effect of process parameters on the growth of n-polar...

Effect of Process Parameters on the Growth of

N-polar GaN on Sapphire by MOCVD

A Thesis

Submitted For the Degree of

Doctor of Philosophy

in the Faculty of Science

by

G R Krishna Yaddanapudi

Department of Materials Engineering

Indian Institute of Science

Bangalore 560 012 INDIA

February 2016

Declaration

I hereby declare that the work reported in this thesis is entirely original. It was

carried out by me in the Department of Materials Engineering & Centre for Nano

Science and Engineering (CeNSE), Indian Institute of Science, Bangalore. I further

declare that it has not formed the basis for the award of any degree, diploma,

membership, associateship or similar title of any university or institution.

Date: (G R Krishna Yaddanapudi)

iii

© G R Krishna Yaddanapudi

February 2016

All rights reserved

To My Loving Parents...

v

Abstract

Group III-Nitrides (GaN, InN & AlN) are considered one of the most important

class of semiconducting materials after Si and GaAs. The excellent optical and

electrical properties of these nitrides result in numerous applications in lighting,

lasers, and high-power/high-frequency devices. Due to the lack of cheap bulk III-

Nitride substrates, GaN based devices have been developed on foreign substrates

like Si, sapphire and SiC. These technologies have been predominantly developed

on the so called Ga-polarity epitaxial stacks with growth in the [0001] direction

of GaN. It is this orientation that grows most easily on sapphire by metal organic

chemical vapor deposition (MOCVD), the most common combination of substrate

and deposition method used thus far. The opposite [0001̄] or N-polar orientation,

very different in properties due to the lack of an inversion centre, offers several ad-

vantages that could be exploited for better electronic and optoelectronic devices.

However, its growth is more challenging and needs better understanding.

The aim of the work reported in this dissertation was a systematic investigation

of the relation between the various growth parameters which control polarity, sur-

face roughness and mosaicity of GaN on non-miscut sapphire (0001) wafers for

power electronics and lighting applications, with emphasis on the realization of

N-polar epitaxial layers. GaN is grown on sapphire (0001) in a two-step process,

which involves the deposition of a thin low temperature GaN nucleation layer (NL)

on surface modified sapphire followed by the growth of high temperature device

quality GaN epitaxial layer. The processing technique used is MOCVD. Various

processing methods for synthesis of GaN layers are described with particular em-

phasis on MOCVD method. The effect of ex situ cleaning followed by an in situ

cleaning on the surface morphology of sapphire (0001) wafers is discussed. The

characterization tools used in this dissertation for studying the chemical bond na-

ture of nitrided sapphire surface and microstructural evolution (morphological and

structural) of GaN layers are described in detail.

The effect of nitridation temperature (TN) on structural transformation of non-

miscut sapphire (0001) surface has been explored. The structural evolution of

nitrided layers at different stages of their process like as grown stage and thermal

annealing stage is investigated systematically. The chemical bond environment

information of the nitrided layers have been examined by x-ray photoelectron

spectroscopy (XPS). It is found that high temperature nitridation (TN ≥ 800oC)

results in an Al-N tetrahedral bond environment on sapphire surface. In con-

trast, low temperature nitridation (TN = 530oC) results in a complex Al-O-N

environment on sapphire surfaces. Microstructural evolution of low temperature

GaN NLs has been studied at every stage of processing by scanning electron mi-

croscopy (SEM) and atomic force microscopy (AFM). Surface roughness evolution

and island size distribution of NLs measured from AFM are discussed. It is found

that NLs processed on sapphire wafers nitrided at (TN ≥ 800oC) showed strong

wurtzite [0002] orientation with sub-nanometre surface roughness. In contrast,

NLs processed at (TN = 530oC) showed zinc blende phase in the as grown stage

with higher surface roughness, but acquired a greater degree of wurtzite [0002]

orientation after thermal annealing prior to high temperature GaN growth.

Polarity, surface quality and crystal quality of subsequently grown high temper-

ature GaN epitaxial layers is described in relation to the structure of the trans-

formed nitrided layers. Higher nitridation temperatures (TN ≥ 800oC) consistently

yield N-polar GaN whereas lower nitridation temperatures (TN = 530oC) yield Ga-

polar GaN. It is found that the relative O atom concentration levels in nitrided

layers control the density of inversion domains in N-polar GaN. The effect of var-

ious growth parameters (NH3 flow rate, growth temperature, NL thickness) on

surface morphology and mosaicity of both Ga & N-polar GaN layers is discussed

in detail. We report device quality N-polar GaN epitaxial layers on non-miscut

sapphire (0001) wafers by careful optimization of growth temperature. It is found

that lower growth temperatures (800oC) are favorable for obtaining smooth N-

polar GaN layers. In contrast, N-polar GaN layers grown at higher temperatures

(1000 to 1080oC) are rough with hexagonal hillocks.

Acknowledgements

It has been a great experience to be part of the successful nitride HEMT research

program at IISc. I am sincerely thankful to my adviser Professor Dipankar Baner-

jee, working under him has indeed been a privilege for me and has also been a very

rewarding experience. Sincere thanks to my second adviser Professor Srinivasan

Raghavan early on gave me the opportunity to work in his world class MOCVD

lab, and to get deeply involved with the hardware. Both of them gave me freedom

to explore my own ideas. During the course of my work I was introduced to several

scientific techniques and processes and in the process also I learned a lot about

microstructural analysis from my advisers.

Many thanks to my colleagues (MOCVD growers): Abheek, Hareesh and Naga-

boopathy, who has taught me almost everything I know about MOCVD reactor.

I would also like to thank them, with whom I have spent late growth nights, ma-

chine trouble shooting and repairs.

I would also like to thank staff at micro and nano characterization facility (MNCF)

at CeNSE, IISc and advanced facility for microscopy and micro analysis (AFMM),

comes under the department of Materials Engineering, IISc, for keeping the facil-

ities in good condition and providing training on the equipments.

I am sincerely thankful to my mother and father for their patience and love. Last

but not least, my other colleagues and friends, who has always inspired, motivated,

entertained me during the good and bad times of my PhD study. To all of you.

Thank you.

ix

Contents

Declaration of Authorship iii

Abstract vii

Acknowledgements ix

List of Figures xv

List of Tables xxi

Abbreviations xxiii

1 Introduction 1

1.1 Gallium Nitride . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 1

1.1.1 Substrates for GaN epitaxy . . . . . . . . . . . . . . . . . . 1

1.2 Structure and polarity of GaN . . . . . . . . . . . . . . . . . . . . . 3

1.3 Polarity and growth of GaN . . . . . . . . . . . . . . . . . . . . . . 3

1.3.1 GaN for lighting and power electronic applications . . . . . . 6

1.3.2 Effect of polarity . . . . . . . . . . . . . . . . . . . . . . . . 7

1.4 Thesis description . . . . . . . . . . . . . . . . . . . . . . . . . . . . 10

2 Experimental Techniques for GaN Synthesis & Characterization 13

2.1 Experimental . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 13

2.1.1 Processing techniques for GaN synthesis . . . . . . . . . . . 14

2.1.2 Metal-organic chemical vapor deposition (MOCVD) . . . . . 15

2.1.3 Time-Temperature (TT) process plot for GaN epitaxy onsapphire by MOCVD . . . . . . . . . . . . . . . . . . . . . . 17

2.2 Characterization . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25

2.2.1 X-ray photoelectron spectroscopy (XPS) . . . . . . . . . . . 26

2.2.2 High resolution x-ray diffractometer . . . . . . . . . . . . . . 27

2.2.3 Atomic force microscopy (AFM) . . . . . . . . . . . . . . . . 29

2.2.4 Differential interference contrast (DIC) light microscopy . . 30

xi

Contents xii

2.2.5 Scanning electron microscopy (SEM) . . . . . . . . . . . . . 32

2.2.6 Transmission electron microscopy (TEM) . . . . . . . . . . . 32

2.2.7 In situ reflectivity and stress monitor analysis tool . . . . . . 33

3 Sapphire Pre-treatment & Microstructural Evolution of LT GaN 37

3.1 Background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 38

3.2 Experimental . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 39

3.3 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 40

3.3.1 In situ thermal treatment of sapphire wafers . . . . . . . . . 41

3.3.2 Nitridation of in situ treated sapphire wafers . . . . . . . . . 43

3.3.2.1 As nitrided sapphire wafers . . . . . . . . . . . . . 43

3.3.2.2 Annealed nitrided wafers . . . . . . . . . . . . . . . 46

3.3.3 LT GaN nucleation layer . . . . . . . . . . . . . . . . . . . . 46

3.3.3.1 Morphological evolution . . . . . . . . . . . . . . . 46

3.3.3.2 Structural evolution . . . . . . . . . . . . . . . . . 50

3.4 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 53

3.4.1 Low temperature nitridation (TN = 530oC) . . . . . . . . . . 53

3.4.2 High temperature nitridation (TN = 1100oC) . . . . . . . . . 54

3.5 Summary & conclusions . . . . . . . . . . . . . . . . . . . . . . . . 55

4 Polarity & Microstructural Evolution of HT GaN 57

4.1 Background . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 57

4.2 Experimental . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 61

4.3 Results . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63

4.3.1 Low temperature nitridation (TN = 530oC) . . . . . . . . . 63

4.3.1.1 V/III ratio . . . . . . . . . . . . . . . . . . . . . . 64

4.3.1.2 Growth temperature . . . . . . . . . . . . . . . . . 67

4.3.1.3 LT GaN NL thickness . . . . . . . . . . . . . . . . 69

4.3.1.4 Polarity of HT GaN . . . . . . . . . . . . . . . . . 70

4.3.2 High temperature nitridation (TN = 1100oC . . . . . . . . . 73

4.3.2.1 V/III ratio . . . . . . . . . . . . . . . . . . . . . . 74

4.3.2.2 Polarity of HT GaN . . . . . . . . . . . . . . . . . 75

4.3.2.3 LT GaN annealing time . . . . . . . . . . . . . . . 76

4.3.2.4 Growth temperature . . . . . . . . . . . . . . . . . 78

4.3.2.5 Carrier gas (H2/N2) . . . . . . . . . . . . . . . . . 82

4.3.3 Summary of results . . . . . . . . . . . . . . . . . . . . . . . 82

4.3.3.1 Low temperature nitridation (TN = 530oC) . . . . 83

4.3.3.2 High temperature nitridation (TN = 1100oC) . . . 83

4.4 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 84

4.4.1 Low temperature nitridation (TN = 530oC) and HT GaN . . 84

4.4.1.1 Polarity . . . . . . . . . . . . . . . . . . . . . . . . 84

4.4.1.2 Crystalline quality of HT GaN . . . . . . . . . . . 85

4.4.1.3 Surface roughness of HT GaN . . . . . . . . . . . . 86

4.4.2 High temperature nitridation (TN = 1100oC) and HT GaN . 88

Contents xiii

4.4.2.1 Polarity . . . . . . . . . . . . . . . . . . . . . . . . 88

4.4.2.2 Surface quality of HT GaN . . . . . . . . . . . . . 88

4.4.2.3 Crystalline quality of HT GaN . . . . . . . . . . . 91

4.5 Summary & conclusions . . . . . . . . . . . . . . . . . . . . . . . . 92

5 Summary, Conclusions and Future Work 93

A Specifications of MOCVD Reactor 97

Bibliography 99

List of Figures

1.1 Non-centro symmetric wurtzite structure of GaN . . . . . . . . . . 4

1.2 Architecture of HEMT device . . . . . . . . . . . . . . . . . . . . . 9

2.1 Block diagram of MOCVD reactor . . . . . . . . . . . . . . . . . . 15

2.2 Aixtron MOCVD Reactor . . . . . . . . . . . . . . . . . . . . . . . 16

2.3 TT process plot . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 17

2.4 Variation in the equilibrium partial pressure of Ga (PGa) with theinput V/III ratio . . . . . . . . . . . . . . . . . . . . . . . . . . . . 19

2.5 Variation in the equilibrium partial pressure of Ga (PGa) with theinput V/III ratio for our MOCVD experimental conditions. . . . . . 20

2.6 Surface diffusion lengths of Ga adatoms . . . . . . . . . . . . . . . . 24

2.7 Mosaic block model: Tilt & Twist angles . . . . . . . . . . . . . . . 27

2.8 HRXRD goniometer rotational angles . . . . . . . . . . . . . . . . 28

2.9 AFM tip-surface force curve . . . . . . . . . . . . . . . . . . . . . . 29

2.10 DIC light microscopy optical lens diagram . . . . . . . . . . . . . . 32

2.11 k -space MOSS in-situ tool . . . . . . . . . . . . . . . . . . . . . . . 34

2.12 k -space MOSS stress measurement . . . . . . . . . . . . . . . . . . 35

3.1 Surface morphology of sapphire wafers after ex situ cleaning fol-lowed by in situ treatment in MOCVD reactor . . . . . . . . . . . . 41

3.2 Surface morphology of sapphire wafers after ex situ cleaning fol-lowed by in situ treatment in MOCVD reactor . . . . . . . . . . . . 42

3.3 N 1s XPS spectrum of nitrided sapphire wafers . . . . . . . . . . . 43

3.4 Normalized O 1s & N 1s XPS intensities from nitrided sapphire wafers 44

3.5 De-convoluted N 1s XPS spectra from nitrided sapphire wafers . . . 45

3.6 De-convoluted N 1s XPS spectra from nitrided sapphire wafers . . . 47

3.7 SEM morphologies of LT GaN . . . . . . . . . . . . . . . . . . . . . 48

3.8 AFM morphologies of LT GaN . . . . . . . . . . . . . . . . . . . . . 49

3.9 AFM surface roughness of LT GaN . . . . . . . . . . . . . . . . . . 50

3.10 High resolution x-ray (0002) ω-scan profiles for LT GaN NLs . . . . 51

3.11 High resolution x-ray 101̄1 φ-scan profiles for LT GaN NLs . . . . . 52

3.12 High resolution x-ray 101̄1 φ-scan profiles for LT GaN NLs . . . . . 52

4.1 Optical reflectivity trace of GaN on sapphire . . . . . . . . . . . . . 59

4.2 Surface morphology of N-polar GaN . . . . . . . . . . . . . . . . . . 59

xvii

List of Figures xviii

4.3 HRXRD rocking curve FWHM values of N-polar GaN in compari-son to conventional Ga-polar GaN . . . . . . . . . . . . . . . . . . . 60

4.4 Surface morphology of HT GaN (V/III = 485) grown for sapphirenitrided at TN = 530oC . . . . . . . . . . . . . . . . . . . . . . . . . 64

4.5 Surface morphology of HT GaN (V/III = 965, 1055 & 1205)) grownfor sapphire nitrided at TN = 530oC . . . . . . . . . . . . . . . . . 65

4.6 RMS surface roughness data of HT GaN (V/III = 965, 1055 &1205)) grown for sapphire nitrided at TN = 530oC . . . . . . . . . . 66

4.7 growth rate and roughening recovery time of HT GaN samples de-posited at different V/III ratios for sapphire nitridation at TN =530oC. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 66

4.8 HRXRD rocking curve FWHM values of HT GaN samples depositedat different V/III ratios for sapphire nitridation at TN = 530oC. . . 67

4.9 Optical micrographs of HT GaN grown at temperatures (a) 1000oC,(b) 1025oC and (c) 1050oC for sapphire nitridation at TN = 530oC. 68

4.10 AFM surface roughness data of HT GaN layers deposited at growthtemperatures: 1000, 1025 and 1050oCC for nitridation at TN =530oC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

4.11 Roughening recovery and growth rate of HT GaN samples grownat different growth temperatures for sapphire nitridation at TN =530oC. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 69

4.12 HRXRD rocking curve FWHM values of HT GaN samples grownat different growth temperatures for sapphire nitridation at TN =530oC. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 70

4.13 AFM surface roughness data of HT GaN as a function of LT GaNNL thickness for sapphire nitridation at TN = 530oC. . . . . . . . . 71

4.14 Roughening recovery time data of HT GaN as a function of LT GaNNL thickness for sapphire nitridation at TN = 530oC. . . . . . . . . 71

4.15 HRXRD rocking curve FWHM values of HT GaN samples grownas a function of LT GaN NL thickness for sapphire nitridation atTN = 530oC. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 72

4.16 SEM morphologies of HT GaN samples before and after KOH etchexperiment for sapphire wafers nitrided at TN = 530oC. . . . . . . . 72

4.17 Nomarski optical microscopy images of HT GaN layers for sapphirewafer nitrided at TN = 1100oC. . . . . . . . . . . . . . . . . . . . . 73

4.18 Nomarski optical microscopy images of HT GaN layers (V/III =965, 1130 & 1205) for sapphire wafer nitrided at TN = 1100oC. . . . 74

4.19 HRXRD rocking curve FWHM values of HT GaN layers (V/III =965, 1130 & 1205) for sapphire wafer nitrided at TN = 1100oC. . . . 75

4.20 SEM morphologies of HT GaN before and after KOH etch experi-ment for sapphire nitridation at TN = 1100oC. . . . . . . . . . . . . 76

4.21 SEM morphologies of HT GaN layers for sapphire nitridation at TN

= 1100oC. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 77

4.22 SEM morphologies after KOH etch experiment of HT GaN layersfor sapphire nitridation at TN = 1100oC. . . . . . . . . . . . . . . . 78

List of Figures xix

4.23 Optical microscopy images of HT GaN layers grown at differentgrowth temperatures for sapphire nitrided at TN = 1100oC. . . . . . 79

4.24 HRXRD rocking curve FWHM values of HT GaN layers grown attemperatures for sapphire nitridation at TN = 1100oC. . . . . . . . 79

4.25 Optical images of HT GaN layers grown at low growth temperaturesfor sapphire nitridation at TN = 1100oC. . . . . . . . . . . . . . . . 80

4.26 Surface roughness of HT GaN layers grown at low growth temper-atures for sapphire nitridation at TN = 1100oC. . . . . . . . . . . . 82

4.27 Optical images of HT GaN layers grown under N2 as carrier gas forsapphire nitridation at TN = 1100oC. . . . . . . . . . . . . . . . . . 83

