Nanowire-Based Light-Emitting Diodes: a New Path Towards High-Speed Visible Light Communication." (2017)

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Fall 9-30-2017 NANOWIRE-BASED LIGHT-EMITTING DIODES: A NEW PATH TOWARDS HIGH- SPEED VISIBLE LIGHT COMMUNICATION Mohsen Nami University of New Mexico

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This Dissertation is brought to you for free and open access by the Electronic Theses and Dissertations at UNM Digital Repository. It has been accepted for inclusion in Physics & Astronomy ETDs by an authorized administrator of UNM Digital Repository. For more information, please contact [email protected]. Mohsen Nami Candidate

Physics and Astronomy Department

This dissertation is approved, and it is acceptable in quality and form for publication:

Approved by the Dissertation Committee:

Professor: Daniel. F. Feezell, Chairperson

Professor: Steven. R. J. Brueck

Professor: Igal Brener

Professor: Sang. M. Han

NANOWIRE-BASED LIGHT-EMITTING DIODES: A NEW PATH TOWARDS HIGH-SPEED VISIBLE LIGHT COMMUNICATION

MOHSEN NAMI

B.S., Physics, University of Zanjan, Zanjan, Iran, 2003 M. Sc., Photonics, Shahid Beheshti University, Tehran, Iran, 2006 M.S., Optical Science Engineering, University of New Mexico, Albuquerque, USA, 2012

DISSERTATION

Submitted in Partial Fulfillment of the Requirements for the Degree of

Doctor of Philosophy Engineering

The University of New Mexico Albuquerque, New Mexico

December, 2017

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Dedication

To my parents, my wife, and my daughter for their endless love, support, and encouragement.

Acknowledgments

I would like to thank all the people who contributed in some way to the work described in this dissertation. I would not be able to finish my dissertation without the guidance, help and support from my adviser, family, and friends. I would like to express my deepest gratitude to my advisor,

Professor Daniel. F. Feezell, for his excellent guidance, caring, patience, and providing me with an excellent atmosphere for doing research. Without him this work would not have been possible. Additionally, I would like to thank my committee members Professor Steven R.J.

Brueck, Professor Sang M. Han, and Professor Igal Brener for their interest in my work and constructive comments. I would like to thank Professor Luke. F. Lester for giving me the honor and opportunity to work in his group and changing my passion toward perusing science and research. I would like to thank Dr. Ashwin. K. Rishinaramangalam for his help and support. I gained a lot from his vast knowledge in epitaxial growth. I would also like to thank Dr. Benjamin

Bryant and Mr. Rhett Ellet for their help in E-Beam lithography. Special thanks go to Dr. Serdal

Okur, and Mr. Isaac Stricklin for their help in micro-photoluminescence measurement. I would like to thank Mr. Mike Fairchild for the timely maintenance of the reactor to enable epitaxial growth using our MOCVD. I would also like to thank Mr. Steve Wawrzyniec and Mr. Douglas

Wozniak for timely maintenance of the cleanroom equipment and also constant source of discussions during process development. Many thanks to my labmates Mr. Saadat Mishkat-Ul-

Masabih, Mr. Arman Rashidi, Mr. Kenneth DaVico, Mr. Andrew Aragon, and Dr. Morteza

Monavarian.

I would like to thank my parents. They were always supporting me and encouraging me with their best wishes. Last but not least, I would like to thank my wife, Zoha. She was always there cheering me up and stood by me through the good times and bad.

NANOWIRE-BASED LIGHT-EMITTING DIODES: A NEW PATH TOWARDS HIGH-SPEED VISIBLE LIGHT COMMUNICATION

MOHSEN NAMI

B.S., Physics, University of Zanjan, Zanjan, Iran, 2002 M. Sc., Photonics, Shahid Beheshti University, Tehran, Iran, 2006 M.S., Optical Science Engineering, University of New Mexico, Albuquerque, USA, 2012

ABSTRACT

Nano-scale optoelectronic devices have gained significant attention in recent years. Among these devices are semiconductor nanowires, whose dimeters range from 100 to 200 nm. Semiconductor nanowires can be utilized in many different applications including light-emitting diodes and laser diodes. Higher surface to volume ratio makes nanowire-based structures potential candidates for the next generation of photodetectors, sensors, and solar cells. Core-shell light-emitting diodes based on selective-area growth of gallium nitride (GaN) nanowires provide a wide range of advantages. Among these advantages are access to non-polar m-plane sidewalls, higher active region area compared to conventional planar structures, and reduction of threading dislocation density.

In this work, metal organic chemical vapor deposition (MOCVD) was employed to grow

GaN nanowires. GaN nanowires were grown selectively using a dielectric mask. The effect of mask geometry and growth conditions on the morphology and dimensions of the GaN nanowires were studied. It is shown that the pitch spacing between each nanowire has a significant effect on

vi the geometry of the GaN nanowires. Growth of quantum wells around the GaN nanowires was also investigated. A feasible approach towards monolithically integrated multi-color LEDs was presented, which is not possible using conventional planar LEDs. The nanowire approach offers a potential way to overcome the current issue of pick-and-place for micro-LED displays, fabricated using planar LEDs. Core-shell nanowire-based LEDs were successfully fabricated and the opto-electrical characteristics of the LEDs were investigated. An equivalent circuit model for the LED was proposed, and a thorough investigation of RF characterizations of the LEDs was presented. This work paves the way to understand the limiting factor for 3-dB modulation bandwidth in nanowire LEDs. Finally, different types of plasmonic nanostructures were proposed to improve the 3-dB modulation bandwidth.

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TABLE OF CONTENTS

LIST OF FIGURES ...... xi LIST OF TABLES ...... xxi Chapter 1 Introduction...... 1 1.1 The properties, advantages and a brief history of III-nitride materials ...... 1 1.2 Crystallographic planes of GaN and their properties ...... 4 1.3 Growth of wurtzite III-nitride materials and origin of threading dislocations ...... 6 1.4 Motivation for the development of nanostructure-based LEDs ...... 8 1.5 A brief history of SAG of GaN/InGaN core-shell nanowire-based LEDs ...... 13 1.6 Motivation for selective area growth of GaN nanowires and advantages ...... 16 1.7 Scope of this dissertation ...... 17 References ...... 19 Chapter 2 Selective-area growth using MOCVD ...... 30 2.1 MOCVD system Veeco-p75 ...... 30 2.2 Processing steps required for nanowire template preparation using electron-beam lithography ...... 31 2.3 Selective-area growth of GaN nanowires ...... 34 2.4 Effect of aperture diameter and the role of capture length ...... 38 2.5 Effect of pitch size on nanowire morphology and the role of growth temperature ...... 40 2.6 Effect of the growth temperature on nanowire dimension and morphology ...... 42 2.7 Study the effect of the TMGa injection time on nanowire dimension ...... 46 2.8 Conclusion ...... 49 References ...... 51

Chapter 3 Growth of GaN/InGaN core-shell nanowires ...... 54 3.1 Effect of the aperture diameter on photoluminescence of GaN/InGaN core-shell nanowires ...... 54 3.2 Study the effect of the pitch size on photoluminescence of GaN/InGaN core-shell nanowires ...... 60 3.3 Study of temperature dependent of photoluminescence and carrier dynamics of GaN/InGaN core-shell nanowires for different pitch sizes ...... 66

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3.3.1 Study of carrier dynamics ...... 71 3.4 Conclusion ...... 75 References ...... 77

Chapter 4 Growth, fabrication, and characterization of electrically injected GaN/InGaN core-shell nanowire light-emitting diodes ...... 80 4.1 Introduction ...... 80 4.2 Growth of GaN/InGaN core-shell nanowire-based LEDs ...... 81 4.3 Electrical characterization and the effect of AlGaN underlayer on device performance ...... 85 4.3.1 Growth of AlGaN underlayer on GaN nanowires ...... 94 4.4 Growth characterization ...... 100 4.5 Fabrication of GaN/InGaN core-shell nanowire light-emitting diodes ...... 103 4.6 Optical characterization ...... 109 4.7 Far field emission pattern ...... 116 4.8 Electrical characterization ...... 118 4.9 RF characterization ...... 122 4.9.1 Introduction ...... 122 4.9.2 RF characteristics of GaN/InGaN core-shell nanowire-based LEDs ...... 123 4.10 Conclusion ...... 131 References ...... 133

Chapter 5 Simulation and design of plasmonic GaN/InGaN core-shell nanowire-based LEDs ...... 140 5.1 Introduction ...... 140 5.2 Purcell effect ...... 141 5.3 Simulation method ...... 144 5.4 Device structures and simulation results ...... 146 5.4.1 Axial structure ...... 146 5.4.2 Core–shell structure with hexagonal cross-section ...... 147 5.4.2.1 Near-field intensity ...... 149 5.4.3 Core–shell structure with hexagonal cross-section and a tapered base ...... 151 5.4.3.1 Near-field intensity ...... 152

5.4.4 Flip-chip core–shell structure with hexagonal cross-section and a tapered base ...... 153 5.4.4.1 Effect of n-GaN thickness ...... 157 5.4.4.2 Embedded dielectric structure ...... 158 5.4.4.3 Near-field intensity ...... 159 5.5 Modulation characteristics ...... 161 5.6 Conclusion ...... 163 References ...... 164

Chapter 6 Conclusions and future work ...... 170 6.1 Conclusions ...... 170 6.2 Future work ...... 173 Appendix A: Processing steps required for fabrication ...... 177

Publications ...... 180

LIST OF FIGURES

Figure 1.1 Mapping of lattice constant vs. band gap of III-nitride semiconductors……………...1

Figure 1.2 Luminous efficiency as a function of operational wavelength, showing a gap in luminous efficiency ranging from 500 nm to 620 nm. (Source OSRAM (chapter 1, reference 23)………………………………………………….…3

Figure 1.3 Schematic of the unit cell for wurtzite GaN showing the low index planes including a-plane, r-plane c-plane, and m-plane………………………………….…...5

Figure 1.4 Simulated energy band diagrams for 3 nm In0.23Ga0.77N for single quantum wells with GaN barrier at a current density of 100 A/cm2 for (a) polar (c- plane), (b) nonpolar (101̅0) (m-plane). The direction of the piezoelectric field and p-n junction built-in field are indicated. The ground-state electron and hole wave functions are also shown. The n-side donor concentration is 1e18 cm-3 and p-side acceptor concentration is 2e19 cm- 3. (Data from chapter 1, reference 38) …………………………………………....……6

Figure 1.5 TEM micrograph of the wurtzite c-plane GaN grown on c-plane sapphire, showing the threading defects originating at the interface between GaN and Sapphire (chapter 1, reference 40)……………………..…………………..…….8

Figure 1.6 XTEM image of a GaN nanowire (a) showing the bending of the threading o dislocation by 90 and termination at the SiNx growth mask. (b) showing the 90o bending of a threading dislocation toward the nonpolar side wall,

and filtering out the threading dislocations by SiNx dielectric growth mask. (chapter 1, reference 40)…….….……………………………..………10

Figure 1.7 The enhancement in the active region area as a function of the nanowire radius with different heights compare to a planar LED on the same substrate footprint area………………………………………………………………11

Figure 1.8 Increase in the number of publications on nanowire related topics from 1996-2016 (Source, ISI; keyword, nanowires). ………………………..……………13

Figure 1.9 A brief history of IQE reported for SAG of GaN/InGaN core-shell nanowire-based LEDs………….……………………………………..………………15

Figure 2.1 (a) Veeco-p75 MOCVD reactor, (b) load lock and transfer arm, (c) and growth chamber…………………………………….……………..…………………31

Figure 2.2 Processing steps for nanowire template using e-beam lithography. (1) 2 µm

GaN template grown on Sapphire, (2) Deposition of 25 nm of SiNx using

plasma-enhanced chemical vapor deposition, (3) Deposition of 90 nm PMMA and 10 nm gold, and e-beam patterning, (4) Gold removal, (5) Develop in cold MIBK:IPA 1:3, (6) Reactive ion etching, (7) Piranha cleaning, (8) E-beam nanowire template….……….………..…33

Figure 2.3 E-beam patterning (a) square shape using higher does, (b) and hexagonal shape using lower dose. Both result in circular aperture………………..………….34

Figure 2.4 PMMA remnant after piranha cleaning. This random pattern, results in a random nanowire growth region with random geometry…….………….…………35

Figure 2.5 E-beam nanowire sample after piranha cleaning under optimized conditions………..…………………………………………….………..…………..35

Figure 2.6 Schematic image of different steps in each cycle for pulsed-mode growth technique, (a) TMGa injection time, (b) TMGa interruption time, (c)

NH3 injection time, (d) and NH3 interruption time………….………………………36

Figure 2.7 Schematic illustration of the processes required to grow GaN nanowires using MOCVD pulsed mode technique, (a) TMGa injection step, (b)

TMGa interruption step, (c) NH3 injection step, (d) and NH3 interruption step………………………….………………………………………………………..36

Figure 2.8 SEM images, at a 45-degree tilt, of GaN nanowires for 1 µm pitch and different aperture diameters (a) 60 nm, (b) 110 nm, (c) 160 nm, (d) and 210 nm…….…….………………………………………..…………….38

Figure 2.9 (a) Average diameter and average height of nanowires as a function of aperture diameter, (b) and average nanowire volume as a function of aperture radius (raw data) and the fitting curve used to determine the capture length…………………….…………………………….…………………….39

Figure 2.10 SEM images, at a 45-degree tilt, containing GaN nanowires showing the differences in growth morphology for pitch (a) 300 nm (b), 500 nm, (c) 1 μm, (d) 2 μm, (e) 3 μm, (f) and 4 μm.………………………………...……………40

Figure 2.11 SEM images, at a 45-degree tilt, of GaN nanowires for 500 nm pitch, showing the differences in growth morphology and nanowire diameter, under different growth temperatures (a) 917oC, (b) 923oC, (c) 930oC. The nanowire morphology evolution as a function of growth temperature is shown in the inset figures...... 42

Figure 2.12 (a) Average nanowire diameter for different pitches at different growth temperatures, (b) lateral overgrowth for different pitches at different

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growth temperatures. The inset images in (b) show the dependence of nanowire morphology on temperature for those grown with a 3 µm pitch, (c) average height of the nanowires for different pitches at different growth temperatures, and (d) aspect ratio for nanowires grown at different temperatures for different pitches….…………..……………………….44

Figure 2.13 Capture length versus pitch at a growth temperature of 926oC …….………………46

Figure 2.14 The difference in diameter of the nanowires for a TMGa injection time of 18 sec and that of 12 sec………………………………………………..……………47

Figure 2.15 (a) Average diameter and average height of nanowires as a function of aperture diameter, (b) average nanowire volume as a function of aperture radius (raw data) and the fitting curve used to determine the capture length………………………………………………………..………………48

Figure 2.16 Average diameter of the nanowire as a function of pitch sizes for different TMGa injection times and different TMGa flow rat, keeping the total amount of flow the same……………………………………………..………………49

Figure 2.17 Average diameter of the nanowire as a function of pitch sizes for different TMGa interruption times…………………………………………….………………50

Figure 3.1 SEM images, at a 45-degree tilt, of GaN nanowires for 1 µm pitch and different aperture diameters (a) 60 nm, (b) 110 nm, (c) 160 nm, (d) 210 nm, before growth of a single quantum well and barrier………..…………………55

Figure 3.2 SEM images, at a 45-degree tilt, of samples of GaN nanowires for 1 µm pitch and different aperture diameters (a) 60 nm, (b) 110 nm, (c) 160 nm, (d) 210 nm, after growth of a single quantum well and barrier……………....56

Figure 3.3 Micro-photoluminescence set-up…………….…………..…...……………………..57

Figure 3.4 Normalized PL spectra for different aperture diameters……….……………………58

Figure 3.5 Estimated thickness of one period, including one QW and barrier……………….…58

Figure 3.6 (a) PL peak wavelength versus aperture diameter and simulated PL peak wavelength versus aperture diameter for constant 27% indium composition (b) Change in PL peak wavelength and estimated indium composition versus aperture diameter. The pitch spacing is 1 m………………….60

Figure 3.7 SEM images of a sample of GaN nanowires showing the differences in growth morphology for pitch = 300 nm (a), 500 nm (b), 1 μm (c), 2 μm (d), 3 μm (e), and 4 μm (f), before QW/barrier growth……….…………..…………61

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Figure 3.8 SEM images of GaN/InGaN core-shell nanowires for pitch = 300 nm (a), 500 nm (b), 1 μm (c), 2 μm (d), 3 μm (e), and 4 μm (f), after QW/barrier growth………………………………….…………………………....62

Figure 3.9 Normalized PL spectra for different pitches………………………...... ……62

Figure 3.10 False colored SEM of arrays of GaN/InGaN core-shell nanowires………….….…63

Figure 3.11 Estimated thickness of one period, including one QW and barrier…...... …64

Figure 3.12 (a) PL peak wavelength versus pitch and simulated PL peak wavelength versus pitch for constant 22% indium composition, (b) Change in PL peak wavelength and estimated change in indium composition versus pitch. The aperture diameter is 40 nm……………………………………..………65

Figure 3.13 Integrated photoluminescence of GaN/InGaN core-shell nanowires for (a) 300 nm (b) 1 µm, (c) and 3 µm pitch size…………...………………..……………67

Figure 3.14 Measured PL peak intensity (normalized at 10 K) versus temperature for (a) an InGaN single QW sample, (b) an InGaN QDs sample (chapter 3, reference 9)………….…………………………………………...…………………68

Figure 3.15 Schematic diagrams indicating the possible mechanism of temperature induced carriers’ redistribution inside shallow and deep localized states inside the active region, (a) 10 K, (b) 25 K, (c) 150 K, (d) and 200 K…….……..…70

Figure 3.16 Time-Resolved Photoluminescence set-up…………....…………..…………..……72

Figure 3.17 Few examples of the room temperature PL transients for excitation powers Of 0.1, 10, and 100 mW for 1 µm pitch size…………………………………...……73

Figure 3.18 PL lifetime as a function of reciprocal temperature for (a) 300 nm, (b) 1 µm, (c) and 3 µm pitch size….………………………………………………..……..74

Figure 3.19 (a) PL lifetime , (b) Potential 3dB modulation bandwidth as a function of excitation power for 300 nm, 500 nm, 1, 2, 3, and 4 µm Pitch sizes……………….………………………….………………………..…….………75

Figure 4.1 Processing steps for nanowire template using interferometric lithography. (1) 2 µm GaN template grown on Sapphire, (2) Deposition of 100 nm of

SiNx using plasma-enhanced chemical vapor deposition, (3) Spin coat of ICON-7 and photoresist, (4) interferometric lithography, (5) Reverse the pattern using Ni deposition and lift-off, (6) Secondary lithography to

define the mesa, and RIE of ARC and SiNx,(7) Piranha cleaning, (8) Growth of GaN nanowires, (9) Microscopic image of nanowire template

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ready for growth.………….…………………………………………………………83

Figure 4.2 Interferometric lithography set-up including laser (Coherent Infinity 40- 100), beam expander, mirror, and stage…………………………..…………………83

Figure 4.3 Schematic of the IL set-up and two different patterning. 1-D patterning results in formation of lines and 2-D patterning results in formation of circular apertures….………………..…………………………….…………………84

Figure 4.4 SEM images of GaN nanowires (a) before (b) after the growth of the quantum wells, barriers, and p-GaN…………..……………………………………..85

Figure 4.5 Room temperature current density-voltage plot of a GaN/InGaN core-shell nanowire light emitting diode under continuous condition. The inset shows J-V plot from -5 V to 0 V…………………………………………………………….86

Figure 4.6 EL spectrum at (a) 10 A/cm2, (b) 75 A/cm2………………………………………….86

Figure 4.7 SEM images of pyramidal nanostripes (a) before (b) after growth of quantum well and p-GaN. ……………………………………………………………88

Figure 4.8 Room temperature current density-voltage plot of a triangular nanostripe light emitting diode under continuous condition……....………………..……………88

Figure 4.9 SIMS proﬁle of the top p-GaN layer of a triangular nanostripe LED vs. depth. The inset shows cross section of triangular nanostripes………………………89

Figure 4.10 (a) SEM image of triangular nanostripe GaN cores, (b) SEM image of triangular naostripe LED with AlGaN underlayer. The inset shows an XTEM image indicating the position of the AlGaN underlayer……………….……91

Figure 4.11 SIMS proﬁle of the top p-GaN layer of a triangular nanostripe LED vs. depth after the growth of n-type AlGaN underlayer. The inset shows cross section of triangular nanostripes………………………………….…………………92

Figure 4.12 Reverse-leakage current density at -5 V measured vs. AlGaN underlayer growth time. The inset shows schematic of triangular nanostripes with and without electron blocking layer (EBL)……….……………………...………………93

Figure 4.13 Reverse-leakage current density at -5 V measured vs. TMAl flow rate……………94

Figure 4.14 Room temperature current density-voltage plot of a triangular nanostripe Light-emitting diode under continuous condition………….……...………………….95

Figure 4.15 SEM images of GaN nanowires (a) before, (b) after the growth of AlGaN underlayer, using continuous mode growth technique………………………………97

Figure 4.16 Room temperature current density-voltage (J-V) plot of a GaN/InGaN core-shell nanowire light emitting diode with AlGaN underlayer grown in continuous mode. The measurement was performed under CW mode. The inset shows J-V plot from -5 V to 0 V…….………………………97

Figure 4.17 Schematic image of different steps in each cycle for pulsed-mode growth of AlGaN underlayer, (a) TMGa and TMAl injection time, (b) TMGa and

TMAl interruption time, (c) NH3 injection time, (d) NH3 interruption time……………………………………………………….…..……….98

Figure 4.18 SEM image of GaN nanowires (a) before, (b) after the growth of AlGaN underlayer using pulsed mode growth technique. ………………….………………98

Figure 4.19 (a) SEM image of GaN/InGaN core-shell nanowire-based LEDs, (b) Schematic image of the device showing sapphire substrate, GaN.

template, SiNx dielectric layer, AlGaN underlayer, MQW, and p-GaN……..…..…99

Figure 4.20 Room temperature current density-voltage (J-V) plot of a GaN/InGaN core-shell nanowire light emitting diode with AlGaN underlayer grown in pulsed mode. The measurement was performed under CW Mode. The inset shows J-V plot from -5 V to 0 V………..…..……………………99

Figure 4.21 STEM images of GaN/InGaN core-shell nanowires in atomic-contrast (ZC) mode showing (a) GaN nanowires with diameter of 600 nm and height of 1.2 µm. (b) Conformal growth of AlGaN around the nanowires. (c) Growth of AlGaN at the top portion of the nanowires. (d, e) Growth of AlGaN layer at the sidewalls of the nanowires. (f) Growth of AlGaN on semipolar facet……………………………………………...…………………101

Figure 4.22 STEM images of a GaN/InGaN core-shell nanowire LED in atomic- contrast (ZC) mode showing AlGaN underlayer, quantum wells, barriers, and EBL (a) at the top, (b) middle, (c) bottom of the GaN/InGaN core- shell nanowire……………………….………..………..…………………………102

Figure 4.23 SEM images of poor quality etched facets using ICP etching……………...…….103

Figure 4.24 SEM images of etched facets using ICP etching under optimized Conditions………………………………………………………...………………104

Figure 4.25 (a) Schematic of the top view of the GaN/InGaN LEDs showing the mesa and SEM image of the grown region, (b) Schematic image of the GaN/InGaN core-shell nanowire LED after ITO deposition……………………..106

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Figure 4.26 Schematic image of the GaN/InGaN core-shell nanowire-based LED (a) after ICP etching,(b) after annealing (c) Schematic image of the top view of the devices showing the mesa and ICP etched region, (d) SEM images of mesa and etched ITO region…………..………………..………………106

Figure 4.27 (a) Schematic image of the LEDs viewed from the top, showing the n- contact, mesa and ITO region (b) cross section of the LED showing the n-contact deposited region, (c) to (d) SEM images of a nanowire-based LED showing the n-contact (yellow color)…………………………107

Figure 4.28 (a) Schematic image of the full fabricated devices viewed from the top, (b) Schematic of the cross section of the LED, (c) and (d) SEM images of fully fabricated devices showing the n-contact, n-pads and p-pads……...………108

Figure 4.29 Schematic images of (a) device cross section and (b) device layout for the GaN/InGaN core-shell nanowire-based LEDs. ……………………………………108

Figure 4.30 (a) PL spectra at different excitation powers (b) PL peak wavelength at different excitation powers (c) PL FWHM at different excitation powers. ….…………………………………...……………………………………110

Figure 4.31 Optical characterization of GaN/InGaN core-shell LED (a) IQE measurement using low temperature/room temperature integrated PL technique (b) Integrated PL vs. excitation power in log-log scale at low temperature. (c) Integrated PL vs. excitation power in log-log scale at room temperature………………..…………...……………………………………112

Figure 4.32 Optical characterization of GaN/InGaN core-shell LED (a) PL lifetime measured at room temperature and low temperature. (b) Radiative lifetime at room temperature. (c) Nonradiative lifetime at room temperature………………………….………..……………………………………115

Figure 4.33 Simulation of a GaN/InGaN core-shell nanowire-based LED. (a) Schematic of the simulated nanowires viewed from the top, and the simulation region. (b) Schematic of the cross section of the nanowire- based LEDs, showing different layers, top monitor to calculate extraction efficiency, and simulation region………………….…..…………………………116

Figure 4.34 (a) Simulated angular distribution of extracted light for a nanowire-based LED along (a) x-axis (b) y-axis……………………………...……………………118

Figure 4.35 (a) J-V plot. Inset image shows an operating LED with 120 µm by 120 µm mesa size (b) L-J-V curve for a device with 120 µm by 120 µm mesa size (c) Estimated EQE…….………………………………………………………120

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Figure 4.36 (a) EL spectra, (b) EL peak wavelength, (c) and FWHM at different current densities……………………………..………………………...…..……121

Figure 4.37 United States frequency allocations in radio spectrum, (reported by NTIA 2016)………………………………………………………………………………123

Figure 4.38 (a) Schematic illustration of the RF set-up, (b) to (d) experimental set-up for the RF measurement. Inset in (d) shows the RF probe, operating LED, and optical fiber………………….…………………………………………124

Figure 4.39 (a) J-V plot for different mesa sizes, the inset show microscopic image of four different mesa sizes, and operating LED with 60 µm mesa size, (b) 3-dB modulation bandwidth for four different mesa sizes………………………125

Figure 4.40 LEDs with different mesa sizes operating in pulsed mode (1% duty cycle with 1 µs pulsed width) at different current densities, showing current crowding form 120 µm mesa size…………………………….…………………..126

Figure 4.41 An arbitrary two-port microwave network system……..………………………….126

Figure 4.42 An arbitrary two-port microwave network system described by ABCD matrix elements……………………………………………………………………127

Figure 4.43 The small-signal equivalent circuit of the GaN/InGaN nanowire-based LED……………………………………………………………………………….128

Figure 4.44 Equivalent two-port network system for the GaN/InGaN nanowire-based LED………………………………………………………………………………129

Figure 4.45 Simultaneous fitting of (a) real and imaginary parts of Eq. (4.11) to experimental data, (b) 20×log (abs(Eq. (4.12)) to the measured modulation response of a 100 µm mesa size device at a current density of 100 A/cm2……………..….……………………………………………………130

Figure 4.46 Carrier recombination time constant and diode time constant………...…………..131

