GALLIUM NITRIDE EPITAXY BY A NOVEL HYBRID VPE TECHNIQUE

A DISSERTATION

SUBMITTED TO THE DEPARTMENT OF MATERIALS SCIENCE AND

ENGINEERING AND THE COMMITTEE ON GRADUATE STUDIES OF

STANFORD UNIVERSITY IN PARTIAL FULFILLMENT OF THE

REQUIREMENTS FOR THE DEGREE OF

DOCTOR OF PHILOSOPHY

David J. Miller

June 2011

http://creativecommons.org/licenses/by-nc/3.0/us/

This dissertation is online at: http://purl.stanford.edu/hz462yv9251

© 2011 by David J. Miller. All Rights Reserved.

Re-distributed by Stanford University under license with the author.

This work is licensed under a Creative Commons Attribution-Noncommercial 3.0 United States License.

ii

I certify that I have read this dissertation and that, in my opinion, it is fully adequatein scope and quality as a dissertation for the degree of Doctor of Philosophy.

James Harris, Primary Adviser

I certify that I have read this dissertation and that, in my opinion, it is fully adequatein scope and quality as a dissertation for the degree of Doctor of Philosophy.

Michael McGehee

I certify that I have read this dissertation and that, in my opinion, it is fully adequatein scope and quality as a dissertation for the degree of Doctor of Philosophy.

GLENN SOLOMON

Approved for the Stanford University Committee on Graduate Studies.

Patricia J. Gumport, Vice Provost Graduate Education

This signature page was generated electronically upon submission of this dissertation in electronic format. An original signed hard copy of the signature page is on file inUniversity Archives.

iii

iv

GALLIUM NITRIDE EPITAXY BY A NOVEL HYBRID VPE TECHNIQUE

David J. Miller, Ph.D.

Stanford University, 2011

Advisors: Glenn S. Solomon & James S. Harris

Gallium nitride is an important material for the production of next-generation

visible and near-UV optical devices, as well as for high temperature electronic

amplifiers and circuits; however there has been no bulk method for the production of

GaN substrates for device layer growth. Instead, thick GaN layers are

heteroepitaxially deposited onto non-native substrates (usually sapphire) by one of two

vapor phase epitaxy (VPE) techniques: MOVPE (metalorganic VPE) or HVPE

(hydride VPE). Each method has its strengths and weaknesses: MOVPE has precise

growth rate and layer thickness control but it is slow and expensive; HVPE is a low-

cost method for high rate deposition of thick GaN, but it lacks the precise control and

heterojunction layer growth required for device structures. Because of the large (14%)

lattice mismatch, GaN grown on sapphire requires the prior deposition of a low

temperature MOVPE nucleation layer using a second growth process in a separate

deposition system. Here we present a novel hybrid VPE system incorporating

elements of both techniques, allowing MOVPE and HVPE in a single growth run. In

this way, a thick GaN layer can be produced directly on sapphire. GaN growth

commences as small (50-100 nm diameter) coherent strained 3-dimensional islands

which coalesce into a continuous film, after which 2-dimensional layer growth

commences. The coalescence of islands imparts significant stress into the growing

film, which increases with the film thickness until catastrophic breakage occurs, in-

situ. Additionally, the mismatch in thermal expansion rates induces compressive

stress upon cooling from the growth temperature of 1025ºC. We demonstrate a

growth technique that mitigates these stresses, by using a 2-step growth sequence: an

initial high growth rate step resulting in a pitted but relaxed film, followed by a low

growth rate smoothing layer. As a result, thick (>50 µm) and freestanding films have

been grown successfully. X-ray rocking curve linewidth of 105 arcseconds and 10K

PL indicating no “yellow” emission indicate that the material quality is higher than

that produced by conventional MOVPE. By further modifying the hybrid system to

include a metallic Mn source, it is possible to grow a doped semi-insulating GaN

template for use in high frequency electronics devices.

v

Acknowledgements:

First and foremost, I would like to thank my advisors Glenn Solomon and

James Harris for their constant support and encouragement to finish. My parents

Barry and Barbara Miller were instrumental in helping me finally attain this goal, and

I am grateful for their love and maintaining their faith in me over the years.

Additional thanks go to Julie Tell who helped me prepare for my thesis defense and

Jody Seltzer who effectively motivated me to put forth the effort to write this

dissertation.