4.28 Summary of x-ray rocking curve FWHM values of HT GaN layersover a wide range of growth conditions for sapphire nitrided at TN

= 530oC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 85

4.29 SEM morphologies of LT GaN annealed under different flow ratesof NH3 for sapphire nitridation at TN = 530oC. . . . . . . . . . . . 87

4.30 Summary of surface roughness values of HT GaN layers over a widerange of growth conditions for sapphire nitrided at TN = 530oC . . 87

4.31 Normalized O 1s intensity of nitrided sapphire wafers at differentstages for sapphire nitridation at TN = 1100oC. . . . . . . . . . . . 89

4.32 Summary of surface roughness values of HT GaN layers over a widerange of growth conditions for sapphire nitrided at TN = 1100oC . . 90

4.33 Summary of x-ray rocking curve FWHM values of HT GaN layersover a wide range of growth conditions for sapphire nitrided at TN

= 1100oC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 91

List of Tables

1.1 Properties of various commonly available foreign substrates avail-able for III-Nitride epitaxy . . . . . . . . . . . . . . . . . . . . . . . 2

1.2 Structural information and bond energy of wurtzite GaN . . . . . . 4

1.3 Band gaps of some typical WBG semiconducting materials in rela-tion to Ge and Si. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 7

2.1 Properties of various commonly available processing techniques forIII-Nitride epitaxy . . . . . . . . . . . . . . . . . . . . . . . . . . . 14

2.2 Group-III & V precursors used for III-Nitride epitaxy . . . . . . . . 17

2.3 Vapor pressure constants a and b for common metalorganic precursors 21

2.4 Activation energy EA for decomposition of GaN, and desorption ofGa and N atoms from GaN. . . . . . . . . . . . . . . . . . . . . . . 23

2.5 Surface diffusion barriers for Ga & N adatoms on Ga-polar andN-polar GaN surfaces. . . . . . . . . . . . . . . . . . . . . . . . . . 25

3.1 Normalized intensities of N 1s deconvoluted peaks from various pos-sible nitrided layer structures from TN = 530, 800 & 1100oC. Nor-malization has been done with respect to Al 2p peak. . . . . . . . 45

4.1 HT GaN growth parameter space for sapphire nitridation at TN =530oC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 62

4.2 HT GaN growth parameter space for sapphire nitridation at TN =1100oC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63

4.3 FWHM values of x-ray rocking curves for the HT GaN samplesgrown directly on LT GaN after ramp up, in relation to the samplesgrown on 4 min annealed LT GaN for sapphire nitridation at TN =1100oC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78

4.4 The surface roughness values of N-polar HT GaN grown at low &high growth temperatures for sapphire nitridation at TN = 1100oC . 81

4.5 HRXRD rocking curve FWHM values of N-polar HT GaN grownat low & high growth temperatures for sapphire nitridation at TN

= 1100oC . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 81

xxi

Abbreviations

MOCVD Metal Organic Chemical Vapor Deposition

HVPE Hydride Vapor Phase Epitaxy

MBE Molecular Beam Epitaxy

LT Low Temperature

HT High Temperature

NL Nucleation Layer

TMGa Tri Methyl Gallium

TMAl Tri Methyl Aluminium

TMIn Tri Methyl Indium

HEMT High Electron Mobility Transistor

2DEG 2 Dimensional Electron Gas

CTE Coefficient of Thermal Expansion

2D 2 Dimensional

3D 3 Dimensional

SEM Scanning Electron Microscopy

AFM Atomic Force Microscopy

RMS Root Mean Square

XPS X-ray Photoelectron Microscopy

TEM Transmission Electron Microscopy

CBED Convergent Beam Electron Diffraction

HRXRD High Resolution X-ray Diffractometer

XRC X-ray Rocking Curve

FWHM Full Width at Half Maximum

ID Inversion Domains

xxiii

Abbreviations xxiv

QW Quantum Well

SQW Single Quantum Well

LED Light Emitting Diode

DFT Density Functional Theory

Chapter 1

Introduction

1.1 Gallium Nitride

Wide band gap semiconductor Gallium Nitride (GaN) and its ternary and quater-

nary alloys along with InN and AlN, have excellent optical and electrical properties

because of which they have found applications in a range of optoelectronic and

high-power/high-frequency electronic applications [1–3]. In the absence of native

bulk III-Nitride crystals for homo-epitaxy, these materials have been deposited

heteroepitaxially on sapphire, Si and SiC substrates.

1.1.1 Substrates for GaN epitaxy

Table 1.1 tabulates the comparison between various foreign substrates for III-

Nitride epitaxy [2, 4]. Sapphire is the first substrate on which GaN technology

was developed and remains the most favored substrate for optoelectronic applica-

tions such as light emitting diodes (LEDs) in the visible and the upcoming UV

range. The cost is reasonable (non-miscut c-plane wafer) at $ 60 - 70 per 2 inch

wafer. The large lattice mismatch and coefficient of thermal expansion (CTE) mis-

match between GaN and sapphire (0001) results in GaN films with defect density

of the order of ∼ 108 cm−2 with residual stresses that are compressive in nature.

1

Chapter 1. Introduction 2

Property Si Sapphire SiC AlN GaN(111) (0001) (0001) (0001) (0001)

Lattice Mismatch % -16.9 16.02 3.48 2.41 0Available Size 12 8 6 4 2Cost per 2 wafer ∼ 25 ∼ 70 ∼ 1500 ∼ 2000 ∼ 2200Thermal Conductivity κ, W K−1 m−1 148 2.3 49 285 130Biaxial Modulus, GPa 165 458 440 510 470CTE (×10−6), K−1 3.6 5 4.46 3.04 3.94Dislocation Density, cm−2 ∼ 109 ∼ 108 ∼ 108 ∼ 105 ∼ 104

Table 1.1: Properties of various commonly available foreign substrates avail-able for III-Nitride epitaxy

The drawback of sapphire is its low thermal conductivity, which makes it difficult

to dissipate heat from devices during operation.

The alternative substrate, SiC, was for long not suitable for RF devices due to

electrical conductivity in the substrate. In 1999, the first semi-insulating SiC

substrates became available, and SiC has since become the leading substrate for

high-power devices, due to its high thermal conductivity. The other advantage

with SiC is its low lattice mismatch with GaN. The high cost is the major draw-

back for SiC substrates; currently semi-insulating 4H-SiC retails for $ 1200 to 1500

per 2 inch wafer, which is nearly 20 times the cost of sapphire.

The advantage of Si is its availability in large sizes and its cost ( $ 20 to 25 per 2

inch Si (111) wafer). Ga-polar GaN based HEMTs growth on Si (111) is well es-

tablished. The tensile nature of residual stress (due to CTE mismatch) in GaN on

Si (111) results in cracking on cool down unless specific stress mitigating strategies

are adopted. In a very commonly used solution, AlGaN transition layers (graded,

step-graded), are inserted between GaN and Si, to induce compressive growth

stress in the GaN layer to overcome the residual tensile stress on cool down [5–8].

The other drawback with Si is that the GaN layers grown are highly defective

(dislocation density is 1 to 2 orders of magnitude higher than GaN films grown on

sapphire and SiC) due to the large lattice and CTE mismatch.

In summary, all three substrates have their advantages and disadvantages as a

result of which they find their niche areas of technological relevance.


This thesis is concerned with growth on sapphire with specific focus on the ability

to control the polarity of GaN. This forms the subject matter of the next section.

1.2 Structure and polarity of GaN

GaN crystallizes in a stable non-centro symmetric wurtzite structure (P63mc) in

addition to a metastable zincblende structure (F4̄3m). The non-centro symmet-

ric wurtzite GaN can either terminate with metal atom (Ga) as the outer layer

along [0001] (+c) direction or anion (N) atom as the outer layer along [0001̄] (-c)

direction (Fig. 1.1 & Table 1.2). For Ga-terminated structure, each Ga atom is

bonded to 3 N atoms below where as for N-terminated structure each N atom is

bonded to 3 Ga atoms below (Fig. 1.1). This apparently simple crystallographic

actually has dramatic consequences for practically every application. As will be

shown in this thesis the growth modes of these two polarities also show very little

resemblance whatsoever to each other. In particular, most current technology has

been developed on Ga-polar GaN based materials as it is the default growth di-

rection. However, the N-polar material which is more difficult to grow, has many

interesting aspects that can enable better power or optoelectronic devices than the

Ga-polar material. Following a brief discussion on the ability of the III-Nitrides to

enable various applications, the high electron mobility transistor, a most common

and simple device is used to illustrate the difference between and the advantage

of the N-polar nitrides in section.

1.3 Polarity and growth of GaN

There are several processing techniques available for III-Nitride materials growth.

Among them metal organic chemical vapor deposition (MOCVD), molecular beam

epitaxy (MBE) and hydride vapor phase epitaxy (HVPE) are commonly used. The

processing technique used in this dissertation is MOCVD available at CeNSE, IISc

Bangalore.


Figure 1.1: Non-centro symmetric wurtzite structured GaN with metal-polarity/Ga-polarity and N-polarity.

Property GaN

Lattice Parameters, nm a = 0.319, c = 0.518Space Group P63mcBond Length 1, nm 0.319Bond Length 2, nm 0.319Bond Length 3, nm 0.195Bond Length 4, nm 0.195Bond Length 5, nm 0.518Bond Energy, eV/atom 8.92

Table 1.2: Structural information and bond energy of wurtzite GaN

The process parameters which are associated with the MOCVD reactor are: tem-

perature, pressure, flow rates of group-III precursors for Ga, Al & In, group-V pre-

cursors for N, and carrier gas flow rates (H2 and/or N2). A sapphire pre-treatment

for polarity selection of nitrides, called sapphire nitridation in this dissertation, in-

volves the transformation of sapphire (0001) surface to a thin complex unknown

AlOxN1−x [9] or AlN [10] layer prior to the growth of subsequent nitride layers.

The nitrided layer reduces the chemical dissimilarity between GaN and sapphire

and reduces the lattice mismatch [11, 12], both of which contribute to a reduction

in interface energy, thereby promoting lateral growth. It can act as a wetting layer

for the subsequently grown nitride layers and it can also aid in obtaining the ori-

entation relationships of these with respect to the underlying sapphire substrate

[12–14].


The two-step growth process, using MOCVD has been used extensively in im-

proving the structural, optical and electrical quality of nitride epitaxial layers in

particular GaN [15, 16]. This process involves the deposition of a low temperature

(LT) GaN/AlN nucleation layer (NL) on sapphire wafers prior to the deposition of

main high temperature (HT) GaN epitaxial layers. Each of these processing steps

are associated with various growth parameters. The LT NL can be controlled

with parameters like growth time, temperature and flow rates of precursors for

NL deposition. Similar parameter optimisation is required for HT GaN deposition

as well. In the two-step growth method, the LT layer helps to provide nuclei of

an optimum density and orientation for subsequently grown HT GaN layer. The

optimum nucleation density is provided by depositing a thin GaN/AlN NL at low

temperature followed by its controlled annealing. The HT GaN layer then grows

by addition of growth species, N and Ga, to the lateral edges of these nuclei.

While the 2-step method is use for N-polar material as well, fine differences exist.

The growth mechanism of N-polar GaN and its alloys is not the same as Ga-polar

GaN due to their differently terminated surfaces. From total energy density func-

tional theory (DFT) calculations it has been predicted that the Ga and N adatom

surface diffusion barriers are different for Ga-polar GaN and N-polar GaN sur-

faces [17]. Thus, the growth parameters which are optimized for Ga-polar GaN

growth on sapphire are not suitable for N-polar GaN growth. In spite of all the

success achieved with Ga-polar GaN and its alloys, N-polar GaN materials grown

under identical conditions yield rough surface morphology with hexagonal hillocks

[18, 19]. Growth parameters which effect the surface morphology and crystalline

quality of N-polar GaN are reactor pressure (P), carrier gas (H2 and/or N2), V/III

ratio and nitridation temperature (TN). The parameter V/III ratio is simply the

ratio of the fluxes of group-III to group-V precursors which are allowed into the

reactor chamber. The hillock density and size of hillocks on N-polar GaN are very

sensitive to V/III ratio and the density increases with the V/III ratio [18]. Longer

nitridation time also seems to affect the surface morphology of N-polar GaN [19].

It has been shown that N2 carrier gas improves the surface morphology of N-

polar GaN whereas H2 carrier gas suppresses the lateral growth of N-polar GaN


[20–22]. The above reports indicate that N-polar GaN growth is very sensitive

to growth parameters, and therefore detailed understanding and careful optimiza-

tion of growth parameters is required to obtain device quality layers. Keller et

al., showed that device quality N-polar GaN epitaxial layers without hexagonal

hillocks can be obtained by using intentionally miscut sapphire (0001) wafers of

2o to 4o along a- and m-directions [23]. However, the cost of highly miscut wafers

are ∼ 4 times the cost of non-miscut sapphire (0001) wafers.

1.3.1 GaN for lighting and power electronic applications

Optoelectronics was the first major application of III-Nitrides that was successfully

commercialized following the pioneering research of Akasaki, Amano and Naka-

mura. Nakamura et al. demonstrated high-brightness blue, green, and yellow light

emitting diodes (LEDs) with InGaN quantum well (QW) structures in 1995 [24].

Ternary InGaN alloys are used as the active layer in GaN-based LEDs and lasers

[24]. The performance of blue and green single quantum well (SQW) structure

LEDs has been improved and at 20 mA, the output power and the external quan-

tum efficiency (EQE) of the blue SQW LEDs were 5 mW and 9.1%, respectively.

Those of green SQW LEDs were 3 mW and 6.3%, respectively [25]. This LED

epitaxial structure is still the basic foundation for all currently commercially avail-

able first-generation blue and green LEDs [25]. It is found that the efficiency of

these LEDs is strongly depends on the In incorporation in the active InGaN layer

[26].

Following optoelectronics, GaN based power electronics and high frequency elec-

tronics is a currently growing area. The GaN based high electron mobility tran-

sistors (HEMTs) has evolved tremendously from their first modest demonstration

in 1993 [27] and is the workhorse of this group of applications. The large band

gaps of GaN and AlGaN provide for large breakdown fields, and thermal stability

of the materials allows for a high temperature of operation [28]. These excellent

properties have led to the demonstrations of devices, with current densities as high

as 2.3 A/mm [29], breakdown voltages around 0.9 kV [30], and power densities of


Material Chemical Symbol Band Gap Energy (eV) Type

Germanium Ge 0.7 IndirectSilicon Si 1.1 IndirectGallium Arsenide GaAs 1.4 DirectSilicon Carbide SiC 3.3 IndirectZinc Oxide ZnO 3.4 DirectGallium Nitride GaN 3.4 DirectDiamond C 5.5 DirectAluminum Nitride AlN 6.02 Direct

Table 1.3: Band gaps of some typical WBG semiconducting materials in re-lation to Ge and Si.

41.4 W/mm at 4 GHz [31].

Table 1.3 shows band gaps of some typical wide ban gap semiconductor materials

in relation to Ge and Si [1–3].

1.3.2 Effect of polarity

The tremendous impact of polarity is demonstrated through its effect on the de-

sign and performance of GaN/AlGaN HEMTs. The success of III-Nitride mate-

rial system is not only due to their direct WBG, intrinsic bulk material transport

properties (the p-material properties do not compare well with n-properties in III-

Nitride materials), but also due to the interface properties.

These polar nitrides behave differently across the interface of the active lay-

ers of devices (Fig. 1.2). For example, the difference in polarization across

the Al0.25Ga0.75N/GaN interface, in combination with discontinuity in conduction

band yields a two dimensional electron gas (2DEG) density across the interface,

well in excess of 1 × 1013 cm−2 [32] even in undoped systems. This is in con-

trast to AlGaAs/GaAs HEMTs, where doping is required to form a 2DEG. The

enhancement in the density, mobility and quantum confinement of 2DEG formed

across the interface are some of the primary requirements for HEMT devices. The

following are the parameters which control the properties of 2DEG formed across

the interface


� 2DEG carrier density can be enhanced by choosing nitride hetero struc-

tures which have higher discontinuity in polarization across the interface

(like GaN/AlN, of course polarization direction matters).

� The parameters which control 2DEG carrier mobility are surface roughness

of base layer, alloy scattering and critical thickness of AlGaN layer for HEMT

stack and dislocation density level in base layers.

� Quantum confinement of 2DEG can be improved by choosing nitride hetero

structures which give high conduction band discontinuity across the inter-

face.

Ga-polar AlxGa1−xN/GaN HEMTs (the device structure is shown in Fig. 1.2)

yields a 2DEG carrier density across the interface that depends on the composition

’x’ of Al in a thin AlxGa1−xN layer (which is under tension) sitting on GaN (Fig.

1.2a). Usually ’x’ ranges from ∼ 25 to 30%, which in turn limits the critical

thickness of AlxGa1−xN (∼ 25 to 20 nm for x ∼ 25 to 30%). The density of 2DEG

increases with the Al content in AlxGa1−xN due to increased spontaneous and

piezo polarizations in the AlxGa1−xN layer. However, higher Al content results

in formation of cracks in the AlxGa1−xN under tension and degrades the device

performance.

In contrast to the above, the design and architecture of HEMT devices for a

highly confined 2DEG density (> 1013 cm−2) involves the deposition of nitride

materials with N-polarity (Fig. 1.2b) [33]. N-polar GaN/AlGaN HEMTs improve

the confinement of 2DEG due to the WBG AlGaN back barrier (Fig. 1.2b). Strong

confinement results in sharp pinch off voltages.

The other advantages with N-polar GaN based HEMTs are low contact resistance

[34], sharp pinch off voltages due to AlGaN back barrier layers [35, 36]. Higher

trans-conductance can be expected with the same gate channel separation used

for Ga-polar AlGaN/GaN HEMTs [37, 38] as well as higher 2DEG density [2,

33, 39]. N-polar GaN alloy devices have shown impressive performance which

are in comparable to Ga-polar AlGaN/GaN HEMTs in the basic figures of merit

[33, 40].