Figure 5.1 Maximum frequency (3-dB modulation bandwidth) as a function of current density for different structures including (left to right) a typical LED (chapter 5, reference 23), a plasmonic LED (SPED) (chapter 5, reference 4), a vertical-cavity surface-emitting laser (VCSEL) (chapter 5, reference 24), and a plasmonic nanolaser (SPASER) (chapter 5, reference 25)……………..………………………………………………………..141

Figure 5.2 Illustration of Purcell enhancement of the spontaneous emission in a cavity. Spontaneous emission rate can be modified by putting a dipole emitter

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inside a cavity………………………………………………..……………………142

Figure 5.3 (a) Contour plot of optimal 3-dB bandwidth for various nanoLED cavities with different modal volumes and quality factors, (b) Optimum 3-dB bandwidth and Q versus cavity volume.(data from (chapter 5, reference 22)………………………………………………………………………….……..144

Figure 5.4 Axial structure: (a) schematic of an axial nanowire structure. (b) Cross- section of the axial nanowire structures as viewed from the top, showing

the location of the p–i–n diode core and SiNx and silver claddings…………………………………………………………………………146

Figure 5.5 Axial structure: (a) calculated Purcell factor versus wavelength. (b) Purcell factor and EXE at a wavelength of 405 nm. (c) Normalized 2D map of the imaginary part of the electric field in the direction of the dipole moment for two dipoles placed at X = ±6 nm……………….…………………...147

Figure 5.6 Core–shell nanowire structure with a hexagonal cross-section: (a) schematic of a core–shell nanowire structure with hexagonal cross- section. (b) Cross-section of the core–shell nanowire structure as viewed from the top, showing the location of the n-GaN core, InGaN active region, p-GaN shell, and silver cladding layer…………………...………..148

Figure 5.7 Core–shell structure: (a) calculated Purcell factor versus wavelength (b) Purcell factor and EXE at a wavelength of 445 nm. (c) 2D map of the normalized imaginary part of the electric field in the direction of the dipole moment for a dipole placed at Y=25 nm…………………………………149

Figure 5.8 Near-field intensity for core–shell structure without a tapered base: (a) schematic of the core–shell nanowire structure with hexagonal cross- section (b) cross-section of the core–shell nanowire structures as viewed from the top. (c) Cross-sectional near-field intensity for a monitor placed at Y = 42 nm. (d) Cross-sectional near-field intensity at XY plane. (e) Cross-sectional near-field intensity at YZ plane…………………..150

Figure 5.9 Core–shell structure with a tapered base: (a) schematic of a core–shell nanowire structure with hexagonal cross-section and a tapered base. (b) Cross-section of the core–shell nanowire structures as viewed from the top, showing the location of the n-GaN core, InGaN active region, p- GaN shell, and silver cladding…………………………………………………….151

xix

Figure 5.10 Core–shell structure with a tapered base: (a) calculated Purcell factor versus wavelength. (b) Purcell factor and EXE at a wavelength of 448 nm. (c) 2D map of the normalized imaginary part of the electric field in the direction of the dipole moment for a dipole placed at Y=25 nm…………....152

Figure 5.11 Near-field intensity for core–shell structure with a tapered base: (a) schematic of the core–shell nanowire structure with hexagonal cross- section and a tapered base. (b) Cross-section of the core–shell nanowire structures as viewed from the top. (c) Cross-sectional near-field intensity for a monitor placed at Y = 42 nm. (d) Cross-sectional near-field intensity at XY plane. (e) Cross-sectional near-field intensity at YZ plane………………………………………………………………………....153

Figure 5.12 (a) Schematic of a core-shell nanowire structure with a hexagonal cross- section and a tapered base. The taper has a 40 nm width at the base and a 14 nm width at the top. (b) Cross-section of the core-shell nanowire structure as viewed from the top, showing the location of the n-GaN core, InGaN active region, p-GaN shell, and silver cladding……………………154

Figure 5.13 Concept of a flip-chip PNLED fabrication process using photoelectrochemical etching to remove the substrate and expose the backside of the NW emitter. Such a structure would allow for fiber coupling or more efficient light extraction into free space than a structure with the substrate still incorporated.Figure 5.14 An arbitrary two-port microwave network system described by ABCD matrix elements………………………………………………………………………….155

Figure 5.14 (a) Calculated Purcell factor vs. wavelength for dipole emitters placed at various vertical locations along the active region. (b) Calculated Purcell factor and EXE into air at 445 nm for dipole emitters placed at various vertical locations along the active region from 5 nm to 45 nm. (c) 2D map of the normalized imaginary part of the electric field in the direction of the dipole moment for a dipole placed at Y=25 nm….……………...156

Figure 5.15 EXE into air at 445 nm for various thicknesses of n-GaN…………………….…..158

Figure 5.16 (a) Schematic of a core-shell nanowire structure with a hexagonal cross- section and a tapered base. A dielectric layer was added in order to increase the EXE. (b) Cross-section of the core-shell nanowire structure as viewed from the top, showing the location of the n-GaN core, InGaN

active region, p-GaN shell, SiNx layer, and silver layer……….…………………159

Figure 5.17 Calculated Purcell factor and EXE into air at 445 nm for a dipole emitter

placed at Y = 25 nm. The thickness of the SiNx layer (s) was swept from 0 to 40 nm and TAg1 is (a) 30 nm and (b) 20 nm……….………………….159

Figure 5.18 (a) Top-view cross-section of the core-shell nanowire structure shown in Fig. 1, the associated coordinate system, and the locations of cross- sectional planes. (b) Cross-sectional near-field intensity within the XZ plane (c) Cross-sectional near-field intensity within the XY plane. (d) Cross-sectional near-field intensity within the YZ plane…….…………………..160

Figure 6.1 (a) SEM image of GaN/InGaN core-shell nanowires (b) zoom in image……….....174

Figure 6.2. Room temperature PL transients a GaN/InGaN core-shell nanowire sample with and without silver nanoparticles………………………………………….....175

Figure 6.3. Room temperature PL transients a GaN/InGaN core-shell nanowire sample with and without silver nanoparticles…………….……………………...………...176

LIST OF TABLES Table 1.1 Comparison of three common substrates used to grow c-plane wurtzite GaN…………………………………………………………………………..7

Table 5.1 Parameters used to find the 3dB modulation bandwidth of the structures discussed in this chapter………………….……………….…………..…163

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Chapter 1

Introduction

1.1 The properties, advantages and a brief history of III-nitride materials

III-nitride materials have a direct band gap ranging from 0.7 ev to 6.2 ev [1-3], which makes these materials a well-suited candidate for a wide range of applications (Figure. 1.1).

Before the advent of III-nitride semiconductors there were no efficient ultraviolet (UV) and blue light-emitting diodes (LEDs). LEDs based on II-VI materials had short lifetimes as they were prone to dislocation generation over time [4]. Gallium nitride (GaN) and its alloys including AlN and InN have enabled deep UV, blue, green and red LEDs and laser diodes (LDs) [5-8].

Furthermore, due to high electron mobilities, high electron velocities, and high breakdown

Figure 1.1: Mapping of lattice constant vs. band gap of III-nitride semiconductors.

1 voltage, III-nitride materials have gained significant attention for high-power and high-frequency electronic devices [9]. Moreover, it is worth mentioning that nitride semiconductors have high mechanical, thermal, and chemical stability, which have enabled them to operate under harsh environments as sensors [10-12]. Finally, noting that the nitride semiconductors are potential candidates for photodetectors over a wide spectral range from UV to IR, and high-efficiency solar cells [13,14].

GaN was synthesized for the first time by Johnson et.al in 1932 by passing ammonia through hot gallium [15]. The first GaN was grown epitaxially in 1969 using halide vapor phase epitaxy on sapphire [16]. Due to significant amount of oxygen incorporation, unintentional background doping of GaN resulted in electron concentration ranging from 1018 to 1020 cm-3

[17]. This high concentrations of electrons in unintentionally doped GaN, prevented researchers from achieving high quality p-type GaN in 1970s. High quality p-GaN is required for high performance GaN p-n diodes. Due to lack of high quality of p-GaN the pace of research on GaN become slow in 1970s. It was not until late 1980s when a big step towards epitaxial growth of

GaN was made. Akasaki’s group proposed insertion of a low temperature GaN or AlN buffer layer before starting the GaN growth [18,19]. In terms of p-GaN quality and dopant, Mg from group II, was long expected to be a suitable acceptor for GaN, but it was almost impossible to detect positive charge carriers in p-GaN at room temperature, and the p-GaN showed a resistivity in the order of 106 Ω.cm. An accidental discovery paved the way to reduce the resistivity of p-

GaN layers from 106 Ω.cm to 2 Ω.cm. Amano et.al [20] showed that low-energy electron beam irradiation activates Mg-doped GaN layers and converts them to a p-type conductive state. A resistivity as low as 2 Ω.cm, was achieved by Nakamura et.al [21] for a Mg-doped p-GaN film, grown in metal-organic chemical vapor deposition (MOCVD) using thermal annealing in N2

Figure 1.2: Luminous efficiency as a function of operational wavelength, showing a gap in luminous efficiency ranging from 500 nm to 620 nm. (Source OSRAM [23]) environment. These achievements led into realization of the first GaN-based p-n junction LED in

1993[22].

To date, blue LEDs with a peak wavelength around 450 nm have been heavily researched. However, there are no high efficiency LEDs operating at the wavelengths between

500 nm to 620 nm, using semiconductor alloys made from AlGaInN materials (Fig 1.2) [23].

This is referred as the “Green Gap” problem for III-nitride based devices. Another issue, related to III-nitride LEDs is attributed to the reduction of the LED efficiency at high current densities which is known as “efficiency droop”. Another significant and enduring challenge facing high- power GaN-based LEDs is efficiency droop, which makes it difficult to utilize the solid-state lighting efficiently [24]. Determining the cause of this is still an active research topic that remains unresolved. Carrier overflow from the quantum wells into regions of efficient nonradiative recombination [25-27], loss of current injection efficiency [26], density-activated

3 defect recombination [28], insufficient hole injection efficiency leading to electron leakage

[24,29], Auger recombination in the quantum wells [30], and carrier delocalization from indium- rich low-defect-density regions at high carrier densities [31,32], are some of the explanations for efficiency droop.

1.2 Crystallographic planes of GaN and their properties

Wurtzite, zincblende, and rocksalt are three common crystalline structures reported for

GaN. Cubic GaN can be grown epitaxially inside silicon (Si) V-groove using MOCVD [33]. The wurtzite structure is the most thermodynamically stable (e.g. lowest energy) structure for III- nitride materials including GaN and its ternary alloys. Wurtzite GaN structure has six fold rotation and two mirror planes. Inside this structure each atom of gallium is surrounded by four nitrogen atoms, and each nitrogen atom is surrounded by four gallium atoms in a tetrahedral orientation. As a result, along the polar axis (c-direction) the relative numbers of Ga and N atoms change moving from one parallel plane to the next. In GaN, the covalent bond between a ‘Ga’ and a ‘N’ atom is in reality, partially ionic [34,35]. Due to differing electronegativities of ‘Ga’ and ‘N’ atoms, each unit cell of GaN is modeled as a charged dipole, with a spontaneous polarization electric field along the c-axis. Growth of heterostructures (e.g. GaN/InxGa1-xN/GaN) results in another type of polarization known as piezoelectric polarization. Piezoelectric polarization is a result of strain applied by lattice mismatch between GaN and its ternary alloys.

Spontaneous polarization accompanied with piezoelectric polarization has a significant effect on the performance of III-nitride semiconductor devices grown on different crystallographic orientations [36]. Schematics of the unit cells for wurtzite GaN and some low index planes are shown in figure 1.3. In the case of quantum wells grown on c-plane, strong internal electric fields caused by spontaneous and piezoelectric polarizations result in significant band bending within

Figure 1.3: Schematic of the unit cell for wurtzite GaN showing the low index planes including a- plane, r-plane c-plane, and m-plane. the quantum wells and spatially separate the electron and holes inside the quantum wells. This separation lowers the electron-hole wave function overlap, also reducing the radiative efficiency of the electron-hole recombination compared to flat-band quantum wells. This phenomenon is known as the quantum-confined stark effect (QCSE). However, for c-plane III-nitride electronic devices, strong polarization is beneficial, because it leads to two-dimensional electron gases

(2DEGs) with large sheet concentrations [37], which are leveraged in electronic devices. The situation is different for other crystallographic planes.

Planes for a hexagonal lattice can be represented by four indices of h,k,i,l as (h,k,i,l) with a relation of h+k+i=0. There are certain relations between these indices, which define general properties of the corresponding planes. For instance, h=k=i=0,l≠0 represents parallel polar c- planes with an inter-planar distance of l. For m-planes and a-planes (Fig. 1.3), due to equal number of atoms of gallium and nitrogen per plane, the strong spontaneous polarization does not exist. These planes are known as nonpolar planes. The absence of a polarization-related electric

Figure 1.4: Simulated energy band diagrams for 3 nm In0.23Ga0.77N for single quantum wells with GaN barrier at a current density of 100 A/cm2 for (a) polar (c-plane), (b) nonpolar (101̅0) (m- plane). The direction of the piezoelectric field and p-n junction built-in field are indicated. The ground-state electron and hole wave functions are also shown. The n-side donor concentration is 1e18 cm-3 and p-side acceptor concentration is 2e19 cm-3. (Data from Ref [38]) field within the quantum wells grown on nonpolar planes, gives rise to higher spatially electron- hole wave function overlap, and higher radiative efficiency (Fig. 1.4) [38]. The polarization- related electric field is somewhat reduced for other categories of planes known as semipolar planes (r-plane in figure 1.3). The term “semipolar plane” refers to any plane that cannot be classified as c-plane, a-plane, or m-plane. In terms of crystallographic terms, a semipolar plane would be any plane that has at least two nonzero h, k, i indices and a nonzero l index.

1.3 Growth of wurtzite III-nitride materials and origin of threading dislocations

Due to the low solubility of nitrogen in gallium and the higher vapor pressure of nitrogen on GaN, this material is not available in nature [39]. Hence, the growth of bulk GaN needs to be performed at high temperature and pressure. The absence of a native substrate for GaN necessitates growth of GaN on non-native substrates with wurtzite crystallographic structure.

Table 1.1 shows the comparison of three common substrates used to grow c-plane wurtzite GaN.

Table 1.1: Comparison of three common substrates used to grow c-plane wurtzite GaN.

Sapphire SiC Silicon

∆풂 (%) -13.9 -3.4 17 풂ퟎ

Electrical n-type SiC is a Insulating Semiconductor properties semiconductor

Relatively Cost inexpensive, up to Expensive Inexpensive 12”

C-plane sapphire is the most common substrate being used to grow commercial III-nitride based LEDs and laser diodes (LDs). In addition to sapphire, silicon carbide and silicon (111) are some other potential substrates to grow GaN. Table 1.1 compares the different type of substrates that are used to grow c-plane GaN. Heteroepitaxially, growth of GaN on lattice mismatched substrates results in GaN being strained (compressive or tensile) depending on the substrates. As the thickness of GaN increases beyond the critical thickness of GaN on these substrates, GaN relaxes and the strain energy is released by formation of threading defects. Growth of wurtzite polar GaN on c-plane sapphire results in threading defect densities ranging from 108 to 1010 cm-2.

Figure 1.5 shows a TEM micrograph of a wurtzite c-plane GaN crystal grown on c-plane sapphire [40]. The threading dislocations originate at the interface, due to the lattice mismatch between sapphire and GaN. Threading defects have adverse effects on the performance of III- nitride materials and devices. Higher threading defect density reduces the lifetime of high power

LEDs and LDs [41]. In terms of electrical properties, higher threading defect increases the reverse leakage current density [42-44], and lowers the electrical breakdown of GaN p-n

Figure 1.5: TEM micrograph of the wurtzite c-plane GaN grown on c-plane sapphire, showing the threading defects originating at the interface between GaN and Sapphire (Data from Ref [40]). junctions [45]. In terms of optical properties, it has been shown that higher threading defects reduces the internal quantum efficiency significantly [46]. Thus, the reduction of threading defects is a potential way to improve the performance of III-nitride devices.

1-4 Motivation for the development of nanostructure-based LEDs

Nanoheteroepitaxy is one of the viable approaches to reduce the threading defects [47].

III-nitride nanostructures can be grown various morphologies like quantum dots (QDs) [48], core-shell structures [49], nanowires (NWs) [50-52], nanopyramids [53, 54], nanorings [55], nanotubes [56], nanoneedles [57], nanopores [58], nanobridge [59] and etc. Among all nanostructures, III-nitride nanowires are one of the most promising structures for the next generation of optoelectronic and electronic devices, such as LEDs [50,60-63], laser diodes [64-

66], high power transistors [67], gas sensors [68,69] visible-blind UV photodetectors [70,71], solar cells [72,73], and plasmonic nanolasers [74,75]. LEDs for lighting based on multi-color

(RGB or RYGB) white light sources could potentially eliminate the inherent optical losses

8 associated with the down conversion of photons in conventional phosphor-based approaches, thereby improving conversion efficiencies. Nanowires offer an approach to monolithically integrated RGB-based white LEDs [76-78]. However, the ability to demonstrate uniform multi- color RGB LEDs on the same wafer using III-nitride nanowires, without the need for complicated fabrication and wafer-bonding techniques, is still in its infancy and requires further development. The multi-color LED solution has the potential to reach luminous efficacies of around 300 lm/W, compared to phosphor-based approaches that saturate around 200 lm/W [79].

In addition, RGB-based white LEDs are preferred for visible light communication (VLC) systems with high data rates,[80] because the data rates in phosphor-converted white LEDs are limited by the slow response of the phosphors [81,82]. Dynamic color tuning can also be achieved via the multi-color approach, which could potentially lead to multiplexing within VLC over several wavelength channels. Researchers have demonstrated polychromatic LEDs on c- plane GaN based on wafer-bonding [83], cascade single-chip using direct wafer bonding [84], and also blue-green color tuning [85]. However, as mentioned in section 1.2, conventional c- plane GaN-based LEDs suffer from internal polarization related electric fields, which hinder their radiative efficiency and modulation speed. These internal electric fields reduce the spatial overlap of the electron-hole wave functions inside the polar c-plane QW, thereby leading to a reduction of the radiative recombination efficiency and thus increasing the radiative recombination lifetime [86]. Nonpolar m-plane InGaN based LEDs have been fabricated by a number of research groups on free-standing GaN [87-89]. The QWs grown on m-plane are devoid of spontaneous and piezoelectric polarization, thereby improving radiative efficiency and reducing efficiency droop [90,91]. Despite promising results, the small area and high cost of the m-plane substrate are two main obstacles for next generation nonpolar LEDs. Alternately,

Figure 1.6: XTEM images of a GaN nanowire (a) showing the bending of the threading dislocation o o by 90 and termination at the SiNx growth mask. (b) showing the 90 bending of a threading

dislocation toward the nonpolar side wall, and filtering out the threading dislocations by SiNx dielectric growth mask. (Data from Ref [40]). devices based on nanostructures such as nanowalls [49,92,93], nanowires [94,95,96], and triangular-nanostripes grown on commercial c-plane sapphire substrates could potentially overcome these obstacles [97,98], and enable RGB LEDs on a single chip.

Nanostructure devices based on bottom-up selective area growth (SAG) have shown a significant reduction in threading dislocation density. The dislocations bend toward the sidewalls at the base of the nanostructures (especially in the case of nanowires) or can be filtered by the dielectric mask used to grow these structures [40] (Fig. 1.6). This results in high qualityGaN at the top of the nanowires.

In addition, GaN nanowires have substantially larger active region surface areas (~4-

12X) compared to their planar foot-print, leading to a lower carrier density for a given drive current, which could mitigate the effects of LED efficiency droop (see figure 1.7). The lateral strain relaxation in nanowires on the nonpolar m-plane sidewalls may also lead to a higher level

Figure 1.7: The enhancement in the active region area as a function of the nanowire radius with different heights compare to a planar LED on the same substrate footprint area.

of indium (In) incorporation during InxGa1-xN/GaN (x= mole fraction of In) core-shell growth as a result of the small dimensions of the nanowires [99].

The small dimension of nanostructure-based emitters also paves the way to improve the radiative efficiency via the Purcell effect. Edward Purcell in 1946 realized that the rate of spontaneous emission could be modified when an emitter is placed inside a cavity [100]. The

Purcell effect states that the enhancement in the emission rate for maximum coupling is directly proportional to the quantity Q/V, as

3 휆 3 푄 퐹 = ( ) ( ) (1.1) 푝 4휋2 푛 푉

Where Fp, is the Purcell enhancement factor, 휆 is the free-space wavelength, n is the refractive index, Q is the quality factor, and V is the mode volume. Nanostructure-based emitters with dimensions in the range of the emission wavelength are potential candidates to shrink the mode volume accompanied by a moderate Q parameter. Reduction of the mode volume results in higher spontaneous emission rate via the Purcell effect, and higher radiative rate of nanostructure-based LEDs.

Due to all mentioned advantages of nanostructures, especially nanowires, vast numbers of papers have been published over the past two decades. The number of papers based on nanowires increased exponentially, with most of the activities and development happening in the last ten years (Fig. 1.8).

Figure 1.8: Increase in the number of publications on nanowire related topics from 1996-2016 (Source, ISI; keyword, nanowires).

Among all nanostructure-based LEDs, GaN/InGaN core-shell nanowire-based LEDs

using bottom-up SAG are the most common, benefiting from lower threading dislocation

densities, higher active region area, and possibility of the growth of quantum wells on nonpolar

m-plane side walls.

1-5 A brief history of selective-area growth of GaN/InGaN core-shell

nanowire-based LEDs

The first demonstration of SAG growth of GaN/InGaN core-shell LEDs, was performed

by W. Bergbauer et.al in 2010 [94]. A continuous-flow MOCVD was used to grow nitrogen

polar GaN nanowires on sapphire, with diameters ranging from 200 nm to 500 nm. Five pairs of

quantum well/barrier were grown around the GaN nanowires cores, and a peak of

13 photoluminescence (PL) of 420 nm was achieved. A 95 nm thick p-type GaN was grown around the active region. Although, a cross-sectional bright-field TEM micrograph of the nanowires showed smooth and perfect side walls, a significant defect-related yellow-band emission for nanowire-based LEDs in the PL spectra is indicative of low-quality, and dim quantum wells. The internal quantum efficiency (IQE) and current-voltage (I-V) plot for the LEDs were not reported for these LEDs. After that, it was in 2013, when S. Albert et.al [101] reported SAG growth of

GaN/InGaN core-shell LEDs. In their work, molecular beam epitaxy was used to grow GaN nanowire with diameter ranging from 200 nm to 500 nm on GaN-buffered Si (111). A peak PL of 540 was reported, and a peak IQE of 37% was achieved as the In/Ga ratio changed during the

QW growth. However, the electrical characteristics including I-V and L-I-V were not reported.

The first demonstration of electrically injected GaN/InGaN core-shell nanowire-based LEDs was performed by C. G. Tu et.al in 2014 [95]. They used MOCVD selective-area growth on c-plane

GaN on c-plane sapphire to grow patterned array of GaN nanowires with diameter of ~500 nm and the height of ~650 nm. Three pairs of quantum well/barrier were grown around the GaN nanowire cores, followed by a 150 nm of p-GaN layer on the side walls. A peak PL of 485 nm was demonstrated from nonpolar QWs on the side walls, and a peak IQE of 21% was achieved.

Although, a high device series resistance of 336 Ω was reported, a turn-on voltage of ~3 V was achieved for the nanowire-based LEDs. M. S. Mohajerani et.al, performed a thorough study of optical properties of SAG grown GaN/InGaN core-shell nanowire-based LEDs in 2016 [102]. In their study, MOCVD SAG was used to grow GaN nanowires with a diameter of 800 nm, and a height of 12 µm on a 5.5 µm thick c-plane GaN on c-plane sapphire. One quantum well was grown around the GaN nanowire, followed by ~500 nm of p-GaN. A peak PL at 423 nm, was shown at 20 K, and an IQE of 8% was reported. A significant change in peak CL was observed

14 at the side walls of the nanowires, ranging from 406 nm to 430 nm, which is a clear indication of different indium incorporation along the side walls of the nanowires. In 2016, T. Schimpkea et.al from OSRAM also reported the highest performing SAG grown GaN/InGaN core-shell nanowire-based LED [96]. They have used MOCVD to grow GaN nanowires on c-plane GaN on c-plane sapphire, with a diameter of 800 nm and a height of ~10 µm. Multiple quantum well were grown around the nanowires, followed by a p-GaN layer. A peak PL wavelength of 445 nm was shown, and an IQE of 58% was reported. This value for IQE was the highest reported IQE for SAG GaN/InGaN core-shell nanowire-based LEDs. I-V and L-J-V (light, current density, voltage) plots were reported and a peak external quantum efficiency of 10% was achieved. A summary of reported IQEs for GaN/InGaN core-shell nanowire-based LEDs is shown in figure

Figure 1.9: A brief history of IQEs reported for SAG of GaN/InGaN core-shell nanowire-based LEDs.

1.9.

1-6 Motivation for selective area growth of GaN nanowires and advantages

Several techniques have been reported in the literature for the growth of GaN nanowires.

They can be categorized into three main approaches: 1) catalyst-assisted, 2) catalyst-free selective area growth (SAG) using in-situ deposition of a dielectric layer, and 3) catalyst-free

SAG using ex-situ deposition of a dielectric layer. In the catalyst-assisted approach, a vapor- liquid-solid (VLS) technique can be used in either MOCVD [103,104] or in the laser assisted catalytic growth technique [105]. In-situ catalyst-free SAG can be done using molecular beam epitaxy (MBE) [106] or MOCVD [107]. Ex-situ catalyst-free SAG may also use either MBE

[108-110] or MOCVD [111-114]. The catalyst-assisted VLS growth technique using MOCVD does not enable well-controlled GaN nanowires and the metal catalyst is often incorporated as a deep level impurity [115]. Among catalyst-free techniques, the ex-situ approach is best suited to increase the homogeneity of the nanowires since the in-situ technique results in poor uniformity in the position and dimension of the nanowires. In the case of the ex-situ technique using MBE, the growth of core-shell QWs is limited by a shadowing effect [63,116]. Using the MOCVD ex- situ catalyst-free technique, the geometry of the nanowires is well controlled with a dielectric mask, the shadowing effect is absent, and no metal catalyst incorporates into the GaN. Therefore, the MOCVD ex-situ catalyst-free technique is promising approach for growing controlled arrays of uniform and high-quality GaN nanowires. To realize the fundamental requirements to ensure high reproducibility and continual optimization of the growth methods, a detailed study of the influence of the MOCVD growth conditions, and mask geometry is required. In addition, precise control on nanostructure properties including position, density, diameter, height and formation of different crystallographic planes are required to achieve uniform arrays of nanowires. This

16 uniformity is essential for obtaining consistent optical and electrical properties required for wafer scale nanowire-based opto-electronic devices.

