doped near -surface region, respectively.

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116 The effect of cryogenic deposition on Schottky barrier height for Au, Ni, Ti, Pt, and Pd contacts to b ulk n ZnO was examined. Ni and Ti samples deposited at both 77K and room temperature show Ohmic behavior. Pt and Pd deposited contacts gave rectifying behavior at both temperatures. Au exhibited rectifying behavior after deposition at 77K, while room temperature deposition resulted in Ohmic conductivity. Barrier heights for Au, Pt and Pd were all = ~0.4 eV. Annealing (up to 300 showed little effect on the barrier heights of Pt and Au. In near ly all cases, the ideality factor for cryogenically deposited contacts was lower than for room temperature deposited contacts, suggesting a more abrupt metal -semiconductor interface for low temperature deposited contacts. A thin -layer interfacial oxide ma y also be present in cryogenically deposited contacts, although the effect of this oxide is a matter of scrutiny. The exact mechanism(s) for barrier height improvement with low temperature deposition is still under consideration and further study will be needed for a complete explanation. The sensitivity for detecting hydrogen using multiple GaN nanowires and InN nanobelts with cluster coating of Pt or Pd on the surface was investigated with different hydrogen concentrations. Both Pt and Pd coatings showe d dramatic improvement with hydrogen detection over their uncoated counterparts for GaN and InN nanostructures. While both coatings increased sensitivity to hydrogen, Pd coatings worked best in all trials, sometimes with sensitivity more than 50% higher t han for Pt under identical conditions. It was noted this result is in contrast to previous data for ZnO hydrogen nano -sensors. In either case, the mechanism for hydrogen gas sensing is believed to arise from the Pt/Pd catalytic dissociation of H2 into at omic hydrogen. Hydrogen sensitivity improved with increasing measurement temperature, up to ~150 contact degradation led to device failure. Our results show metal -coated GaN nanowires and

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117 InN nanobelts are promising for hydrogen gas sensors with fast response, high sensitivity, and low -power consumption. Additionally, because of the marked improvement in sensitivity from metal -coatings, these sensors may be applicable for low temperature spacecraft devices.

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128 BIOGRAPHICAL SKETCH Jonathan Wright was born in West Lafayette, Indiana in 1983. He grew in Midland, Michigan, graduating from Midland High School in 2001. In June 2005, Jonathan earned a Bachelor of Science in m aterials s ci ence and e ngineering and a minor in English from t he Ohio State University. While at Ohio State, Jonathan was involved in undergraduate superconductivity research with Prof. Mike Sumption and steel process engineering technology with Prof. Yogesh Sahai. A fter completing his undergraduate education, Jonathan completed his doctorate study in m aterials s cience and e ngineering at the University of Florida under the direction and guidance of Prof. Stephen J. Pearton. His research there focused in large part on process engineering for novel compound semiconductors, particularly GaN and ZnO. Jonathan worked on internship during summer 2007 at NASA Glenn Research Center in Cleveland, Ohio under the mentorship of Principle Research Scientists Gus Fralick and John Wrbanek Jonathan completed a second internship during summer 2008 at Sandia National Laboratories in Albuquerque, New Mexico under the mentorship of Dr. Randy J. Shul.