Figure 1.2: Architectures of HEMT devices: (a) conventional Ga-polar Al-GaN/GaN HEMT, (c) shows the corresponding band structure , (b) & (d)N-polar GaN/AlGaN/GaN HEMT and its corresponding band diagram.

It has been also shown that N-polar GaN enhances the In incorporation in the

subsequently grown InGaN QWs on the GaN base layer for LED applications,

which results in LEDs with greater efficiency [41].


1.4 Thesis description

The theme of this dissertation is the development and study of N-polar GaN epi-

taxial layers on non-miscut sapphire (0001) wafers for power electronics and light-

ing applications. The processing technique used is MOCVD, available at CeNSE,

IISc Bangalore. The main scope of this thesis is to understand the relation be-

tween various growth parameters which control the polarity, surface roughness

and mosaicity of GaN and epitaxial layers on non-miscut sapphire (0001) wafers.

The various steps of growth such as sapphire nitridation conditions and LT GaN

growth conditions is extensively studied in this work. Growth parameters such as

NH3 flow rate, growth temperature, and NL thickness have been systematically

varied. Particular emphasis is placed on the correlation between the structure of

these precursor layers (nitrided layer and LT GaN NL) on the quality (surface

morphology and mosaicity) and polarity of the subsequently grown nitride semi-

conductor layers. We report for the first time device quality N-polar GaN epitaxial

layers on non-miscut sapphire (0001) wafers.

In chapter 2, various processing methods for synthesis of GaN layers are described

with particular emphasis on MOCVD method. The characterization tools used in

this dissertation for studying the chemical bond nature of nitrided sapphire surface

and micro-structural evolution (morphological and structural) of LT GaN NL &

HT GaN layers are described in detail.

Chapter 3, starts with the effect of ex situ cleaning followed by an in situ cleaning

on the surface morphology of sapphire (0001) wafers. The effect of nitridation

temperature (TN) on structural transformation of non-miscut sapphire (0001) sur-

face has been studied systematically. The structural evolution of nitrided layers

at different stages of their processing like as grown stage and thermal annealing

stage is investigated systematically. The chemical bond environment information

of the nitrided layers has been examined by x-ray photoelectron spectroscopy. It is


found that high temperature nitridation (TN ≥ 800◦C) results in a Al-N tetrahe-

dral bond environment on sapphire surface. In contrast, low temperature nitrida-

tion (TN = 530◦C) results in a complex Al-O-N environment on sapphire surfaces.

Micro-structural evolution of LT GaN NLs have been studied at every stage of

processing by scanning electron microscopy and atomic force microscopy. Surface

roughness evolution is measured from atomic force microscopy is described. It is

found that NLs processed on sapphire wafers nitrided at (TN ≥ 800◦C) showed

strong [0002] orientation with sub-nanometer surface roughness. In contrast, NLs

processed at (TN = 530◦C) showed the presence zincblende phase in the as grown

step with higher surface roughness, but acquired a greater degree of wurtzite [0002]

orientation after thermal annealing prior to high temperature GaN growth.

In chapter 4, polarity, surface quality and crystal quality of subsequently grown

HT GaN epitaxial layers is described in relation to the structure of the trans-

formed nitrided layers. High nitridation temperatures (TN ≥ 800◦C) consistently

yield N-polar GaN material whereas low nitridation temperatures (TN = 530◦C)

yield Ga-polar GaN. It is found that the relative oxygen atom concentration levels

in nitrided layers control the density of inversion domains in N-polar GaN layers.

The effect of various growth parameters such as (NH3 flow rate, growth temper-

ature, NL thickness) on surface morphology and mosaicity of both Ga & N-polar

GaN layers is discussed in detail. First time we report device quality N-polar GaN

layers at low growth temperatures 800◦ non-miscut sapphire (0001) wafers.

Chapter 5 contains the summary of the thesis results followed by directions and

suggestions for the future work.

Chapter 2

Experimental Techniques for GaN

Synthesis & Characterization

Group III-Nitride epitaxy on sapphire is a well-established platform and vari-

ety of experimental techniques such as metal organic chemical vapor deposition

(MOCVD), molecular beam epitaxy (MBE) and hydride vapor phase epitaxy

(HVPE) are available to deposit these semiconducting nitride materials. This

chapter contains a brief description of various processing techniques for GaN syn-

thesis with particular emphasis on MOCVD method followed by a brief introduc-

tion to various characterization tools which were used in this dissertation.

2.1 Experimental

The non-miscut 2 inch sapphire (0001) wafers were brought from the following ven-

dors EPISTONE, MONOCRYSTAL and KYOCERA. The unintentional miscuts

of the wafers provided by vendors were in the range of 0 ± 0.3o. The wafers were

diced into 1 cm2 size pieces by a MTI Corporation EC400 wafer dicing saw, which

is furnished with diamond blades of different thicknesses varying from 300 to 100

µm. The saw is computerized with a positional accuracy of 0.01 mm. The diced

wafers were then ultra-sonicated and cleaned ex-situ with acetone, isopropanol

13

Chapter 2. Experimental Techniques & Characterization 14

Property HVPE MBE MOCVD

Growth Rate ∼ 50 µm/h ∼ 50 nm/h ∼ 2 µm/h

Strength Large scale production Sharp interface Large scale productionGood quality film In-situ monitor Sharp interfaceVery high growth rate High purity High purity

H2 free ambient No ultra high vacuumPlasma assisted growth In-situ monitorLaser assisted growth High growth rateUniformity

Weakness Complex process Expensive Expensive sourcesExtreme temperature Need ultra-high vacuum Hazardous sourcesconditionsHazardous sources Low growth rate Large quantities of

NH3 required

Table 2.1: Properties of various commonly available processing techniques forIII-Nitride epitaxy

and de-ionized water. It is observed that the ex situ cleaning procedure is very

critical and it effects the surface morphology of sapphire wafers, as described in

subsequent chapter.

2.1.1 Processing techniques for GaN synthesis

Polycrystalline wurtzite GaN was originally synthesized using HVPE on sapphire

wafers in 1969 [42]. Subsequently several technical breakthroughs enabled single

crystal wurtzite GaN, low residual background carrier concentration in undoped

GaN, conductivity control of p-type GaN, epitaxial layer stacks for LEDs, LDs

and HEMTs, and these led to the first modest LED and HEMT devices being in-

troduced in 1993 [27, 43]. Table 2.1 shows the comparisons between the commonly

available processing techniques used for III-Nitride epitaxy.


2.1.2 Metal-organic chemical vapor deposition (MOCVD)

MOCVD is a process for the deposition of materials that utilizes volatile metal

organic compounds to transport metallic atoms that are relatively non-volatile

at the convenient deposition temperature. The organometallic compounds are

usually mixed with other source materials such as Hydrides that react to form

compound semiconductor films. Fig. 2.1 shows the simple block diagram of the

horizontal flow MOCVD reactor. Fig. 2.2 shows the AIXTRON 200/4 RF-S

MOCVD reactor which is installed at CeNSE, IISc.

Figure 2.1: Simple block diagram of horizontal flow MOCVD reactor

The group-III precursors (TMGa, TMAl & TMIn, Table 2.2) are injected into

the reactor chamber with the aid of carrier gas (H2/N2). Group-III & group-V

precursors are injected into the reactor chamber through separate nozzles. The

separation plate inside the chamber prevents the pre-reaction of these precursors.

Eventually, all these precursors are allowed to react onto the substrate surface

which is held at typical growth temperatures. Unwanted products formed will be

pumped to the scrubber for dilution before they are released to the outer atmo-

sphere. The substrate sits on a rotating susceptor to maintain the uniformity of

the deposited film. NH3 is used as the group-V hydride source and H2 used as the


Figure 2.2: Aixtron AIX 200/4 RF horizontal flow MOCVD reactor availableat CeNSE, IISc Bangalore.

carrier gas. All the group-III & group-V precursor bubblers and gas lines are fur-

nished with the necessary mass flow controllers (MFCs) and pressure controllers

(PCs) to control the flow rates of gases during the deposition of materials. The

maximum temperature limit of the reactor is 1500oC. The minimum reactor pres-

sure is 10 mbar and the maximum is 1000 mbar. The reactor chamber is equipped

with the removable quartz ware to prevent the reactor walls from contamination

during the growth of materials. The reactor is also equipped with an in-situ real

time thickness and stress monitor tool to understand the stress evolution and

growth behavior of films.


Precursor Chemical Formula

Trimethyl Gallium (TMGa) (CH3)3GaTrimethyl Aluminium (TMAl) (CH3)3AlTrimethyl Indium (TMIn) (CH3)3InAmmonia NH3

Carrier Gas H2, N2

Table 2.2: Group-III & V precursors used for III-Nitride epitaxy

2.1.3 Time-Temperature (TT) process plot for GaN epi-

taxy on sapphire by MOCVD

Fig. 2.3 shows the standard time-temperature (TT) process diagram for two-step

GaN epitaxy on sapphire wafers at two different nitridation temperatures TN =

530 and 1100oC.

Figure 2.3: Typical temperature Vs time (TT) process plot for two-step GaNepitaxy on sapphire wafers

The process starts with in situ thermal cleaning of sapphire wafers under the flow

of purified H2 at a temperature of 1100oC. The other process steps which are

involved in GaN epitaxy are

� Nitridation of sapphire at ∼ TN = 500 to 1100oC to enable a structural trans-

formation of sapphire surface to reduce the chemical dissimilarity between

sapphire and subsequently grown GaN layers


� LT GaN NL deposition at temperature ∼ 500 to 600oC to provide optimum

nucleation density for the subsequently grown HT GaN epitaxial layers

� Annealing of LT GaN NL at temperature ∼ 1000 to 1080oC involves decom-

position of NLs, surface migration of decomposed atoms, incorporation and

growth of new GaN nuclei.

� Growth of HT GaN epitaxial layer at temperature ∼ 1000 to 1080oC involves

epitaxial layer growth on annealed NLs, with adequate kinetics to grow the

necessary thickness.

The synthesis of GaN by using MOCVD technique is well known [44]. The for-

mation of GaN involves the complex chemical reactions between TMGa, NH3 and

their intermediate gas phase adducts (TMGa:NH3). The corresponding formation

energies for each step in the reaction involves are not documented well. The for-

mation energy of GaN can be estimated by assuming a simple model, where Ga

reacts with the NH3 according to the following reaction [44]

Ga (g) + NH3 (g) ⇔ GaN (s) +3

2H2 (2.1)

The estimated equilibrium constant K for the reaction

K =aGaN .P

3/2H2

PGa .PNH3

(2.2)

and

log10(K) = −12.2 + 1.78 ×(

104

T

)+ 1.79 log10(T) (2.3)

Where aGaN is the activity of GaN, P is the partial pressure of the reactants

which are participating in the above reaction. Fig. 2.4 shows the variation in

equilibrium partial pressures of Ga (PGa) with the corresponding input V/III ratio

[44]. V/III ratio is the ratio of flow rates of input group-V (NH3) and group-III

(Ga, Al & In precursors) sources. The input V/III ratio is varied by changing


the partial pressure PNH3 and by keeping PoGa at a constant value. The red color

points indicate the corresponding equilibrium partial pressures of Ga (PGa) for our

MOCVD experimental conditions at two different V/III ratios 260 and 1200.

Figure 2.4: Variation in the equilibrium partial pressure of Ga (PGa) with theinput V/III ratio. Copyright: Koukitu et al., JJAP, Vol.36, L1136, 1997. TheJapan Society of Applied Physics.

Fig. 2.5 shows the variation in equilibrium partial pressure of Ga (PGa) calculated

from equation (2.2) with the V/III ratio for our MOCVD experimental conditions,

which we have used in this dissertation.

It is found that the equilibrium partial pressure of Ga (PGa) decreases with the

increase in V/III ratio.

The driving force for the deposition of GaN is given by equation (2.4) and is

controlled by the partial pressure of NH3 (PNH3) (Fig. 2.4).

∆PGaN = PoGa − PGa (2.4)

Where PGa is the equilibrium partial pressure of Ga inside the reactor chamber,

PoGa is the input partial pressure of Ga which is kept at a constant value. For


Figure 2.5: Variation in the equilibrium partial pressure of Ga (PGa) withV/III ratio for our MOCVD experimental conditions.

example in our case, we have kept the corresponding input flow rate of TMGa at

4.1 sccm for GaN growth, and PvGa is the vapor pressure of Ga. Partial pressures

of reactants are varied by changing the corresponding flow rates. We do keep the

TMGa (which is in liquid form) bubbler at a particular temperature where TMGa

is in equilibrium with its vapor. In the subsequent process the carrier gas H2 is

passed through the TMGa bubbler to evaporate liquid TMGa, before it is fed into

the reactor chamber. The input flow rate of TMGa (or PoGa) is controlled by the

flow rate of carrier gas H2 which is passing through the bubbler. The vapor partial

pressure of metal organics (TMGa, TMAl etc.) depends upon the temperature.

The relation can be expressed as [45]

PvIII = 10(a− b

T) × 1013.5

760mbar (2.5)

The vapor pressure for group-III precursors and their corresponding temperatures

are shown in Table 2.3 [45].

The driving force decreases with the decrease in V/III ratio. Three kinds of modes

of deposition is possible [44] from a consideration of Fig. 2.4.


Precursor a b (K) Vapor pressure P vIII (mbar)

TMGa 8.07 1703 90 (0oC)TMAl 8.22 2134 9.60 (17oC)TMIn 10.52 3014 1.78 (17oC)

Table 2.3: Vapor pressure constants a and b for common metalorganic pre-cursors

PGa > PoGa: Etching

This mode is achieved by decreasing the V/III ratio or by reducing the partial

pressure of NH3 (PNH3). During this mode GaN starts decomposing and the

decomposition rate is higher than the incorporation rate of Ga and N adatoms

into the growing crystal. GaN decomposes at above 800oC at a pressure of 1 atm

and at lower temperature in vacuum [46, 47]. It was obtained that from mass

spectroscopy and thermogravimetric experiments, GaN decomposes into Ga and

NH3. There is huge scatter in the data reported for the activation energy barrier

for GaN decomposition, which spans from 0.34 to 3.1 eV and depends strongly

on conditions of the ambient (H2 and/or N2) in which the decomposition takes

place [46, 47]. The unit processes involved in GaN decomposition and growth are

explained in the subsequent section.

PGa > PvGa: Droplet formation

As discussed earlier in this context, the equilibrium vapor is the pressure where

the liquid TMGa is in equilibrium with its own vapor. This mode occurs when the

equilibrium partial pressure of Ga (PGa) inside the reactor chamber exceeds the

vapor pressure of Ga (PvGa). During this mode the Ga atoms tend to condense on

the growing surface and eventually form liquid metallic Ga droplets on the surface.


PGa < PoGa & PGa < Pv

Ga: Growth

This condition defines the growth window for GaN deposition. During this mode

of growth, the incorporation rate of Ga and N atoms into the growing bulk crystal

is higher than the decomposition rate of GaN. However, the incorporation rate

of adatoms in to the growing crystal is controlled by several factors such as the

life time of adatoms on the growing surface and the diffusion lengths of adatoms

which in turn controlled by the corresponding surface diffusion barriers at typi-

cal MOCVD growth temperatures. The driving force for this growth window is

increases with the increase in partial pressure of NH3 (PNH3), which is consistent

with the experimental MOCVD trends.

The above results indicate that for GaN growth to occur the equilibrium partial

pressure of Ga (PGa) is always has to be lower than the vapor pressure PvGa and

the input partial pressure PoGa.

GaN decomposition has been studied extensively and several mechanisms were

proposed to explain it. The following reactions are reported for GaN decomposi-

tion [48, 49].

2GaN (s) → 2Ga (g) + N2 (g) (decomposition) (2.6)

2GaN (s) → 2Ga (l) + N2 (g) → 2Ga (g) + N2 (g) (desorption) (2.7)

2GaN (s) → GaN (g) or [GaN]x (g) (sublimation) (2.8)

It was reported that H2 could assist the GaN decomposition by reformation of

NH3 via [47].

2GaN (s) + 3H2 (g) → 2Ga (l) + 2NH3 (g) (2.9)

Table 2.4 shows the activation energy barriers measured for the above reactions

under the ambient of H2 [47–49]. The activation energies are measured by sev-

eral techniques such as thermogravimetry, mass spectrometry and reflection high

energy electron diffraction (RHEED).


Process EA eV Technique

GaN decomposition 3.1 Thermogravimetry, Mass spectrometryGa desorption from GaN 2.2 - 2.76 RHEED, Growth rate Vs Growth temperatureN desorption from GaN 6.1 RHEED

Table 2.4: Activation energy EA for decomposition of GaN, and desorption ofGa and N atoms from GaN.

The other unit step which takes place in GaN decomposition and growth is surface

diffusion. The factor surface diffusion is critical and it controls the surface quality

of the growing surface.

The mode of growth is classified into two types based on surface diffusion lengths

of the adatoms on the growing surface. The mode step-flow (2D mode) occurs

when the diffusion length of adatoms is longer than the terrace width or step

width involved during growth. In contrast, island mode (3D mode) occurs when

the diffusion length is smaller than the corresponding terrace width [50]. The

diffusion lengths are controlled by the corresponding surface diffusion barriers

at typical process temperatures. The diffusion length can be estimated by the

following equation [50]

λs =√

Dsτs =

(a2 ν= exp

(−Esd

kT

)τs

) 12

(2.10)

Where, Ds is the diffusion coefficient, ν= is vibrational frequency of adatoms paral-

lel to the surface, a is mean distance between the adsorption sites, Esd correspond-

ing surface diffusion activation energy and τs can be classified into two cases, which

are life time before desorption (τD), which depends on the desorption rate and the

life time before lattice incorporation (τI), which depends on growth rate. The

mean life time of residence of adatoms (τD) on the growing surface before being

re-evaporated and is given by

τD =1

ν⊥exp

(Edes

kT

)(2.11)

Where ν⊥ is the vibrational frequency of adatoms normal to the surface and Edes is


the desorption activation energy of an adatom from the growing surface. Assuming

ν = ν⊥ = ν=, we have the following expression for surface diffusion length [50]

λs,Ga = a exp

(Edes − Esd

2kT

)(2.12)

Koleske et al. [51] measured λs,Ga based on τD using equation (2.12) and is shown

in Fig. 2.6 as a solid line, which indicates λs,Ga decreases slightly as the temper-

ature increases. This is due to the increase in desorption rate, which results in

reduced τD. But, at typical MOCVD growth pressures Ga desorption appears to

be suppressed and eventually yields longer τD. The limiting life time in such cases

is the incorporation life time (τI) into the growing lattice. The value of τI depends

inversely on growth rate, as slower the growth rate the larger τI.