1-7 Scope of this dissertation

This dissertation mainly focuses on the design, selective-area growth, fabrication, and characterization of GaN/InGaN core-shell nanowire-based LEDs. Particular emphasis is placed on selective-area growth and fabrication techniques, understanding the effect of different growth parameters on the nanowire’s morphology and electro-optical and RF characterization. The dissertation is divided into six chapters:

Chapter I: The present chapter has given a general introduction on advantages and issues related to III-nitride materials. A brief historical view of the development of III-nitride semiconductors is given, and the obstacles pertaining to the conventional planar structures are discussed. Based on particular advantages of nanostructures, a new path to overcome challenges related to planar structures is given. A brief history of selectively grown GaN/InGaN core-shell nanowire-based LEDs is provided, which shows the performance of these types of devices has improved significantly, but still more detailed studies need to be done to improve the performance of these devices further.

Chapter II: In this chapter MOCVD-SAG of nanostructures is discussed, and a systematic study of the influence of the different growth conditions on the geometry of the nanostructures is reported. This study opens a new window towards precise control of the nanostructure morphology, which is an important key for optimization of such devices.

Chapter III: The third chapter discusses the growth of core-shell InGaN quantum wells around the nanostructures, and a study of indium incorporation for the active region. In this

17 chapter a thorough study of the influence of the mask geometry and nanostructures morphology on indium incorporation is presented. This section provides a path towards multi-color monolithically integrated LEDs based on nanostructures, which is extremely difficult for planar structures.

Chapter IV: Electrically injected GaN/InGaN core-shell nanowire-based LEDs are presented in this chapter. Scanning transmission electron microscopy (STEM) was performed to characterize the growth of thin films, including quantum wells, barriers, and electron blocking layers (EBL). Micro-photoluminescence (µ-PL), internal quantum efficiency, low temperature, and room temperature time-resolved PL, and far-field emission pattern of the device are given.

Radiative and nonradiative lifetimes are decoupled and presented. A detailed study of the carrier dynamics of devices is reported in this chapter. Electrical properties including J-V, L-J-V, and external quantum efficiency are reported. RF characteristics of the device are given and an equivalent circuit model is proposed for the device.

Chapter V: This chapter presents a theoretical study of nanoemitters, and their interaction with metals. Several plasmonic structures are proposed and their optical properties are given.

This chapter also provides a thorough study of the Purcell effect, and its calculation method. The finite-difference time domain method is applied to calculate the Purcell enhancement factor, for a various types of nanowire LEDs.

Chapter VI: This chapter talks about the most important findings of this dissertation, and provides a path towards future work in this field.

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Chapter 2

Selective-area growth using MOCVD

2.1 MOCVD system Veeco-p75

This section provides some information about our metalorganic chemical vapor deposition (MOCVD) equipment. MOCVD vertical rotating disc reactors (RDR) are widely used for large-scale production of GaN-based semiconductor devices including light-emitting diodes

(LEDs) and field-effect transistors (FETs). In RDRs, wafer rotation provides an effective averaging of the deposition rate distribution, this results in highly uniform growth of epitaxial layers. The details of RDR technology can be found elsewhere [1]. We use a Veeco-p75 system, which is a commercial Veeco vertical turbo disc reactor, with a rotation speed of up to 1500 rpm.

This reactor is a vertical high-speed RDRs with a 75 mm disc diameter, where ammonia and alkyls diluted in nitrogen and/or hydrogen carrier gas are delivered into the growth chamber by a flow flange gas injector. The flow flange has three different zones: one hydride zone and two alkyl zones. This allows separate reactant delivery, which prevents a premature reaction between the precursors. The two alkyl supply zones enable a more efficient way to control flow and layer uniformity [2]. Using a Veeco-p75, it is possible to grow epitaxial layers with pressures ranging from 20 Torr to 500 Torr, and growth temperatures up to 1200oC. This MOCVD reactor was utilized to grow all III-nitride materials in this dissertation including GaN templates, GaN nanostructures, InxGa1-xN/GaN multiple quantum well/barrier, AlxGa1-xN, and p-GaN layers for planar and 3D structures. Typical group III precursors including trimethylgallium (TMGa), triethylgallium (TEGa), trimethylindium (TMIn), and trimethylaluminum (TMAl) have been used with a combination of group V sources -ammonia (NH3) -to grow GaN templates, quantum

Figure 2.1 (a) Veeco-p75 MOCVD reactor, (b) load lock and transfer arm, (c) and growth chamber. wells/barriers, and GaN nanowires. Silane and biscyclo pentadienyl magnesium (CP2Mg) have been used as n-dopants and p-dopants respectively. Figure 2.1, shows the reactor used for material growth, including load lock, transfer arm, susceptor stage, and the growth chamber. The load lock configuration enables samples to be moved toward the growth chamber using a transfer arm. This prevents contamination of the reactor by atmospheric particles during the loading process. This is an easy and effective way to prevent O2 contamination of the chamber and maintain a stable humidity inside the chamber.

2.2 Processing steps required for nanowire template preparation using electron-beam lithography

To leverage the Purcell effect and enhance the spontaneous emission rate, nanoemitters with a small mode volume and a moderate Q parameter are required (Eq. 1.1). Electron-beam lithography is a feasible way to pattern dielectric thin films with a resolution of 10 to 20 nm

[3,4]. Electron-beam (e-beam) patterning, accompanied with selective-area growth (SAG), is a potential way to reduce the size of the nanoemitters below 100 nm. This section provides experimental techniques required for e-beam patterning and SAG of GaN nanowires. The growth of GaN nanowires was carried out on a selectively patterned c-plane GaN template on c-plane

31 sapphire. Processing steps required for e-beam patterning and SAG of GaN nanowires are given in this section. Initially, a 2-µm-thick GaN layer doped with silicon (n-GaN) was grown on a 2 inch, single-side-polished sapphire substrate. A thin (less than 10 nm) low temperature GaN buffer layer was grown on sapphire before the growth of the GaN template. A conventional method of continuous mode MOCVD and typical group III (TMGa) and group V (NH3) precursors was used to grow the GaN template. The GaN template growth was carried out at the pressure of 500 Torr at the growth temperature of 1090oC, with a group V to group III ratio

(V/III ratio) of 3200. The rotation speed was held at 1500 rpm during the growth. The wafer was then cleaned in piranha solution (mixture of sulfuric acid and hydrogen peroxide 4:1) and 25 nm

o of SiNx was deposited using chemical vapor deposition (CVD) at 300 C (Samco PD-10). The sample was cleaned and 90 nm of PMMA (495PMMA-C2) was deposited on it followed by 10 nm of gold. This layer plays an important role in the performance and reproducibility of e-beam patterning. After e-beam patterning (hexagonal symmetry) the gold layer is removed using gold etchant-TFA. The sample is then developed in cold methyl isobutyl ketone (MIBK): isopropyl alcohol (IPA) (1:3). Subsequently, Cl2-based reactive-ion etching (RIE) is used to pattern arrays of circular apertures. During these processing steps, many challenges had to be overcome.

Patterning SiNx for apertures with 40 nm diameter was the main challenge. A thin layer of

PMMA is required to reduce the aperture sizes, however using RIE, the etch rate of PMMA is 3 times faster than SiNx. This results in removing the whole PMMA before etching the dielectric.

Hence, 90 nm thick PMMA is required to be able to etch 30 nm of SiNx. Using 90 nm of PMMA and 30 nm of SiNx, we were able to make apertures in the SiNx with 40 nm in diameter. E-beam irradiation shape and dose were other parameters that needed to be optimized to achieve consistent and repetitive apertures. Figure 2.3 shows two different levels of irradiation doses

Figure 2.2: Processing steps for nanowire template using e-beam lithography. (1) 2 µm GaN template

grown on Sapphire, (2) Deposition of 25 nm of SiNx using plasma-enhanced chemical vapor deposition, (3) Deposition of 90 nm PMMA and 10 nm gold, and e-beam patterning, (4) Gold removal, (5) Develop in cold MIBK:IPA 1:3, (6) Reactive ion etching, (7) Piranha cleaning, (8) E- beam nanowire template. with targeted and resulted aperture shapes. It is shown that, to get circular apertures with diameters ranging from 25-30 nm, a rectangular shape pattern with a side length of 16 nm or a hexagonal shape pattern with a diagonal length of 30 nm are required accompanied with a higher

(4000 µc/cm2) or lower (3000 µc/cm2) e-beam irradiation dose, respectively. Removing the

PMMA using piranha was another challenge (see figure 2.3). This issue was solved using piranha cleaning in a sonication machine, with only one sample in each cleaning process (see figure 2.4). Finally, we were able to optimize the e-beam parameters, PMMA thickness, and

Figure 2.3: E-beam patterning (a) square shape using higher dose, (b) and hexagonal shape using lower dose. Both result in circular aperture. piranha cleaning. By doing so, we achieved clean e-beam nanowire templates that were ready for growth.

2.3 Selective-area growth of GaN nanowires

Once a clean e-beam nanowire template was ready, it was loaded into the reactor for the nanowire growth. The GaN nanowires were grown using pulsed MOCVD in an H2-N2 atmosphere without any nucleation steps. Figure 2.6 shows a schematic illustration of the pulsed

MOCVD growth technique, used to grow GaN nanowires. In a pulsed mode technique, H2 and

N2 gases are flowing continuously while the group V (N) and group III (Ga) precursors are temporally pulsed. Selective-area pulsed MOCVD growth technique was introduced by Hersee et.al to control the geometry and position of the GaN nanowire growth [5]. The pulsed MOCVD mode technique consists of four different sections in a growth cycle: TMGa injection time,

Figure 2.4: PMMA remnants after piranha cleaning. This random pattern results in a random nanowire growth region with random geometry.

Figure 2.5: E-beam nanowire sample after piranha cleaning under optimized conditions.

TMGa interruption time, NH3 injection time, and NH3 interruption time. During the first section,

TMGa is flowing into the chamber, which results in an excess Ga adatoms in the chamber. In the second section, Ga adatoms have time to migrate along the dielectric growth

Figure 2.6: Schematic image of different steps in each cycle for pulsed-mode growth

technique, (a) TMGa injection time, (b) TMGa interruption time, (c) NH3 injection time, (d)

and NH3 interruption time.

Figure 2.7: Schematic illustration of the processes required to grow GaN nanowires using a MOCVD pulsed mode technique, (a) TMGa injection step, (b) TMGa interruption step, (c)

NH3 injection step, (d) and NH3 interruption step. mask and reach the apertures to start the nucleation of GaN nanowires. Also, the adatoms may reach forming nanowires and contribute in the growth of these nanowires. During the TMGa

36 interruption time, Ga adatoms that are absorbed on the m-plane sidewalls may diffuse to the top c-plane or desorb from the m-plane. Since the sticking coefficient of the Ga adatoms on the c- plane is higher than that on the m-plane, Ga adatoms can stay on top of the c-plane longer until the NH3 injection starts. Once the NH3 injection starts, excess amounts of N adatoms interact with the Ga adatoms on the top c-plane and contribute in the vertical growth of nanowires [6]. In addition, it is shown that under Ga rich conditions, the Ga atoms are able to replace the hydrogen atoms that passivate the semipolar planes and provide growth on those planes as well [7,8]. The

Ga interruption time is a critical parameter to the growth of the GaN nanowires in a pulsed mode technique. An appropriate selection of the Ga interruption time gives rise to a suppression of the growth rate on the m-plane side walls. Yet the growth rate on the c-plane and semipolar planes remain high. This helps to grow nanowires more vertically than laterally. Conventional continuous growth mode might occur if TMGa is injected right after the NH3. This is due to the residual amount of NH3 in the chamber which results from the high partial pressure of NH3 typically used to grow GaN. Continuous mode growth typically results in a formation of nanopyramidal structures [9]. Thus, a NH3 interruption is required to avoid a continuous growth mode. The GaN nanowires were grown with a pressure of 90 Torr, rotation speed of 1500 rpm,

H2 flow of 3000 sccm, and N2 flow of 1000 sccm. The TMGa injection time, TMGa interruption time, NH3 injection time, and NH3 interruption time were 18 sec, 27 sec, 6 sec, and 15 sec,

-1 -1 respectively. The TMGa flow rate and NH3 flow rate were 26.7 µmol min and 8 mmol min , respectively. The V/III ratio was 100. It is shown that a low V/III ratio (Ga rich condition) is required to grow nanowires vertically [10,11]. The pulsed mode technique is one way to reduce the V/III ratio, locally at the position of the nanowire growth [11]. Recently, Choi et. al [12], and

(a) (b) (c) (d)

Figure 2.8: SEM images, at a 45-degree tilt, of GaN nanowires for 1 µm pitch and different aperture diameters (a) 60 nm, (b) 110 nm, (c) 160 nm, (d) and 210 nm.

Coulon et. al [13] also successfully demonstrated GaN nanowire growth using the continuous growth mode at a very low V/III ratio.

2.4 Effect of aperture diameter and the role of capture length

Since different crystallographic orientations of nanostructures have different optical and electrical properties, the nanostructure geometry has a significant effect on the performance of nanostructure-based electro-optical devices. In this section we study the effect of aperture diameter on the nanowire dimension and morphology. The pitch spacing between each nanowire was kept constant and equal at 1 µm, and the aperture diameter was changed form 60 nm to 210 nm. 10 pulsed-mode cycles (Fig. 2.6) were employed to grow the GaN nanowires. The other growth parameters were the same as previously discussed, and the growth temperature was kept at 926oC. Figure 2.8 shows the SEM images of the GaN nanowires grown using different aperture diameters with a pitch of 1 µm. The morphology of the nanowires did not significantly change even though the aperture diameter was varied. The unchanged morphology of the nanowires grown in different sized apertures with constant pitch is attributed to a fixed annular surface around a nanowire within which the growth nutrient (TMGa and NH3) can be captured to participate in the growth. This annular distance is defined as the capture length [14]. The capture length can be calculated as follows:

Figure 2.9: (a) Average diameter and average height of nanowires as a function of aperture diameter, (b) and average nanowire volume as a function of aperture radius (raw data) and the fitting curve used to determine the capture length.

푉 ∝ 퐴푐푎푝푡푢푟푒 (2.1)

푉 = 퐾 × 휋(푅 + 푎)2 (2.2)

The term 푉 in equation 2.1 is the average structure volume and K in equation 2.2 is a proportionality constant which is only related to growth conditions and is assumed to be independent of the pitch size. Furthermore, 푎 is the aperture radius and R is the capture length which is unique for a given pitch size.

Figure 2.9(a) shows the average diameter and average height of the nanowires as a function of the aperture diameter. Figure 2.9(b), shows the average structure volume, V, as a function of the aperture radius (a). Statistical analysis of the nanowire size and morphology was performed using SEM images of nanowires. For nanowire arrays with different aperture diameters, the statistics were taken from 100 nanowires. Thus, the reported nanowire sizes are the statistical average obtained for arrays of nanowires with 1 µm pitch and different aperture diameters. The standard deviation is reported as the error for each set of measurements. Using

(a) (b) (c)

(d) (e) (f)

Figure 2.10. SEM images, at a 45-degree tilt, of a sample containing GaN nanowires showing the differences in growth morphology for pitch (a) 300 nm, (b) 500 nm, (c) 1 μm, (d) 2 μm, (e) 3 μm, (f) and 4 μm. equation (2.2) and non-linear curve fitting based on chi-square minimization in a commercial plotting software, the best fit curve is obtained, giving 퐾휋 = 89.7 nm and 푅 = 470 nm. The excellent fit suggests that the capture length is independent of the size of the initial aperture diameter for the given growth conditions.

2.5 Effect of pitch size on nanowire morphology and the role of growth temperature

In the previous section the effect of the aperture size on nanowire morphology was studied. In this section we study the effect of the nanowire spacing pitch between each nanowire on the nanowire dimension and morphology. The pitch was changed from 300 nm to 4 µm. 10 pulsed-mode cycles (Fig. 2.6) were employed to grow the GaN nanowires with aperture diameter of 40 nm. The growth parameters were kept the same as previously discussed, and the growth

40 temperature was kept at 926oC. Figure 2.10 shows SEM images of the GaN nanowires grown out of 40 nm openings with different pitches. Clearly, the nanowires with the smallest spacing between them exhibit the smallest volume indicating an increased competition for the growth nutrients between adjacent nanowires as the pitch is reduced. The diameters of the nanowires increased with increasing pitch, indicating a reduction in the competition for the growth nutrients with larger pitch. A difference in growth morphology between the nanowires on different pitches was also observed. The explanation lies in the higher surface mobility of TMGa compared to

NH3 [15,16]. The 300 nm and 500 nm pitch nanowire arrays were found to exhibit mostly nonpolar m-plane facets and angled semipolar {101̅1} facets. By increasing the pitch to 1 μm and 2 μm, the nanowire arrays exhibit predominantly flat c-plane facets and m-plane sidewall facets, while the proportion of the angled semipolar {101̅1} facets is significantly reduced. This results from a locally enhanced gallium adatom density for the larger pitches due to the lower density of nanowires and the higher TMGa surface mobility. This leads to a localized reduction of the V/III ratio. Under these conditions, the growth rate of the semipolar plane is likely enhanced relative to the more stable c-plane and m-plane. This causes the nanowire to grow out and lead to the prominence of the top (0001) and side wall {101̅0} facets [10]. With a further increase in the pitch, the extent of the m-plane facets is significantly reduced and the c-plane and the semipolar plane become the dominant facets in the nanowires. This was previously explained by an increase in the density of hydrogen atoms available from the decomposition of NH3 per unit area of GaN growth. The excess hydrogen atoms bond to the nitrogen on the {101̅1} surface. This prevents the incoming Ga adatoms from bonding to the nitrogen atoms, which leads to a reduction in the growth rate [17,18]. Thus, the semipolar facet becomes the dominant crystallographic orientation observed when the pitch is greater than 3 µm.

(a) (b) (c)

Figure 2.11: SEM images, at a 45-degree tilt, of GaN nanowires for 500 nm pitch, showing the differences in growth morphology and nanowire diameter, under different growth temperatures (a) 917oC, (b) 923oC, (c) and 930oC. The nanowire morphology evolution as a function of growth temperature is shown in the inset figures.

2.6 Effect of the growth temperature on nanowire dimension and morphology

Due to the strong dependence of the adsorption and desorption rates of precursors at different growth temperatures, the effect of growth temperature on nanowire morphology was studied in this section. This study was carried out by varying the growth temperature from 917oC to 930oC while the remaining growth parameters were kept constant, and are the same as discussed in section 2.3. To be consistent, 10 pulsed-mode cycles (Fig. 2.6) were employed to grow the GaN nanowires. Figure 2.11, shows SEM images of the GaN nanowires grown at different growth temperatures. At low growth temperatures, the density of parasitic polycrystalline GaN on the dielectric mask is considerable and reduces with increasing growth temperature. The higher parasitic growth at low temperature results from the lower surface mobility of the growth nutrients. With increasing growth temperature, the growth nutrients gain sufficient kinetic energy to traverse the growth dielectric to initiate nucleation inside the apertures. Thus, the increased growth temperature leads to a reduction in the parasitic polycrystalline GaN grown on the dielectric. The inset images in figure 2.11 show the dependence of nanowire size and morphology on growth temperature for the 500 nm pitch

42 arrays. Statistical analysis of the nanowire size and morphology was performed using SEM to elucidate the optimal growth temperatures needed to achieve particular nanowire characteristics.

The statistics were taken from 100 nanowires for pitch sizes less than 1 µm and from 50 nanowires for pitch sizes more than 1 µm. Thus, the reported nanowire sizes are the statistical average obtained for each pitch. The standard deviation is reported as the error for each set of measurements. The average diameter of the nanowires and the lateral overgrowth (nanowire diameter – aperture dimeter) versus periodic spacing for different growth temperatures are shown in figures 2.12 (a) and 2.12(b), respectively. As expected, the nanowire diameter and lateral overgrowth increases with increasing pitch due to a reduction in the competition with adjacent nanowires for growth nutrients. It can be observed in figure 2.12(a) and 2.12(b) that increasing growth temperature leads to a reduction in the nanowire diameter and lateral overgrowth. This is due to Ga desorption from the nanowires with an increase in temperature which reduces the growth rate. The nanowire diameter did not significantly increase at a growth temperature of

930oC for pitches greater than 2 µm which is attributed to an equilibrium established between the lateral growth rate and the thermal etching rate of GaN in hydrogen ambient at high temperatures

[19]. SEM images showing the changes in growth morphology as a function of temperature for

3-µm-pitch nanowires are shown as insets in figure 2.12(b). As the temperature increases, the Ga desorption from all facets increases, resulting in a reduction of the nanowire size. However, due to the higher Ga adatom desorption rate on the m-plane compared to the c-plane [20], the m- plane sidewalls were more significantly reduced than the top c-plane. As the temperature is increased further, the stability of the semipolar {101̅1} facets is enhanced by hydrogen passivation of the semipolar plane. This leads to structures dominated by semipolar facets at high temperature. As shown in figure 2.12(c), the average height of the nanowires increases by

(a) (b)

(d) (c(c) ) (d)

Figure 2.12: (a) Average nanowire diameter for different pitches at different growth temperatures, (b) lateral overgrowth for different pitches at different growth temperatures. The inset images in (b) show the dependence of nanowire morphology on temperature for those grown with a 3 µm pitch, (c) average height of the nanowires for different pitches at different growth temperatures, and (d) aspect ratio for nanowires grown at different temperatures for different pitches. increasing the pitch from 300 nm to 4 µm for growth temperatures ranging from 917oC to 926oC.

At a growth temperature of 930oC, the nanowire height increases for pitches ranging from 300 nm to 2 µm, while the height is slightly reduced for pitches beyond 2 µm. The reduction of

44 height for pitches beyond 2 µm is a clear indication of thermal etching at 930oC. As the growth temperature increases from 917oC to 930oC, both the nanowire height and aspect ratio (average height/average diameter) are reduced, as shown in figure 2.12(d). The maximum aspect ratio in figure 2.12(d) is 2.2. The aspect ratio using selective-area pulsed-mode MOCVD is limited by transport of adatoms to the tops of the nanowires and is therefore dependent upon the temperature. Jung et. al [11] reported an aspect ratio of 5.9, and we have previously observed aspect ratios up to 11. Jung et. al [11], and Lin et. al [6] have shown an increase in aspect ratio at higher growth temperatures, due to higher surface diffusion of Ga adatoms from the side-wall m- planes to the top c-plane at higher temperatures [21]. As discussed before the sticking coefficient of the Ga adatoms on the c-plane is higher than that of the m-plane [6, 22, 23]. The diffused Ga adatoms can stay longer on the top c-plane until the NH3 injection begins, and Ga adatoms contribute to the growth of the nanowires vertically rather than laterally. However, if the growth temperature is high enough to increase the desorption rate of Ga adatoms from the top c-plane

(thermal etching), the vertical growth will be reduced which leads to a reduction of the nanowire height and aspect ratio, as we observe in figure 2.12(c).

Having obtained the average volume of the nanowires from the data in figure 2.12 and the K parameter from section 2.3, the capture lengths for different pitch sizes at 924oC were calculated using equation 2.2 and are shown in figure 2.13. The capture length increases as the pitch increases and begins to saturate for the largest pitches. The pitch is approximately twice the capture length at a pitch of 1 µm. Below 1 µm, the pitch is less than twice the capture length.

Figure 2.13. Capture length versus pitch at a growth temperature of 926oC .

Above 1 µm, the pitch is more than twice the capture length and the capture length begins to saturate because the nanowires share fewer adatoms with adjacent nanowires [24].

2.7 Study the effect of the TMGa injection time on nanowire dimension

In order to take advantage of the Purcell effect, a reduction of the nanowire diameter is required. To reduce the nanowire dimeter below 100 nm, we performed a study on the effect of the TMGa injection time on the nanowire diameter. The TMGa injection times of 18, 15, and 12 sec were used while other growth parameters were kept the same as section 2.3. 10 pulsed-mode cycles were employed to grow the GaN nanowires at the growth temperature of 926oC. The average diameter of the GaN nanowires as a function of pitch size are shown in figure 2.14. The diameter of the GaN nanowires increases as the pitch size increases, which could be explained by increasing the capture length as a function of pitch size. The GaN nanowires with average diameter of ~95 nm was achieved for a TMGa injection time of 12 sec. The difference in the

Figure 2.14: The difference in diameter of the nanowires for a TMGa injection time of 18 sec and that of 12 sec. average diameter of the nanowires for a TMGa injection time of 18 sec and that of 12 sec is shown in figure 2.15. The slope of this difference is higher for smaller pitch sizes (less than

1µm) and reduces significantly as the pitch size goes over 1µm and saturates at 3µm. This means the diameter of the nanowires is more sensitive to the TMGa injection time for lower pitches than larger pitches. As shown in figure 2.13, a 1µm pitch size is where the nanowires stop sharing the adatoms since the capture length is close to half of the pitch. As long as the capture length is more than half of the pitch size, by increasing the TMGa injection times, a higher amount of TMGa adatoms participate in the growth of the nanowires compared to larger pitches in which the capture length is less than half of the pitch.

For the next experiment, we studied the effect of the TMGa injection time and TMGa flow rate while keeping the parameter of TMGa flow × TMGa injection time the same. For the first experiment, we have used the TMGa flow rate of 26.7 µmol min-1 with an injection time of

Figure 2.15: Average diameter of the nanowire as a function of pitch sizes for different TMGa injection times. 18 sec and for the second one a TMGa flow rate of 32.04 µmol min-1 with an injection time of 15 sec was used. Both of these conditions provide the value of 8.01 µmol of TMGa flowing into the chamber to grow the nanowires. Due to higher TMGa flow rate for a shorter lifetime, the Ga adatom desorption is lower in the first experiment compared to the second one. The results are shown in figure 2.16. As shown in this figure, for pitch sizes less than 1µm, these two conditions result in same size of nanowires. For pitch sizes less than 1µm, nanowires share the adatoms.

This means Ga adatoms that desorb from one nanowire can be absorbed be the adjacent nanowires and contribute in the growth of those nanowires. However for pitch sizes over 1µm, since the nanowires do not share the adatoms the condition with higher desorption rate resulted in smaller nanowires. Based on this finding, one way to reduce the GaN nanowire diameter is to increase the Ga adatom desorption rate. Increasing the TMGa interruption time, is an easy and effective way to increase the Ga adatom desorption rate. For the next set of experiments, the

Figure 2.16: Average diameter of the nanowire as a function of pitch sizes for different TMGa injection times and different TMGa flow rate, keeping the total amount of flow the same.

TMGa interruption time increased from 27 sec to 33 sec, and the average of the nanowire diameters for different pitch sizes were measured using SEM images. The results are shown in figure 2.17. By increasing the TMGa interruption time from 27 sec to 33 sec, nanowires with diameter of 74 nm were achieved [25].

2.8 Conclusion

In this chapter the pulsed mode MOCVD growth technique has been used to grow GaN nanowire arrays in a hexagonal symmetry. The effect of mask geometry, growth temperature,

TMGa injection time, and TMGa interruption time on growth, morphology and dimensions of

GaN nanowires were studied. We have shown that mask geometry has an important role in dimension and morphology of GaN nanowires. These findings pave the way for a deeper and a clear understanding of the growth of GaN nanowires using pulsed mode SAG. The effects of the

Figure 2.17: Average diameter of the nanowire as a function of pitch sizes for different TMGa interruption times. aperture diameter and pitch size on the growth of GaN nanowires were investigated, and we have shown that the pitch size plays an important role on the growth and morphology of GaN nanowires. The theory of capture length was proposed and the changes in the GaN nanowire dimensions and morphology were explained based on this theory. The effect of TMGa injection time on the nanowires diameter was studied. We show that the nanowires diameters reduce as the

TMGa injection time reduces. Using a 12 sec TMGa injection time, a GaN nanowire with a diameter of 95 nm was achieved. It is shown that the nanowire diameter is more sensitive to the

TMGa injection time in the regime where there is a strong competition for the growth nutrients between adjacent nanowires. This regime occurs where the capture length is more than half of the pitch sizes. The effect of the TMGa interruption time on GaN nanowire diameter was investigated. We have shown that by increasing the TMGa interruption time from 27 sec to 33 sec, GaN nanowires with a dimeter of 74 nm can be achieved.