The research for this dissertation was done at CBL Technologies, Inc. in

Redwood City, CA. I would like to thank Glenn again for his role as CEO and

director of the company in helping me define and shape this project. Our company’s

technicians Randy Carston and Rodney Worth greatly assisted me in performing

numerous crystal growth runs and constant system maintenance. The Matsushita

Electric Company of Japan gave generous financial support in the form of a joint

development agreement with CBL and provided Tetsuzo Ueda and Tadao Hashimoto,

two outstanding engineers who served in numerous roles at our facility. Finally, I

would like to thank Manfred Ramsteiner, Oliver Brandt, Achim Trampert, and Klaus

Ploog at the Paul Drude Institute for Solid State Electronics in Berlin for their

microstructural and optical characterization of our GaN material.

vi

Table of Contents

Chapter 1: The Need for Gallium Nitride Substrates ...................................... 1

1.1 Introduction ........................................................................................... 1

1.2 Properties of GaN, AlN and InN ............................................................. 6

Chapter 2: A New Hybrid VPE Method for GaN ............................................ 9

2.1 Towards a gallium nitride substrate ......................................................... 9

2.2 Substrates for GaN heteroepitaxial growth ............................................ 14

2.2.1 Silicon ............................................................................................. 14

2.2.2 Silicon carbide ................................................................................ 16

2.2.3 Sapphire .......................................................................................... 16

2.2.4 Lithium gallate ................................................................................ 17

2.3 Summary of heteroepitaxial substrate choices ...................................... 17

2.4 Vapor-phase epitaxy of GaN ................................................................. 19

2.4.1 MOVPE .......................................................................................... 19

2.4.2 HVPE .............................................................................................. 22

2.4.3 Hot and Cold Walled Reactors ....................................................... 26

2.4.4 Near-equilibrium vs. far from equilibrium processes ..................... 28

2.4.5 Low Temperature Nucleation layer ................................................ 30

2.4.6 The Hybrid VPE system ................................................................. 33

Chapter 3: Hybrid MOVPE/HVPE GaN process optimization ..................... 35

3.1 Stress in heteroepitaxial GaN ................................................................ 35

3.1.1 Lattice mismatch stress ................................................................... 36

3.1.2 Coalescence stress in GaN on sapphire .......................................... 38

3.1.3 Thermal mismatch stress ................................................................ 40

3.2 Effects of stress ...................................................................................... 42

3.2.1 Cracking ......................................................................................... 43

3.2.2 Peeling and delamination ............................................................... 47

3.3 The surface morphology of HVPE GaN films ....................................... 48

3.3.1 Hillocks ........................................................................................... 49

3.3.2 Pits .................................................................................................. 52

3.3.3 Hexagonal pits ................................................................................ 53

3.3.4 Irregular pits .................................................................................... 59

3.3.5 Quantifiable roughness measurement .............................................. 61

3.4 Effects of substrate temperature, growth rate and V/III ratio ............. 64

3.5 smoothing layer growth .......................................................................... 67

3.6 The 2-step growth process ...................................................................... 70

vii

3.7 Summary of GaN deposition process optimization techniques .............. 72

Chapter 4: Microstructural Characterization of VPE-grown GaN ................ 74

4.1 Structure of thin GaN layers ............................................................... 74

4.2 A 3-zone layered growth model ............................................................. 77

4.3 Structural improvement with increasing thickness ................................. 79

4.4 X-ray methods for determining approximate dislocation density .......... 80

4.41 Accounting for the effect of bowing on XRD linewidth .................. 81

4.4.2 The relationship between FWHM and dislocation density ............. 84

4.43 Reduction in the dislocation density with increased thickness ........ 85

4.5 Summary of microstructural characterization ....................................... 88

Chapter 5: Photoluminescence characterization ............................................ 90

5.1 Photoluminescence characterization of GaN .......................................... 90

5.1 10K Photoluminescence ..................................................................... 90

5.1.2 Photoluminescence setup ................................................................. 92

5.2 Photoluminescence characterization of a 12 µm sample ....................... 93

5.3 Photoluminescence of 28 micron layer .................................................. 95

5.4 Emission from Freestanding 60 µm GaN .............................................. 96

5.5 Two-electron transitions ......................................................................... 97

5.6 Summary of photoluminescence results ................................................. 98

Chapter 6: A semi-insulating GaN Alloy: GaMnN ..................................... 100