Figure 2.6: Measured surface diffusion length λs,Ga of Ga adatoms at differenttemperatures. The solid line corresponds to λs,Ga based on τD and dashed linescorrespond to λs,Ga based on τI. Koleske et al., JAP, 84, 1998, 1998. Reprintedwith permission.

Fig. 2.6 shows the dependence of τI on the growth rate and is λs,Ga plotted as

dashed lines for different growth rates ranging from 3 to 100 nm/min. τI was

measured by the thickness per monolayer (actually bilayer, which is 0.258 nm and

represents one half of the lattice parameter along c-direction) by the growth rate. It

was also shown that as growth rate decreases, both and increases, which eventually


Adatom Ga-polar GaN Ga-polar GaN N-polar GaN N-polar GaNN-terminated N-terminated

(eV) (eV) (eV) (eV)

Ga 0.4 1.8 0.2 1.0N 1.4 - 0.9 -

Table 2.5: Surface diffusion barriers for Ga & N adatoms on Ga-polar andN-polar GaN surfaces.

yields a more ordered lattice because the number of adatoms that incorporate is

increased. It was found that the diffusion length λs,N for N adatoms is much

smaller than Ga adatoms and is due to the high surface diffusion barrier for N

adatoms when compared to Ga adatoms [17]. At the same time the other possible

reason for lower λs,N values, is due to high vapor pressure of N2, once N migrates

to next to an adjacent N, N2 forms and desorbs.

From total energy density functional theory (DFT) it is predicted that the surface

diffusion barriers for Ga & N adatoms on (0001) and (0001̄) surfaces are not same

[17]. Table 2.5 shows the estimated surface diffusion barriers for Ga & N adatoms

at typical growth temperatures.

From Table 2.4 it can be suggested that the estimated diffusion lengths for Ga &

N adatoms on N-polar GaN surface must be higher than on Ga-polar GaN surface.

As we described in Sec. 1.4.2 of Chapter 1, the surface morphology of N-polar GaN

grown under identical conditions is rough and is associated with inversion domains

(IDs). In such cases the estimation of corresponding surface diffusion lengths of

Ga & N adatoms on both domains (IDs) is complicated, which eventually makes

difficult to understand the behavior of adatoms on the growth surface and hence

the growth mechanism of N-polar GaN.

2.2 Characterization

This section describes some of the standard characterization tools that have been

used to analyze our samples. The primary characterization tools employed in this


study are high resolution x-ray diffractometry, x-ray photoelectron spectroscopy,

atomic force microscopy, scanning electron microscopy, differential contrast in-

terference optical microscopy (Nomarski) and the in situ reflectivity and stress

measurement tool. In some cases transmission electron microscopy has been used

to examine the polarity of our samples. All these are available at the Centre for

Nano Science & Engineering (CeNSE), IISc Bangalore. The transmission electron

microscopy studies were done at the Advanced Facility for Microscopy and Micro

analysis (AFMM), IISc Bangalore and Defense Metallurgical Research Laboratory,

Hyderabad.

2.2.1 X-ray photoelectron spectroscopy (XPS)

This technique has been extensively used to study the chemical nature of sap-

phire wafers nitrided at different temperatures. This ex situ tool was supplied

by Kratos Analytical AXIS Ultra DLD, Manchester. It is a surface sensitive tech-

nique, which analyzes the kinetic energy (KE) of photoelectrons which are emitted

from the sample surface via photoemission process. The spectrometer measures

the binding energy (BE) of the ejected electrons from the measured KE. The BE

of the photoelectrons which are emitted from the sample surface is given by the

following relation.

BE = hν − KE − φ (2.13)

Where φ represents the combined electron spectrometer and sample work functions

and is an instrument dependent factor normally derived for each instrument as part

of a calibration procedure. The technique derives its chemical sensitivity from the

fact that nearest neighbor atoms will have a direct effect upon the binding energy

of the core level electrons. Therefore any change in the chemical environment such

as oxidation state will lead to a modification of the KE.

The instrument has the following key components: X-ray source, electron transfer

lens, electron energy analyzer, and detection system. All of these components are

contained within an ultra-high vacuum envelope (10−9 Torr). The spectrometer


is equipped with monochromatic x-rays as the primary source. The x-ray gun is

used in combination with a focusing monochromator due to which only the Al

Kα component is diffracted from the quartz crystal. The natural line width of

this component is < 0.26 eV. A charge balance option is available for insulators.

An Ar ion beam was used to clean the sample surface from contamination and/or

for depth profile information. Charge corrections were done with respect to C

1s (Carbon) peak located at 284.8 eV. Elemental quantification of all the survey

spectra and high resolution scans were carried out by CASA XPS software.

2.2.2 High resolution x-ray diffractometer

HRXRD has been used to examine the crystal mosaicity of our samples. The

epitaxial film is assumed to consist of single crystallites called mosaic blocks with

misorientations with respect to each other [52, 53]. The out-of-plane rotation

of the blocks perpendicular to the surface normal is the mosaic tilt, and the in-

plane rotation around the surface normal is the mosaic twist (Fig. 2.7). The

Figure 2.7: Mosaic blocks in a crystal showing tilt and twist with respectto each other. Tilt corresponds to out-of-plane rotation whereas twist corre-sponds to in-plane rotation of mosaic blocks. Srikant et al., JAP, 82, 4286,1997. Reprinted with permission.


average absolute values of tilt and twist are directly related to the full width

at half maximum (FWHM) of the corresponding distributions of crystallographic

orientations [52, 53]. The rocking curve measurement (ω-scan) method is the

most frequently employed method to analyze the mosaicity of epitaxial layered

materials. In this geometry, the detector remains stationary and the sample is

rotated about ω-axis (Fig. 2.8).

Figure 2.8: X-ray diffractometer with different rotational angles (ω, ψ, φ) inrelation to the sample reference frame.

ω-scan (000l) reflections are used to measure tilt of mosaic blocks [52, 53]. Tilt is

sensitive to screw and mixed dislocations but not to edge dislocations because edge

dislocations do not distort (000l) planes as their burger vectors (1/3 < 112̄0 >)

lie within those planes. Twist is caused by edge and mixed dislocations and is

usually measured with ω-scans of off-axis reflections (h or k 0). Off-axis reflections

occurring at higher ψ values are used, so that FWHM is dominated by in-plane

twist. Thus, lower FWHM values for (000l) and off-axis reflections indicate that

epitaxial films of better crystalline quality with low mosaicity.

The instrument that has been used for XRC analysis was supplied by RIGAKU

SmartLab, Japan. This tool was equipped with the 2 bounce and 4 bounce Ge

(220) incident beam monochromators. The Cu (Kα1: with λ = 0.15405 nm and

Kα2: with λ = 0.15443 nm) x-ray source is collimated and monochromated by a

4 bounce Ge (220) monochromator to achieve high angular resolution.


2.2.3 Atomic force microscopy (AFM)

The AFM used in this dissertation is Dimension Icon ScanAsyst Bruker. The

AFM has been used to study the surface features of LT GaN NLs and the sur-

face roughness evolution of HT GaN epitaxial layers.AFM can be operated in two

modes such as contact mode and non-contact mode. Tapping mode falls in some-

where between the contact and non-contact mode. All AFM measurements were

done in this dissertation using tapping mode technique. The classification of AFM

working modes is described in Fig. 2.9.

Figure 2.9: Typical inter atomic force curve between the cantilever tip ofAFM and sample surface

The contact mode falls in the repulsive interaction region whereas the non-contact

mode falls in the attractive Van der Waals interaction region. The tapping mode

falls somewhere in between the two interaction regions. In the contact mode the

deflection of cantilever is kept constant. In contrast, the tip is oscillated at it’s

resonant frequency in non-contact mode and the amplitude of the oscillation is

kept constant. Contact mode imaging is heavily influenced by frictional and adhe-

sive forces, and can damage samples and distort image data. Non-contact imaging

generally provides low resolution and can also be hampered by the contaminant

(e.g., water) layer which can interfere with oscillation.

Tapping mode imaging takes the positive features of both these modes and over-

comes problems associated with friction, adhesion, electrostatic forces, and other


difficulties that plague conventional AFM scanning methods, by alternately plac-

ing the tip in contact with the surface to provide high resolution and then lifting

the tip off the surface to avoid dragging the tip across the surface. Tapping mode

imaging is implemented in ambient air by oscillating the cantilever assembly at or

near the cantilever’s resonant frequency using a piezoelectric crystal. The piezo

motion causes the cantilever to oscillate with a high amplitude( typically greater

than 20nm) when the tip is not in contact with the surface. The oscillating tip is

then moved toward the surface until it begins to lightly touch, or tap the surface.

During scanning, the vertically oscillating tip alternately contacts the surface and

lifts off, generally at a frequency of 50,000 to 500,000 cycles per second. As the

oscillating cantilever begins to intermittently contact the surface, the cantilever

oscillation is necessarily reduced due to energy loss caused by the tip contacting

the surface. The reduction in oscillation amplitude is used to identify and measure

surface features.

When the oscillating cantilever approaches the surface the amplitude of oscillation

decreases. This is due to the additional restoring force works on the cantilever and

this can be seen as an increase in the spring constant of the cantilever. This drop

in amplitude can be used as the feedback parameter for AFM imaging, just like

the cantilever deflection in contact mode.

Tapping mode inherently prevents the tip from sticking to the surface and causing

damage during scanning. Unlike contact and non-contact modes, when the tip

contacts the surface, it has sufficient oscillation amplitude to overcome the tip-

sample adhesion forces. Also, the surface material is not pulled sideways by shear

forces since the applied force is always vertical.

2.2.4 Differential interference contrast (DIC) light microscopy

DIC light microscopy (Nomarski) is a technique which produces impressive 3D-

like images of unstained specimens. The shadowing effects of the technique are

remarkable. DIC light microscopy was used in this dissertation to provide an initial

idea about surface quality of the HT GaN epitaxial layers. Nomarski microscopy


utilizes a system of dual beam interference optics that transforms local gradients

in optical path length in a specimen into regions of contrast in an image. The

working principle of the Nomraski light microscope is described in Fig. 2.10

� Light passes through a standard polarizer before entering the condenser,

producing plane-polarized light

� This light enters a Wollaston prism located in the front focal plane of the

condenser. The prism interacts with the polarized light to produce two sep-

arate wavefronts which are parallel and polarized perpendicularly to each

other. These are termed the ordinary (O) and extraordinary (E) rays. Fur-

thermore, these two wavefronts are separated by a very small difference (less

than the resolution of the system) generally ranges from 0.2 - 2 µm and the

separation is depends upon the prism used

� The two wavefronts pass through the specimen, phase of one beam may be

differentially shifted with respect to the other if there is a local gradient in

optical path length.

� The light now enters a second Wollaston prism set-up which recombines the

wavefronts. If there has been a phase shift between the two rays as they

pass through areas of different refractive index then elliptically polarised

light is the result. DIC optics are designed to convert these phase differences

between the two beams of light into amplitude differences which can be

visualized by the human eye

� Finally the light enters a second polarizing filter, termed an analyzer. The

initial polarizer and this analyzer form crossed polars. The analyzer will

permit the passage of some of the elliptically polarized light to form the

final image. All the remaining light will be blocked by the analyzer.


Figure 2.10: Typical optical lens diagram of differential interference contrast(DIC) light microscopy. A. Lasslett, Microscopy Division, Olympus UK Ltd,Southall, Middlesex, UK. Reprinted with permission

2.2.5 Scanning electron microscopy (SEM)

The SEM used in this study was Carl Zeiss Ultra 55 FESEM. This tool has been

used to study the morphology of our samples at different stages of their process.

The probe current range: 15 to 20 nA, aperture size range: 7.5 to 120 µm and

accelerating voltage range: 5 to 30 kV.

2.2.6 Transmission electron microscopy (TEM)

TEM was used in this dissertation to examine the polarity of the grown epitax-

ial layers. Convergent beam electron diffraction (CBED) technique was used in

combination with java electron microscope simulator (JEMS) simulated patterns

to examine the polarity of our samples at different thicknesses. A 300 kV Technai


TM G2 F30 S-TWIN TEM with FEG at Advanced Facility for Microscopy & Mi-

croanalysis (AFMM) IISc and a FEI Tecnai G2-20T 200KV TEM at Defense Met-

allurgical Research Laboratory (DMRL) Hyderabad were used to examine GaN

layers in this dissertation.

Cross sectional TEM foils were made using the regular sandwich technique, in

which two pieces of the sample are cut and glued together keeping the film sides

of the two pieces facing each other. The sandwich assembly is inserted and stuck

inside a slotted rod and this complete assembly is pushed and glued inside a hollow

tube with 3 mm outer diameter. Thin slices from this combination is cut by using

Buehler ISOMET low speed saw and thinned down to ∼ 100 µm using MULTI-

PREPTM equipment manufactured by Allied Instruments. TEM foils were further

thinned down to electron transparency by using conventional ion-beam thinning

technique using Gatan PIPS (Precision Ion Polishing System) with Cold Stage.

2.2.7 In situ reflectivity and stress monitor analysis tool

The in situ tool that was used in this study is the k -space Associates multi-

beam optical stress sensor (MOSS). This technique has been used to monitor the

growth behavior of GaN films on sapphire. The primary information that can be

extracted is the real time film thickness, real time stress and growth rate. This in

situ capability enables the monitoring of film stress and thickness as it develops.

In the MOSS technique (Fig. 2.11a), multiple parallel laser beams illuminate

the surface simultaneously and the beam positions are measured with a CCD

detector. The measurement of film thickness is based on interference between the

light beams reflected from the top surface of the film and from the film/substrate

interface (Fig. 2.11b). If the film is not too absorbing, interference between the

reflected beams results in a modulation of the reflected intensity. The interference

of beams depends upon the path difference of the rays which in turn depends on

growing film thickness. The period of the intensity oscillations is used to determine

the growth rate of the film. In addition, the amplitude of the intensity oscillations

depends on the reflectivity of the interfaces, and therefore also a probe of surface


morphology during growth. Fig. 2.11c shows an example of GaN two-step growth

on sapphire.

Figure 2.11: A schematic diagram of k -space MOSS set up (a) multiple par-allel laser beams fall on the substrate surface and will be reflected and detectedby CCD camera, (b) interference effects for the rays which are reflected fromthe surface and the film/substrate interface, and (c) a sample reflectivity tracefor GaN epitaxy on sapphire, which is recorded by in-situ k -space MOSS toolattached to MOCVD reactor.

The growth rate is measured by

λ

2η= Period of one oscillation (2.14)


Where, λ is wavelength of light used = 660 nm.

η is refractive index of the growing film and is 2.3 for GaN.

The in situ tool can also monitor the real time stress evolution in films during their

growth. MOSS stress monitoring is based on measuring the curvature induced in

the substrate by the stressed film (Fig. 2.12).

Figure 2.12: The radius of curvature of the stressed film effects the spacingbetween laser beams which are reflecting from the surface of the film.

The relationship between the film stress, σ , and the radius of curvature, R, is

given by the following equation developed originally by Stoney [54]

σ =Msh

2s

6hfR(2.15)

Where hf is the film thickness, Ms is the biaxial modulus of the substrate, hs is the

substrate thickness, and R is radius of curvature. For flat films R is infinite, so that

parallel beams of light will be reflected from the surface and the beam separation

will be same as the incident beam separation. If a surface is curved, parallel beams

of light that strike it at different positions will be reflected at different angles (Fig.

2.12). The amount of deflection of the light beam is related to the curvature by

1

R=δd

dcosα / 2L (2.16)

Where δd is the measured change in the spacing of the beams, d is the initial

spacing, L is the distance from the sample to the detector and α is the angle

between the incident beams and the surface normal.


This tool is used extensively to understand the growth behavior of GaN epitaxial

layers on sapphire (0001) wafers.

Chapter 3

Sapphire Pre-treatment &

Microstructural Evolution of LT

GaN

It is well known that the high energy interface, lattice mismatch and CTE mis-

match between HT GaN and sapphire prevents the HT GaN in growing laterally

and promotes 3-dimiensional growth of GaN with poor crystalline quality. The

structural quality of HT GaN layers is dramatically improved, when a thin LT

GaN/LT AlN NL is deposited on sapphire prior to the growth of HT GaN. The

function of the subsequently grown LT NL could be described as [55]

� The supply of nucleation sites of low orientational fluctuation and

� The promotion of lateral growth of HT GaN

This chapter addresses the following two things. First one is to understand the

structural transformation of sapphire surface at three different nitridation tem-

peratures TN = 530, 800 and 1100oC. Second one is to understand the relation

between the transformed nitrided layers on sapphire surface and the microstruc-

tural evolution of the subsequently grown LT GaN NLs.

37

Chapter 3. Sapphire Pre-treatment & Microstructural Evolution of LT GaN 38

3.1 Background

Nitridation is the foundation of N-polar nitride materials deposited on sapphire

wafers. It is well known that GaN growth on non-nitrided sapphire (0001) surface

results in Ga-polar GaN while that grown on sapphire (0001) surface nitrided in

the temperature range above 950oC results in N-polar GaN [19, 56, 57]. The nitri-

dation of sapphire surface prior to the growth of LT GaN NL is known to reduce

the chemical dissimilarity and lattice mismatch between GaN and sapphire [9–12].

Literature suggests that the nitrided layer is complex unknown AlOxN1−x [9], AlN

and/or AlN in O atom rich environment [10]. Sapphire surface gets damaged if it

is exposed to NH3 and H2 ambient for longer times during nitridation at higher

temperatures [9, 19]. The possible reason behind surface damage is due to the

strong chemical reaction between NH3 and sapphire [19]. Sun et al., showed that

higher nitridation temperatures result in protrusions on sapphire surface during

nitridation, which in turn affects the surface morphology of subsequently grown

N-polar GaN epitaxial layers [19]. The chemical nature and structure of these

protrusions are unknown.