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Chapter 3

Growth of GaN/InGaN core-shell nanowires

3.1 Effect of the aperture diameter on photoluminescence of GaN/InGaN core- shell nanowires

In this section the growth of InGaN quantum wells (QW) is presented, and the conditions required for the growth of QWs are given. After developing the growth conditions for the GaN nanowire cores (chapter 2), the coaxial growth of a single InGaN QW and a GaN barrier was added to the nanowire cores. For the first experiment, a single QW and barrier was grown around the nanowires with a constant pitch size of 1 µm and different aperture diameters. The aperture diameters were varied from 60 to 210 nm, with a step of 50 nm (see section 2.4). These nanowires were grown based on the growth conditions explained in chapter 2. A single QW and barrier were grown on the nanowire arrays, using the continuous mode MOCVD growth technique. The QW and barrier were grown without any GaN buffer layer on the core at 730oC

o and 790 C, respectively, with nitrogen (N2) as the carrier gas as well as triethylgallium (TEGa) and Ammonia (NH3) as the sources. The V/III ratio was maintained at 11,600 during the growth of QW and barrier, and the chamber pressure was maintained at ~200 Torr. The SEM images of the nanowires before and after growth of the QW/barrier are given in figures 3.1 and 3.2 respectively.

Micro-Photoluminescence (µ-PL) measurements were performed on the nanowire arrays using a frequency-doubled Ti:sapphire laser with an excitation wavelength of 405 nm. The samples were excited normal to the surface from the tops of the core-shell nanowires. A long

Figure 3.1. SEM images, at a 45-degree tilt, of GaN nanowires for 1 µm pitch and different aperture diameters (a) 60 nm, (b) 110 nm, (c) 160 nm, (d) 210 nm, before growth of a single quantum well and barrier. working distance (50x) micro-objective was used to excite a circular area with a diameter of ~20

µm. The emission was collected normal to the surface from the top using the micro-objective.

Figure 3.3 shows the experimental setup for conducting the µ-PL measurements of the

GaN/InGaN core-shell nanowires. A Coherent Chameleon titanium-sapphire pulsed laser, tuned to 810 nm with pulse widths of 140 fs and a repetition rate of 80 MHz, was used to excite the nanowires. The tunable range of the laser is 680 nm to 1080 nm, therefore frequency doubling by

Figure 3.2. SEM images, at a 45-degree tilt, of GaN nanowires for 1 µm pitch and different aperture diameters (a) 60 nm, (b) 110 nm, (c) 160 nm, (d) 210 nm, after growth of a single quantum well and barrier. a Coherent harmonics generator frequency doubler was required to obtain the excitation wavelength of 405 nm. From the frequency doubler, the laser light was passed through a continuous neutral density filter to select different excitation powers. A Semrock high-pass 405 nm dichroic mirror directed this reduced beam towards a Montana Instruments cryostation which housed the nanowires. The cyrostation was used to control the temperature of the nanowires, which ranged from 293 K to 10 K during the measurements. The photoluminescence emitted from the nanowire was composed of wavelengths longer than 405 nm which could pass through

Figure 3.3. Micro-photoluminescence set-up. the dichroic mirror. The PL was focused through a lens into an optical fiber, and coupled into an

Acton SP2300 Spectrometer. An Andor Tech Newton CCD camera then recorded the intensity spectrum from the spectrometer.

Figure 3.4 shows the normalized PL spectra for nanowires shown in figure 3.2, while figure 3.5 shows the estimated thickness of one period, which included one QW and barrier. The thickness of one period (QW/barrier) was determined using SEM, with a total thickness comprised of 40% QW and 60% barrier. The QW/barrier thicknesses are thicker for smaller aperture diameters. As the aperture diameter is increased, the QW/barrier thickness is reduced.

As shown in figure 2.9(b), the nanowire volume increases as the aperture diameter is increased.

This can explain the reduction of one period thickness by increasing the aperture diameter.

Figure 3.6(a) shows the measured PL peak wavelength, which decreases as the aperture diameter

Figure 3.4. Normalized PL spectra for different aperture diameters.

Figure 3.5. Estimated thickness of one period, including one QW and barrier.

58 increases. Surface band bending also has an effect, but band bending should increase for larger aperture diameter (thinner barriers), resulting in a red shift. The effects due to QW thickness and band bending on the PL wavelength were decoupled from the effects of indium composition using a commercial Schrӧdinger-Poisson solver (SiLENSe) to model the expected transition wavelength as a function of QW thickness at constant indium composition. With this approach, any difference remaining between the measured and simulated peak wavelengths is due to an increase in indium composition. The simulation parameters were set to establish 0.55 eV band bending at the surface of the n-type GaN nanowires to account for band bending in the QWs [1].

The expected peak wavelengths including only the effects of QW thickness and band bending

(i.e., for constant indium composition for all aperture diameters) were determined. For the largest aperture diameter (210 nm) an indium composition of 27% is needed to match the simulated peak wavelength with the measured PL peak wavelength. The indium composition was held constant and the QW and barrier thicknesses were then increased based on figure 3.5 to obtain the expected effect on the peak wavelength by changing the QW thickness and band bending on the peak wavelength if there is a constant indium composition. The simulated emission wavelengths with constant indium composition are shown in figure 3.6(a) (blue triangles). The difference in measured and simulated peak wavelength is shown in figure 3.6(b). For each aperture diameter, the indium composition in the simulation was then increased until the simulated peak wavelength matched the measured peak wavelength. The required increases in indium composition needed to match the simulation to the measured peak wavelength are shown in figure 3.6(b). The result shows that the estimated indium composition for the 60 nm diameter wires is 4% higher than that of the 210 nm diameter nanowires. The highest indium composition is observed for the nanowire with the smallest diameter, where there is not significant

Figure 3.6. (a) PL peak wavelength versus aperture diameter and simulated PL peak wavelength versus aperture diameter for constant 27% indium composition (b) Change in PL peak wavelength and estimated indium composition versus aperture diameter. The pitch spacing is 1 m. competition for adatoms. This occurs for given growth conditions when the sum of the nanowire radius and capture length is less than half the pitch. For larger aperture diameters, the nanowires compete for adatoms, which slows the growth rate and reduces the indium composition.

3.2 Study the effect of the pitch size on photoluminescence of GaN/InGaN core-shell nanowires

In this section the effect of pitch size on the optical properties of the GaN/InGaN core- shell nanowires is investigated. The pitch spacing between nanowires was varied from 300 nm to

4 µm, with a fixed aperture diameter of 40 nm. Six different nanowire arrays with pitches of 300 nm, 500 nm, 1, 2, 3, and 4 µm were grown. SEM images of the nanowires are given in figure

3.7. The nanowires were grown based on the growth conditions explained in chapter 2. A QW and barrier were grown on the nanowires using the continuous mode MOCVD growth technique.

The QW/barrier were grown with the growth conditions explained in section 3.1. SEM images

Figure 3.7. SEM images of GaN nanowires which illustrate the differences in growth morphology for pitch sizes = 300 nm (a), 500 nm (b), 1 μm (c), 2 μm (d), 3 μm (e), and 4 μm (f), before QW/barrier growth. of the nanowires after the growth of QW/barrier are given in figure 3.8. The target indium composition and QW thickness for the 1 µm pitch were ~25% and 3 nm, respectively, based on our planar c-plane GaN growth conditions.

µ-PL measurements were performed on the nanowire arrays using a frequency-doubled

Ti:sapphire laser with an excitation wavelength of 405 nm, as explained in section 3.1. The normalized PL intensity versus wavelength for each pitch is shown in figure 3.9. The spacing between the oscillation peak intensities in the PL spectra is consistent with a Fabry-Perot cavity formed between the nanowires and the ~1.9 m GaN template/sapphire interface. The peak near

560 nm for the 0.3 m pitch is attributed to yellow-band emission. The PL broadening can be attributed to non-uniform indium content on the nanowire sidewalls, indium-rich regions at the m-plane/semipolar plane junctions, or different indium incorporation rates on different planes [2-

4]. A false colored SEM of arrays of GaN/InGaN core-shell nanowires is shown in figure 3.10.

Figure 3.8. SEM images of GaN/InGaN core-shell nanowires for pitch = 300 nm (a), 500 nm (b), 1 μm (c), 2 μm (d), 3 μm (e), and 4 μm (f), after QW/barrier growth.

Figure 3.9. Normalized PL spectra for different pitches

Figure 3.10. False colored SEM of arrays of GaN/InGaN core-shell nanowires.

In order to estimate the thickness of the QW and barrier, the diameters of the nanowires were measured using SEM before and after the QW/barrier growth. An average was taken from

100 nanowires for pitch sizes less than 1 μm and from 50 nanowires for pitch sizes more than

1 μm. Thus, the reported one period thickness is the statistical average obtained for each pitch.

The standard deviation is reported as the error for each set of measurements. The thickness of one period is shown in figure 3.11. There is a clear increase in the thickness of the shell layers as the pitch is increased, leading to an increase in emission wavelength for larger pitches. Based on the growth rates of the QW and barrier, the QW comprises 40% and the barrier comprises 60% of the total thickness of one period, respectively. Thus, the QW thickness ranges from 2 to 10.4 nm and the barrier thickness ranges from 3 to 15.6 nm. The measured PL peak wavelength is shown in figure 3.12(a) (red squares). As the pitch increases, a significant red-shift is observed, which in general could be related to increasing QW thickness and/or increasing indium

Figure 3.11. Estimated thickness of one period, including one QW and barrier. composition. Surface band bending also has an effect, but band bending should decrease for larger pitches (thicker barriers), resulting in a blue shift. The effects of QW thickness and band bending on the PL wavelength were decoupled from the effects of indium composition using the technique explained in section 3.1. With this approach, any difference remaining between the measured and simulated peak wavelengths is due to an increase in indium composition. The simulated peak wavelength and the measured peak wavelength for the 300 nm pitch (2 nm QW) were then matched using an indium percentage of 22% in the simulation. The indium percentage was held constant and the QW and barrier thicknesses were then increased based on figure 3.11, to obtain the effect of changing QW thickness and band bending on the peak wavelength for a constant indium composition. The simulated emission wavelengths with constant indium composition are shown in figure 3.12(a) (blue triangles). The difference in the measured and simulated peak wavelengths are shown in figure 3.12(b). The simulated and measured peak

Figure 3.12. (a) PL peak wavelength versus pitch and simulated PL peak wavelength versus pitch for constant 22% indium composition, (b) Change in PL peak wavelength and estimated change in indium composition versus pitch. The aperture diameter is 40 nm. wavelengths are well aligned for the first two pitches, but differ significantly for pitches of 1 m and above, indicating that the actual indium compositions in the simulation for the larger pitches should be higher to achieve the measured peak wavelengths. For each pitch, the indium composition in the simulation was then increased until the simulated peak wavelength matched the measured peak wavelength. The required increases in indium composition needed to match the simulation to the measured peak wavelength are shown in figure 3.12(b). The result shows that the estimated indium composition in the 4 m pitch is nearly 10% higher than that of the

300 nm pitch. By increasing the pitch from 300 nm to 4 m, the capture length of the adatoms

(Fig. 2.13) increases, thus leading to an increase in the growth rate that is associated with achieving higher indium incorporation [5]. Similar to the trend in capture length, the indium composition increases sharply until a pitch of 2 µm and then saturates since the capture length becomes less than half of the pitch size and the nanowires share fewer indium adatoms.

3.3 Study of temperature dependent of photoluminescence and carrier dynamics of GaN/InGaN core-shell nanowires for different pitch sizes

Measuring integrated photoluminescence intensity (IPLI) as a function of temperature is an effective way to investigate localization effects and the quantum efficiency of the GaN/InGaN core-shell nanostructures. Temperature dependent IPLI for pitch sizes of 300 nm, 1µm, and 3

µm, was performed. PL spectra of the nanostructures were measured using the µ-PL set up, explained in section 3.1 (Fig. 3.3). Excitation power was kept at 25 mW, with an excitation wavelength of 405 nm, and a laser spot size of 20 µm. The normalized IPLI as a function of reciprocal temperature for three different pitch sizes of 300 nm, 1 µm, and 3 µm are shown in figure 3.13. An anomalous behavior of temperature-dependent photoluminescence of

GaN/InGaN core-shell nanowires with a pitch size of 300 nm is shown in figure 3.13(a). Figure

3.13(b) shows the normal behavior of the temperature dependent IPLI which is related to thermal quenching [6,7], and figure 3.13(c) shows a rise of IPLI at 112 K. As shown in figure 3.13(a), the

IPLI has a peak at close to 200 K instead of 10 K. This behavior is in contrast with thermal quenching effect. Thermal quenching behavior is explained by thermal escape mechanism or thermally-activated nonradiative recombination mechanisms [6,7], which is the typical behavior for planar structures with InGaN/GaN [6] and InGaAs/GaAs [7] quantum wells/barriers. These two mechanisms have similar expressions, thus it is assumed that the nonradiative recombination is principally contributed by the carriers at the high energy levels [8].

Anomalous behavior of the IPLI as a function of temperature has been previously observed for GaN/InGaN quantum dots (QDs) and single QWs structures [9] (see Fig. 3.14).

This phenomenon is attributed to temperature induced carrier redistribution inside shallow and deep localized states within the active region [9,10]. The localization states for InGaN active

Figure 3.13. Integrated photoluminescence of GaN/InGaN core-shell nanowires for (a) 300 nm (b) 1 µm, (c) and 3 µm pitch size. regions are mostly caused by indium inhomogeneity, quantum well thickness variation, and interface fluctuations. For a structure consisting of QDs with the lateral dimension of 10 nm and the height of 1.5 to 2 nm, a significant enhancement in IPLI was observed as the temperature increased from 150 K to 215 K [10]. The increase in IPLI was related to carrier redistribution inside the potential minima caused by localization effect inside the active region [10,11]. As shown in figure 3.11, for GaN/InGaN core-shell nanowire arrays for a pitch size of 300 nm, the measured thickness of one period (one quantum well and barrier) is close to 5 nm, with the total thickness comprised of 40% QW and 60% barrier. This results in QW thickness of 2 nm. It’s expected that the QW thickness variation on the m-plane sidewalls of the nanowires results in formation of the QDs on the sidewall. Formation of the QDs on the sidewalls of the nanowires gives rise to this abnormal behavior of the IPLI. A schematic mechanism of the possible carrier redistribution inside shallow and deep localized states within the QDs is shown in figure 3.15. At a low temperature of 10 K, carriers are possibly distributed among both deep and shallow potential minima caused by potential fluctuations and the radiative recombination process is dominant. As the temperature increases to 25 K, some of the shallow localized carriers are thermally activated and relax down into other deep localizations via hopping, leading to the

Figure 3.14. Measured PL peak intensity (normalized at 10 K) versus temperature for (a) an InGaN single QW sample, (b) an InGaN QDs sample [9]. initial increase in IPLI. Typically, as the temperature increases, the defect-related nonradiative recombination becomes more dominant when compared to radiative recombination, this results in thermal quenching and reduction of IPLI [6,7]. The unusual increase in the IPLI -as the temperature varied from 100 K to 200 K can be understood by assuming that the deep localized carriers are responsible for the dominant recombination mechanisms in these QDs structures

[10,11]. Due to higher effective crystalline quality in deep localization minima [11-13] and isolation from nonradiative centers, it can be assumed that the radiation efficiency of carriers in deep localization areas is higher. Higher radiative efficiency gives rise to a more rapid increase in emission intensity. As the temperature increases, the carriers escape from the shallow localizations and then converge to fill the deep localization centers by relaxation and recapture processes. Hence, the carriers consumed by radiative recombination in deep localizations can be

68 compensated rapidly. As a consequence, a majority of carriers would recombine in the deep localizations with higher radiation efficiency and eventually lead to an anomalously enhanced emission [11]. This is shown in figure 3.15(c). For temperature above 200 K, the defect assisted nonradiative recombination cannot be ignored anymore, and thermal escape from the deep localizations becomes dominant. During the relaxation and recapture processes, some portion of carriers would be captured by the nonradiative centers, which suppress the compensation of carriers that are consumed by radiative recombination, which results in a decreasing emission intensity. For temperatures over 200 K, the carriers may have sufficient energy to repopulate the shallow localizations, and interact more with the nonradiative recombination centers. This process, along with thermal escape of the carriers from deep localization minima, results in a rapid thermal quenching of the IPLI.

For a pitch size of 1 µm the IPLI shows the normal thermal quenching behavior, which has been reported for InGaN/GaN quantum well/barrier structures [6]. With this pitch size, the thickness of one period (one quantum well and barrier) is close to 8.5 nm (Fig. 3.11), where the total thickness is comprised of 40% QW and 60% barrier. The QW thickness is calculated to be

3.4 nm. It’s believed that quantum well uniformity should be higher for a 3.4 nm thick QW as compared to 2 nm QW in an array with a 300 nm pitch. Higher uniformity within the QW results in a reduction of localized minima inside the quantum well. So the abnormal behavior of the

IPLI as a function of temperature is less dominant compared to deep localization effects observed in arrays with a 300 nm pitch size. However a small increase in IPLI has been observed as the temperature increases from 50 K to 65 K, which can be due to minor localization effect inside the QW.

Figure 3.15. Schematic diagrams indicating the possible mechanism of temperature induced carriers’ redistribution inside shallow and deep localized states inside the active region, (a) 10 K, (b) 25 K, (c) 150 K, (d) and 200 K.

For a pitch size of 3 µm, an increase in IPLI has been observed as the temperature increases from 50 K to 112 K. A PL peak wavelength of 525 nm has been observed and a quantum well thickness of 9.2 nm has been calculated for an array with this pitch size of GaN/InGaN core-shell nanowires. The combination of this quantum well thickness and peak PL intensity, results in indium incorporation of 31% (Fig.3.12(b)). Higher quantum well thickness results in higher degree of uniformity within the active region. However, it is shown that higher indium incorporation in the InGaN quantum wells results in higher indium segregation inside the wells

[14-16]. Indium segregation gives rise to formation of localized centers with potential minima inside the quantum well [12]. We suspect that the increase in IPLI for temperature rising from 50

K to 112 K is due to formation of this type of localized center.

3.3.1 Study of carrier dynamics

In order to further elucidate the kinetics of the carrier recombination, time-resolved photoluminescence (TRPL) measurements were performed as a function of temperature for 300 nm, 1 µm, and 3 µm pitch sizes. Figure 3.16, shows the time-resolved PL setup used to measure the carrier lifetime of the GaN/InGaN core-shell nanowires at different temperatures and excitation powers. The excitation source was a Coherent Chameleon Ti:Sapphire pulsed laser with a repetition rate of 80 MHz, which was tuned to 810 nm, and the output frequency was doubled to have a corresponding wavelength of 405 nm. The 405 nm emission was directed into a micro-objective lens by reflecting off a dichroic mirror, and was focused onto the nanowire- based LED. After the excitation of the nanowire-based LED, the photoluminescence passed through the dichroic mirror, next through a lens, and into the Acton SP2300 Spectrometer, where it was eventually fed into a Hamamatsu C5680 streak camera. The streak camera was equipped with the Hamamatsu M5678 blank unit to measure the time variation of the intensity of the PL

Figure 3.16. Time-Resolved Photoluminescence set-up.

over a time range of 1 nanosecond. To obtain the instrument response function (IRF) of the streak camera to the excitation source, a TRPL measurement was performed on scattered light originating from the 405 nm laser beam. Figure 3.17 shows a few examples of the room temperature PL transients for excitation powers of 0.1, 10, and 100 mW on nanowires within a 1

µm pitch array. The instrument response function (IRF) is also shown and verifies it is much shorter than the measured lifetimes. As shown in figure 3.17, the normalized TRPL decay curves with varied excitation powers exhibit two obvious decay stages, which are relatively fast in the early stage and slower in the extended range. The decay curves can be well fitted with a bi- exponential function [18-21]

Figure 3.17. Few examples of the room temperature PL transients for excitation powers of 0.1, 10, and 100 mW for 1 µm pitch size.

−푡 −푡 ⁄휏 ⁄휏 퐼(푇) = 퐴1푒 1 + 퐴2푒 2 (3.1)

where A1 and A2 are the amplitudes and τ1 and τ2 are the time constants of the fast decay and slow decay components, respectively. The slow decay component is considered to be the PL lifetime [19,20,22,23]. Figure 3.18 shows the PL lifetimes (extracted from the slow decay component) as a function of reciprocal temperature for 300 nm, 1 µm, and 3 µm pitch sizes. As shown in figure 3.18, the PL lifetime reduces as the temperature increases. The onset of reduction of the PL lifetime, for all three different pitches, corresponds to the point at which IPLI shows a significant change. For 300 nm, and 3 µm pitch sizes this point corresponds to the point at which the IPLI start increasing, and the point which the IPLI starts reducing for a 1 µm pitch.

It is suspected that the reduction of the PL lifetime for a pitch size of 300 nm as the temperature is increased from 100 K to 200 K, can be attributed to reduction of the radiative lifetime, which

Figure 3.18. PL lifetime as a function of reciprocal temperature for (a) 300 nm, (b) 1 µm, (c) and 3 µm pitch size. can be due to accumulation of the carriers inside deep localized centers. As the temperature increases over 200 K, the nonradiative centers become more dominant in PL lifetime reduction.

For pitch size of 1 µm pitch size, it is believed that nonradiative recombination is the main mechanism for reducing the PL lifetime, for temperatures above 100 K. For a 3 µm pitch size we suspect PL lifetime begins reducing at 50 K, which corresponds to the point where IPLI starts to increase slightly. This can be attributed to a reduction of the radiative lifetime due to increasing the number of deep localized carriers. The PL lifetime and potential 3dB modulation bandwidth

(equation 3.2) for 0.3, 0.5, 1, 2, 3, and 4 µm pitch sizes and different excitaion powers are shown in figure 3.19 (a), 3.19(b) respectively.

1 푓(3푑퐵) ≈ (3.2) 2휋휏푝푙

Where 휏푝푙 is PL lifetime. As shown in figure 3.19(a) the PL lifetime reduces as the excitation power increases. By increasing the excitation power, the carrier density inside the active region increases, which results in a reduction of the PL lifetime. As the nanowire diameter increases the active region thickness also increases (Fig. 3.11), which gives rise to a reduction of the carrier density inside the active region for the same excitation powers. This all leads to an increase in the PL lifetime.

Figure 3.19. (a) PL lifetime, (b) Potential 3dB modulation bandwidth as a function of excitation power for 300 nm, 500 nm, 1, 2, 3, and 4 µm pitch sizes.

3.4 Conclusion

In this chapter the growth of InGaN quantum wells on the GaN cores were studied. One quantum well and barrier were grown around the GaN nanowire cores, and µ-PL measurements were obtained and analyzed for a variety of nanowire dimensions, array pitches, and aperture diameters. An 80 nm redshift in the peak PL wavelength was observed as the nanowire pitch spacing was increased from 300 nm to 4 µm, while increasing the aperture diameter from 60 nm to 210 nm resulted in a blueshift of ~35 nm. The thickness of one period (QW/barrier) as a function of pitch and aperture diameter were inferred using SEMs, which were thicker as the pitch was increased. The effects of pitch and aperture diameter on indium composition were decoupled from the measured PL by simulating the expected peak wavelength for a constant indium composition. Significant increases in indium composition were predicted for larger pitches and smaller aperture diameters. Temperature dependent IPLI and TRPL were performed on GaN/InGaN core-shell nanowire arrays. For 300 nm and 2 µm pitch seize array an anomalous

75 behavior of the IPLI as a function of temperature was observed. This phenomenon is attributed to temperature induced carrier redistribution inside shallow and deep localized states inside the active region. The source of these types of localization effects can be related to formation of QDs on the sidewalls of the GaN nanowires during the growth of thin (less than 2nm) quantum wells, or indium segregation for quantum wells with high amounts of indium incorporation. This detailed study provides significant insight into the effects of mask configuration and growth conditions on the GaN/InGaN core-shell nanowire properties and is applicable to the fabrication of monolithically integrated, multi-color GaN-based light-emitters on a single chip, which can potentially be used in multichannel VLC systems with high data rates.

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Chapter 4

Growth, fabrication, and characterization of electrically injected GaN/InGaN core-shell nanowire light-emitting diodes

4.1 Introduction

In this chapter the growth, fabrication and characterization of electrically injected

GaN/InGaN core-shell nanowire micro-light-emitting diodes (µ-LEDs) is presented. µ-LEDs are expected to become the next generation of pixel-level emitters in display technology. These displays include indoor/outdoor video walls, smart phones, tablets, televisions, smart watches, and head-mounted displays. µ-LEDs offer potential advantages compared to conventional organic LEDs (OLEDs) and liquid crystal displays (LCDs). Some advantages include higher brightness, higher transparency, robustness, lower power consumption, and shorter response times [1-4]. III-nitride based µ-LEDs exhibit a luminance of 105 cd/m2, while the luminescence of LCDs and OLEDs are 3000 cd/m2 and 1500 cd/m2, respectively [4]. Additionally, µ-LEDs exhibit nanosecond response times, in contrast to the millisecond and microsecond response times of LCDs and OLEDs, respectively [4]. The short response times of µ-LEDs pave the way for the next generation of indoor LED-based data communication, known as light-fidelity (Li-Fi)

[5]. The brightness, etendue, and acceptable light efficiency of LED-based projection systems are also strongly dependent on the far-field emission pattern of the individual devices [6]. A small far-field radiation angle of ±15-20° is critical in full-color LED displays specially LED projection systems [6,7]. The photonic crystal (PhC) effect has been used to engineer the far- field emission pattern of conventional planar LEDs, although the processing of photonic crystal structures poses challenges [8-12]. Alternatively, µ-LEDs based on nanowire structures

(nanowire-based LEDs) enable engineering of the far-field emission pattern using the periodic nature of the nanowires themselves. Indeed, a directional far-field emission pattern has been achieved experimentally using nanowire-based LEDs [7,13]. In addition, in chapter three it is shown that nanowire-based emitters offer an approach to monolithically integrated RGB-based white LEDs and lasers [14]. These monolithically integrated multi-color µ-LEDs can be achieved in a single growth by varying the pitch spacing between the nanowire pattern, potentially providing significant cost reduction due to fewer processing steps and consumed sources compared to pick-and-place approaches. In addition to display technology, RGB-based white µ-LEDs are preferred over phosphor-converted approaches for visible-light communication (VLC) systems with high data rates [15] due to the slow response of the phosphors in conventional phosphor-converted white LEDs [16,17] Dynamic color tuning is also achievable using a multi-color approach [18], which could potentially lead to multiplexing within VLC over several wavelength channels. Interferometric lithography [19], or nanoimprint lithography can be used to pattern the dielectric mask. These two approaches are more industrial compatible compared to e-beam lithography.

The techniques, explained in chapters two and three were utilized to grow GaN nanowires and quantum wells respectively. The work presented in this chapter addresses the issues regarding the fabrication and characterization of GaN/InGaN core-shell micro-LEDs (µ-

LEDs). A detailed study of optical and electrical characterization is performed.