6.1 Semi-insulating GaN for use in microwave amplifiers ........................ 100

6.2 Semi-insulating GaN through the incorporation of Mn using HVPE .. 101

6.3 Initial Characterization of GaMnN ....................................................... 104

6.4 TEM analysis: a second phase and a new crystal structure ................ 107

6.4.1 Mn-rich second phase in GaMnN ................................................. 108

6.4.2 Sublattice ordering – a new crystal structure ................................ 110

6.5 Summary: semi-insulating GaMnN ...................................................... 112

Chapter 7: Conclusions and future directions .............................................. 114

7.1 Conclusions ......................................................................................... 114

7.2 Future directions and research ............................................................. 116

Chapter 8: List of References ...................................................................... 118

viii

List of Tables

Table 1.1. Room temperature band gap and lattice constants for GaN, InN and

AlN ................................................................................................................................. 7

Table 2.1. A comparison of lattice and thermal mismatch, chemical

compatibility issues, and cost for GaN heteroepitaxial substrates. .............................. 18

Table 6.1. Characterization summary of 5 µm thick GaMnN layers grown with

various HCl flow ratios. ............................................................................................. 105

ix

List of Figures

Figure 2.1. Schematic diagram showing the bonding arrangement and unit cell

for GaN. ........................................................................................................................ 10

Figure 2.2. Bandgap vs. lattice constant for group III-nitrides ....................... 11

Figure 2.3. The Ga - N2 phase diagram. ........................................................... 13

Figure 2.4. Temperature variation of the thermodynamic driving force for

MOVPE and HVPE. ................................................................................................... 26

Figure 2.5. Schematic diagram showing the different heating schemes used for

“hot-” and “cold-” walled reactors.. ............................................................................. 27

Figure 2.6. Atomic Force Microscope images of MOVPE nucleation layer

before and after recrystallization. ................................................................................. 32

Figure 2.7. Model of a novel hybrid MOVPE/HVPE deposition system

featuring hot-walled and cold-walled heating systems. ............................................... 34

Figure 3.1. A schematic drawing of a GaN (0001) unit cell overlaid onto the

(0001) sapphire unit cell. .............................................................................................. 37

Figure 3.2. A schematic representation of the evolution of coalescence stress in

heteroepitaxial GaN on sapphire. ................................................................................. 40

Figure 3.3. Illustration of the effects of the 33% mismatch in thermal

expansion coefficient between sapphire and GaN ........................................................ 41

Figure 3.4. A comparison of tensile and compressive stress cracking

mechanisms. ................................................................................................................. 45

x

Figure 3.5. Plan-view micrograph of a 10 µm thick GaN film that has cracked

and buckled during the cool down process ................................................................... 46

Figure 3.6. 100x optical micrograph of a 2 µm film grown at 1050ºC, showing

presence of numerous hexagonal hillock-shaped prominences. ................................... 50

Figure 3.7. Cross-section optical micrograph of a hexagonal-shaped pit in a 24

µm HVPE GaN film. .................................................................................................... 53

Figure 3.8. The basic tetrahedral bonding arrangement between Ga and N

atoms in GaN.. .............................................................................................................. 54

Figure 3.9. The GaN unit cell as viewed along the [1010] azimuth. ............... 55

Figure 3.10. Schematic diagram of the distortion of hexagonal pits as a

mechanism for strain relief.. ......................................................................................... 58

Figure 3.11. The effect of surface thermal pretreatment prior to deposition of

the low temperature MOVPE layer. ............................................................................. 61

Figure 3.12. Plan-view and cross-sectional micrographs of pitted and smooth

HVPE GaN films. ......................................................................................................... 63

Figure 3.13. The effect of substrate temperature on resulting surface

morphology of 2 µm HVPE films. ............................................................................... 65

Figure 3.14. The effect of the growth rate on 2 µm thick HVPE GaN films

grown at 1025ºC.. ......................................................................................................... 65

Figure 3.15. The effect of the V/III ratio on the surface morphology of 2 µm

HVPE GaN films grown at 1025ºC. ............................................................................. 66

Figure 3.16. The effect of smoothing layer regrowth on a pitted film. ........... 68

xi

Figure 3.17. The two-step growth process to achieve GaN film thickness in

excess of 15 µm. ........................................................................................................... 70