The structural transformation of sapphire surface critically depends upon the pro-

cess conditions such as flow rates of NH3, H2 and/or N2, nitridation temperature

(TN) and process time. It is well known that NH3 decomposes at above 500oC and

the complete decomposition of NH3 is assumed to occur easily at 1100oC [58]. The

mechanism of nitridation involves the counter diffusion of N (inward diffusion) and

O (outward diffusion) atoms from the sapphire surface [58]. Diffusion data for N

and O atoms in sapphire lattice is well documented [58, 59].

It has been reported that the as-grown LT GaN NL on sapphire contains cubic

(zincblende) phase in addition to stable wurtzite phase [59, 60]. The usual tem-

perature range for LT GaN growth falls in between 500 to 600oC [59, 60]. It has

been reported well in the literature that the morphology, distribution, polarity

and orientation of LT GaN critically depends upon the sapphire pre-treatment

called sapphire nitridation in this dissertation. The surface morphology of the as

deposited LT GaN on non-nitrided sapphire wafers is very rough with 3D discrete


islands, in contrast the surface morphology is relatively smooth with complete

coverage for the LT GaN deposited on nitrided sapphire wafers [11, 61]. It was

also found that the as grown LT GaN deposited on nitrided sapphire results in

Ga-polar LT GaN and the polarity is transformed to N-polar in the subsequent

annealing steps [57], the polarity of LT GaN is determined by potassium hydroxide

(KOH) etching methods [57]. The annealing of LT GaN NLs take place at typical

GaN epitaxial layer growth temperatures ∼ 1000 to 1080oC under the ambient of

NH3, H2 and/or N2 and it was found that the structural quality of subsequently

grown HT GaN epitaxial layer is a strong function of annealing conditions of LT

GaN [62–66]. N-polar GaN is more sensitive to the ambient present during its

annealing process and it was found that N-polar GaN etches at faster rate than

Ga-polar GaN during annealing period under the ambient of NH3 and H2 at typ-

ical MOCVD annealing temperatures [67]. The crystalline quality and electrical

properties of HT GaN have also been studied as a function of growth rate/thick-

ness of LT GaN NL [63, 68]. The distribution and density of LT GaN nuclei for

the subsequent HT GaN growth can be controlled by the growth parameters such

as V/III ratio and reactor pressure during the growth of LT GaN [62, 69].

This chapter therefore explores the modification of the non-miscut sapphire (0001)

wafers at three nitridation temperatures TN = 530, 800 and 1100oC to provide a

link between GaN two-step growth on non-nitrided sapphire and the conventional

high temperature nitridation.

3.2 Experimental

Experiments were carried out in an Aixtron 200/4 RF-S horizontal flow MOCVD

reactor with removable quartz ware. The in situ thermal treatment of wafers were

carried out at 1100oC under the flow of purified H2 for 10 min before nitridation.

The nitridation step was performed at three different temperatures TN = 530,

800 and 1100oC and 200 mbar reactor pressure. Wafers were nitrided under the

ambient of NH3 and H2 for 1 min. NH3 flow rate was kept constant at 1500 sccm

throughout the process. The nitrided wafers were taken out from the MOCVD


reactor chamber for ex situ XPS characterization.

After sapphire nitridation, LT GaN was grown to a 60 nm thick layer to under-

stand the effect of nitridation on microstructural evolution of NL. The thickness

calibration was done by the in situ optical tool as discussed in Sec. 2.2.7 of Chap-

ter 2. The effect of LTGaN thickness on HT GaN growth is described in Chapter

4. LT GaN NL was deposited on nitrided wafers at 530oC for 8 min (to deposit 60

nm thick layer), 200 mbar, at a V/III ratio of 2535, TMGa flow rate of 0.6 sccm

and H2 flow rate of 6500 sccm. The wafers were then heated up from the LT GaN

NL growth temperature to the LT GaN NL annealing condition in 5 minutes. The

NLs were annealed thermally for 4 minutes at 1000oC and a reactor pressure of

400 mbar in a mixture of 2000 sccm of NH3 and 4700 sccm of H2.

During the annealing process of NL, it is possible that the nitrided sapphire surface

regions which are not covered by LT GaN NLs, are exposed to NH3 and nitrided

further. We have investigated these changes by subjecting the nitrided sapphire

surface to identical ramp up and hold cycles as they would encounter during the

deposition of LT GaN NL and its subsequent process. Therefore, after the nitrida-

tion step, the temperature is ramped to the NL deposition temperature of 530oC

and held at 530oC for 8 min followed by temperature ramp up to NL annealing

temperature of 1000oC in 5 minutes. Once the temperature reaches the annealing

condition, it is held for 4 min.

The main tool used to characterize nitrided layers was XPS as described in Sec.

2.2.1 of Chapter 2. SEM & AFM are used for morphological studies, and HRXRD

is for structural analysis of LT GaN NL.

3.3 Results

This section starts with a few results on the effect of ex situ and in situ cleaning on

surface morphology of sapphire (0001) wafers followed by subsequent nitridation

and LT GaN NL deposition steps.


3.3.1 In situ thermal treatment of sapphire wafers

As described in Sec. 2.1.3 of Chapter 2, the diced sapphire wafers undergo ex

situ cleaning with acetone, isopropanol and de-ionized water and are dried with

a stream of N2 gas to dry solvents from the surface until no droplets are found.

In many cases, the surface may contain microscopic droplets which might leave

organic residues on the surface upon drying with N2. Fig. 3.1a shows the surface

of sapphire which has been subjected to the ex situ cleaning process with acetone,

isopropanol, de-ionized water and dried with N2 gas prior to the in-situ thermal

cleaning in the MOCVD reactor.

Figure 3.1: Surface morphology of sapphire (0001) wafer which is cleaned exsitu, (a) microscopic organic remnants on sapphire surface after drying with N2

gas, and (b) morphology of sapphire surface after in-situ treatment with H2 at1100oC for 10 min in MOCVD reactor. Fig. (c) & (d) are magnified views of(b).

Fig. 3.1b shows the surface morphology of sapphire after in situ treatment under

the flow of purified H2 at 1100oC. The etched morphology is due to the reaction


between H2 and the organic residues which are left after ex situ cleaning [70]. Fig.

3.1c & d are magnified images of in situ treated sapphire surface. To resolve this

issue, we have optimized the ex situ cleaning procedure and it is found that the

cleaning of wafers many times with de-ionized water after acetone and isopropanol

helps to remove all the organic residues from the surface. Fig. 3.2 shows the

surface of sapphire which is cleaned ex situ with acetone, isopropanol and multiple

times with de-ionized water followed by an in situ H2 treatment at 1100oC in the

MOCVD reactor.

Figure 3.2: Surface micrographs of sapphire wafer which have undergone anoptimized ex situ cleaning procedure, (a) & (b) are SEM micrographs at twodifferent magnifications, and (c) & (d) are corresponding AFM morphologieswith sub-nanometre RMS roughness ∼ 0.35 nm.

AFM image shows (Fig. 3.2c & d) atomic steps on the sapphire surface with rms

surface roughness ∼ 0.35 nm after in situ treatment with H2 at 1100oC in the

MOCVD reactor.


3.3.2 Nitridation of in situ treated sapphire wafers

3.3.2.1 As nitrided sapphire wafers

Fig. 3.3 shows the high resolution x-ray N 1s photoelectron spectrum from sap-

phire wafers nitrided at three different temperatures TN = 530, 800 and 1100oC.

Figure 3.3: The N 1s x-ray photoelectron peak from sapphire wafers nitridedat TN = 530, 800 and 1100oC

The distinct N 1s binding energy photoelectron peak indicates that there is incor-

poration of N atoms into the sapphire lattice during nitridation. Three effects of

nitridation temperature are discernible: the change in intensity of the N1s peak,

the shift in peak positions and a peak broadening is observed with the decrease

in nitridation temperature. The variation in N1s peak intensity indicates that the

increase in incorporation of N atoms into the sapphire surface with the nitridation

temperature. The shift in peak position and peak broadening indicates that the

environment of the N atoms in the modified sapphire undergoes changes with ni-

tridation temperature. Since surface modification is expected to occur by inward

diffusion of N into sapphire to replace O atoms. Fig. 3.4 shows the normalized


intensities of N 1s and O 1s peaks and are plotted against nitridation tempera-

tures. The normalization is done with respect to the Al 2p peak intensity. Increase

in nitridation temperature results in greater N/O ratios in the modified sapphire

surface.

Figure 3.4: Normalized intensities of (a) O 1s and (b) N 1s peaks from sapphirewafers nitrided at TN = 530, 800, and 1100oC. Normalization was done withrespect to Al 2p peak

The normalized O 1s intensity is compared with the O 1s intensity from non-

nitrided sapphire surface as indicated by blue color line in Fig. 3.4b. It is found

that the non-nitrided sapphire surface has high O atom content when compared

to nitrided sapphire surfaces. Peak broadening and shift in peak positions have

been explored by deconvoluting the N1s peaks. Fig. 3.5 shows the deconvoluted

spectra of N 1s photoelectron peaks at different nitridation temperatures.

Systematic changes in both the intensity and peak positions are observed with the

change in nitridation temperature. A strong peak at ∼ 396.7 eV is present for TN

=800 and 1100oC (Fig. 3.5b & c). This coincides with that of the binding energy

of Al-N bond in the bulk AlN of 396.7 eV [58, 71] and indicates the structural

transformation of sapphire surface to AlN at TN ≥ 800oC. There are additional

peaks at ∼ 398.5 eV and ∼ 395.5 eV. The intensity of the former peak decreases

and the latter increases with the change in nitridation temperature from TN =

800 to 1100oC. The peaks at ∼ 395 eV correspond to sub-stoichiometric AlNx<1

with Al-Al bonds [72]. The peak at ∼ 398 eV has been attributed to incomplete


Figure 3.5: De-convoluted spectra of N 1s photoelectron peaks from sapphirewafers nitrided at (a) TN = 530oC, (b) TN = 800oC, and (c) TN = 1100oC.

TN AlN (Al-N bond) AlON (Spinel) AlOxN1−x Sub-stoichiometricAlN (Al-Al bond)

(oC) IN1s (a. u.) IN1s (a. u.) IN1s (a. u.) IN1s (a. u.)

530 0.30 0.13 1.81 -800 1.2 - 0.8 0.251100 1.34 - 0.35 0.55

Table 3.1: Normalized intensities of N 1s deconvoluted peaks from variouspossible nitrided layer structures from TN = 530, 800 & 1100oC. Normalizationhas been done with respect to Al 2p peak.

substitution of O by N such that tetrahedral A-N-O bonds are present [58, 73].

With decrease in nitridation temperature, peak positions shift to higher energy

levels. At TN = 530oC, the binding energy of 401.41 eV indicates that N is in O

atom rich environment [59, 74–77] and the peak at 405.31 eV in Fig. 3.5a arises

from AlON cubic spinel structure [76]. Table. 3.1 summarizes these results in the

form of normalized intensities of N 1s deconvoluted peaks from various possible

nitrided layer structures from TN = 530oC to 1100oC. Normalization has been

done with respect to Al 2p peak.


3.3.2.2 Annealed nitrided wafers

The XPS plots shown in Fig. 3.6 suggest that the as-nitrided layers undergo

further nitridation during ramp and anneal conditions. For TN = 530oC, the N 1s

peak just after ramp up to the annealing temperature shifts toward lower energy

levels indicating a transitional bonding environment with decreasing amount of

O atom content, and tetrahedral Al-N bond formation (Fig. 3.6a). The peak

corresponds to 405 eV (AlON cubic spinel) is no longer observed after ramp up

and anneal step.

It is found that after 4 min anneal step the dominant peak located at 397.4 and

is corresponds to the Al-N bond energy levels in O atom rich environment (Fig.

3.6c). In contrast, no significant changes are observed for the samples that have

been nitrided at TN = 1100oC as shown in Fig. 3.6b & 3.6d. The dominant peak

is always found to be at ∼ 396.7 eV that corresponds to tetrahedral Al-N bonds

in bulk AlN [58, 71]. The minor peak located at ∼ 398 eV in as-nitrided sample

is found to remain even after ramp up step (Fig. 3.6b), but has shifted to ∼ 397

eV after 4 min annealing step (Fig. 3.6d). Fig. 3.6e shows the compiled binding

energy scale for the N 1s photoelectron peak with their corresponding chemical

bond environments from the literature [58, 59, 71–77] .

3.3.3 LT GaN nucleation layer

3.3.3.1 Morphological evolution

Fig. 3.7 shows the surface morphology of LT GaN NL after the three different

stages: deposition, ramp up to LT GaN annealing temperature of 1000oC in 5

minutes (0 min annealing) and hold at 1000oC for 4 minutes (4 min annealing).

Fig. 3.8 shows the corresponding AFM morphologies. Fig. 3.9 shows the RMS

surface roughness values of the NLs derived from AFM data.

The most striking aspect of the LT GaN NLs evolution summarized in Fig. 3.7 to

3.9 is the significant difference in surface morphology and AFM roughness between


Figure 3.6: N 1s x-ray photoelectron peaks from nitrided sapphire wafers atdifferent stages of their process: (a) & (c) are for TN = 530oC, after ramp up and4 min annealing, and (b) & (d) are for TN = 1100oC, after ramp up and 4 minannealing at 1000oC. The compilation of N 1s binding energy scale with theircorresponding chemical bonding environments are shown in (e) [58, 59, 71–77]

NLs deposited on surfaces nitrided at TN = 530oC versus those deposited on

surfaces nitrided at TN ≥ 800oC. A levelling effect of the ramp up and anneal

at higher temperatures can be seen from Fig. 3.7 & 3.9. On sapphire nitrided

at TN = 530oC, Fig. 3.7a, the nuclei appears to be arranged along the atomic

steps on sapphire and evolve in this arrangement even during the ramping step.


Figure 3.7: SEM surface morphologies of LT GaN NLs: (a), (b) & (c) are forthe NL at different stages of their processing on sapphire wafers nitrided at TN

= 530oC, (d), (e) & (f) are for the NL processed on sapphire wafers nitride atTN = 800oC, (g), (h) & (i) are for the NL on sapphire wafers nitrided at TN =1100oC.


Figure 3.8: AFM surface morphologies of LT GaN NLs: (a), (b) & (c) are forthe NL at different stages of their processing on sapphire wafers nitrided at TN

= 530oC, (d), (e) & (f) are for the NL processed on sapphire wafers nitride atTN = 800oC, (g), (h) & (i) are for the NL on sapphire wafers nitrided at TN =1100oC.


Figure 3.9: AFM surface roughness evolution of LT GaN NLs deposited onsapphire wafers nitrided at TN = 530, 800 and 1100oC at different stages oftheir processing.

However, further exposure to NH3 at higher temperatures 1000oC during the hold

step annihilates these preferential orientations of the islands. In the case of the

samples nitrided at TN = 800 and 1100oC the ramp up significantly smoothens the

LT GaN surface. However annealing appears to lead to an element of preferential

sublimation leaving behind features that appear in light contrast.

3.3.3.2 Structural evolution

The effect of nitridation on the crystallography of the LT GaN NLs is revealed

through the HRXRD scans of Fig. 3.10 to 3.12. Fig. 3.10. Shows (0002) HRXRD

ω-scan profiles of LT GaN NLs at different stages of their processing. It is clearly

seen that the NL deposited on surfaces nitrided at TN ≥ 800oC show strong (0002)

peaks as compared to that deposited on the surface nitrided at TN = 530oC. The

LT GaN processed on sapphire nitrided at TN ≥ 800oC is found to be wurtzite

LT GaN with a strong [0002] orientation as indicated by six {101̄1} asymmetric

peaks shown in Fig. 3.11c and d. The sample is tilted to 61.96o about ψ-axis and

is rotated 360o about φ-axis to get diffraction from six {101̄1} planes. NLs at the

end of their 4 min annealing step have significantly reduced in thickness and it is


difficult to pick up an XRD signal from them.

In contrast, the NLs grown on sapphire nitrided at TN = 530oC does not show

this strong [0002] orientation in the as-grown conditions as indicted by the φ-

scan in Fig. 3.11a. This is due to the presence of zincblende phase. Fig. 3.12

shows the phi scan measurements for cubic LT GaN {220} peaks. Six {220} peaks

which are separated by 60o angle indicates the presence of zincblende phase in

the as grown LT GaN. The sample is tilted to 35.26o about the ψ-axis and the

detector is fixed at 2θ of 57.91o to get the diffraction from {220} planes. However

{101̄1} peaks appear after the 4 min annealing step at 1000oC (Fig. 3.11b),

indicating that decomposition and reformation of GaN during this process results

in transformation of LT GaN from zincblende phase to wurtzite phase with [0002]

oriented crystallography.

Figure 3.10: High resolution x-ray (0002) ω-scan profiles for NL at differentstages of their processing: (a) is for as grown LT GaN NL on sapphire surfacesnitrided at TN = 530, 800 and 1100oC, (b) is for NL after the 0 min annealingconditions and (c) is for NL after 4 min annealing.

The results indicate that HT nitridation of sapphire yields LT GaN NL with

strong [0002] orientation. In contrast LT nitridation yields LT GaN NL with the

zincblende in the as grown step and transformed to wurtzite phase with [0002]

orientation in the subsequent anneal stages.


Figure 3.11: High resolution x-ray {101̄1} diffraction φ-scan for NLs: (a), (b)are for NLs processed on sapphire wafers nitrided at TN = 530oC, (c), (d) arefor NLs processed on sapphire wafers nitrided at TN = 1100oC.

Figure 3.12: High resolution x-ray diffraction φ-scan measurements forzincblende {022} peak of as-grown LT GaN on sapphire nitrided at TN = 530oC.