4.2 Growth of GaN/InGaN core-shell nanowire-based LEDs

The growth of GaN nanowires was carried out on a selectively patterned c-plane GaN template on c-plane sapphire. The processing steps required for patterning of dielectric mask

81 using interferometric lithography (IL) [19], and selective-area growth (SAG) of GaN nanowires are given in this section. The process began with the growth of a 2-µm-thick GaN layer doped with silicon (n-GaN) on a 2 inch, single-side-polished sapphire substrate. A thin (less than 10 nm) low temperature (550oC) GaN buffer layer was grown on sapphire before the growth of the

GaN template. A conventional method of continuous mode MOCVD and typical group III

(TMGa) and group V (NH3) precursors had been used to grow the GaN template. GaN template growth carried out at the pressure of 500 Torr, and the growth temperature of 1090oC, with a group V to group III ratio (V/III ratio) of 3200. The rotation speed was held at 1500 rpm during the growth. The wafer was then cleaned in piranha solution (mixture of sulfuric acid and hydrogen peroxide 4:1) and 100 nm of SiNx was deposited using plasma-enhanced chemical vapor deposition (PECVD) at 300oC. ICON-7 was sputtered as an anti-reflecting layer (ARC), followed by Ultra I 123.8 photoresist (PR). Then IL was performed on the sample to pattern the

PR. To reverse the pattern 15 nm of Ni was deposited on the sample, followed by lift-off.

Secondary contact lithography was performed on the sample to define the mesa. Reactive ion etch (RIE) was used to etch the dielectric and ARC, then the sample was cleaned in Piranha and loaded into the MOCVD reactor to grow GaN nanowires. The processing steps required to make nanowire template are shown in figure 4.1. The IL set-up including the laser (Coherent 4-100

Trippled ND:YAG at 355nm), beam expander, mirror, and stage are shown in figure 4.2. A schematic of the IL set-up, examples of 1-D and 2-D patterning are shown in figure 4.3. To perform the 2-D patterning two exposures are required one at zero degrees (compare to the edge of the mirror), and the other one at 60 degrees. 1-D patterning results in formation of lines in the

PR while 2-D patterning results in formation of elliptical apertures in the PR.

Figure 4.1: Processing steps for nanowire template using interferometric lithography. (1) 2 µm GaN template grown on Sapphire, (2) Deposition of 100 nm of SiNx using plasma-enhanced chemical vapor deposition, (3) Spin coat of ICON-7 and photoresist, (4) interferometric lithography, (5) Reverse the pattern using Ni deposition and lift-off, (6) Secondary lithography to define the mesa, and RIE of ARC and SiNx,(7) Piranha cleaning, (8) Growth of GaN nanowires, (9) Microscopic image of nanowire template ready for growth.

Figure 4.2: Interferometric lithography set-up including laser (Coherent Infinity 40-100), beam expander, mirror, and stage.

Figure 4.3: Schematic of the IL set-up and two different patterning. 1-D patterning results in formation of lines and 2-D patterning results in formation of quasi circular apertures.

Once the nanowire template was ready, it was loaded into our MOCVD reactor. Pulsed mode

growth technique (section 2.3) was used to grow the GaN nanowires. Growth conditions were

kept the same as section 2.3. Figure 4.4 shows the GaN nanowires grown with 150 cycles under

the growth conditions explained in section 2.3. After the growth of GaN nanowires, four pairs of

quantum wells and barriers were grown around the nanowires. The growth conditions used to

grow quantum wells and barrier were similar to what was explained in section 3.1. After the

growth of the active region, a p-GaN layer was grown around the nanowires. Biscyclo

pentadienyl magnesium (CP2Mg) was used as a p-dopant. The p-GaN layer was grown at the

o growth temperature of 920 C, pressure of 200 Torr in N2 environment with a V/III ratio of 7550.

The SEM images of the GaN nanowires are shown in figure 4.4(a) and figure 4.4(b) shows the

SEM images of the GaN nanowires after the growth of active region and the p-GaN layer. As

shown in figure 4.4(b), the p-GaN layer coalesced at the base of the nanowires, and the final

structure looks more like truncated pyramids, rather than nanowires.

Figure 4.4: SEM images of GaN nanowires (a) before (b) after the growth of the quantum wells, barriers, and p-GaN.

4.3 Electrical characterization and the effect of AlGaN underlayer on device performance

Shadow masking is an easy and effective technique which provides access to the electrical characteristics of samples without needing to go through complicated and long fabrication processes. A mask with circular apertures and diameters of 75, 100, 200 and 300 µm was used to deposit 20 nm of palladium followed by 300 nm of gold as p-contacts. A scratched region at one corner of the sample was used as an n-contact. The current density-voltage (J-V) plot of our first generation of electrically injected GaN/InGaN core-shell nanowire-based LED is shown in figure 4.5. As shown in this figure the J-V plot shows a significant reverse-leakage current density at −5 (푉), and a series resistance of ~748 훺. High reverse-leakage current density and high series resistance reduce the performance of the LED. The electroluminescence

(EL) spectra at 10 A/cm2 and 75 A/cm2 are shown in figures 4.6(a) and 4.6(b) respectively. A

Figure 4.5: Room temperature current density-voltage plot of a GaN/InGaN core-shell nanowire LED under continuous condition. Inset shows the J-V plot from -5 V to 0 V.

Figure 4.6: EL spectrum at (a) 10 A/cm2, (b) 75 A/cm2.

86 peak EL of 512 nm and 479 nm has been observed at 10 A/cm2 and 75 A/cm2. The shift in EL spectra for nanostructure-based LEDs is previously reported, and attributed to different indium incorporation at the top and the bottom portion of the nanostructures [20, 21]. Reduction of the reverse-leakage current density is the first task that was addressed, to improve the performance of the LEDs. To have an estimate of the magnesium (Mg) doping level in top p-GaN layer secondary-ion mass spectroscopy (SIMS) was tried on GaN/InGaN core-shell nanowires LEDs.

However, due to complexity of the structures (2-D periodic structure) we were not able to achieve accurate SIMS results. Therefore, a more simple structure (triangular nanostripe), with

1-D periodic structures was used. SEM images of the triangular nanostripe before and after the growth of quantum well, barrier, and p-GaN layer are shown in figures 4.7(a) and 4.7(b) respectively. Triangular nanostripes have been grown selectively in continuous growth mode using SiNx as a dielectric growth mask [20, 21]. Processing steps with 1-D patterning (section

4.2) have been used to make nanostripes in SiNx dielectric layer. Triangular nanostripe cores were grown using trimethylgallium (TMGa) and ammonia (NH3) as the precursors at a V/III ratio of 2000. The core was also doped with silicon to achieve n-type conduction and was grown in a mixed H2–N2 atmosphere (H2:N2=2:1). Four pairs of InGaN (3 nm)/GaN (9 nm) quantum wells/barriers were then grown using trimethylindium (TMIn), triethylgallium (TEGa), and

NH3 under a high V/III ratio (12,000). The quantum wells and barriers were grown at the growth temperature of 770oC and 810oC respectively. This was followed by the growth of 200 nm of magnesium-doped GaN (p-GaN) for device characterization and 500 nm of p-GaN to perform

SIMS. A thick layer of p-GaN has been used to coalesce the top p-GaN layer, and make flat surface which is preferred for SIMS performance. A room temperature current density-voltage plot of triangular nanostripe in continuous mode is shown in figure 4.8. As shown in this figure

Figure 4.7: SEM images of pyramidal nanostripes (a) before (b) after the growth of the quantum well and p-GaN.

Figure 4.8: Room temperature current density-voltage plot of a triangular nanostripe LED under continuous condition. the triangular nanostripe also exhibit a high reverse-leakage current density at −5 (푉). SIMS measurements were performed to quantify the magnesium and other impurity concentrations including oxygen and silicon in the p-GaN layer. Previously, micro-SIMS was utilized to study impurity profiles in multiple layers on similar triangular stripes [22]. The Figure 4.9 shows the

Figure 4.9: SIMS proﬁle of the top p-GaN layer of a triangular nanostripe LED vs. depth. The inset shows cross section of triangular nanostripes.

SIMS measurement profile of the top p-type GaN grown on the triangular nanostripes. The SIMS profile indicates a very large concentration of silicon and oxygen in the top 0.5 mm of the film, despite no intentional flow of Silane (SiH4) or oxygen-containing precursors during the growth of the p-GaN. The likely source of these type impurities is the ex-situ deposited silicon nitride

(SiNx) dielectric layer, which is patterned and used as a mask for the bottom-up selective-area nanostructure growth. As explained in section 4.2 PECVD was employed to deposit the dielectric layer for the SAG nanostructures at 300°C and at a pressure of 1 Torr. The refractive index of this SiNx was determined by ellipsometry to be ~1.78. This is significantly lower than the anticipated refractive index of SiNx (~2.01) [23]. Claassen et al. have attributed the lower

89 refractive index to a significant concentration of Si atoms bonded with hydrogen, instead of nitrogen, within the Si3N4 network. This arises when NH3 is used as the primary source of nitrogen [24]. Under high temperature MOCVD growth conditions, the weak Si-H bonds [25,

26] in the SiNx growth dielectric dissociate and serve as the source of Si. Si may readily incorporate onto gallium sites and serve as a donor-type impurity in GaN, making p-GaN growth challenging because the intended p-type dopant (Mg) also prefers to incorporate onto gallium sites.

Furthermore, the formation energy of Si donors on Ga sites in the GaN crystal lattice is lower than that of magnesium [27]. This formation energy is even lower when the Fermi level is closer to the valence band (Mg-doped GaN). Thus, the Mg incorporation into a vacant Ga site may also be reduced by the presence of silicon. The measured level of ~4 − 6 × 1019 cm-3 is lower than expected based on the CP2Mg flow rate during the growth. Another reason for the lower refractive index of our PECVD SiNx could be incorporation of O during the PECVD arising from unintentional leaks in the deposition chamber [28]. This is also evident from the

SIMS data, which indicates a very high O concentration within the first 0.5 μm of the device. O atoms prefer to incorporate on nitrogen sites [29] and also behave as n-type dopants. The collective incorporation of Si and O in the Mg-doped p-GaN results in a large reverse leakage current density and poor p-n junction characteristics, as shown in the inset of Fig. 4.8. The large leakage current prevents the fabrication of efficient electrically injected devices. This effect is also magnified in SAG nanostructured LEDs whose p-GaN layer is not coalesced like our test structures.

One way to passivate the dielectric layer and prevent the dissociation of Si and O into the p-GaN, is inclusion of a thin layer of aluminum gallium nitride (AlGaN) [30]. Growth of a

Figure 4.10: (a) SEM image of triangular nanostripe GaN cores, (b) SEM image of triangular naostripe LEDs with AlGaN underlayer. The inset shows an XTEM image indicating the position of the AlGaN underlayer.

thin layer of n-type AlGaN before the growth of quantum wells and barriers also results in reduction of Si and O impurities during the growth of the active region. This resulted in improving the quality of the quantum wells, and the LED performance. It is well known that the reaction mechanism between TMAl and NH3 (leading to AlN) involves a number of reaction intermediates (concerted mechanism) [31], which result in pre-reactions and the formation of

AlN without the need for a conducive surface (e.g., single crystal GaN). In the case of bottom-up

SAG, a polycrystalline AlN layer is formed on the growth mask. This is in contrast with the reaction mechanism between TMGa and NH3 (leading to GaN) that proceeds via homogeneous pyrolysis and radical reaction intermediates, predominantly involving the gallium or nitrogen atoms already bonded to the GaN surface. The growth of polycrystalline AlN is shown in figure

Figure 4.11: SIMS proﬁle of the top p-GaN layer of a triangular nanostripe LED vs. depth after the growth of n-type AlGaN underlayer. The inset shows cross section of the triangular nanostripes.

4.10. Figure 4.10(a) shows SEM image of GaN triangular nanostripe cores, and figure 4.10(b) shows SEM image of triangular nanostripe LEDs with n-type AlGaN underlayer. The AlGaN underlayer was grown after the growth of nanostripe cores, with the growth temperature of

o 940 C, pressure of 200 Torr, and H2, N2 and NH3 flow rates of 2020,1010, and 1010 sccm. The

V/III ratio was kept constant at 3800 during the growth of the AlGaN underlayer with TMGa,

TMAl, and Silane flow rates of 5.64 and 6.04 µmol/min, and 20 nmol/min, respectively. This amount of Silane results in a dopant level of 8 × 1017cm-3 in gas phase for the n-type AlGaN layer. Figure 4.10(b) clearly shows the formation of polycrystalline AlN on dielectric mask. This polycrystalline AlN layer not only resists desorption of silicon and oxygen from the SiNx by

Figure 4.12: Reverse leakage current density at -5 V measured vs. AlGaN underlayer growth time. The inset shows schematic of triangular nanostripe with and without electron blocking layer (EBL). forming a masking layer, it also getters oxygen atoms due to the high affinity of aluminum to oxygen. This is evident from figure 4.11 that shows the SIMS profile of the top p-GaN layer for triangular stripe LEDs with an AlGaN underlayer. The measured silicon concentration in the top p-GaN layer is at least 3-4 orders of magnitude lower than that in figure 4.9 (LED without

AlGaN underlayer). The measured oxygen concentration in the LED with the AlGaN underlayer is also at least 50 times lower compared to the LED without the underlayer. The effect of reduction of Si and O impurities on reverse-leakage current density was studied using triangular nanostripe LEDs with 200 nm of p-GaN. The reduction in Si and O concentrations in the p-type

GaN directly correlated with reduced reverse-leakage current density measured at -5 V, as shown in figure 4.12. A substantial reduction in the reverse-leakage current density is observed with the addition of just a 30 second n-AlGaN growth step, which corresponds to ~20 nm AlGaN

Figure 4.13: Reverse-leakage current density at -5 V measured vs. TMAl flow rate.

thickness on the semipolar (101̅1) triangular stripes. The leakage current density at -5 V was further reduced with an increase in the growth time (thickness) of the AlGaN underlayer. The effect of a p-type AlGaN electron-blocking layer (EBL) was also studied. The p-type AlGaN was grown after the quantum wells/barriers and before the p-GaN layer. As shown, in figure 4.12,

EBL to be less effective in reducing the leakage current, in the absence of the underlayer, The presence of Si and O impurities in the sample without the n-AlGaN underlayer leads to doping compensation in the EBL, thereby leading to a less effective corresponding Mg-doped p-GaN growth that follows immediately. In contrast, the samples with AlGaN underlayers exhibit even lower reverse leakage current density, which also further reduced with increasing TMAl flow during the growth of the AlGaN underlayer, as shown in 4.13. This results from an increase in the pre-reaction rate of the TMAl and NH3 at higher TMAl flow rates, which produces a thicker polycrystalline AlN on dielectric mask. The current density-voltage (J-V) plot of a triangular

Figure 4.14: Room temperature current density-voltage plot of a triangular nanostripe light emitting diode under continuous condition. nanostripe LED with 20 nm of n-type AlGaN under layer is shown in figure 4.14. As shown in this figure the reverse-leakage current density at -5 V is reduced significantly.

4.3.1 Growth of AlGaN underlayer on GaN nanowires

In this section the effect of n-type AlGaN underlayer on performance of GaN/InGaN core-shell nanowire-based LEDs is studied. Figure 4.15(a) shows the GaN nanowire grown under pulsed mode condition explained in section 2.3. 150 cycles of pulsed mode growth were used to grow nanowires with diameter of 500-600 nm and the height of 1.1-1.2 µm. After the growth of GaN nanowires in pulsed mode growth, n-type AlGaN underlayer was grown under continuous growth conditions explained in previous section for 7.5 minutes. Figure 4.15(b) shows the nanowires after the growth of AlGaN underlayer. After the growth of the AlGaN under layer four pairs of quantum wells and barriers were grown around the nanowire followed

95 by a ~200 nm of p-GaN layer. The J-V plot for the device is shown in figure 4.16. As shown in this figure, the reverse-leakage current density reduced significantly compared to figure 4.5. The reverse-leakage current density reduced from 8.2 A/cm2 to 2.2 A/cm2, and the series resistance reduced from 748 Ω to 358 Ω. Despite, 4 times reduction in reverse-leakage current density and

2 times reduction in series resistance the obtained values for reverse-leakage current density and series resistance are far from those for high performance planar LEDs with reverse-leakage current density and series resistance in the range of 0.05 A/cm2 and 25 Ω. So, more improvement is needed to put the GaN/InGaN core-shell nanowire-based LEDs at the level of high performance planar LEDs.

As shown in figure 4.15(b), the semipolar (101̅1) surface area has been increased significantly after the growth of n-type AlGaN underlayer. Due to H2 passivation, the growth rate of semipolar (101̅1) facet is much lower than that of m-plane sidewalls and top c-plane [32-34].

This results in a very thin layer of p-GaN grown on semipolar facets. Thin p-GaN layer gives rise to higher leakage current for electrically injected devices. One way to overcome this problem is passivation of semipolar facets using a thin dielectric layer to reduce the reverse-leakage current

[35, 36], which requires some extra processing steps. Alternatively, a more conformal growth of

AlGaN underlayer might be a potential way to reduce the semipolar facets. Reduction of semipolar facet surface area might result in reduction of reverse-leakage current, and series resistance. Due to locally lowered of V/III ratio and a better control on the growth of 3D structures, pulsed mode growth technique is a feasible way to grow a more conformal AlGaN underlayer around the nanowires. A schematic image of the pulsed mode growth of n-type

AlGaN underlayer is shown in figure 4.17. During pulsed mode growth of AlGaN underlayer, H2 and N2 flow continuously while TMGa, TMAl, and NH3 are pulsed. The AlGaN underlayer was

Figure 4.15: SEM images of GaN nanowires (a) before, (b) after the growth of AlGaN underlayer, using continuous mode growth technique.

Figure 4.16: Room temperature current density-voltage (J-V) plot of a GaN/InGaN core-shell nanowire light emitting diode with AlGaN underlayer grown in continuous mode. The measurement was performed under CW mode. The inset shows J-V plot from -5 V to 0 V.

grown in pressure of 90 Torr, rotation speed of 1500 rpm, H2 flow of 3000 sccm, and N2 flow of

1000 sccm. The TMGa and TMAl injection time, TMGa and TMAl interruption time, NH3

Figure 4.17: Schematic image of different steps in each cycle for pulsed-mode growth of AlGaN

underlayer, (a) TMGa and TMAl injection time, (b) TMGa and TMAl interruption time, (c) NH3

injection time, (d) NH3 interruption time.

Figure 4.18: SEM images of GaN nanowires (a) before, (b) after the growth of AlGaN underlayer using pulsed-mode growth technique. injection time, and NH3 interruption time were 18 sec, 27 sec, 6 sec, and 15 sec, respectively.

-1 -1 The TMGa flow rate, TMAl, and NH3 flow rate were 26.7 µmol min , 6.04 µmol min , and 8 mmol min-1, respectively with a V/III ratio of 81. 25 cycles have been used to grow AlGaN

Figure 4.19: (a) SEM image of GaN/InGaN core-shell nanowire-based LEDs, (b) Schematic

image of the device showing sapphire substrate, GaN template, SiNx dielectric layer, AlGaN underlayer, MQW, and p-GaN.

Figure 4.20: Room temperature current density-voltage (J-V) plot of a GaN/InGaN core-shell nanowire light emitting diode with AlGaN underlayer grown in pulsed mode. The measurement was performed under CW mode. The inset shows J-V plot from -5 V to 0 V. underlayer. Figure 4.18(a) shows the GaN nanowires, and figure 4.18(b) shows the GaN nanowires after the growth of AlGaN underlayer. Both GaN nanowires and AlGaN underlayer were grown in pulsed mode. As shown in figure 4.18(b) a more conformal growth of AlGaN has

99 been achieved using the pulsed mode growth technique. Four pairs of quantum wells and barriers were grown on the nanowires shown in figure 4.18(b) under the growth condition explained in section 3.1. A thin layer of p-type AlGaN layer was grown after the growth of the quantum wells as an electron blocking layer (EBL), followed by ~ 200 nm of p-GaN layer to make GaN/InGaN core-shell nanowire light emitting diodes. SEM image of the GaN/InGaN core-shell nanowire- based LED and schematic of the device are shown in figures 4.19(a) and 4.19(b) respectively.

Electrical characterization is performed and J-V plot of the device is shown in figure 4.20. Using pulsed mode growth of AlGaN underlayer the reverse-leakage current density was reduced to

0.06 A/cm2 and the series resistance was reduced to 289Ω.

4.4 Growth characterization

Scanning transmission electron microscopy (STEM) is a powerful technique that enables characterization and study of thin-film materials. STEM images of GaN/InGaN core-shell nanowires in atomic-contrast (ZC) mode, are shown in figure 4.21. As shown in figures 4.21(a) and (b), GaN nanowires with diameter of 600 nm and height of 1.2 µm were achieved and a fairly uniform AlGaN underlayer was grown around the GaN nanowires. Figure 4.21(c) shows the growth of the AlGaN underlayer on top portion of the GaN nanowire including top c-plane and semipolar planes. As shown in figure 4.21(d) the AlGaN underlayer thickness increases from bottom to top of the nanowire. The thickness of the AlGaN underlayer was measured using

STEM images. 28.5, and 13.7 nm of AlGaN underlayer were measured on top c-plane and semipolar planes of the GaN nanowires. The thickness of the AlGaN underlayer changes from

8.5 nm at the bottom to 19.5 nm at top part of the GaN nanowire. The thickness of the EBL, also follows the same trend and changes from 2.4 nm at the bottom to 6.6 nm at the top of the GaN nanowires. Figure 4.22(a) to (f) show high resolution STEM images in ZC mode of the AlGaN

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Figure 4.21: STEM images of GaN/InGaN core-shell nanowires in atomic-contrast (ZC) mode showing (a) GaN nanowires with diameter of 600 nm and height of 1.2 µm. (b) Conformal growth of AlGaN around the nanowires. (c) Growth of AlGaN at the top portion of the nanowires, (d, e) Growth of AlGaN layer at the sidewalls of the nanowires, (f) Growth of AlGaN on semipolar facet. underlayer, quantum wells, and EBL for top (a and b), middle (c and d), and bottom (e and f) of the GaN/InGaN core-shell nanowire LED. These figures clearly show the change in the thickness of the AlGaN underlayer, quantum well, and EBL moving from top to bottom of the structure.

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Figure 4.22: STEM images of a GaN/InGaN core-shell nanowire LED in atomic-contrast (ZC) mode showing AlGaN underlayer, quantum wells, barriers, and EBL (a) at the top, (b) middle, (c) bottom of the GaN/InGaN core-shell nanowire.

The thickness of the quantum well changes from 2.0 nm at the bottom to 2.54 nm at the top of the nanowires. The images show a quantum well thickness variation of only 0.54 nm along 1.2

µm of the GaN nanowire height.

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4.5 Fabrication of GaN/InGaN core-shell nanowire light-emitting diodes

This section addresses the fabrication of GaN/InGaN core-shell nanowire-based LEDs.

Device fabrication requires different processing steps including, indium tin oxide (ITO) deposition, inductively coupled plasma (ICP) etching, metal deposition, and different steps of photolithography and cleaning. ITO is required to have a more uniform current spreading along the mesa. ITO deposition is the first step of the fabrication. However before ITO deposition pre- cleaning step is required to remove any residue left over on the sample. Pre-cleaning has been done using dilute hydrochloric acid (HCL) and deionized water (DI) (1:3) for 3 minutes and then

DI rinse for another 3 minutes. The sample was then loaded in a CHA dielectric deposition machine to deposit 150 nm to 200 nm of ITO. Conventional contact photolithography was performed to cover the mesa with photoresist (PR) followed by ICP etching to etch the ITO,

SiNx, and 100 to 200 nm of the GaN template. In the absence of H2 cooling using ICP etching, high gradient of temperature from the substrates to top PR layer, results in PR reflow. PR reflow gives rise to reduction the verticality and smoothness of the etched facet. This leads to reduction of the etched facet angle from 80-90 degree to 30-40 degree, and rough facets. SEM images of some damaged facets using ICP etching are shown in figure 4.23. As shown in this figure the etched facets were damaged and a very poor quality etching profile was achieved with low verticality. In addition, we noticed that using mung to adhere the sample to the substrate, results in non-uniform etching profile along the sample. To overcome issues regarding the ICP etching we used an H2 cooling system, and replaced mung with photoresist AZ 4330 to have a more uniform etching profile. SEM images of the etched facets under optimized conditions of ICP etching are shown in figure 4.24. As shown in these figures damage-free facets with a more vertical sidewalls were achieved, under optimized ICP etching conditions. After ICP etching

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Figure 4.23: SEM images of poor quality etched facets using ICP etching. the sample was cleaned in 1165 developer to remove remaining PR. The sample was annealed at

o 550 C for 10 minutes in N2 environment, using rapid thermal annealer machine (RTA). Figure

4.25(a) shows a schematic image of the sample (top view) and a SEM image showing the mesa region with grown nanowire-based LEDs and a size of 120 µm by 120 µm. Figure 4.25(b) shows cross section of the device and different layers including blanket deposit of ITO, p-GaN, multiple quantum wells (MQW), AlGaN underlayer and GaN cap layer. Figure 4.26(a) and (b) show schematic image of the cross section of the device after ICP etching, and annealing respectively.

After annealing, the ITO becomes more transparent. Transparent ITO is required for high light extraction efficiency. Figure 4.26(c), and (d) show schematic image (top view) of the mesa and etched region of the ITO after annealing, and SEM image of the sample showing the mesa and

ITO region. Conventional contact lithography was performed to define regions for n-contact deposition. Titanium, aluminium, nickel, gold, has been used as n-contact with thickness of

10,100,100, and 300-500 nm respectively. The thickness of the last layer (gold) was designed to

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Figure 4.24: SEM images of etched facets using ICP etching under optimized conditions.

o level the height of the n-contact with the ITO deposited on the SiNx. 1165 solution at 80 C has been used for lift-off process, and clean the sample. Figures 4.25(a) and (b) show schematic illustration of the sample (top view) after n-contact deposition, and cross section of the device showing the n-contact and other elements. Figures 4.25(c) through (d) show false colored SEM images of the device and the yellow highlighted n-contact. Conventional contact lithography was used to define regions for p-pads, and n-pads. 10 nm of chromium and 500 nm of gold have been used as p-pad and n-pad. Figures 4.26 (a) and (b) show the schematic illustration (top view) of the device and cross section of the nanowire-based LED after p-pad deposition. Figures 4.26(c) and (d) show false-colored SEM images of the full fabricated showing the n-pads, p-pads, n- contact, and ITO. More details about the fabrication processes are given in appendix A. Figure

4.27 shows schematic illustration of the cross section of a device and a full fabricated

GaN/InGaN core-shell nanowire-based LED.

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Figure 4.25: (a) Schematic of the GaN/InGaN LEDs viewed from the top, showing the mesa and SEM image of the grown region, (b) Schematic image of the GaN/InGaN core-shell nanowire LED after ITO deposition.

Figure 4.26: Schematic image of the GaN/InGaN core-shell nanowire-based LED (a) after ICP etching, (b) after annealing (c) Schematic image of the devices viewed from the top, showing the mesa and ICP etched region, (d) SEM images of mesa and etched ITO region.