3.4 Discussion

3.4.1 Low temperature nitridation (TN = 530oC)

The as nitrided surface at low nitridation temperatures shows a relatively high O

atom content with a strong N1s peak at ∼ 401 eV accompanied by less intense

peaks at ∼ 405 eV and ∼ 397 eV as shown in Fig. 3.5a. These peaks have been

identified with various levels of O in AlN [58, 59, 71–77] (Fig. 3.6e). It appears

plausible to state that the surface is dominated by Al2O3 based structural motifs

on nitridation at TN = 530oC accompanied by small amounts of AlON cubic spinel

associated with the peak at ∼ 405 eV [76].

The deposition of LT GaN on such a surface results in zincblende GaN with (111)

orientation and an incomplete coverage of the surface (Fig 3.7a). The possible

reason for incomplete coverage of LT GaN is due to incomplete nitridation of

sapphire at low temperatures. It has been shown earlier that GaN deposited

on unnitrided sapphire is rough with 3D islands and with incomplete coverage

[11]. This is due to the high energy interface between GaN and sapphire [11].

Therefore we believe that the high energy interface between GaN and sapphire

is not significantly affected by low temperature nitrided layer and we continue to

obtain 3D islands of GaN with incomplete coverage. The possible reason for the

directionality of LT GaN (Fig. 3.7a) is due to the nucleation of LT GaN along the

atomic steps of the sapphire surface (Fig. 3.2). However the ramp up and anneal

process provides a more complete coverage along with distinct texturing along the

[0002] axis of wurtzite GaN (Fig 3.7c & 3.11b). The ramp up and anneal steps

are accompanied by N1s peak shifts from the underlying nitrided surface (Fig. 3.5

and 3.6), the uncovered portions of which are exposed to H2 and NH3 ambient

at higher temperatures of 1000oC. As we described in Sec. 2.1.3 of Chapter 2,

during ramp up and anneal stages LT GaN starts decompose and redeposits on

sapphire surface, which is in the ambient of NH3 and H2. The decomposition

and redeposition of LT GaN involves the surface migration of decomposed Ga and

N adatoms, and incorporation of these adatoms into the newly formed growing


nuclei and/or on the existing nuclei, which eventually yields LT GaN with different

surface morphology when compared to the as grown stage (Fig. 3.7a, b & c).

It appears that the decomposition and redeposition of GaN occurs on further

nitrided region complexes which are now rich in N atom content, given that the

high intensity N1s photoelectron peak has a value of ∼ 397 eV which corresponds

to Al-N bond in O atom environment (Fig. 3.6). Therefore, it is plausible to

state that the increase in coverage of LT GaN during ramp up and anneal steps is

due to reduction in interface energy between GaN and sapphire via changes in the

transformed structure of underlying nitrided layer from AlOxN1−x to dominant

Al-N bond in O atom environment. It also appears that this process leads to the

formation of wurtzite LT GaN oriented along [0002] direction (Fig. 3.11b).

3.4.2 High temperature nitridation (TN = 1100oC)

At higher nitridation temperatures (TN = 800 and 1100oC) there is adequate ni-

trogen from the decomposed NH3 source to completely replace the O atoms in

sapphire surface [58], and stable Al-N bonds will be formed due to complete ex-

change of N and O atoms in sapphire surface [78]. This is indicated by the high

intensity XPS N1s peak at 396.7 eV, which corresponds to that of Al-N binding

energy in bulk AlN [58, 71]. There is no significant change in this structure of ni-

trided layer on annealing after ramp up as shown in Fig 3.6. As a consequence LT

GaN deposited on such surfaces after high temperature nitridation results in strong

[0002] oriented wurtzite LT GaN with complete surface coverage in as grown and

annealed stages (Fig. 3.7d & g and 3.11c & d). The complete coverage of LT GaN

is due to the transformed nitrided layer (Al-N bond environment), which results in

reduced interface energy between GaN and sapphire, and eventually promotes 2D

growth of LT GaN. It has been shown that nitrided layer acts as wetting layer and

it reduces the interface energy between GaN and sapphire [61]. Thus, we believe

that the complete coverage of LT GaN on sapphire nitrided at temperatures TN ≥

800oC (Fig. 3.7d & g) is due to the wetting nature of transformed nitrided layer.

The surface coverage of LT GaN is complete even after ramp up stage also, but


the morphology contains pits of various size. The possible reason for the pits of

various size is due to the polarity of LT GaN layer. It was reported that N-polar

GaN decomposes rapidly in the ambient of NH3 & H2 than Ga-polar GaN [67].

We suspect that the as grown LT GaN contains GaN of both polarities and in the

subsequent ramp up process the N-polar domains etch at faster rate than Ga-polar

domains, which eventually leads to pits in the LT GaN of various size (Fig. 3.7e &

h). The surface morphology at the end of the anneal step contains islands of small

in size with bright contrast (Fig. 3.7f & i), are could be remnants of Ga-polar

LT GaN, which have been left after etch of N-polar LT GaN. The strong [0002]

orientation of wurtzite LT GaN is also due to the transformed Al-N bond nitrided

layer which reduces the lattice mismatch between GaN and sapphire, and aids the

NL in forming with the [0002] orientation with respect to the underlying substrate

[58, 59, 71–73].

Therefore, we state that the structural evolution of nitrided layer plays an im-

portant role in determining the morphology, surface coverage and orientation of

subsequently grown LT GaN NL with respect to the underlying sapphire substrate.

3.5 Summary & conclusions

Chemical bond evolution of transformed nitrided layers on sapphire surfaces have

been investigated systematically at different stages of their processing. It is found

that HT nitridation (TN = 1100oC) yields modified sapphire surface with Al-N

bond rich environment. In contrast LT nitridation yields (TN = 530oC) sapphire

surface modified with the dominant Al2O3 based structural motifs in the as grown

stage and transformed to Al-N bond with O atom content. The effect of these

seed layers on the microstructural evolution of subsequently grown LT GaN NLs

have been studied systematically. It is observed that the morphology, surface

coverage and orientation of LT GaN NL strongly depends on the evolution of

the transformed nitrided layer on sapphire surface. LT GaN NLs processed on

HT nitrided sapphire are found to have strong [0002] orientation with complete

surface coverage. In contrast, LT GaN processed on LT nitrided sapphire showed


zincblende phase with incomplete coverage in the as grown stage and later it got

transformed to wurtzite phase with [0002] orientation and with enhanced surface

coverage after the subsequent ramp up and anneal steps.

Chapter 4

Polarity & Microstructural

Evolution of HT GaN

This chapter focuses on the following two points, the first one is to understand

the relation between the structure of the transformed nitrided layer on sapphire

and polarity of the subsequently grown two-step HT GaN epitaxial layers. The

second one is to understand the relation between various growth parameters such

as growth temperature, V/III ratio, LT GaN parameters and carrier gas etc., on

the surface morphology and crystalline quality of HT GaN layers.

4.1 Background

The influence of various growth parameters on microstructural evolution of HT

GaN layers can be easily understand by monitoring the in-situ optical reflectivity

traces as discussed in subsequent Sec. 2.2.7 of Chapter 2. Fig. 4.1 shows one such

reflectivity trace of GaN two-step epitaxy on sapphire. Different processing stages

are indicated in relation to the two-step process.

� Stage 1 corresponds to the pre-treatment of sapphire wafers. During this

period wafers are thermally annealed at 1100oC under the ambient of H2

57

Chapter 4. Polarity & Microstructural Evolution of HT GaN 58

followed by sapphire nitridation step. During this stage the reflectivity stays

constant at sapphire reflectivity (0.07595). The optical reflectivity trace is

normalized to bare sapphire reflectivity.

� Stage 2, starts with the deposition of LT GaN NL at 530oC. The rise in

reflectivity is due to change in refractive index from sapphire to GaN and

the reflectivity continues to increase due to interference effects of optical rays

from LT GaN and sapphire.

� In stage 3, the as deposited LT GaN layer is annealed thermally at 1000oC

under the ambient of H2 and NH3. This stage involves several steps such

as decomposition of LT GaN, surface migration of decomposed Ga and N

adatoms and redeposition of GaN nuclei. All these steps eventually lead to

roughening of LT GaN surface and reduction in the thickness of LT GaN

layer [62, 79]. The reason for the drop in reflectivity shown in Fig. 4.1 is due

to roughness as well as reduction in thickness of annealed LT GaN [62, 79].

� Stage 4, starts with the deposition of HT GaN at 1000oC, 400 mbar. During

this stage HT GaN starts growing in island mode (3D) and the reflectivity

drops to a minimum due to high surface roughness of the sample [69, 80].

The reflectivity continues to stay at a minimum until the coalescence of HT

GaN islands occur. The coalescence of islands lead to recovery in optical re-

flectivity and after this point the HT GaN starts grow in layer by layer mode

(2D) [69, 80]. The time duration between the point where the reflectivity

hits a minimum and where the reflectivity recovers is called roughening re-

covery period. This period is a direct measure of crystalline quality of HT

GaN and is sensitive to the growth parameters.

� In stage 5, HT GaN starts grow in layer by layer mode (2D mode). The

oscillations in the reflectivity trace is due to interference effects of the optical

rays which are reflected from GaN surface and GaN/sapphire interface. The

period of oscillation gives the growth rate of the film as discussed in Sec.

2.2.7 of Chapter 2.


Figure 4.1: In-situ optical reflectivity trace of two-step GaN epitaxy on sap-phire (0001) wafers.

In spite of all the success achieved with Ga-polar GaN and its alloys, the surface

morphology of N-polar GaN films grown under identical conditions, are rough

with hexagonal facets and are not suitable for device fabrication [18] (Fig. 4.2).

Intentionally miscut sapphire (0001) wafers (2o to 4o) along a- and m-directions

were used to obtain device quality N-polar GaN layers as shown in Fig. 4.2b

[23]. Fig. 4.3 shows the HRXRD rocking curve FWHM values of N-polar GaN

layers grown on miscut and non-miscut nitrided sapphire wafers in comparison to

conventional Ga-polar GaN layers [19].

Figure 4.2: Surface morphology of N-polar GaN layers grown on (a) non-miscut and (b) highly miscut nitrided sapphire wafers. Sumiya et al., JAP, 88,1158, 2000. Keller et al., JAP, 102, 083546, 2007. Reprinted with permission.


Figure 4.3: High resolution x-ray diffraction rocking curve FWHM values forN-polar GaN in comparison to conventional Ga-polar GaN. Sun et al., JCG,311, 2948, 2009. Reprinted with permission.

Although the growth mechanism of Ga-polar GaN has been clarified to a great

extent, this knowledge can not be applied to N-polar GaN directly, because the

surface structures of Ga-polar GaN and N-polar GaN are not the same. The sur-

face diffusion barriers for Ga and N adatoms are different on Ga-polar GaN and

N-polar GaN surfaces [17]. It is suggested that the rough surface morphology of

N-polar GaN is due to mixed polar domains and the hillock formation is due to

difference in growth rates of Ga-polar and N-polar domains [81]. The centre of the

hillock is found to be terminate with Ga-polar domain whereas the surrounding

matrix is termonated with N-polar GaN [81]. These inversion domains (IDs) are

found to originate from the regions which are rich in O atom content [82, 83].

They are also found to originate from GaN/sapphire film interface [82, 83] and

also from the steps on the sapphire surface [84]. However, by careful optimization

of growth parameters such as nitridation temperature Sun et al. obtained rela-

tively smoother N-polar GaN films [19]. Device quality N-polar GaN layers were

obtained on intentionally miscut (2o to 4o) sapphire (0001) wafers [19, 23]. Inten-

tionally miscut wafers (2o to 4o) enhances the mobility of adatoms significantly

and when the step density is high the probability of incorporation Ga adatoms at

the step edges is high and it results in a step flow growth mode. At lower step

densities (0.5o to 1o) the distance between adjacent steps is large and it results in


nucleation of islands on terraces between steps [23].

Therefore, it is plausible to state that the surface quality of N-polar GaN is con-

trolled by the following two factors

� Surface mobility of adatoms on the growth surface, which is controlled by

the miscut of sapphire wafers [23].

� Density of IDs, which is controlled by both the O atom content on the

nitrided surface and miscut of sapphire wafers [82–84].

However, to date there are no reports on systematic analysis of polarity selection

and microstructural evolution of LT GaN NLs and HT GaN layers in relation to

the structure of the transformed nitrided layers on sapphire wafers. This chap-

ter addresses the factors which control the surface morphology of N-polar GaN

through surface mobility of adatoms and density of IDs on non-miscut sapphire

(0001) wafers.

4.2 Experimental

The in situ thermal treatment of wafers were carried out at 1100oC under the flow

of purified H2 for 10 min before nitridation. The nitridation step was performed

at three different temperatures TN = 530, 800 and 1100oC and 200 mbar reactor

pressure. Wafers were nitrided under the ambient of NH3 and H2 for 1 min. NH3

flow rate was kept constant at 1500 sccm throughout the process.

After sapphire nitridation, LT GaN was grown to a 60 nm thick layer to under-

stand the effect of nitridation on microstructural evolution of NL. The thickness

calibration was done by the in situ optical tool as discussed in Sec. 2.2.7 of Chap-

ter 2. LT GaN NL was deposited on nitrided wafers at TN = 530oC for 8 min

(to deposit 60 nm thick layer), 200 mbar, at a V/III ratio of 2535, TMGa flow

rate of 0.6 sccm and H2 flow rate of 6500 sccm. The wafers were then heated up

from the LT GaN NL growth temperature to the LT GaN NL annealing condition

in 5 minutes. The NLs were annealed thermally for 4 minutes at 1000oC and a


Parameter A1 A2 A3 A4 A5

LT GaN Thickness, nm 60 60 60 60 to 605

LT GaN Anneal Time, min 4 4 4 4 4

LT GaN Anneal, NH3 flow, sccm 2000 3000 to 4000 4000 47005000

HT GaN Growth Temperature, oC 1000 1000 1000 to 1000 10001050

HT GaN Growth Pressure, mbar 400 400 400 400 400

HT GaN TMG flow, sccm 4.1 4.1 4.1 4.1 4.1

HT GaN NH3 flow, sccm 2000 3000 to 4000 4000 47005000

HT GaN V/III ratio 485 725 to 1205 965 965 1135

HT GaN Thickness, µm ∼ 1.5 ∼ 1.5 ∼ 1.5 ∼ 1.5

HT GaN Carrier gas H2 H2 H2 H2 H2

Table 4.1: HT GaN growth parameter space for sapphire nitridation at TN =530oC

reactor pressure of 400 mbar in a mixture of 2000 sccm of NH3 and 4700 sccm of H2.

Several growth parameters are varied to understand the polarity and micro-structural

evolution of HT GaN layers. These are tabulated in Table 4.1 & 4.2. Table 4.1

corresponds to growth parameters for sapphire nitridation at TN = 530oC and

Table 4.2 is for growth parameters at TN = 1100oC.


Parameter B1 B2 B3 B4 B5 B6 B7

LT GaN Thickness, nm 60 60 60 60 60 30 60

LT GaN Anneal Time, min 4 4 4 0 4 0 4

LT GaN Anneal 2000 3000 to 4700 4700 4700 1500 4700NH3 flow, sccm 5000

HT GaN Growth 1000 1000 1000 1000 1000 to 900 1000Temperature, oC 1050 800

HT GaN Growth 400 400 400 400 400 400 400Pressure, mbar

HT GaN TMG flow, sccm 4.1 4.1 4.1 4.1 4.1 2.0 4.1

HT GaN 2000 3000 to 4700 4700 4700 1500 4700NH3 flow, sccm 5000

HT GaN V/III ratio 485 725 to 1135 1135 1135 750 11351205

HT GaN Thickness, µm ∼ 1 ∼ 1 ∼ 1 ∼ 1 ∼ 1 ∼ 1

HT GaN Carrier gas H2 H2 H2 H2 H2 H2 N2

Table 4.2: HT GaN growth parameter space for sapphire nitridation at TN =1100oC

4.3 Results

4.3.1 Low temperature nitridation (TN = 530oC)

Fig. 4.4a shows the surface morphology of HT GaN layer deposited on 4 min

annealed LT GaN for sapphire wafer nitrided at TN = 530oC (Fig. 3.7c of Chapter

3), as per growth conditions given in column A1 of Table 4.1, that is at a V/III

ratio of 485. The morphology is quite rough and is due to incomplete coalescence

of HT GaN islands. The corresponding reflectivity trace is shown in Fig. 4.4b.

There is no recovery in reflectivity due to high surface roughness of the film. Red

color point on the trace indicates the starting point of HT GaN growth.


Figure 4.4: Surface morphology of HT GaN grown for sapphire nitrided atTN = 530oC, at a V/III ratio of 485 and its corresponding optical reflectivitytrace.

To enhance the island coalescence and to obtain device quality HT GaN epitaxial

layers for the nitridation temperature of TN = 530oC we have systematically varied

several growth parameters such as V/III ratio, growth temperature of HT GaN

and LT GaN thickness/growth time. The effect of each of these growth parameters

on surface and microstructural evolution of HT GaN will be described in the

subsequent sections.

4.3.1.1 V/III ratio

V/III ratio is the ratio of input flow rates of group-V (NH3) and group-III pre-

cursors (TMGa, TMAl & TMIn). V/III ratio is varied in this dissertation by

keeping the TMGa flow rate constantly at 4.1 sccm. NH3 flow rate is varied from

2000 to 5000 sccm. Fig. 4.5 shows the surface morphologies of HT GaN layers

grown at different V/III ratios with their corresponding reflectivity traces. The

corresponding growth conditions are given in column A2 of Table 4.1.

Fig. 4.6 shows the corresponding RMS surface roughness data of these samples.

The surface roughness is found to be increase to 0.75 nm (V/III = 1205) from 0.5

nm (V/III = 965).


Figure 4.5: Nomarski optical micrographs of HT GaN surfaces and their cor-responding in-situ optical reflectivity traces, grown at V/III ratios (a) 965, (b)1055 and (c) 1205 for a nitridation temperature of TN = 530oC. In all tracesthe red points indicate starting point of HT GaN growth.

The information that we extract from the reflectivity traces shown in Fig. 4.5 are

the corresponding growth rates and the roughening recovery times of HT GaN. It

is found that the growth rate of HT GaN decreases with the increase in V/III ratio.

The growth rate can be obtained from the period of the reflectivity oscillations as

described in the Sec. 2.2.7 of Chapter 2.

It is also found that the roughening recovery time decreases with the increase

in V/III ratio. Fig. 4.7 shows the variation in the growth rate and roughening

recovery time of HT GaN layers plotted against the V/III ratios.

The mosaicity of all these HT GaN samples were characterized by HRXRD. Fig.


Figure 4.6: AFM surface roughness data for the HT GaN samples grown atdifferent V/III ratios for sapphire nitridation at TN = 530oC.

Figure 4.7: Variation in the growth rate and roughening recovery time of HTGaN samples deposited at different V/III ratios for sapphire nitridation at TN

= 530oC.

4.8 shows the high resolution XRD rocking curve FWHM values for both symmet-

ric (0002) and asymmetric {101̄1} peaks of GaN layers. FWHM values are found

to increase with the V/III ratio.


Figure 4.8: High resolution XRD rocking curve measurements for HT GaNsamples grown at different V/III ratios for sapphire nitrided at TN = 530oC.

The results suggest that the surface quality and crystalline quality of HT GaN

layers depends critically on the V/III ratio.

4.3.1.2 Growth temperature

The other growth parameter which controls the surface morphology of HT GaN

is the growth temperature. We have varied HT GaN growth temperature from

1000oC to 1050oC using the growth parameters given in column A3 of Table 4.1.

Fig. 4.9 shows the surface morphology of HT GaN layers with their corresponding

reflectivity traces for sapphire nitridation of TN = 530oC. Fig. 4.10 shows the

corresponding surface roughness evolution of HT GaN layers. The roughness is

found to be decrease with the increase in growth temperature.

Fig. 4.11 shows the corresponding changes in growth rate and roughening recovery

time of the HT GaN layers. A slight increase in the growth rate of HT GaN is

observed. In contrast, it is found that the roughening recovery time decreases to

∼ 1 min from ∼ 25min and this is due to the faster coalescence of HT GaN islands

at higher growth temperatures.


Figure 4.9: Nomarski optical micrographs of HT GaN surfaces and their cor-responding in-situ optical reflectivity traces, grown at temperatures (a) 1000oC,(b) 1025oC and (c) 1050oC for sapphire nitridation at TN = 530oC. In all tracesthe red points indicate starting point of HT GaN growth.

Interestingly, we did not find any significant change in the crystalline quality

(mosaicity) of HT GaN samples grown at different growth temperatures. Fig.

4.12 shows the HRXRD rocking curve FWHM values for symmetric (0002) and

asymmetric {101̄1} peaks.


Figure 4.10: AFM surface roughness evolution of HT GaN layers deposited atgrowth temperatures: 1000, 1025 and 1050oCC for nitridation at TN = 530oC.

Figure 4.11: Variation in roughening recovery and growth rate of HT GaNsamples grown at different growth temperatures for sapphire nitridation at TN

= 530oC.

4.3.1.3 LT GaN NL thickness

In order to improve the crystalline quality of HT GaN layers further we have

systematically varied the thickness/growth time of LT GaN NLs to control the


Figure 4.12: High resolution XRD rocking curve FWHM values for symmetric(0002) and asymmetric {101̄1} peaks of HT GaN samples grown at differentgrowth temperatures for sapphire nitridation at TN = 530oC.

nucleation density (process conditions: column A4 of Table 4.1). Fig. 4.13 shows

the RMS roughness as a function of LT GaN thickness. The roughness is insensitive

to thickness except at very low thicknesses where it rises sharply. Fig. 4.14 shows

the roughening recovery rate which decreases with the increase in thickness. The

FWHM values are found to be lowest for 25 nm thickness of LT GaN NL (Fig.

4.15). The optimum thickness of LT GaN and the optimum process parameters

are may not be same for all the reactors and it may vary for different reactor

geometries.

4.3.1.4 Polarity of HT GaN

It is well known from the literature that the HT GaN samples grown on non-

nitrided sapphire consistently yield Ga-polar GaN. Potassium hydroxide (KOH)

etching and convergent electron beam diffraction (CBED) experiments were car-

ried out on these HT GaN samples for polarity measurement. It is known from the

literature that KOH etches the N-polar GaN whereas Ga-polar GaN is resistant to

KOH etch [85–88]. The HT GaN samples are etched in 0.2 mol of KOH solution


Figure 4.13: AFM surface roughness data of HT GaN as a function of LTGaN NL thickness for sapphire nitridation at TN = 530oC.

Figure 4.14: Roughening recovery time data of HT GaN as a function of LTGaN NL thickness for sapphire nitridation at TN = 530oC.

at 60oC for 10 min. Fig. 4.16 shows the surface morphology of HT GaN samples

before and after KOH etching experiment.

It can be seen that HT GaN samples are not significantly affected by the KOH

treatment as shown in Fig. 4.16a and 4.16b. The experimental CBED patterns

were taken along < 101̄0 > zone axis of GaN and the patterns were compared

with JEMS simulated patterns. The JEMS patterns were simulated for different


Figure 4.15: High resolution XRD rocking curve FWHM values for symmetric(0002) and asymmetric {101̄1} peaks of HT GaN samples as a function of LTGaN NL thickness for sapphire nitridation at TN = 530oC.

Figure 4.16: Surface morphologies (a) before and (b) after KOH etch experi-ments of HT GaN layers grown on annealed NLs for sapphire wafers nitrided atTN = 530oC. The corresponding convergent beam electron diffraction patternsalong with their cross-sectional images of GaN films are shown in (c).


thicknesses 110 and 145 nm from the GaN/sapphire interface. The 0002 disc

indicates a bright band at the center whereas 0002̄ disc indicates a dark band

at the center for Ga-polar GaN film. These etching results and CBED patterns

shown in Fig. 4.16c indicate that the two-step HT GaN layers grown on surfaces

nitrided at TN = 530oC are Ga-polar GaN [19, 85–89].

4.3.2 High temperature nitridation (TN = 1100oC

)

Fig. 4.17 shows the surface morphology of HT GaN layer deposited on 4 min

annealed LT GaN NL for sapphire wafers nitrided at TN = 1100oC (column B1

of Table 4.2 ). The morphology is quite rough and contains hexagonal faceted

islands. The corresponding reflectivity trace is shown in Fig. 4.17b. There is no

recovery in the reflectivity due to high surface roughness of the film.

Figure 4.17: Nomarski optical microscopy images of HT GaN layers for sap-phire wafer nitrided at TN = 1100oC. The corresponding reflectivity trace isalso shown in the figure. The red point on the trace indicates the starting pointof HT GaN growth.

Immediately after HT GaN growth reflectivity rises substantially and then drops

down to a minimum after two small oscillations. This indicates that HT GaN starts

with a 2D growth mode at the beginning of the growth and then in the subsequent

process the growth mode is transformed to a 3D mode. We have varied several


growth parameters to enhance the surface quality of HT GaN layers. These studies

are described below.

4.3.2.1 V/III ratio

Fig. 4.18 shows the Nomarski optical microscopy images of HT GaN layers grown

at different V/III ratios of 965, 1130 and 1205. The corresponding growth param-

eters are given in column B2 of Table 4.2

Figure 4.18: Nomarski optical microscopy images of HT GaN layers depositedat V/III ratios (a) 965, (b) 1130 and (c) 1205 for sapphire nitridation at TN =1100oC.

In all cases the surface morphology is rough with faceted hexagonal islands. The

corresponding reflectivity traces are not shown here because there was no recovery

in the reflectivity traces due to high surface roughness. It has been shown earlier

that higher V/III ratio can enhance the surface quality of HT GaN for sapphire ni-

tridation at TN = 530oC. In contrast, for HT GaN samples for sapphire nitridation

at TN = 1100oC, the surface morphology remains rough with faceted hexagonal

islands and with small changes in the density and size of the islands for the given

V/III ratio range.

Fig. 4.19 shows the FWHM values of x-ray rocking curves of grown HT GaN

samples at different V/III ratios for sapphire nitridation at TN = 1100oC. FWHM

values of (0002) and {101̄1} peaks are found to increase with the V/III ratio.


Figure 4.19: High resolution XRD rocking curve FWHM values for HT GaNlayer: (a) (0002) and (b) {101̄1} peaks for sapphire nitridation at TN = 1100oC.

4.3.2.2 Polarity of HT GaN

The surface morphology of HT GaN layers for high nitridation temperatures is

rough with many hexagonal hillocks as shown in Fig. 4.17 & 4.18. Prior to

further optimization of HT GaN layers, potassium hydroxide (KOH) etching and

convergent electron beam diffraction (CBED) experiments were carried out on

these HT GaN samples for polarity measurement. HT GaN samples are etched

in 0.2 mol of KOH solution at 60oC for 10 min. Fig. 4.20 shows the surface

morphology of HT GaN samples before and after KOH etching experiment. The

growth conditions for the samples are given in column B3 of Table 4.2

The etching and CBED studies indicate that the samples grown on annealed NL

for sapphire wafers nitrided at TN = 1100oC are N-polar GaN. As can be seen

from Fig. 4.20a and 4.20b, they are substantially etched away when subjected

to the same KOH etch treatment [85–88]. CBED patterns shown in Fig. 4.20c

confirm that they are N-polar [19, 89]. The CBED patterns were taken along

< 101̄0 > zone axis of GaN and the patterns were compared with the JEMS

simulated patterns. The JEMS patterns were simulated for different thicknesses

100, 105 and 115 nm from the GaN/sapphire interface. In this case both 0002

disc and 0002̄ discs are interchanged in contrast due to N-polarity of GaN film as


Figure 4.20: Surface morphologies (a) before and (b) after KOH etch experi-ments of HT GaN layers grown on annealed NLs for sapphire wafers nitrided atTN = 1100oC. The corresponding convergent beam electron diffraction patternsalong with their cross-sectional images for GaN films are shown in (c).

compared to the TN = 530oC case (Fig. 4.16). The hillock remnants which are

left after KOH etch are Ga-polar since they do not etch.

4.3.2.3 LT GaN annealing time

The rough surface morphology with faceted hexagonal islands seen in the previous

section may originate from mixed polarities of HT GaN. It has been suggested in

the literature that the faceted hexagonal hillocks originate from inversion domains

in the structure that is from Ga polar domains that grow faster than the sur-

rounding N-polar domains [81]. It has also been suggested that Ga-polar domains

form on O atom rich surfaces [82, 83]. It was thought that the annealing period

at the high temperature nitridation may enhance the O atom level on the nitrided

surface as a consequence of diffusion from the sapphire substrate into the nitrided

layer. To enhance the surface quality of HT GaN layers we have grown these layers

on samples that were ramped up to the growth temperature but not annealed.


Fig. 4.21 shows the Nomarski optical microscopy images of HT GaN layers grown

under identical conditions except the LT GaN annealing time (column B4 of Table

4.2). The sample shown in Fig. 4.21a is deposited on 4 min annealed LT GaN as

described in previous sections. In contrast, HT GaN grown on samples directly af-

ter ramp up without annealing results in a substantial reduction of the hexagonal

hillock density (Fig. 4.21b).

Figure 4.21: Surface morphologies of HT GaN layers deposited on (a) 4 minannealed LT GaN layer and (b) as ramped up LT GaN layers for sapphirenitridation at TN = 1100oC.

To confirm the polarity in the HT GaN that was deposited in the as-ramped up

condition, the above samples are subjected to same KOH etch treatment. Fig. 4.22

shows the side view of SEM surface morphologies of HT GaN epitaxial layers shown

in Fig. 4.21a and 4.21b, after KOH etch experiment. Both process conditions lead

to N-polar GaN. However, the density of the remnant hillocks after the KOH

etching were significantly less for the HT GaN grown directly after ramp up (Fig.

4.22b), that is the smoother HT GaN was associated with a lower density of

remnant hillocks (Fig. 4.21b). The results suggest that device quality N-polar

GaN layers can be obtained on non-miscut sapphire (0001) wafers.

Table 4.3 shows the HRXRD rocking curve FWHM values for the sample grown

directly on LT GaN after ramp up, in relation to the samples grown on 4 min

annealed LT GaN for sapphire nitridation at TN = 1100oC.


Figure 4.22: Surface morphologies of HT GaN layers after KOH treatmentdeposited on (a) 4 min annealed LT GaN layer and (b) 0 min annealed (rampedup) LT GaN layers for sapphire nitrided at TN = 1100oC.

Layer LT GaN LT GaN (0002) {101̄1}Thickness Anneal Time FWHM FWHM

(nm) (min) (arc-sec) (arc-sec)

HT GaN 60 4 702 2523HT GaN 60 0 540 1130

Table 4.3: FWHM values of x-ray rocking curves for the HT GaN samplesgrown directly on LT GaN after ramp up, in relation to the samples grown on4 min annealed LT GaN for sapphire nitridation at TN = 1100oC

4.3.2.4 Growth temperature

The other process parameter which we have varied to control the surface mor-

phology of HT GaN is the growth temperature. We have varied HT GaN growth

temperature as described in the previous section from 1000oC to 1050oC. The cor-

responding growth parameters are listed in column B5 of Table 4.2.

It is found that the size and density of hexagonal islands are very sensitive to

growth temperature (Fig. 4.23). The density of islands is found to increase while

the size of the hexagonal islands is found to decrease with the increase in growth

temperature.

Fig. 4.24 shows the corresponding HRXRD rocking curve FWHM values. The

FWHM values of (0002) and {101̄1} are found to be increase with the growth


Figure 4.23: Nomarski optical microscopy images of HT GaN layers grownat (a) 1000oC, (b) 1025oC and (c) 1050oC growth temperatures for sapphirenitrided at TN = 1100oC.

temperature.

Figure 4.24: HRXRD rocking curve FWHM values of HT GaN layers grownat temperatures: (a) 1000oC, (b) 1025oC and (c) 1050oC growth temperaturesfor sapphire nitridation at TN = 1100oC.

Since density and size of hexagonal islands critically depends upon the growth

temperature and is found to increase with the growth temperature. In order to

understand further the surface evolution of N-polar GaN layers, we have also grown

HT GaN samples at lower growth temperatures < 1000oC. Fig. 4.25 shows the

Nomarski optical images of HT GaN layers grown at two different temperatures

of 900 and 800oC as per the growth conditions given in column B6 of Table 4.2.

Interestingly, the reflectivity oscillations were observed for the sample grown at


Figure 4.25: Optical surface images of HT GaN layers grown at differenttemperatures (a) 900oC, (b) 800oC (LT GaN annealed at 800oC) and (c) 800oC(LT GaN annealed at 1000oC). The corresponding in-situ optical reflectivitytraces are also shown.

temperature 800oC (Fig. 4.25b). The sample surface looks mirror like. In con-

trast, in the sample grown at 900oC, decay in the reflectivity oscillations has been

observed (Fig. 4.25a) and the corresponding surface morphology of the sample

looks relatively rougher than the sample grown at 800oC. The results indicate

that lower growth temperatures < 1000oC seems to be favorable for obtaining N-

polar GaN with good surface quality. These samples are subjected to the same


Layer LT GaN LT GaN LT GaN HT GaN SurfaceThickness Anneal Anneal Growth Roughness

Time Temperature Temperature(nm) (min) (oC) (oC) (nm)

HT GaN 60 4 1000 1000 RoughHT GaN 60 0 1000 1000 ∼ 3.5HT GaN 30 0 800 800 ∼ 1.5HT GaN 30 0 1000 800 ∼ 1.2HT GaN 30 0 900 900 ∼ 7.5

Table 4.4: The surface roughness values of N-polar HT GaN grown at low &high growth temperatures for sapphire nitridation at TN = 1100oC

Layer LT GaN LT GaN HT GaN (0002) {101̄1}Thickness Anneal Growth FWHM FWHM

Temperature Temperature(nm) (oC) (oC) (arc-sec) (arc-sec)

HT GaN 60 1000 1000 702 2523HT GaN 60 1000 1000 820 2182HT GaN 30 800 800 1080 2160HT GaN 30 1000 800 540 1800

Table 4.5: HRXRD rocking curve FWHM values of N-polar HT GaN grownat low & high growth temperatures for sapphire nitridation at TN = 1100oC

KOH treatment described in the earlier sections and are confirmed to be N-polar.

The surface roughness values are indicated in Table 4.4 corresponding HRXRD

rocking curve FWHM values are tabulated in Table 4.5.

To enhance the crystalline quality of the sample, we have annealed the LT GaN

NL at 1000oC and reduced the temperature to 800oC for the subsequent HT GaN

growth (column B6 of Table 4.2). The FWHM values are reduced drastically

(Table 4.4) and the surface morphology of the corresponding sample along with

its reflectivity trace is shown in Fig. 4.25c. The corresponding AFM morphologies

of samples (Fig. 4.25b & c) are shown in Fig. 4.26. The RMS surface roughness

of the samples is found to lie in between 1 to 1.5 nm. The results suggest that

device quality N-polar HT GaN layers can be obtained on non-miscut sapphire

(0001) wafers by careful optimization of growth temperature.


Figure 4.26: AFM morphologies of HT GaN layers grown at temperature800oC for high temperature nitridation: (a) LT GaN is annealed at 800oC (HTGaN RMS roughness ∼ 1.5 nm) and (b) LT GaN is annealed at 1000oC (HTGaN RMS roughness ∼ 1.2 nm).

4.3.2.5 Carrier gas (H2/N2)

It has been reported well in the literature that the carrier gas N2 enhances the

lateral growth of N-polar GaN. In contrast, the carrier gas H2 suppresses the lateral

growth of N-polar GaN [20–22]. In order to investigate the effect of carrier gas on

surface morphology of N-polar GaN. We have deposited N-polar HT GaN layers

under identical conditions while altering the carrier gas from H2 to N2 (column

B7 of Table 4.2).

Fig. 4.27 shows the surface morphology of HT GaN layers. The density and size

of the hexagonal islands are drastically reduced for the sample grown under N2 as

carrier gas (Fig. 4.27b).

4.3.3 Summary of results

This section describes the summary of key results of HT GaN layers obtained at

nitridation temperatures TN = 530 & 1100oC.