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Figure 4.27: (a) Schematic image of LEDs viewed from the top showing the n-contact, mesa and ITO region (b) cross section of the LED showing the n-contact deposited region, (c) to (d) SEM images of a nanowire-based LED showing the n-contact (yellow color).

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Figure 4.28: (a) Schematic image of the full fabricated devices viewed from the top, (b) Schematic of the cross section of an LED, (c) and (d) SEM images of fully fabricated devices showing the n-contact, n-pads and p-pads.

Figure 4.29: Schematic images of (a) device cross section and (b) device layout for a GaN/InGaN core-shell nanowire-based LED.

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4.6 Optical characterization

The PL characteristics of the nanowire-based LEDs were measured to evaluate the fundamental active region quality in the absence of electrical injection effects. Micro- photoluminescence (µPL) measurements were performed using 405 nm excitation from a frequency-doubled Ti:Sapphire laser ( see section 3.1). A long working distance (50x) micro- objective was used to excite a circular area with a diameter of ~15 µm. The pumping wavelength

405 nm ensures photo-generation of carriers only within the active region and enables uniform pumping of the quantum wells. The room-temperature (RT) µPL as a function of laser excitation power is shown in figure 4.30(a). The µPL peak wavelength and full-width at half maximum

(FWHM) are shown in figures 4.30 (b) and 4.30 (c), respectively. The absence of defect related yellow-band emission in the µPL spectra is indicative of high-quality, bright quantum wells. As the excitation power increases from 0.1 to 75 mW, the µPL peak wavelength shifts from 483 nm to 462 nm and the FWHM decreases from 53 nm to 43 nm. For nonpolar active regions in nanowires, the blueshift in the PL peak wavelength is attributed to the band-filling effect and non-uniform indium distribution along the nanowires [37, 38].

To further understand the quality of the quantum wells, the internal quantum efficiency

(IQE) of the LEDs was measured. We applied the conventional method of dividing the integrated

µPL at room temperature (RT) by that at low temperature (LT) -10oK- to measure the IQE at different excitation levels [39]. Figure 4.31(a) shows the measured IQE versus excitation power, and the inset shows the µPL measurements at an excitation power of 10 mW at RT and LT. The

IQE versus excitation power plot provides important information about the efficiency droop for the nanowire-based LED. Previously, other groups only reported the peak IQE, which ranged from 8% to 58% in various studies [38, 40-42]. The peak IQE here is 62%, which is among the

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Figure 4.30: (a) PL spectra at different excitation powers (b) PL peak wavelength at different excitation powers (c) PL FWHM at different excitation powers.

110 highest reported for GaN/InGaN core-shell nanowire-based LEDs. Figures 4.31(b) and 4.31(c) show the integrated PL versus excitation power at LT and RT, respectively. At LT (10oK) the slope of the integrated PL vs. excitation power is close to 1 for excitation powers less than 1 mW, which indicates radiative recombination is dominant [43]. The slope decreases to 0.61 at higher excitation powers. Even at LT, a sub-linear slope at higher excitation powers has been observed [44] and can be attributed to either Auger recombination [45,46], absorption saturation

[47-49], or generation of hot carriers [50]. The integrated PL versus excitation power at RT has a slope of 1.27 for excitation powers below 1 mW. This clearly indicates the presence of a combination of radiative and non-radiative recombination [43]. The slope approaches 1 for excitation powers ranging from 1-10 mW, which suggests the radiative recombination rate has become dominant over the non-radiative rate [43]. The slope decreases to 0.65 for excitation powers greater than 10 mW, indicating Auger recombination or carrier leakage has become dominant [43, 51]. The change in the slope of the integrated PL (see Fig. 4.31(c)) at 10 mW directly corresponds with the peak of the IQE. The IQE is maximum at 10 mW and exhibits efficiency droop at higher excitation powers.

In addition to high quantum efficiency, the response time of the LEDs is critical for their implementation in displays, VLC, and Li-Fi technologies. The carrier recombination rate is a key parameter in determining the response time. Time-resolved photoluminescence (TRPL) was used to measure the carrier lifetime of the nanowire-based LEDs at RT and LT for excitation powers ranging from 0.1 mW to 75 mW (see section 3.2). The PL transients were fit using the bi-

−푡 −푡 ⁄휏 ⁄휏 exponential decay function, 퐴1푒 1 + 퐴2푒 2, where A1 and A2 are amplitudes and τ1 and τ2 are the time constants of the fast decay and slow decay components, respectively. The slow decay component is considered the PL lifetime (see section 3.3). Figure 4.32(a) shows the PL

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Figure 4.31: Optical characterization of GaN/InGaN core-shell LED (a) IQE measurement using low temperature/room temperature integrated PL technique (b) Integrated PL vs. excitation power in log-log scale at low temperature. (c) Integrated PL vs. excitation power in log-log scale at room temperature.

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lifetimes (extracted from the slow decay component) versus excitation power at RT and

LT. The inset of figure 4.32(a) shows a few examples of the RT PL transients for excitation powers of 1, 5, and 25 mW. The instrument response time (IRF) is also shown and verifies it is much shorter than the measured lifetimes. The PL lifetime measured for the nanowire-based

LED shows a minimum of 1.3 ns. This lifetime is at least 3 times shorter than that of typical planar c-plane LEDs, which are in the range of 4-20 ns at high excitation powers [52, 53]. The shorter lifetime in the nanowire-based LEDs is mostly attributed to the higher electron-hole wave function overlap for QWs grown on the m-plane side walls of the nanowires, rather than nonradiative surface recombination since the p-GaN is ~ 200 nm thick. Shorter carrier lifetimes provide the possibility of higher 3dB modulation bandwidth.

Figure 4.32 (a) shows the PL lifetime decreases as the excitation power increases both at

RT and LT. Also, the PL lifetime at LT is longer than that at RT for all excitation powers. The difference between the PL lifetime at LT and the PL lifetime at RT is highest at low excitation powers. At low excitation powers and LT, radiative recombination is dominant (slope  1 in Fig.

4.31(b)), while at RT for the same excitation powers both radiative and non-radiative recombination exist (slope > 1 in Fig. 4.31(c)). The combination of radiative and non-radiative recombination lowers the PL lifetime at RT.

Having obtained the IQE and PL lifetime data at RT, the radiative and non-radiative lifetimes can be decoupled using equations (4.1) and (4.2),

휏푅 = 휏푃퐿⁄퐼푄퐸 (4.1)

휏푁푅 = 휏푃퐿⁄(1 − 퐼푄퐸) (4.2)

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where 휏푃퐿 is the PL lifetime, 휏푅 is the radiative lifetime, and 휏푁푅 is the non-radiative lifetime. Figures 4.32 (b) and 4.32 (c) show the radiative and non-radiative lifetimes versus excitation power respectively. As the excitation power increases from 0.1-10 mW, the radiative lifetime decreases. For excitation powers beyond 10 mW, the radiative lifetime remains fairly constant. The onset of the constant radiative lifetime above 10 mW in Figure 4.32 (b) corresponds to the point at which the slope of the integrated PL versus excitation power changes to 0.65 in figure 4.31 (c). The non-radiative lifetime in Figure 4.32 (c) continues to decrease for excitation powers above 10 mW. At high excitation powers, both Auger recombination and carrier leakage are present based on the slope = 0.65 in figure 4.31(c). Auger recombination is the dominant process in reducing the non-radiative lifetime in figure 4.32 (c). Carrier leakage reduces the carrier density inside the active region. This results in increasing the no-radiative lifetime, and opposes the reduction in non-radiative lifetime and eventually leads to the saturation of the non-radiative lifetime at higher excitation powers.

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Figure 4.32: Optical characterization of GaN/InGaN core-shell LED (a) PL lifetime measured at room temperature and low temperature, (b) Radiative lifetime at room temperature, (c) Nonradiative lifetime at room temperature

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Figure 4.33 Simulation of a GaN/InGaN core-shell nanowire-based LED. (a) Schematic of the simulated nanowires as viewed from the top and the simulation region. (b) Schematic of the cross section of nanowire-based LEDs showing different layers, top monitor to calculate extraction efficiency, and simulation region.

4.7 Far field emission pattern

The far-field emission pattern of µ-LEDs has an important role in their performance. A relatively directional output power, ranging from ±15-20o, is preferred for full-color displays

[6,7]. As mentioned in the introduction such far-field emission patterns have been achieved by utilizing the photonic crystal effect for conventional planar structures [10-12]. A commercial- grade simulator based on the finite-difference time-domain (FDTD) method (Lumerical FDTD

Solutions) was used to simulate the far-field emission pattern and light extraction efficiency

(EXE) of the nanowire-based LED structure shown in figure 4.19(a). The FDTD method is a fully vectorial approach that naturally gives both the time domain and frequency domain information [54]. The geometry of the nanowire-based LEDs was extracted from the SEM and

STEM images. A cross section of the nanowires, and the simulation region used, are shown in figures 4.33(a) and 4.33(b). A simulation region of 2 µm by 1.7 µm was selected as a unit cell.

Perfectly matched layer (PML) boundary conditions were used for the top z-axis boundary condition, periodic boundary conditions were used for x and y axes, and metal boundary conditions were used for the bottom z-axis. To represent the quantum well emission, we used the

116 common approach of incorporating a radiating electric dipole source at the location of the active region [55]. The emission from nonpolar InGaN quantum wells is well modeled by a single dipole source directed along the a-direction of the wurtzite crystal [56, 57]. Six simulations were performed for each dipole at each side of the nanowire-based LED. The results were added incoherently to find the extraction efficiency and far-field emission pattern of the LED using a monitor at the top of the structure.

The angular distributions of extracted light along the x and y-axes are shown in figures

4.34(a) and 4.34(b). These angular distributions exhibit multiple lobes and a relatively directional emission pattern compared to the Lambertian-type far-field pattern of conventional planar LEDs [58, 59]. The higher directionality is attributed to the periodic nanowire structure, which enhances the diffracted power normal to the LED surface. A directional far field emission pattern ranging between ±15o is predicted for our nanowire-based LED. Directional far-field emission patterns in the range of ±30o and below have been achieved by other research groups using nanowire-based LEDs [7,13]. This unique property makes nanowire-based LEDs a good candidate for the next generation of high brightness displays. We note that no extra PhC patterning or micro lenses are needed to obtain this relatively directional emission pattern from nanowire-based LEDs. EXE from the top surface of the nanowire-based LED was also calculated. An EXE of 13.2% from the top surface was calculated for the nanowire-based LED, which is higher than the simulated EXE (8.1%) from the top surface of a planar LED.

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Figure 4.34: (a) Simulated angular distribution of extracted light for a nanowire-based LED along (a) x-axis (b) y-axis.

4.8 Electrical characterization

While optical characterization techniques provide useful information for understanding the quantum well quality, these techniques do not provide information on the properties of the

LED under electrical injection. These properties play a critical role in the performance of nanowire-based LEDs and include carrier transport, reverse-leakage current, series resistance, and turn-on voltage. Figure 4.35(a) shows the continuous-wave J-V plot for a nanowire-based

LED with a 120-µm mesa size. The current density (J) is calculated based on the planar footprint. If the effective nanowire surface area is used to calculate the current density, the current density would be reduced by a factor of ~2 in figures 4.35 and 4.36. Figures 4.35(b) and

4.35(c) show the light–current density–voltage (L–J–V) characteristics and estimated EQE under room-temperature pulsed operation (2% duty cycle, 2 µs pulse width), respectively. The J-V plot does not exhibit any reverse leakage current. The LED shows a turn-on voltage of 2.9 V.

Deposition of 150 nm of ITO resulted in reduction of series resistance to 25 Ω. The combination of a low turn-on voltage and low series resistance places this device among the highest performing of those reported for GaN/InGaN core-shell nanowire-based LEDs in terms of

118 current-voltage characteristics and internal quantum efficiency [36, 40-42]. With advanced LED packaging techniques unavailable for this work, the light extraction efficiency (EXE) was simulated to enable an estimate of the total output power and external quantum efficiency (EQE) of the device (section 4.6). Assuming an injection efficiency (IE) of 1, the total extracted power from the top surface of the LED (P) was calculated using equation (4.3). Here, h is Plank’s constant, υ is the LED electroluminescence (EL) emission frequency, q is electron charge, and I is current. In this calculation, we assume that the peak IQE coincides with the peak of the EQE

(i.e., EXE does not depend upon injection current)

P=EQE hυ/q I (4.3) where,

EQE=EXE×IQE. (4.4)

Figures 4.36(a), 4.36(b), and 4.36(c) show the electroluminescence (EL) spectra, peak wavelength, and FWHM at different current densities, respectively. As the current density increases from 80 A/cm2 to 1.9 kA/cm2, the EL peak wavelength exhibits a blueshift from 452 nm to 444 nm. The shift is attributed to non-uniformities in the indium incorporation in different regions of the nanowires [36-38,41]. Large blueshifts in the EL peak wavelength of nanowire- based LEDs, ranging from 62 nm to 180 nm, have been observed by other research groups [33,

36]. As the current density increases from 80 to 500 A/cm2 the FWHM of the EL spectra decreases from 52 nm to 38 nm. The FWHM remains fairly constant from 500 A/cm2 to 1.25 kA/cm2 and increases to 40 nm for higher current densities.

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Figure 4.35. (a) J-V plot. Inset image shows an operating LED with 120 µm by 120 µm mesa size (b) L-J-V curve for a device with 120 µm by 120 µm mesa size (c) Estimated EQE.

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Figure 4.36: (a) EL spectra, (b) EL peak wavelength, (c) and FWHM at different current densities.

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4.9 RF characterization

4.9.1 Introduction

Wireless data traffic is growing exponentially due to the evolution of computing, consumer electronics, and mobile communication technologies. A recent forecast by Cisco indicates that mobile communication traffic doubles every two years [61]. This means if the global traffic volume increases at the current rate, by the end of 2022 the traffic volume will be

10 times more than current traffic. Since the available radio frequency (RF) communication spectrum is very limited, there will soon be a ‘spectrum crisis’ as RF technology cannot keep pace with the demand [62]. The available RF spectra are efficiently utilized by using advanced signal processing concepts such as massive multiple-input multiple output (MIMO) systems and shrinking cell size leading to femtocells [62]. Figure 4.37 shows the United States frequency allocations in radio spectrum, which clearly indicates the radio frequency spectrum limitation for the next generations of RF communication systems. A potential solution to the looming spectrum crisis lies in the migration of wireless communication into the visible light spectrum [63]. The availability of hundreds of THz of the unregulated spectrum makes visible light communication

(VLC) attractive for wireless communications. VLC based on green and blue LEDs through plastic optical fiber (POF) [64], free-space, and under water wireless optical communication

(UWOC) [65] have gained significant attention in the last years. 3dB modulation bandwidth of

LEDs is a key parameter in performance of LED-based VLC technology. D. Tsonev et.al have employed orthogonal frequency division multiplexing (OFDM) to reach data rate of 3-Gb/s using micro-LEDs with 3-dB modulation bandwidth of 60 MHz [63]. X. R. Ferreira et.al have utilized

DC biased optical OFDM to reach data rate of 5-Gb/s in free space using micro-LED arrays with

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Figure 4.37: United States frequency allocations in radio spectrum, (reported by NTIA 2016).

3-dB modulation bandwidth of 800 MHz [66]. Wavelength division multiplexing has been used to reach data rates of 8 and 10 Gb/s using micro-LEDs [67,68].

4.9.2 RF characteristics of GaN/InGaN core-shell nanowire-based LEDs

In this section, RF characterization of GaN/InGaN core-shell nanowire-based LEDs is presented. To study the effect of LED size on RF characteristics, nanowire-based LEDs were fabricated on four different square mesas sizes with 60, 80, 100, and 120 µm sides. Figure

4.38(a) shows schematic illustration of the RF set-up including network analyzer (NA) (Agilent

PNA-X N5247A), bias-tee (ZFBT-4R2GW-FT+), micro-RF ground-signal-ground probe

(ACP40-GSG-150, Cascade Microtech, Inc.), high-speed photodetector (PD) (DET025AFC,

ThorLabs, Inc.), low noise amplifier (LNA) (PE15A1009, Pasternack Enterprises), and optical

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Figure 4.38: (a) Schematic illustration of the RF set-up, (b) to (d) experimental set-up for the RF measurement. Inset in (d) shows the RF probe, operating LED, and optical fiber. fiber. Figures 4.38(b) to (d) show the experimental set-up of the RF characterization. The inset in figure 4.38(d) shows a microscopic image of the LED mesa, RF probe, and optical fiber.

The frequency of the small signal was swept from 10 MHz to 1 GHz. The nanowire- based LEDs were probed using the micro-RF probe (Fig. 3.38 (d)) and the modulated output light was fiber coupled and collected into a high-speed photodetector. The LNA was used to amplify the signal received by the photodetector, which was then coupled into port 2 of the NA to determine the modulation response (S21). The real and imaginary parts of the frequency response were extracted from the measured reflection coefficient (S11), which was collected over the same frequency range as S21.

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Figure 4.39: (a) J-V plot for different mesa sizes, the inset show microscopic image of four different mesa sizes and operating LED with 60 µm mesa size, (b) 3-dB modulation bandwidth for four different mesa sizes. The inset shows few examples of normalized frequency response for a device with 80 µm mesa size.

The J-V plots for four different mesa sizes are shown in figure 4.39(a) and 3-dB modulation bandwidth of the devices are shown in figure 4.39 (b). The inset in 4.39 (b) shows few examples of normalized frequency response for a device with 80 µm mesa size. As shown in figure 4.39 (b), the 3-dB modulation bandwidth saturates at ~ 140 MHz independent of device sizes. However, before reaching a saturation point, bigger devices have shown higher 3-dB bandwidth modulation. Higher 3-dB bandwidth modulation for bigger device sizes might be due to current crowding. Figure 4.40 shows operating LEDs in a pulsed mode with 1% duty cycle and 1 µs pulse width, with different mesa sizes at different current densities. As shown in this figure, current crowding is clearer for a device with 120 µm mesa size.

To better understand the saturation of the 3-dB modulation bandwidth, a small-signal equivalent circuit model was derived to have an excellent fit of real and imaginary part of S11 and absolute value of S21 to the experimental data [69, 70]. Scattering matrix element method has

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Figure 4.40: LEDs with different mesa sizes operating in pulsed mode (1% duty cycle with 1 µs pulsed width) at different current densities, showing current crowding for 120 µm mesa size. been employed to have a complete description of the two-port network as seen at each port.

While the impedance and admittance relate the total voltage and currents at the ports, the scattering matrix relates the voltage waves incident on the ports to those reflected from the ports

[71]. For a two-port network system shown in figure 4.41, the scattering matrix [S], is defined as follow:

Figure 4.41: An arbitrary two-port microwave network system.

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− + 푉1 푆11 푆12 푉1 [ −] = [ ] [ +] (4.5) 푉2 푆21 푆22 푉2

+ - Where Vn is the amplitude of the voltage wave incident on port n and Vn is the amplitude of the voltage wave reflected from port n. A specific element of the scattering matrix can be determined as:

− 푉푖 푆푖푗 = + |푉+=0 푓표푟 퐾≠푗 (4.6) 푣푗 푘

In other words, equation 4.6 says that Sij is found by driving port j with an incident wave

+ - of voltage Vj and measuring the reflected wave amplitude Vi coming out of port i. The incident waves on all ports except the j th port are set at zero, which means that all ports should be terminated in matched loads to avoid reflectance [71]. S parameters can be used to fully characterize a microwave network system. In the case of two-port network system it is convenient to define a 2 × 2 transmission, or ABCD matrix for each two-port network. Figure

4.42 shows a simple model of a two port-network system, where:

푉 퐴 퐵 푉 [ 1] = [ ] [ 2] (4.7) 퐼1 퐶 퐷 퐼2

Figure 4.42: An arbitrary two-port microwave network system described by ABCD matrix elements.

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Figure 4.43: The small-signal equivalent circuit of the GaN/InGaN nanowire-based LED

Having obtained the ABCD matrix elements, S parameters can be easily calculated.

Figure 4.43 shows an equivalent circuit model for the GaN/InGaN core-shell nanowire- based LED [69, 70], this model is equivalent to a two port-network structure shown in figure

4.44, where:

′ 푍2푍3 푍1 = 푅푠 + 푍2 + 푍3 + 푍4

′ 푍4푍3 푍2 = (4.8) 푍2 + 푍3 + 푍4

′ 푍2푍4 푍3 = 푍2 + 푍3 + 푍4

And:

1 푍2 = 푗휔푐푑

푍3 = 푅푑 (4.9)

푅푄푊 푍4 = 1 + 푗휔푐푄푊 푅푄푊

And ABCD matrix elements are:

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Figure 4.44: Equivalent two-port network system for the GaN/InGaN nanowire-based LED.

′ ′ ′ 푍1 ′ ′ 푍1푍2 퐴 = 1 + ′ 퐵 = 푍1 + 푍2 + ′ 푍3 푍3 (4.10) ′ 1 푍2 퐶 = ′ 퐷 = 1 + ′ 푍3 푍3

S11 and S21 can be calculated from ABCD elements:

퐵 퐴 + ⁄푍 − 퐶푍0 − 퐷 푆 = 0 (4.11) 11 퐵 퐴 + ⁄ + 퐶푍0 + 퐷 푍0

2 (4.12) 푆 = 21 퐵 퐴 + ⁄ + 퐶푍0 + 퐷 푍0

Where Z0 is 50 Ω.

To achieve a unique solution for the shared unknown parameters in the circuit model, the real and imaginary part of S11 (equation 4.11) and modulation response (20 × 퐿표푔(퐴푏푠(푆21)) of the circuit are simultaneously fit to the same experimental data sets from the GaN/InGaN core- shell nanowire-based LED, and the results for one data set are shown in figure 4.45. Nonlinear regression is used to achieve excellent fittings to the RF data without pre-defining or assuming

129

Figure 4.45: Simultaneous fitting of (a) real and imaginary parts of Eq. (4.11) to experimental data, (b) 20×log (abs(Eq. (4.12)) to the measured modulation response of a 100 µm mesa size device at a current density of 100 A/cm2. any parameters based on DC measurements [69, 70]. Excellent agreement between the theory and experiment justifies that the proposed model (figure 4.43) is an accurate representative of the

GaN/InGaN core-shell nanowire-based LED. Carrier recombination time constant and diode time constant have been defined as 휏푟푒푐 = 푅푄푊퐶푄푊 and 휏푑 = 푅푑퐶푑 respectively [69, 70]. Both

휏푟푒푐 and 휏푑 have been shown in figure 4.46. As shown in this figure at high current densities,

휏푟푒푐 is in the range of 1.3 to 2 ns while 휏푑 is in the range of 0.2 to 0.5 ns. This means the 3-dB modulation bandwidth of the nanowire-based LED is probably limited by carrier lifetime rather than diode time constant. Thus to increase the 3-dB modulation speed a novel active region design is required to lower the carrier recombination lifetime. Plasmonics is a potential way to enhance the radiative recombination rate through Purcell effect. Enhancement of the radiative rate leads to reduction of recombination lifetime, and 3-dB modulation bandwidth enhancement

[72]. In the next chapter we will discuss about the simulation and design of plasmonic

GaN/InGaN nanowire LEDs, with a potential 3-dB modulation speed in GHz range.

130

Figure 4.46: Carrier recombination time constant and diode time constant.

4.10 Conclusion

In chapter four, electrically injected GaN/InGaN core-shell nanowire-based LEDs were presented. Interferometric lithography was used to fabricate nanowire templates. Pulsed mode

SAG was employed to grow the GaN nanowires. It was shown that an AlGaN underlayer is required to prevent the desorption of Si and O from dielectric mask. Dissociation of Si and O from SiNx dielectric mask gives rise to compensation of the p-GaN layer, reduction of the efficiency of this layer, and a significant reverse-leakage current density. We have shown that growth of an AlGaN underlayer leads to the growth of polycrystalline AlN layer on the dielectric mask and passivation of this layer. To achieve conformal growth of the AlGaN underlayer, a pulsed mode growth technique was used. GaN/InGaN core-shell nanowire-based LEDs were fabricated and electro-optical and RF characteristics were presented. Peak IQE of 62% and series resistance of 25 Ω were achieved. These results are among the highest performance levels for

131 nanowire-based LEDs thus far. TRPL was used to study the carrier dynamics of the LEDs at room temperature and low temperature. In addition, FDTD modeling revealed that the nanowire- based LEDs have a strongly directional far-field emission pattern. Properties such as high IQE, short carrier lifetime, and emission directionality are attractive for solid-state lighting, visible- light communication, and -LED displays, respectively. The RF characteristics of GaN/InGaN core-shell nanowire-based LEDs were presented, and an equivalent circuit model was proposed for the LED. Using scattering matrix elements and the ABCD transfer matrix method, the S11 and

S21 parameters were evaluated and fit to the experimental data. Using the fitting parameters we decoupled the carrier recombination lifetime from the diode time constant. It was shown that the

3-dB modulation bandwidth of the LEDs is limited by carrier lifetime rather than diode time constant.

132

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139

Chapter 5

Simulation and design of plasmonic GaN/InGaN core-shell nanowire-based

LEDs

5.1 Introduction

In this chapter, plasmonic GaN/InGaN core-shell nanowire-based LEDs are discussed, and the potential 3-dB modulation bandwidth of each design is presented. Metal-clad semiconductor nanocavities have attracted significant attention as a means to enable continued size scaling of photonic devices beyond the diffraction limit. The primary motivations for these devices are their potential for thresholdless lasing [1], high-speed modulation [2-4], low power consumption [5,6], and dense integration. These features may be attractive for optical communications, chemical and biological sensing, and fluorescence imaging. Among devices based on metal-clad nanocavities, plasmonic nanolasers (and spasers) have been heavily researched as the nanoscale counterparts to conventional dielectric diode lasers [7]. A key consideration in the operation of plasmonic nanolasers is overcoming the high optical losses associated with the metal cladding layers. Although optically pumped [8,9] and electrically injected [10-15] plasmonic nanolasers have been theoretically predicted and experimentally realized, the practicality of electrically injected devices with strong mode confinement in three dimensions remains an open question [16,17]. More specifically, the threshold current density of a plasmonic nanolaser with three-dimensional mode confinement (i.e., a spaser) was predicted to be ~ 1 MA/cm2 [18]. Thus, the requirement of low power consumption may be difficult to attain in room temperature electrically injected plasmonic nanolasers. Furthermore, concerns have been raised regarding the linewidth and speed of plasmonic nanolasers [18]. As an alternative,

140

Figure 5.1: Maximum frequency (3-dB modulation bandwidth) as a function of current density for different structures including (left to right) a typical LED [23], a plasmonic LED (SPED) [4], a vertical- cavity surface-emitting laser (VCSEL) [24], and a plasmonic nanolaser (SPASER) [25].

plasmonic nanoscale light-emitting diodes (PNLEDs), or surface plasmon emitting diodes

(SPEDs), utilizing incoherent spontaneous emission may provide many of the attractive

characteristics of plasmonic nanolasers without requiring large operating current densities [17-

21]. PNLEDs have the potential for excellent high-speed performance [22] and can operate with

lower power consumption. These devices are designed to operate below the lasing threshold and

utilize the enhanced spontaneous emission rates achievable in highly resonant cavities.

Figure 5.1 [18] shows maximum frequency (3-dB modulation bandwidth) as a function of

current density for different structures including (left to right) a typical LED [23], a plasmonic

LED (SPED) [4], a vertical-cavity surface-emitting laser (VCSEL) [24], and a plasmonic

nanolaser (SPASER) [25]. As shown in this figure for pump current density in the range of 1

kA/cm2, typical LED structures have the maximum frequency in the range of couple of hundreds

of MHz, while plasmonic nanoLEDs have the potential to achieve ultra-high frequency (> 100

GHz) utilizing the Purcell effect [4,22,26].

5.2 Purcell effect

141

Figure 5.2: Illustration of Purcell enhancement of the spontaneous emission in a cavity. Spontaneous emission rate can be modified by putting a dipole emitter inside a cavity.

Edward Purcell in 1946 discovered that the rate of spontaneous emission could be modified when an emitter is placed inside a cavity [27]. He showed that for maximum coupling, the spontaneous emission rate can be enhanced by a factor of Fp, which is called Purcell enhancement factor:

3 휆 3 푄 퐹 = ( ) ( ) (5.1) 푝 4휋2 푛 푉

Where Fp, is the Purcell enhancement factor, 휆 is the free-space wavelength, n is the refractive index, Q is the quality factor, and V is the mode volume. A schematic of the Purcell enhancement is illustrated in figure 5.2. A dipole emitter in free space has a continuum of optical states available to couple into. When the emitter is placed in a resonator cavity, which supports some modes due to the boundaries, the rate of spontaneous emission can be modified by coupling into the cavity modes. The rate enhancement maximizes, when the emission wavelength matches with the wavelength of a cavity mode and when the emitter is positioned at the field maximum

[28]. Considering equation 5.1, a high Q parameter and small mode volume is preferred to have a

142 high Purcell factor. However, for a high Q parameter results in increase in photon lifetime, leading to slow photon decay time and low modulation speed. Therefore, a balance between Q parameter and 3-dB modulation speed is required. It is shown that the 3-dB bandwidth modulation as a function of Q parameter, and Purcell enhancement factor (Fp) can be derived as

[22]:

1 1 푓3푑퐵,푚푎푥 ≈ (5.2) 2휋 2 2 √휏푝+휏푠푝

푄 휏푠푝0 Where 휏푝 = ⁄ is the photon lifetime, 휔0 is the optical cavity frequency, 휏푠푝 = ⁄ is the 휔0 퐹푝

effective lifetime, and 휏푠푝0 is the bulk spontaneous lifetime. The 푓3푑퐵,푚푎푥 will be optimized when the condition 푑(푓3푑퐵,푚푎푥)/푑푄 = 0 is satisfied. This optimal condition occurs when the photon lifetime and effective lifetime are equal. This results in optimum value for Q (Qopt) and 3- dB modulation bandwidth (f3db,opt) :

2 휋 휏푠푝0휔0푉푛 푄 = √ 표푝푡 6 (5.3)

1 3휔0 푓3푑퐵,표푝푡 ≈ √ 2 2휋 휋 휏푠푝0푉푛

Where Vn is effective modal volume and is equal to [22]:

3 푉푛 = 푉⁄(휆⁄2푛) (5.4)

Figure 5.3 shows the counter plot of optimal 3-dB bandwidth modulation as a function of effective modal volume and Q parameter, with a dotted line showing Qopt. The dotted line separates devices dominated by strong (above dotted line) and weak (below dotted line) coupling regimes [22,29,30]. Strong coupling regime is shaded, as it demonstrates a more complicated

143

Figure 5.3: (a) Contour plot of optimal 3-dB bandwidth for various nanoLED cavities with different modal volumes and quality factors, (b) Optimum 3-dB bandwidth and Q versus cavity volume.(data from [22]) dynamics and cannot be simply described by bandwidth. The strong coupling regime (Q > Qopt) will exhibit Rabi oscillations that decay at the photon lifetime. This regime has gained significant interest for quantum information processing. Due to lower Q parameter, the weak coupling regime (Q < Qopt) allows for faster emission of photons and is attractive for high-speed optical communications. In the strong coupling regime only quantum electrodynamics (QED) can give an accurate description of atom-field interactions. On the other hand, in the weak-coupling regime the QED and classical theory give the same results for modification of the spontaneous emission decay rates [30].

5.3 Simulation method

In this section, the method used to simulate plasmonic III-nitride based nanostructures is presented. A commercial-grade simulator based on the FDTD method (Lumerical FDTD

Solutions) was used to perform the calculations. The FDTD method is a fully vectorial approach that naturally gives both the time domain and frequency domain information [31]. The Finite difference time domain (FDTD) method has become the state-of-the-art method for solving

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Maxwell’s equations in complex geometries. To simulate the plasmonic III-nitride structures the dispersion relations of n-type gallium nitride (n-GaN), p-type GaN (p-GaN), and indium gallium nitride (InGaN) is required. The n-GaN, p-GaN, and InGaN dispersion relations were determined from [32], while that for the silver cladding layer was determined from [33]. To represent the

QW emission, we used the common approach of incorporating a radiating electric dipole source at the location of the active region [34]. The emission from InGaN QWs is well modeled by a single dipole source which is swept along the active region. In the weak coupling regime (dipole approximation), where the Q parameter is low, the Purcell factor is defined as the ratio of the classical radiation power for a dipole in a cavity to a dipole emission power in a bulk dielectric material [30, 35]:

휏푏푢푙푘 푃푐푎푣푖푡푦 퐹 = 푠푝 = 푐푙푎푠푠푖푐푎푙 (5.5) 푝 푐푎푣푖푡푦 푃푏푢푙푘 휏푠푝 퐶푙푎푠푠푖푐푎푙

푏푢푙푘 푐푎푣푖푡푦 Where, 휏푠푝 is the spontaneous emission lifetime in bulk material and 휏푠푝 is the spontaneous emission lifetime in a cavity. The Purcell factor is directly related to local density of optical states (LDOS). LDOS describes the electromagnetic environment and can be calculated from the imaginary part of the dyadic Green’s function in the direction of the dipole moment at the dipole position [30, 36]. Thus, Purcell enhancement factor of the spontaneous emission is a local effect, to account for this, the location of the dipole emitter was swept along the active region for plasmonic nanostructures to find the Purcell factor at each point. Equation (5.5) has been used to find the Purcell enhancement factor. The light extraction efficiency (EXE) is defined as the ratio of the power exiting the structure (bigger box in figure 5.4(a)) to the total power generated by the dipole inside the active region (smaller box in figure 5.4(a)). The near field intensity is excited using arrays of dipoles inside the active region. The quality factor (Q) is defined as υ F/Δυ FWHM where υ F is the frequency at the peak of the Purcell factor and Δυ FWHM

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Figure 5.4: Axial structure: (a) schematic of an axial nanowire structure. (b) Cross-section of the axial nanowire structures as viewed from the top, showing the location of the p–i–n diode core and

SiNx and silver claddings. is the linewidth. Lumerical FDTD Solutions contains pre-defined scripts to calculate the quality factor.

5.4 Device structures and simulation results

5.4.1. Axial structure

A schematic of an axial structure with a hexagonal cross section is shown in figure 5.4.

The structure consists of a 3 nm thick c-plane QW sandwiched between 23.5 nm of n-GaN on the bottom and 23.5 nm of p-GaN on the top. The total height of the nanowire (NW) is 50 nm. The

SiNx layer on the GaN substrate is 25 nm thick and the SiNx layer surrounding the p-GaN, QW, and n-GaN is 10 nm thick. These layers are necessary to prevent shorting of the junction. The width (W) shown in figure 5.4(a) is 20 nm. In this geometry the QW emission for the active region is represented by a single dipole, with the dipole moment in the Z direction. The location of the dipole emitter was swept along the horizontal dimension (X, width) of the active region from −8 to 8 nm. The Purcell factor, EXE, and normalized imaginary part of the electric field inside the cavity in the direction of dipole moment for two dipoles placed at X = ±6 nm, are

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Figure 5.5: Axial structure: (a) calculated Purcell factor versus wavelength. (b) Purcell factor and EXE at a wavelength of 405 nm. (c) Normalized 2D map of the imaginary part of the electric field in the direction of the dipole moment for two dipoles placed at X = ±6 nm. shown in figures 5.5(a), 5.5(b), and 5.5(c) respectively. Figure 5.5(a) shows the Purcell factor versus wavelength for dipole emitters placed at various horizontal locations along the active region. Figure 5.5(b) shows the Purcell factor and EXE at wavelength of 405 nm. As shown in figure 5.5(b), the Purcell factor increases when the dipole emitter moves toward the silver cladding layer, however the EXE decreases. In this structure, there is a tradeoff between the

Purcell factor and EXE. The peak of the Purcell factor is at 405 nm. The average Purcell factor and EXE at 405 nm are 26 and 4%, respectively. A Purcell factor equal to 26 is insufficient to obtain high speed operation, because the axial design contains polar c-plane QWs, which have an inherently higher carrier lifetime than nonpolar QWs. Nonpolar QWs are practically attainable using a core–shell design and have shorter natural carrier lifetimes. The axial structure therefore requires a 5–10X higher Purcell factor than the core–shell structure to reduce the carrier lifetime to the same value [37].

5.4.2 Core–shell structure with hexagonal cross-section

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Figure 5.6: Core–shell nanowire structure with a hexagonal cross-section: (a) schematic of a core– shell nanowire structure with hexagonal cross-section. (b) Cross-section of the core–shell nanowire structure as viewed from the top, showing the location of the n-GaN core, InGaN active region, p- GaN shell, and silver cladding layer. If fabricated, the horizontally oriented c-plane InGaN QW in the axial geometry would have inherent polarization-related electric fields that lower the electron–hole wave function overlap and radiative rate. As an alternative, we examine a core–shell configuration containing vertically oriented nonpolar m-plane QWs, which are free from polarization-related electric fields [38] (see chapter one). Indeed, low-temperature time resolved photoluminescence measurements on nonpolar m-plane QWs has shown >10X decrease in the radiative lifetime compared to polar c-plane QWs [37]. Thus, nonpolar QWs should provide advantages for increasing the device efficiency and achieving higher modulation bandwidths. A schematic of a core–shell NW structure with a hexagonal cross-section is shown in figure 5.6. The n-GaN core has a width of 14 nm and a height of 75 nm and is surrounded by a 3 nm thick InGaN QW with a

50 nm height. The QW is surrounded by a 10 nm thick p-GaN layer and a 40 nm thick silver cladding layer. A 25 nm thick layer of SiNx was included to prevent electrically shorting the pn junction. The QW emission for the active region is represented by a single dipole, with the dipole moment in the z direction. The location of the dipole emitter was swept along the vertical dimension (Y, height) of the active region from 5 to 45 nm. The emission from nonpolar InGaN

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Figure 5.7: Core–shell structure: (a) calculated Purcell factor versus wavelength (b) Purcell factor and EXE at a wavelength of 445 nm. (c) 2D map of the normalized imaginary part of the electric field in the direction of the dipole moment for a dipole placed at Y=25 nm.

QWs is well modeled by a single dipole source directed along the a-direction of the wurtzite crystal (z-direction in figure 5.6(a)) [39,40]. The Purcell factor, EXE, and normalized imaginary part of the electric field inside the cavity in the direction of the dipole moment for a dipole placed at Y=25 nm are shown in figure 5.7. As shown in figure 5.7(a), the peak of the Purcell factor changes from 90 at 590 nm for a dipole placed at Y = 5 nm to 58 at 445 nm for a dipole placed at Y = 45 nm. Figure 5.7(b) shows how the Purcell factor and EXE change at 445 nm at different dipole positions inside the structure. The average Purcell factor and EXE at 445 nm are

45 and 5%, respectively. As shown in figure 5.7(b), the Purcell factor is non-uniform along the active region. The non-uniform Purcell factor is due to the fact that dipoles at the bottom of the cavity interact more with the 25 nm thick SiNx and dipole near the top of the cavity interact more with the metal cladding. Thus, if the active region is moved away from the SiNx, a nearly uniform Purcell factor along the cavity can be achieved. To explore this, we investigated core– shell structures with a tapered base.

5.4.2.1 Near-field intensity

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Figure 5.8: Near-field intensity for core–shell structure without a tapered base: (a) schematic of the core–shell nanowire structure with hexagonal cross-section (b) cross-section of the core–shell nanowire structures as viewed from the top. (c) Cross-sectional near-field intensity for a monitor placed at Y = 42 nm. (d) Cross-sectional near-field intensity at XY plane. (e) Cross-sectional near-field intensity at YZ plane.

To understand the distribution of energy within the cavity and the metal cladding, we investigated the near-field intensity of the structures in figures 5.6. Figure 5.8 shows the near- field electric field intensity for the three orthogonal cross-sections of the core–shell structure with a hexagonal cross-section at 445 nm. Figure 5.8(a) shows the side view cross-section of the core–shell NW structure and the coordinate system. Figure 5.8(b) shows the top-view cross- section of the core–shell NW structure, the coordinate system, and the locations of the cross- sectional planes that were examined. Figure 5.8(c) shows the cross-sectional near-field intensity viewed from the top of the NW within the XZ plane at Y = 42 nm, revealing that the electric field is strongly confined near the interface between the p-GaN and the silver cladding. Figure

5.8(d) shows the cross-sectional near-field intensity for the XY plane at a position illustrated in

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Figure 5.9: Core–shell structure with a tapered base: (a) schematic of a core–shell nanowire structure with hexagonal cross-section and a tapered base. (b) Cross-section of the core–shell nanowire structures as viewed from the top, showing the location of the n-GaN core, InGaN active region, p-GaN shell, and silver cladding. figure 5.8(b). Figure 5.8(e) shows the cross-sectional near-field intensity for the YZ plane at a position illustrated in figure 5.8(b).

5.4.3 Core–shell structure with hexagonal cross-section and a tapered base

This structure is similar to that in section 5.4.2 but contains a tapered base with a 50 nm height as shown in figure 5.9. The tapered NW structure can be experimentally realized under particular growth conditions in which the semipolar {101̅1} family of planes is thermodynamically stable during only the initial stages of growth [41, 42]. The taper has a 40 nm width at the base and a 14 nm width on top. The n-GaN core has a 14 nm width and a 50 nm height and is surrounded by a 3 nm thick InGaN nonpolar QW. The QW emitter is represented by a dipole with the dipole moment along the z direction. The location of the dipole emitter was swept along the vertical dimension (Y, height) of the active region from 5 to 45 nm. The QW is surrounded by a 10 nm thick p-GaN and a 40 nm thick silver cladding layer. The taper added to the bottom of the structure might also make the structure easier to fabricate, as the lithography

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Figure 5.10: Core–shell structure with a tapered base: (a) calculated Purcell factor versus wavelength. (b) Purcell factor and EXE at a wavelength of 448 nm. (c) 2D map of the normalized imaginary part of the electric field in the direction of the dipole moment for a dipole placed at Y=25 nm . resolution requirements would be relaxed. The Purcell factor, EXE, and normalized imaginary part of the electric field inside the cavity in the direction of the dipole moment for a dipole placed at Y=25 nm are shown in figure 5.10. Figure 5.10(a) shows that the Purcell factor versus wavelength has the same behavior along the vertical dimension of the cavity height. The peak of the Purcell factor is at 448 nm. As shown in figure 5.10(b), the peak of Purcell factor is more uniform in comparison with the non-tapered base structure shown in figure 5.7(b). The average

Purcell factor and EXE for this structure at 448 nm are 57 and 3%, respectively. In comparison with the non-tapered base structure, the tapered base structure shows a higher Purcell factor and lower EXE, which shows a trade-off between Purcell factor and EXE.

5.4.3.1 Near-field intensity

Figure 5.11(a) shows the side-view cross-section of the core–shell NW structure with a tapered base and the coordinate system. Figure 5.11(b) shows the top-view cross-section of the core–shell NW structure, the coordinate system, and the locations of the cross-sectional planes that were examined. Figure 5.11(c) shows the cross-sectional near-field intensity viewed from the top of the NW at 448 nm within the XZ plane at Y = 42 nm, it clearly shows that the electric

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Figure 5.11: Near-field intensity for core–shell structure with a tapered base: (a) schematic of the core–shell nanowire structure with hexagonal cross-section and a tapered base. (b) Cross-section of the core–shell nanowire structures as viewed from the top. (c) Cross-sectional near-field intensity for a monitor placed at Y = 42 nm. (d) Cross-sectional near-field intensity at XY plane. (e) Cross-sectional near-field intensity at YZ plane. field is strongly confined near the interface between the p-GaN and the silver cladding. Figure

5.11(d) shows the cross-sectional near-field intensity for the XY plane at a position illustrated in figure 5.11(b). Figure 5.11(e) shows the cross-sectional near-field intensity for the YZ plane at a position illustrated in figure 5.11(b). In both structures, most of the near-field intensity is concentrated near the Ag/semiconductor interfaces, leading to the relatively high Purcell factors and low EXEs.

5.4.4 Flip-chip core–shell structure with hexagonal cross-section and a tapered base

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Figure 5.12: (a) Schematic of a core-shell nanowire structure with a hexagonal cross-section and a tapered base. The taper has a 40 nm width at the base and a 14 nm width at the top. (b) Cross-section of the core-shell nanowire structure as viewed from the top, showing the location of the n-GaN core, InGaN active region, p-GaN shell, and silver cladding.

Extraction of light from the investigated structures (5.4.1 to 5.4.3) primarily occurred through a ~300 µm thick substrate, which severely reduced the light power density at the extraction surface, thus limiting the amount of useful light that could be coupled to a fiber or directed into free space. Considering these limitations, it is desirable to develop a flip-chip configuration in which the substrate has been removed and the light extraction surface is in close proximity (~100 nm) to the nanowire emitter. Figure 5.12 shows a schematic of the proposed

PNLED structure, which consists of a flip-chip Ag-clad GaN/InGaN core-shell nanowire. The n-

GaN core is 50 nm tall and the diameter is 14 nm. The active region is circumscribed around the core and consists of a single 3-nm-thick nonpolar InGaN quantum well. The quantum well (QW) is surrounded by a 10-nm-thick p-GaN layer and an optically thick (~300 nm) Ag cladding layer

(TAg). The thicknesses of the n-GaN (TGaN), SiNx, Au (TAu), and Al2O3 (sapphire) (Ts) submount are 100 nm, 25 nm, 1 µm and 300 µm, respectively. The width of the Ag cladding (L) and Al2O3 submount (D) are 2 µm and 4 µm, respectively. At the bottom of the NW core, a tapered base is included to enable a higher and more uniform Purcell factor. The light extraction efficiency

(EXE) is defined as the ratio of the power exiting the structure (top monitor figure 5.12(a)) to the

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Figure 5.13: Concept of a flip-chip PNLED fabrication process using photoelectrochemical etching to remove the substrate and expose the backside of the NW emitter. Such a structure would allow for fiber coupling or more efficient light extraction into free space than a structure with the substrate still incorporated. total power generated by the dipole inside the active region (smaller box in figure 5.12(a)). The tapered base has a height of 25 nm and a taper angle of 62o.

Figure 5.13 illustrates the fabrication steps that could potentially be used to produce an electrically injected flip-chip PNLED structure. The proposed process leverages band-gap selective photoelectrochemical (PEC) etching [43, 44] of an InGaN sacrificial layer embedded beneath the NW to enable substrate removal. The first step in the fabrication process would utilize electron beam lithography and selective-area epitaxy to pattern and grow the core-shell nanowire structure shown in figure 5.13(a). The NW would then be coated in thick Ag for the plasmonic effect and thick Au for subsequent flip-chip bonding (figure. 5.13(b)). Wet and dry chemical etching would then be used to define a device mesa and expose the InGaN sacrificial etch layer (figure. 5.13(c)). The structure would then be flip-chip bonded using Au-Au bonding

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Figure 5.14: (a) Calculated Purcell factor vs. wavelength for dipole emitters placed at various vertical locations along the active region. (b) Calculated Purcell factor and EXE into air at 445 nm for dipole emitters placed at various vertical locations along the active region from 5 nm to 45 nm. (c) 2D map of the normalized imaginary part of the electric field in the direction of the dipole moment for a dipole placed at Y=25 nm. techniques to an Al2O3 host submount coated in Au. Band-gap selective PEC etching would then be used to separate the device mesa from the substrate (figure. 5.13(d-e)), effectively removing the substrate and leaving the device bonded to the host submount [45]. Finally, a ring n-contact would be applied to the backside using conventional photolithography (figure. 5.13(e)). The proposed process describes a practical approach to achieve electrically injected flip-chip

PNLEDs and could be transferred to other III-V materials.

Figure 5.14(a) shows the Purcell factor vs. wavelength for a dipole placed at various vertical locations along the active region of the tapered nanowire structure shown in figure 5.12.

As shown in 5.14(a), the Purcell factor vs. wavelength is relatively uniform as a function of vertical position within the cavity and the peak Purcell factor occurs at 445 nm. Figure 5.14(b) shows the Purcell factor and EXE into air at 445 nm for a dipole placed at various vertical locations along the active region. When the dipole is swept vertically from 5 to 45 nm, the

Purcell factor is relatively constant, increasing from a minimum of 53 to a maximum of 60.

Conversely, the EXE is strongly non-uniform along the vertical direction of the cavity, decreasing from 0.7% to 0.16% as the dipole moves away from the open side of the nanowire.

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The average EXE into air is 0.44%, which is nearly the same as the EXE for a dipole placed at the point Y = 25 nm. The average of the Purcell factor at 445 nm for all of the dipole positions is

57. Figure 5.14(c) shows the normalized imaginary part of the electric field inside the cavity in the direction of the dipole moment for a dipole placed at Y=25 nm. As expected, the field intensity is also peaked near 445 nm, is relatively uniform in the vertical direction within the nanowire, and drops to nearly zero within the GaN taper. This shows that the field is well contained within the wire despite the wire being open and unconfined on the substrate side. The high average Purcell factor indicates that a significant portion of the energy in the mode resides within the metal cladding or near the semiconductor-metal interfaces. Consequently, absorption losses will be high and the EXE into air is expected to be very low.

5.4.4.1 Effect of n-GaN thickness

In addition to the geometry and dimensions of the nanowire itself, the thickness of the n-

GaN layer (TGaN) has an important effect on the EXE into air. Figure 5.15 shows the EXE into air at 445 nm vs. n-GaN thickness for a dipole placed at Y = 25 nm. As shown in figure 5.15, the

EXE into air vs. n-GaN thickness has a sinusoidal behavior, showing a transmittance dependence similar to that of a Fabry-Pérot etalon. This behavior is due to the interference effect between forward and backward travelling photons. Similar Fabry-Pérot effects on EXE have been reported in GaN-based LEDs for solid-state lighting [46]. For a substrate thickness of 100 nm, the maximum EXE into air is obtained. figure 5.13 illustrates the importance of accurately controlling the n-GaN thickness to achieve the highest possible EXE. The proposed flip-chip device affords this control, as it utilizes a highly selective PEC etching process.

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Figure 5.15: EXE into air at 445 nm for various thicknesses of n-GaN.

5.4.4.2 Embedded dielectric structure

The relatively high Purcell factor obtained in the structure shown in figure 5.12 also results in low EXE. It is desirable to explore additional configuration that might result in higher

EXE. To achieve a high Purcell factor, a very thick Ag layer is unnecessary and a thinner layer on the order of 20-30 nm can be used. This enables an additional device geometry that utilizes an embedded dielectric layer to significantly increase the EXE without strongly decreasing the

Purcell factor. The structure is shown in figure 5.16 and is the same as that shown in figure 5.12 except a SiNx layer is embedded within the silver cladding. The thickness of this layer (S) was swept from 0 to 40 nm for two different thicknesses of the first silver layer (TAg1 = 20 nm and

TAg1 = 30 nm). The Purcell factor and EXE into air are shown in figure 5.17. Figure 5.17 illustrates the clear trade-off between EXE and Purcell factor, whereby increasing the thickness of the SiNx layer the EXE increases and the Purcell factor decreases. Figure 5.17(a) shows the case for TAg1 = 30 nm, where increasing the thickness of the SiNx layer from 0 to 40 nm results in the EXE increasing from 0.43% to 0.86% and the Purcell factor decreasing from 59 to 51.

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Figure 5.16: (a) Schematic of a core-shell nanowire structure with a hexagonal cross-section and a tapered base. A dielectric layer was added in order to increase the EXE. (b) Cross- section of the core-shell nanowire structure as viewed from the top, showing the location of

the n-GaN core, InGaN active region, p-GaN shell, SiNx layer, and silver layer.

Figure 5.17: Calculated Purcell factor and EXE into air at 445 nm for a dipole emitter

placed at Y = 25 nm. The thickness of the SiNx layer (s) was swept from 0 to 40 nm and

TAg1 is (a) 30 nm and (b) 20 nm.

Although the Purcell factor is reduced by 13%, the EXE increases by a factor of two. Figure

5.17(b) shows the case for TAg1 = 20 nm, where increasing the thickness of the SiNx layer from 0 to 40 nm results in the EXE increasing from 0.43% to 1.29% and the Purcell factor decreasing from 59 to 38. Although the EXE increases by a factor of three in this case, the Purcell factor is also significantly lowered. Schemes such as embedded dielectric layers could potentially be used to increase EXE without significantly lowering the Purcell factor.

5.4.4.3 Near-field intensity

159

Figure 5.18: (a) Top-view cross-section of the core-shell nanowire structure shown in Fig. 1, the associated coordinate system, and the locations of cross-sectional planes. (b) Cross- sectional near-field intensity within the XZ plane (c) Cross-sectional near-field intensity within the XY plane. (d) Cross-sectional near-field intensity within the YZ plane.

To understand the distribution of energy within the cavity and the metal cladding, we investigated the near-field intensity of the structure in figure 5.12 over the three possible cross sections of the nanowire. Figure 5.18(a) shows the top-view cross-section of the core-shell nanowire structure, the coordinate system, and the locations of the cross-sectional planes that were examined. Figure 5.18(b) shows the cross-sectional near-field intensity viewed from the top of the nanowire within the XZ plane at Y = 42 nm and wavelength of 445 nm, revealing that the electric field is strongly confined near the interface between the p-GaN and the silver cladding.

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Figure 5.18(c) shows the cross-sectional near-field intensity for the XY plane at a position illustrated in figure 5.18(a), revealing that the electric field within the XY plane is also primary confined near the interface between the p-GaN and the silver cladding and that the field intensity is several times higher than that of the XZ cross-section. Figure 5.18(d) shows the cross-sectional near-field intensity for the YZ plane at a position illustrated in figure 5.18(a). For the YZ cross- section, the peak of the electric field resides within the core of the semiconductor structure.

However, the peak field intensity is lower for this cross section than for the XZ or XY cross- sections.

5.5. Modulation characteristics

A primary motivation for research on PNLEDs is their potential for high-speed modulation. High speeds are enabled when a large Purcell factor results in a significant reduction of the natural spontaneous emission lifetime. In this section, we predict the maximum carrier- lifetime-limited modulation bandwidth for all PNLEDs examined in this work. We first consider the natural spontaneous emission lifetimes for GaN/InGaN QWs. For polar c-plane InGaN QWs

2 at a current density of 100 A/cm , we assume a radiative lifetime (τr) of 3 ns and a non-radiative lifetime (τnr) of 5 ns, as reported by David and Grundmann [47].