Figure 4.27: Nomarski optical microscopy images of HT GaN layers grownunder (a) H2 as carrier gas, and (b) N2 as carrier gas for sapphire nitrided atTN = 1100oC.

4.3.3.1 Low temperature nitridation (TN = 530oC)

� Low temperature nitridation (TN = 530oC) yields Ga-polar GaN.

� The surface quality of Ga-polar GaN critically depends upon the parameters

such as V/III ratio and growth temperature. The quality is found to increase

with the growth temperature and decrease at higher V/III ratios due to

statistical roughening.

� The crystalline quality is found to decrease with the increase in V/III ra-

tio and LT GaN thickness, and is not significantly affected by the growth

temperature.

4.3.3.2 High temperature nitridation (TN = 1100oC)

� High temperature nitridation (TN = 1100oC) yields N-polar GaN.

� The surface quality of N-polar GaN is more sensitive to the HT GaN growth

temperature. The quality is found to decrease with the increase in growth

temperature, and is not significantly affected by the parameter V/III ratio.

� Low growth temperatures (800oC) are found to be favorable for obtaining

smooth N-polar GaN materials.


� The crystalline quality is found to decrease with the increase in V/III ratio

and growth temperature.

4.4 Discussion

The experiments described in the previous section explore a wide range of process

parameters in an attempt to provide optimum combinations of polarity, surface

roughness and crystalline quality of HT GaN layers. This section provides a ratio-

nale for our observations first for HT GaN growth with low temperature nitridation

(TN = 530oC) and then for HT GaN growth with high temperature nitridation

(TN = 1000oC).

The experience on optimization of growth parameters for Ga-polar GaN in our

MOCVD reactor has been taken as a bench mark for understanding the growth

mechanism of N-polar GaN layers.

4.4.1 Low temperature nitridation (TN = 530oC) and HT

GaN

4.4.1.1 Polarity

Our results suggest a central role for the O atom content in the nitrided layer, and

the related surface structure of the nitrided sapphire surface, in determining the

polarity of the HT GaN epitaxial layers. It has been shown that HT GaN grown

on non-nitrided sapphire inevitably has Ga polarity [56] and we continue to obtain

Ga-polar GaN even with lower nitridation temperatures TN = 530oC.

As shown in Table 3.1 and Fig. 3.4 of Chapter 3, the as-nitrided surface at low

nitridation temperatures shows a relatively high O content. Ab-initio calculations

of surface stability indicate that the existence of various Al2O3 based structural

motifs associated with different degrees of replacement of O by N at these tem-

peratures [78]. The as-grown LT GaN has a cubic structure (Fig. 3.11 & 3.12


of Chapter 3). On annealing the LT GaN prior to HT GaN deposition the un-

derlying nitrided layer appears to change in character to a AlN based structural

motifs, albeit with still high O levels. LT GaN during the anneal decomposes and

redeposits, presumably on the altered nitrided layers (Fig. 3.6 of Chapter 3) with

more uniform coverage (Fig. 3.7 of Chapter 3) and with a hexagaonal structure

(Fig. 3.11 of Chapter 3). It is suggested that this redeposited hexagonal LT GaN

is Ga-polar since the nitrided layer has still a high O atom content shown in Table

3.1 of Chapter 3, in similarity with the Ga-polar HT GaN grown on non-nitrided

sapphire. As a consequence the subsequent HT GaN epitaxial layer is Ga-polar.

Thus our results suggest Ga-polar GaN is continues to form on surfaces that have

been nitrided as long as the nitridation process (at low temperatures) results in O

containing AlN structural motifs.

4.4.1.2 Crystalline quality of HT GaN

The crystalline quality of Ga-polar HT GaN has been investigated through a

characterization of the HRXRD rocking curve FWHM values of (0002) and {101̄1}

peaks over a wide range of parametric conditions that are summarized in Fig 4.28.

Both these peaks show similar trends. Fig. 4.28 therefore provides values of just

the (0002) peaks.

Figure 4.28: Summary of (0002) x-ray rocking curve FWHM values of HTGaN layers over a wide range of growth conditions for sapphire nitrided at TN

= 530oC.


The values increase with V/III ratio, go through a minimum with NL thickness

and are not sensitive to the growth temperature. A comparison with roughening

recovery time in Fig. 4.7 as function of V/III ratio shows that the FWHM values

decrease with the increase in roughening recovery time. Since the roughening

recovery time is an indicator of LT GaN nucleus density, this suggests that FWHM

values improve with lower nucleus density as might be expected if coalescence of

island during the recovery stage control the defect density in HT GaN [90–93].

In order to understand the effect of V/III ratio on nucleation density, we have

examined the LT GaN layer formed with different V/III ratios before the growth

of HT GaN was initiated. The results shown in Fig. 4.29 demonstrate quite

convincingly that LT GaN nucleation density decreases with the decrease in V/III

ratio. In analogy, it may be expected that a decrease in the LT GaN thickness

(other parameters remaining constant) would imply a lower density of nucleation

sites (and increase roughening recovery time as observed in Fig. 4.7 & 4.14) and

a consequent decrease the defect density and thus the FWHM values, and this

is what is observed in Fig. 4.28. However at very low nucleation density, the

FWHM values rise steeply again. It is suggested this occurs because of poor

or substantially incomplete coverage of the substrate by LT GaN below a certain

critical thickness so that LT GaN no longer provides effective nucleation sites (Fig.

4.15). The FWHM values are insensitive to growth temperature for a given LT

GaN thickness and V/III ratio. This is expected since the reduced recovery times

and increasing growth rates in this case (Fig. 4.12) arise from diffusional effects

due to higher growth temperatures and do not reflect changes in the underlying

LT GaN nuclei size.

4.4.1.3 Surface roughness of HT GaN

The surface roughness of HT GaN increases with V/III ratio, decreases with

growth temperature and is insensitive to NL layer thickness of 25 nm below which

it increases sharply, as summarized in Fig. 4.30.


Figure 4.29: Surface morphologies of 4 min annealed LT GaN NLs at twodifferent NH3 flow rates: (a) 2000 sccm and (b) 5000 sccm

Figure 4.30: Summary of surface roughness values of HT GaN layers over awide range of growth conditions for sapphire nitrided at TN = 530oC.

The data suggests that surface roughness originates from a statistical roughening

process that occurs when the diffusion length is shorter than the mean distance

between the binding sites [50]. The increase in surface roughness with increasing

in V/III ratio (Fig. 4.6) results from a decreased mobility of Ga atoms on a surface

that is increasingly dominated by excess N atoms [17] as discussed in Sec. 2.1.3

of Chapter 2. The increase in growth temperature on the other hand increases

mobility in general and therefore reduces statistical roughening (Fig. 4.10). The

roughening is expected to be insensitive to LT GaN nucleation layer thickness,

as is observed (Fig. 4.13), at constant V/III ratio and growth temperature. The

sharp increase in roughness below a critical thickness of LT GaN arises from an


entirely different effect that is the LT GaN now provides incomplete coverage of

the nitrided sapphire and is no longer effective in providing nucleation sites.

4.4.2 High temperature nitridation (TN = 1100oC) and HT

GaN

4.4.2.1 Polarity

At high nitridation temperature the nitride layer is dominated by AlN based struc-

tural motifs with significant lower oxygen content as compared to the nitrided layer

for low temperature nitridation (Table 3.1 and Fig. 3.4 of Chapter 3). There is

no significant change in the nitrided layer structure on annealing after ramp up

as shown in Fig 3.6 of Chapter 3. As has been shown earlier [19, 56, 57] HT GaN

deposited under these conditions is dominantly N-polar. The origin of N-polarity

under these conditions is still not clear. However the surface quality of the N-polar

HT GaN is poor as shown in Fig. 4.17 & 4.18. We have therefore attempted to

explore process parameters that would improve surface quality as described in the

next section.

4.4.2.2 Surface quality of HT GaN

When HT GaN is deposited after the annealing step, a high density of hexagonal

hillocks are observed (Fig. 4.21a). These can be largely eliminated if HT GaN

is deposited directly after ramp-up to the HT GaN deposition temperature (Fig.

4.21b). The hillock surface morphology of N-polar GaN shown (Fig. 4.21a & b)

has been attributed to the presence of Ga-polar domains forming from regions

rich in O atom content [81–83]. Our results show that the density of these Ga-

polar domains is significantly lower when HT GaN is grown directly after ramp-up

instead of after the annealing step (Fig. 4.22 a & b). We believe that these

Ga-polar domains that define the hillock density form from O-rich Al-N bond

complexes in AlN based structural motifs associated with the minority N1s ∼ 398


and 397 peaks present after ramp up and anneal as shown in Fig. 3.5 & 3.6 of

Chapter 3. It is likely that the sublimation and redeposition of LT GaN during

ramp up and anneal gives rise to a mixed polarity of LT GaN in which the density

of Ga-rich domains depends on the O content of the underlying nitrided layer. The

polaity of the HT GaN will then depend on the polarity of the LT GaN nuclei. The

correlation between the effect of relative O atom content in the nitrided layers and

the surface morphology of N-polar GaN is supported by examining the normalized

XPS intensity of O 1s peak from O-N and Al bond environments at different stages

of nitrided layer processing (Fig. 4.31). Fig. 4.22 supports this hypothesis in that

a lower density of Ga-polar hillocks is left behind after KOH etching under identical

etching conditions in the HT GaN deposited directly after ramp up. It is known

that O atoms diffuse out from sapphire to GaN at typical growth temperatures.

The higher O content in the nitrided layer after the annealing step must be due

to increased diffusion of O atoms from the sapphire, which eventually leads to

formation of more Ga-polar domains in the growing film.

Figure 4.31: XPS normalized O 1s intensity from O-N & Al chemical bondenvironments from sapphire nitride at TN = 1100oC, at different stages.

This result suggested that lower growth temperature of GaN may result in de-

creased oxygen diffusion from sapphire substrate into the nitrided layer and vice

versa. Therefore we have explored the effect of growth temperature on surface


morphology of N-polar GaN and these results shown in Fig. 4.23 and Fig. 4.25

demonstrate that decreasing HT GaN growth temperatures improve surface qual-

ity. Indeed the morphology of the film grown at 800oC is mirror like. The cor-

responding reflectivity traces shown in Fig. 4.25 indicates that 3D to 2D growth

mode transition is no longer observed for N-polar GaN films. The reflectivity rises

immediately after HT GaN growth and the oscillations were continued till the end

of the growth. This is in contrast to Ga-polar GaN films where 3D to 2D transition

is commonly observed [89].

For conventional MOCVD GaN growth, H2 is used as carrier gas. It was found

by several groups that H2 can passivate the N-terminated surface during growth

[20–22], which explains the lower growth rate on the N-polar surface compared

to Ga-polar surface. Fig. 4.27 shows that improvement in surface quality can be

achieved if N2 is used as carrier gas instead of H2 gas.

The results on surface quality in the process parameter space explored in this

thesis are summarized in Fig. 4.32.

Figure 4.32: Summary of surface roughness values of HT GaN layers over awide range of growth conditions for sapphire nitrided at TN = 1100oC.


4.4.2.3 Crystalline quality of HT GaN

The data for crystalline quality is summarized in Fig. 4.33 in the processing

parameter space explored in this thesis.

Figure 4.33: Summary of (0002) x-ray rocking curve FWHM values of HTGaN layers over a wide range of growth conditions for sapphire nitrided at TN

= 1100oC.

For dominantly N-polar GaN containing inversion domains, the crystalline quality

FWHM values would be influenced both by inversion domain density as well as

dislocation density arising out coalescence effects. For given growth temperature

and NL thickness the HRXRD rocking curve FWHM for (0002) and {101̄1} peaks

increases with V/III ratio and for a given V/III ratio and NL thickness, the FWHM

values increase with growth temperature. The increase in FWHM with V/III ratio

can be attributed to the higher hillock density and therefore a higher inversion

domain density. The effect of inversion domain density on FWHM of samples is

also evident in HT GaN grown after an anneal or directly after ramp up. However

since all these conditions have rough surface morphology, these conditions are

not of much interest. Of particular interest the effect of lower HT GaN growth


temperatures at 800 and 900oC. Decreasing the growth temperature substantially

improves the surface quality but increases the FWHM of rocking curves. Assuming

that the FWHM values are affected by LT GaN nucleation density and therefore

the coalescence density as explained in the preceding section, HT GaN was grown

at 800oC after annealing the LT GaN at 1000oC to reduce the LT GaN nucleation

density. A substantial decrease in FWHM is observed as a consequence. Detailed

analysis on the effect of inversion domain density and other defects in N-polar

LT GaN is required to understand fully the effects of growth parameters on the

crystalline quality of HT N-polar LT GaN.

4.5 Summary & conclusions

� A large range of growth parameter space that includes LT GaN thickness,

V/III ratio and growth temperature of HT GaN growth have been explored

for varying nitridation temperatures.

� Low temperature nitridation upto TN = 530oC results in Ga-polar GaN while

nitridation above TN = 530oC results in N-polar GaN.

� Growth parameter conditions have been optimized for both surface quality

and crystalline quality of Ga-polar GaN and the rationale for processing

effects developed.

� Growth parameter conditions have also been arrived at, which allow the re-

alization of device quality N-polar GaN without the necessity of deliberately

miscut sapphire (0001) wafers.

Chapter 5

Summary, Conclusions and

Future Work

We have explored a large domain of growth parameter space including nitrida-

tion temperature, LT GaN nucleation layer control, HT GaN growth with varying

V/III ratios and growth temperature to understand the dependence on these pa-

rameters of polarity and surface and crystalline quality of GaN grown on basal

plane sapphire.

� It has been shown that a modified sapphire (nitrided sapphire) surface plays a

key role in determining the polarity of subsequently grown GaN layers. The

modified surface is characterized by its chemical nature and is controlled

by nitridation temperature (TN). The presence of Al-N bond environment

formed at HT nitridation (TN ≥ 800oC) yields strong c-oriented wurtzite

LT GaN films. In contrast, presence of AlOxN1−x (Al2O3 based structural

motifs) formed at (TN = 530oC) yields (111) oriented cubic LT GaN films in

the as grown stage. The orientation of LT GaN in the subsequent processes is

controlled by the structural transformation of the underlying nitrided layer.

� The polarity of two-step based HT GaN layers grown on such a modified

sapphire surfaces is not the same. N-polar grows preferentially on modified

93

Chapter 5. Summary, Conclusions and Future Work 94

sapphire surface with an Al-N chemical bond environment with AlN struc-

tural motifs, while Ga-polar GaN grows on AlOxN1−x sapphire surfaces. The

former is observed at nitridation temperatures down to TN = 800oC while

the latter is found at nitridation temperatures to TN = 530oC.

� The effect of various growth parameters on HT GaN eipilayers discussed in

this dissertation shows that the surface quality, and mosaicity of HT GaN

layers critically depends upon the process parameters such as growth tem-

perature, V/III ratio, LT GaN growth and annealing time. A domain of

processing was identified for high surface quality Ga-polar GaN with accept-

able crystalline quality, and a rationale for the observed effects has been

presented.

� It was shown that the surface quality of N-polar GaN epitaxial layers grown

on non-miscut sapphire (0001) wafers, strongly depends upon the growth

temperature. The relative amount of O atom content presented in the pre-

cursor nitrided layers is a critical factor to obtain better surface quality of N-

polar GaN base layers and device quality N-polar GaN layers were obtained

for the first time on non-miscut sapphire wafers at low growth temperatures

(TN = 800oC) layers.

Several directions for additional work emerge from the thesis:

� Many observations of the thesis and the proposed rationale have depended

on analysis by techniques such as XPS and HRXRD coupled with SEM and

AFM characterization. Direct evidence of posed structures should be ob-

tained by TEM using recent aberration corrected techniques to elucidate

the details of structure and chemistry of the nanoscale interfaces and pre-

cursor layers that are vital to the quality of the GaN epilayers.

� Similarly details of defect structures in Ga-polar and N-polar epilayers grown

in this work require characterization by TEM to relate these defect structures

to the observed FWHM data obtained in this work.

Chapter 5. Summary, Conclusions and Future Work 95

� A better understanding of the mechanisms underlying the effect of growth

temperature on the quality of N-polar GaN would be desirable.

� The precise origin of polarity remains unknown. The results of this thesis

suggest that the O atom content on AlN based structural motifs of the ni-

trided layer determine the polarity. However the termination of the underly-

ing sapphire at different temperatures and its stable surface structure at any

temperature may play a role. The results of this thesis coupled with TEM

will provide a framework in which ab-initio calculations of stable structure

could be performed to understand the role of these surfaces and interfaces

in determining polarity.

� It has been determined that nitridation temperatures upto TN = 530oC

results in Ga-polarity of HT GaN epilayers while nitridation temperatures

above TN = 800oC result in N-polar GaN epilayers. A more precise definition

in the temperature domains would be useful in terms of device processing.

� As discussed in Sec. 1.3.2 of Chapter 1, there are certain advantages with

the N-polar based HEMT devices when compared to the conventional Ga-

polar GaN based HEMTs. The results of this thesis suggest that it should

be possible to optimize the Al0.25Ga0.75N layer growth on the N-polar GaN

base layers for HEMT device applications.

Appendix A

Specifications of MOCVD

Reactor

System: Aixtron make research scale single wafer Metal Organic Chemical Vapor

deposition (MOCVD).

Model:AIX 200/4 RF.

Type: Double walled water cooled cold wall reactor.

Hot zone: Detachable quartz liner tube with SiC and TaC thermal insulations.

Gas flow mechanism: Horizontal laminar gas flow with group III and group

V species carried in to the growth chamber through separate gas line.

Substrate compatibility: 1 cm2, 2 inch, 3 inch wafers.

Substrate rotation mechanism: Gas foil rotation by lift off mechanism.

Substrate heating mechanism: Radiative RF heating.

97

Appendix A. Appendix A 98

Working temperature range: From above room temperature to 1500oC.

Working pressure range: 10 to 1000 mbar.

Wafer handling: Through dry nitrogen conditioned glove box free from O2 and

H2O below 10 ppm to avoid oxygen incorporation during growth.

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