For nonpolar m-plane InGaN QWs (core-shell structures), the lifetimes have been shown to be approximately 10X lower than those on c-plane [37], so for the nonpolar InGaN QWs of core-shell nanowires we assume the natural lifetimes without cavity effects are τr ≈ 0.3 ns and τnr

≈ 0.5 ns. The effective carrier lifetime (τeff) with cavity effects (i.e, Purcell enhancement) is given by [48, 49, 50]:

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1 퐹 1 = + (5.6) 휏푒푓푓 휏푟 휏푛푟

Where F is the Purcell factor and τeff is the Purcell enhanced carrier lifetime. The 3dB bandwidth is given by [22]:

1 1 푓3푑퐵 ≈ (5.7) 2휋 2 2 √휏푝+휏푒푓푓

Where τp is the photon lifetime (see section 5.2). For cavities with quality factors of less than several hundred, τp is a small portion of the overall lifetime in Eq. (5.6) and f3dB depends primarily on τeff . The 3dB bandwidth in this case becomes approximately:

1 1 푓3푑퐵 ≈ (5.8) 2휋 휏푒푓푓

Using the calculated Purcell factors, equations (5.6) and (5.7), and the natural carrier lifetimes from [47], we estimate the maximum achievable modulation bandwidth for the various structures discussed above. Table 1, shows the wavelength () at the peak of the Purcell factor, the average Purcell factor (F), the radiative lifetime (τr), the non-radiative lifetime (τnr), the calculated cavity quality factor (Q), and the resulting 3dB modulation bandwidth for all of the structures discussed in this paper. Based on table 5.1, the core-shell structure with the hexagonal cross-section and a tapered base would have the highest modulation bandwidth (30.2 GHz). It is important to note that the axial structure would have significantly lower maximum modulation bandwidth because the natural carrier lifetime is much longer for polar c-plane QWs and because the Purcell factor for the axial structure is lower than those of the core-shell structures.

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Table 5.1. Parameters used to find the 3dB modulation bandwidth of the structures discussed in this chapter.

τ τ τ f Structure 휆(nm) F r nr eff Q 3dB (ns) (ns) (ps) (GHz)

Axial structure with hexagonal cross- 405 26 3.0 5.0 115.0 10 1.4 section

Core-shell structure with hexagonal 445 45 0.3 0.5 6.6 12 23.9 cross-section

Core-shell structure

with hexagonal 448 57 0.3 0.5 5.2 12 30.2 cross-section and a tapered base

Flip-ship core-shell

structure with hexagonal cross- section and a 445 57 0.3 0.5 5.2 14 30.2 tapered base

5.6 Conclusions

The main focus of this chapter was on design and simulation of GaN/InGaN plasmonic nanostructures. The theory of Purcell effect was presented and methods to calculate the Purcell enhancement factor are given. We utilized the FDTD method to investigate the Purcell factor,

EXE, and cavity quality parameter (Q) of axial and core-shell III-nitride based nanostructures. It was shown that for axial structures due to polar c-plane quantum wells, the inherent polarization-

163 related electric fields lower the electron–hole wave function overlap and radiative rate. However for core-shell structures, nonpolar active region can be achieved by the growth of quantum wells on GaN nanowire sidewalls. We have shown that a more uniform Purcell enhancement factor can be achieved using a tapered-based structure. A flip-chip structure was proposed to remove the sapphire substrate and improve the EXE into air. To increase the EXE into air even further an embedded structure was proposed. The potential modulation response of the devices was presented and maximum 3-dB modulation of 30.2 GHz was predicted.

164

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[43] Tamboli, A.C., Schmidt, M.C., Hirai, A., DenBaars, S.P. and Hu, E.L., 2009. Photoelectrochemical undercut etching of m-plane GaN for microdisk applications. Journal of The Electrochemical Society, 156(10), pp.H767-H771.

[44] Kim, D.U., Chang, H., Cha, H., Jeon, H. and Jeon, S.R., 2013. Selective lateral electrochemical etching of a GaN-based superlattice layer for thin film device application. Applied Physics Letters, 102(15), p.152112.

[45] Holder, C., Speck, J.S., DenBaars, S.P., Nakamura, S. and Feezell, D., 2012. Demonstration of nonpolar GaN-based vertical-cavity surface-emitting lasers. Applied Physics Express, 5(9), p.092104.

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[47] David, A. and Grundmann, M.J., 2010. Droop in InGaN light-emitting diodes: A differential carrier lifetime analysis. Applied Physics Letters, 96(10), p.103504.

[48] Nami, M. and Feezell, D.F., 2014. Optical properties of plasmonic light-emitting diodes based on flip-chip III-nitride core-shell nanowires. Optics express, 22(24), pp.29445- 29455.

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[50] Nami, M., Wright, J. and Feezell, D.F., 2014, June. Investigation of purcell factor and light extraction efficiency in Ag-Coated GaN/InGaN core-shell nanowires. In Lasers and Electro-Optics (CLEO), 2014 Conference on (pp. 1-2). IEEE.

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Chapter 6

6.1Conclusions

Simulation and design, growth, fabrication and characterization of GaN/InGaN core- shell nanowire-based LEDs were presented in this dissertation. Chapter one addressed the motivation of the nanostructure-based optoelectronic devices and a brief history of electrically injected GaN/InGaN core-shell nanowire-based LEDs was given.

In chapter two, the use of pulsed mode selective-area growth to grow GaN nanowires was discussed. The effect of mask geometry, growth temperature, TMGa injection time, and TMGa interruption time on growth, morphology and dimensions of GaN nanowires were studied. We have shown that the mask geometry has an important role in the dimension and morphology of

GaN nanowires. These findings pave the way for a deeper and a clearer understanding of the growth of GaN nanowires using pulsed mode SAG. The effects of the aperture diameter and pitch size on the growth of the GaN nanowires were investigated, and we showed that the pitch size plays an important role on the growth and morphology of GaN nanowires. The theory of capture length was proposed, and changes in the GaN nanowire dimensions and morphology were explained based on this theory. The effect of TMGa injection time on the nanowire diameter has been studied. It was shown that as the TMGa injection time reduces nanowires diameters reduces. Using 12 sec of TMGa injection time, GaN nanowires with a diameter of 95 nm was achieved. It was shown that the nanowire diameter is more sensitive to TMGa injection time in the regime where there is a strong competition for the growth nutrients between adjacent nanowires. This regime occurs where the capture length is more than half of the pitch sizes. The effect of the TMGa interruption time on GaN nanowires diameter was investigated. We have

170 shown that by increasing the TMGa interruption time from 27 sec to 33 sec, GaN nanowires with a dimeter of 74 nm can be achieved.

In chapter three, growth of an InGaN quantum well around the GaN nanowire cores was studied. One quantum well and one barrier were grown around the GaN nanowires cores, and µ-

PL was obtained and analyzed for a variety of nanowire dimensions, array pitch spacings, and aperture diameters. An 80 nm redshift in peak PL wavelength was observed as the nanowire pitch spacing was increased from 300 nm to 4 µm, while increasing the aperture diameter from

60 nm to 210 nm resulted in a ~35 nm blueshift. The thickness of one QW/barrier period as a function of pitch and aperture diameter was inferred using SEM, and increased as the pitch increased. The effects of pitch and aperture diameter on indium composition were decoupled from the measured PL by simulating the expected peak wavelength for a constant indium composition. Significant increases in indium composition were predicted for larger pitches and smaller aperture diameters. Temperature dependent measurements of the IPLI and TRPL were performed on GaN/InGaN core-shell nanowire arrays. For 300 nm and 2 µm pitch size arrays, an anomalous behavior of the IPLI as a function of temperature was observed. This phenomenon is attributed to temperature induced carrier redistribution inside shallow and deep localized states inside the active region. The source of these types of localization effects can be related to formation of QDs on the sidewalls of the GaN nanowires during the growth of thin (less than 2 nm) quantum wells, or indium segregation for quantum wells with high amount of indium incorporation. This detailed study provides significant insight into the effects of mask configuration and growth conditions on the GaN/InGaN core-shell nanowire properties and is applicable to the fabrication of monolithically integrated, multi-color GaN-based light-emitters

171 on a single chip, which can potentially be used in multichannel VLC systems with high data rates.

In chapter four, electrically injected GaN/InGaN core-shell nanowire-based LEDs were presented. Interferometric lithography was used to make nanowire template. Pulsed mode SAG was employed to grow the GaN nanowires. It was shown that an AlGaN underlayer is required to prevent desorption of Si and O from dielectric mask. Dissociation of Si and O from SiNx dielectric mask gives rise to compensation of the p-GaN layer, reduction the efficiency of this layer, and a significant reverse-leakage current density. We have shown that growth of AlGaN underlayer leads to the growth of polycrystalline AlN layer on the dielectric mask and passivation of this layer. To enable conformal growth of AlGaN underlayer, a pulsed mode growth technique was used. GaN/InGaN core-shell nanowire-based LEDs were fabricated and electro-optical and RF characteristics was presented. Peak IQE of 62% and series resistance of

25 Ω were achieved. These results are among the highest performance levels for nanowire-based

LEDs thus far. TRPL was used to study the carrier dynamics of the LEDs at room temperature and low temperature. In addition, FDTD modeling revealed that the nanowire-based LEDs have a strongly directional far-field emission pattern. Properties such as high IQE, short carrier lifetime, and emission directionality are attractive for solid-state lighting, visible-light communication, and -LED displays, respectively. RF characteristics of GaN/InGaN core-shell nanowire-based LEDs was also presented, and an equivalent circuit model was proposed for the

LED. Using scattering matrix elements and the ABCD transfer matrix method, the S11 and S21 parameters were measured and fit to the experimental data. Using the fitting parameters we decoupled the carrier recombination lifetime from the diode time constant. It was shown that the

172

3-dB modulation bandwidth of the LED is limited by carrier lifetime rather than the diode time constant.

The main focus of the fifth chapter was on design and simulation of GaN/InGaN plasmonic nanostructures. The theory of Purcell effect was presented and methods to calculate the Purcell enhancement factor were given. We utilized the FDTD method to investigate the

Purcell factor, EXE, and cavity quality parameter (Q) of axial and core-shell III-nitride based nanostructures. It was shown that for axial structures due to polar c-plane quantum wells, the inherent polarization-related electric fields lower the electron–hole wave function overlap and radiative rate. However for core-shell structures, nonpolar active region can be achieved by the growth of quantum wells on GaN nanowire sidewalls. We showed that a more uniform Purcell enhancement factor can be achieved using a tapered-based structure. A flip-chip structure was proposed to remove the sapphire substrate and improve the EXE into air. To increase the EXE into air even further, an embedded structure was proposed. The modulation response of the devices was presented and a maximum 3-dB modulation of 30.2 GHz was predicted.

6.2 Future work

The initial results of RF characterization are promising, however the 3-dB modulation bandwidth needs to be improved to reach the state-of-the-art devices with 3-dB modulation of

0.8 to 1 GHz. The equivalent circuit model and the fitting approach provided us with the possibility to decouple the carrier lifetime and diode time constant. Since the 3-dB modulation bandwidth of the LED is limited by carrier lifetime rather than diode time constant a novel active region design is required to reduce the carrier lifetime. The plasmonic nanostructures originally proposed here offer a promising approach to reduce the carrier lifetime by taking the advantage

173

Figure 6.1: (a) SEM image of GaN/InGaN core-shell nanowires (b) zoom in image. of Purcell enhancement factor. However, implementation of such structures is extremely difficult. Throughout the course of this dissertation, we have developed electrically injected nanowire LEDs from the ground up, study various designs capable of high-speed operation, and elucidated the detailed growth mechanisms for bottom up nanostructures. Furthermore, initial experiments have shown a reduction in PL lifetime for GaN/InGaN core-shell nanowires coated in silver nanoparticles. Reduction of the PL lifetime from 149 to 87 ps has been observed for the nanowires using silver nanoparticles with diameters of 30 nm. Figure 6.1 shows GaN/InGaN core-shell nanowires with four pairs of quantum wells/barriers and a 12 nm of cap layer. Figure

6.2 shows reduction in PL lifetime for the GaN/InGaN core-shell nanowire after silver deposition.

Reduction of the carrier lifetime can be also be achieved by leveraging the Purcell effect and reduction of modal volume. Growth the GaN nanowire cores with diameters less than 100 nm might result in reduction of the modal volume and increase of the Purcell enhancement factor. Higher Purcell enhancement factor gives rise to higher carrier recombination rate and higher 3-dB modulation speed. One way to reduce the size of the GaN nanowire cores is to use

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Figure 6.2: Room temperature PL transients a GaN/InGaN core-shell nanowire sample with and without silver nanoparticles. e-beam lithography to make the nanowire template with aperture size of 40 to 50 nm. We previously demonstrated nanowires with 30 nm diameter.

Another promising topic for future work is phosphor-free white light sources by leveraging the results from chapter three. Having obtained GaN nanowire cores with three different pitch sizes of 500 nm, 1 and 2 µm, multi-color LEDs can be achieved in one device mesa using a single epitaxial growth. Figure 6.3 shows an example of a phosphor-free white light source using three different regions with GaN nanowires with 500 nm (region 1), 1 µm (region

2), and 2 µm pitch size (region 3). The surface area of each region can be designed to tune the dominant wavelength.

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Figure 6.3: A proposed approach to make Phosphor-free white light source using three different regions emitting violet, blue and green.

176

Appendix A

Processing steps required for fabrication

Solvent Acetone 3min, Isopropanol 3min, DI rinse Cleaning No ultrasonic Bench 3min p-Electrode ITO Deposition 1:3 HCl:H2O 1 min, DI rinse 30s x 3 rinse BHF Dip Acid Bench and dump E-Beam 2500Å, 1-2Å/sec, O2 flow rate: 4.2 sccm, ITO Deposition Tool used is EVAP03 CHA ~2×10-5Torr ITO patterning Positive Tone, spin speed = 3krpm, soft bake at 112C for 90 Mesa Contact PR: SPR220 min, 5.6sec exposure at Photolithography Aligner cp=220, followed by a 60 sec development in AZ300MIF ITO etch rate = 30 Manual mode etching: Cl2/Ar = 20/5 sccm, Mesa Etch PE01 nm/min; Typically etch ICP/RF = 350/50W (~0.3 um/min) for 8 min Solvent Inspect: check for Strip PR 1165 80 C, followed by DI rinse. Bench complete PR removal. Check Etch Depth SEM

ITO Anneal RTA 600°C, 10min, N2 NA Mesa Formation Positive Tone, spin speed = 3krpm, soft bake at 112C for 90 Mesa Contact PR: SPR220 min, 5.6sec exposure at Photolithography Aligner cp=220, followed by a 60 sec development in AZ300MIF GaN etch rate = 200 Manual mode etching: Cl2/Ar = 20/5 sccm, Mesa Etch PE01 nm/min; Typically etch ICP/RF = 350/50W (~0.3 um/min) for 45 sec Solvent 1165 80 C (105 C on hotplate), followed by Inspect: check for Strip PR Bench DI rinse. complete PR removal Check Etch Depth SEM n-contact deposition

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Solvent Acetone 3min, Isopropanol 3min, DI rinse Clean NO ultrasonic Bench 3min Negative Tone, exp = 4 sec gives better sidewall profile for liftoff. Spin @ 3K rpm, followed by soft bake at 112C for 90 sec. CP n-Electrode Contact AZnLOF2020 mode on aligner, set Photolithography Aligner power to 201. Post exposure bake at 112 for 60 sec, followed by developing in AZ300:MIF for 60-70 sec. Clean PR and developer 1:3 HCl:H2O 1 min, DI rinse 30s x 3 rinse Acid cleaning Acid Bench left over before contact and dump deposition Ti/Al/Ni/Au, 100/1000/1000/2000Å n-Electrode Metal EVAP01/0 (Depends on thickness needed for the Deposition 2 leveling of p-pads and n-contact), 2Å/sec Inspect: check for Solvent 1165 80 C (105 C on hotplate), followed by complete PR removal. Liftoff Bench DI rinse. Place sample in hot 1165 and pipette n-Electrode Check RTA Calibration RTA 450°C, 3min, N2, 5sccm Anneal First Pads Solvent Acetone 3min, Isopropanol 3min, DI rinse Clean NO ultrasonic Bench 3min

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Negative Tone, exp = 4 sec gives better sidewall profile for liftoff. Spin @ 3K rpm, followed by soft bake at 112C for 90 sec. CP n-Electrode Contact AZnLOF2020 mode on aligner, set Photolithography Aligner power to 201. Post exposure bake at 112 for 60 sec, followed by developing in AZ300:MIF for 60-70 sec. Cr/Au, 100/5000Å (Depends on thickness n-Electrode Metal EVAP01/0 needed for the leveling of p-pads and n- Deposition 2 contact), 2Å/sec Solvent 1165 80 C (105 C on hotplate), followed by Liftoff Bench DI rinse.

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Peer Reviewed Journal Publications

1. M. Nami, Isaac E. Stricklin, Kenneth M. DaVico, Ashwin K. Rishinaramangalam, S. R. J. Brueck, Igal Brener, and Daniel F. Feezell “Carrier Dynamics and Electro- Optical Characterization of High-Performance GaN/InGaN Core-Shell Nanowire Light-Emitting Diodes” submitted to Scientific Reports (August 2017)

2. M. Monavarian, A. Rashidi, A. Aragon, S. H. Oh, M. Nami, S. P. DenBaars, and D. Feezell “Explanation of low efficiency droop in semipolar (202̅1̅) InGaN/GaN LEDs through evaluation of carrier recombination coefficients” Optics Express 25(16), 19343-19353. (August 2017)

3. A. Rashidi, M. Nami, M. Monavarian, A. Aragon, K. DaVico, F. Ayoub, S. Mishkat- Ul-Masabih, A. Rishinaramangalam, and D. Feezell “Differential Carrier Lifetime and Transport Effects in Electrically Injected III-Nitride Light-Emitting Diodes” Journal of applied Physics,122, 035706 (July 2017).

4. M. Behzadirad, M. Nami, A. Rishinaramagalam, D. Feezell, T. Busani, “GaN nanowire tips for nanoscale atomic force microscopy” Nanotechnology, 28, 20LT01, (April. 2017).

5. A. Rashidi, M. Monavarian, A. Aragon, S. Okur, M. Nami, A. Rishinaramangalam, S. Mishkat-Ul-Masabih, and D. Feezell, “High-Speed Nonpolar InGaN/GaN LEDs for Visible-Light Communication” Photon. Technol. Lett. 29, 4, 381-384, (January. 2017).

6. S. Okur, M. Nami, A. K. Rishinaramangalam1, S. H. Oh, S. P. DenBaars, S. Liu, I. Brener, and D. F. Feezell, “Internal quantum efficiency and carrier dynamics in semipolar (202̅1̅) InGaN/GaN light-emitting diodes” Optics Express 25, 3, 2178- 2186, (January. 2017).

7. M. Nami, R. F. Eller, S. Okur, A. K. Rishinaramangalam, S. Liu, I. Brener, and D. F. Feezell, “Tailoring the morphology and luminescence of GaN/InGaN core-shell nanowires using bottom-up selective-area epitaxy” Nanotechnology, 28, 2, 025202 (January. 2017).

8. A. K. Rishinaramangalam, M. Nami, D. M. Shima, G. Balakrishnan, S. R. J. Brueck and D. F. Feezell, “Reduction of reverse-leakage current in selective-area-grown GaN-based core-shell nanostructure LEDs using AlGaN layers” Phys. Status Solidi A, 1–5 (December. 2016).

9. A. Rishinaramangalam, M. Nami, M. Fairchild, D. Shima, G. Balakrishnan, S. Brueck, and D. Feezell, “Semipolar InGaN/GaN Nanostructure Light-Emitting

180

Diodes on c-Plane Sapphire,” Appl. Phys. Express, 9, 032101(1-4), (February. 2016).

10. A. Rishinaramangalam, M. Nami, B. Bryant, R. Eller, D. Shima, G. Balakrishnan, S. Brueck, and D. Feezell, “Ordered arrays of bottom-up III-nitride core-shell nanostructures,” SPIE Optics and Photonics Conference, San Diego, (August. 2015).

11. A. Kazemi, B. Klein, M. Nami, J. Kim, M. Allen, J. Allen, B. Wenner, D. Brown, A. Urbas, D. Feezell, B. Mitchell, S. Krishna, “Investigation of plasmonic enhancement in a quantum dot-in-a-well structure,” 2015 SPIE Photonics West (OPTO) Conference, 937019-937019-10, San Francisco, CA, (February. 2015).

12. M. Nami and D. Feezell, “Optical Properties of Ag-Coated GaN/InGaN Axial and Core-Shell Nanowire Light-Emitting Diodes,” J. Opt, 17, 025004(1-9) (January. 2015).

13. M. Nami and D. Feezell, “Optical Properties of Plasmonic Light-Emitting Diodes Based on Flip-Chip III-Nitride Core-Shell Nanowires,” Opt. Express, 24, 29445- 29455 (November. 2014).

14. M. Nami, A. Rishinaramangalam, and D. Feezell, “Analysis of Light Extraction Efficiency for Gallium Nitride-Based Coaxial Microwall Light-Emitting Diodes,” Phys. Status Solidi C, 11,766– 770 (January. 2014).

15. A. Hurtado, M. Nami, I. D. Henning, M. J. Adams, and L. F. Lester, “Two- Wavelength Switching With a 1310-nm Quantum Dot Distributed Feedback Laser” JSTQE, 19, 4 (July/August 2013).

16. A. Hurtado, J. Mee, M. Nami, I. D. Henning, M. J. Adams, and L. F. Lester, “Tunable microwave signal generator with an optically-injected 1310nm QD-DFB laser” Optics. Express, 21, 9, 10772-10778 (May. 2013).

17. A. Hurtado, M. Nami, I. D. Henning, M. J. Adams, and L. F. Lester, “Bistability patterns and nonlinear switching with very high contrast ratio in a 1550 nm quantum dash semiconductor laser” Appl. Phys. Lett. 101, 161117 (October. 2012).

Conference Presentations

1. M. Nami, I. E. Stricklin, A. K. Rishinaramangalam, S. R. J. Brueck, I. Brener, and D. F. Feezell “Electro-Optical Characterization of Selectively-Grown GaN/InGaN Core-Shell Nanowire LEDs” accepted for oral presentation at ISSLED 2017, Banff, Canada. (October 2017)

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2. M. Monavarian, A. Rashidi, A. Aragon, M. Nami, and Daniel F. Feezell “Small- Signal Modulation Characteristics of Nonpolar and Semipolar InGaN/GaN Light- Emitting Diodes” ISSLED 2017, Banff, Canada. (October 2017)

3. M. Nami, A. K. Rishinaramangalam, D. M. Shima, G. Balakrishnan, S. R. J. Brueck, and D. F. Feezell. “Effect of AlGaN Underlayer on Reverse-Leakage Current Reduction in GaN/InGaN Core-Shell Nanostructure Light-Emitting Diodes” 18th US Workshop on Organometallic Vapor Phase Epitaxy (OMVPE-18). Santa Fe, NM (July 2017).

4. M. Nami, A. K. Rishinaramangalam, I. E. Stricklin, S. R. J. Brueck, Igal Brener, and D. F. Feezell, “Optical and Electrical Characterization of GaN/InGaN Core-Shell Nanowire Light-emitting diodes” Electronic Materials Conference (EMC), Notre Dame, IN, (June. 2017).

5. M. Nami, R. F. Eller, S. Okur, A. K. Rishinaramangalam, S. Liu, I. Brener, and D. F. Feezell, “Anomalous Behavior of Temperature-Dependent Photoluminescence in GaN/InGaN Core-Shell Nanostructures” International Workshop on Nitride Semiconductors (IWN), Orlando, (October. 2016).

6. M. Behzadirad, N. Dawson, M. Nami, A. Rishinaramangalam, D. Feezell, and T. Busani, “GaN Nanowire Tip for High-Aspect-Ratio Nanoscale AFM Metrology,” SPIE Optics and Photonics Conference, San Diego, (August. 2016).

7. A. Rishinaramangalam, M. Fairchild, M. Nami, O. Johnson, D. Shima, G. Balakrishnan, S. Brueck and D. Feezell, “Semipolar (10-11) GaN-Based Core-Shell Nanostructure LEDs on c-Plane Sapphire Using Selective-Area MOCVD,” Electronic Materials Conference (EMC), Newark, DE, (June. 2016).

8. M. Nami, R. Eller, S Okur, A. Rishinaramangalam, S. Liu, I. Brener, and D. Feezell, “Effect of Growth Temperature and Mask Geometry on Morphology and Photoluminescence of GaN/InGaN Core-Shell Nanowires,” Electronic Materials Conference (EMC), Newark, DE, (June. 2016).

9. S. Okur, M. Nami, A. Rishinaramangalam, S. H. Oh, S. Liu, I. Brener, S.P. DenBaars, and D. Feezell, “Effect of active region design on carrier dynamics in semipolar (20-2-1) InGaN/GaN light-emitting diodes,” 2016 Electronic Materials Conference (EMC), Newark, DE, (June 2016).

10. M. Nami and D. Feezell, “Investigation of Purcell Factor and Light Extraction Efficiency in Ag Coated GaN/InGaN Core-Shell and Axial Hexagonal Nanowires,” ICN+T Conference, Vail, CO, (July 2014).

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11. M. Nami, J. Wright, and D. Feezell, “Investigation of Purcell Factor and Light Extraction Efficiency in Ag-Coated GaN/InGaN Core-Shell Nanowires,” Conference on Lasers and Electro-Optics (CLEO), San Jose, CA, (May 2014).

12. A. Hurtado, J. Mee, M. Nami, I.D. Henning, M.J. Adams and L.F. Lester, “Tunable Microwave Signal Generation with an Optically Injected 1310nm Quantum Dot Distributed-Feedback Laser” Conference on Lasers and Electro-Optics (CLEO), San Jose, CA, (June 2013).

13. M. Nami, A. Rishinaramangalam, and D. Feezell, “Analysis of Light Extraction Efficiency for GaNBased Coaxial Microwall Light-Emitting Diodes,” International Conference on Nitride Semiconductors, Washington D.C. (August. 2013).

14. A. Hurtado, J. Mee, M. Nami, I. D. Henning, M. J. Adams, and Luke F. Lester, “Analysis and Applications of an Optically-Injected 1310 nm Quantum-Dot Distributed Feedback Laser” 15th International Conference on Transparent Optical Networks (ICTON) (June. 2013).

15. A. Hurtado, M. Nami, I.D. Henning , M.J. Adams and L. F. Lester, “Bistability and switching with very high contrast ratio in an optically-injected 1550nm-QDash Fabry-Perot laser” International Semiconductor Laser Conference (ISLC) (October. 2012)

16. Hurtado, J. Mee, M. Nami, L.F. Lester, I.D. Henning and M.J. Adams, “Tunable microwave, millimeter-wave and THz signal generation with a 1310nm quantum dot laser” IEEE International Topical Meeting on Microwave Photonics (MWP) (October. 2013)

17. S. Saghafi, A. Heydari, M. Nami, “Influence of laser irradiation on wheat growth” International multi-conference IEEE/LEOS CAOL 2005 (Yalta 12-17 September. 2005) on Advanced Optoelectronics and Lasers.

18. S. Saghafi, M. Nami, A. Heydari, “Improving micromachining process using a new mode converter system” International multi-conference IEEE/LEOS CAOL 2005. (Yalta 12-17 September. 2005) on Advanced Optoelectronics and Lasers.

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