Figure 3.18. 500x cross-sectional optical micrograph of a 40 µm film grown by

the 2-step method, with comb overlay showing thickness variation. ........................... 71

Figure 4.1. High-resolution X–ray diffraction (XRD) ω-scan of a 1.5 µm thick

GaN film ....................................................................................................................... 75

Figure 4.2. High resolution (0002) XRD Gaussian curve fits to the data shown

in Figure 4.1. ................................................................................................................. 76

Figure 4.3. Schematic and TEM bright field image of the three-zone GaN

growth on a sapphire substrate. .................................................................................... 78

Figure 4.4. Mosaic of images of a series of TEM micrographs taken of a 12 µm

thick hybrid VPE-grown GaN film. ............................................................................. 79

Figure 4.5. Schematic diagram of the X-ray diffraction geometry for a bowed

substrate with bending radius r.. ................................................................................... 81

Figure 4.6. The variation in full-width half maximum (FWHM) X-ray

diffraction linewidth with X-ray source slit width for a 12 µm thick GaN layer on

sapphire. ........................................................................................................................ 83

Figure 4.7. The variation of the square of the magnitude of the zero-slit

extrapolation FWHM with GaN layer thickness. ......................................................... 86

Figure 5.1. The calculated band structure around the Γ point in wurtzite GaN.

...................................................................................................................................... 91

Figure 5.2. 10K PL spectra of a 12 µm thick HVPE-grown GaN film. .......... 93

Figure 5.3. 10K PL spectrum of a 28 µm thick HVPE GaN layer. ................. 95

xii

Figure 5.4. 10K PL spectra from a piece of 60 µm freestanding GaN. ........... 96

Figure 6.1. Schematic diagram of the Ga(Mn)N deposition system. ............. 102

Figure 6.2. The calculated MnCl2 vapor pressure as a function of temperature

.................................................................................................................................... 103

Figure 6.3. The linear relationship between the ratio of HCl flow over Mn to

HCl flow over Ga and the resulting GaMnN film’s Mn content ................................ 104

Figure 6.4. Plan-view 500x optical micrographs of 5 µm thick GaN and

GaMnN alloys. ........................................................................................................... 106

Figure 6.5. The relationship between ω-scan rocking curve FWHM and Mn

concentration in 5 µm thick GaMnN HVPE films. .................................................... 107

Figure 6.6. Transmission-electron microscopy (TEM) images of a GaMnN

sample with estimated 0.16% Mn.. ............................................................................ 108

Figure 6.7. Energy Dispersive X-ray (EDX) analysis on a TEM sample of

GaMnN ....................................................................................................................... 109

Figure 6.8. High resolution TEM image of GaMnN deposited on GaN on

sapphire. ...................................................................................................................... 111

1

Chapter 1: The Need for Gallium Nitride Substrates

1.1 Introduction

Within the last twenty years, applications using gallium nitride (GaN) have

evolved from use as an obscure industrial ceramic into an economically significant

semiconductor. At room temperature GaN has a direct electronic bandgap in the near

ultraviolet (360 nm or 3.4 eV), and it can be alloyed with the smaller-gap indium

nitride (InN) or the larger-gap aluminum nitride (AlN) to produce a material capable

of emitting or absorbing light from the infrared part of the spectrum through the mid-

ultraviolet. In addition to its direct bandgap, GaN also has a moderate intrinsic carrier

concentration and strong resistance to thermal and radiation degradation, properties

which potentially have use in a wide range of applications. These include high power,

high frequency, low noise microwave amplifiers for avionics and communications

systems, LEDs that can be made to emit from red to ultraviolet, useful in video

displays for computers, televisions and the solid-state white lighting industry, the blue

lasers used in HD DVD storage systems, and solid-state detectors for ultraviolet

radiation, including water purification and the militarily significant solar-blind region

of the spectrum (4-6 eV).

While the promise of GaN-based materials is great, there remain significant

technical challenges that must be remedied before this materials system can be fully

commercialized. Specifically, there are currently no native substrates of moderate size

(>2”) available for device layer growth. Bulk methods for GaN growth are still

somewhat experimental and difficult to reproduce on an industrial scale; this results in

2

the necessity for heteroepitaxial growth methods to deposit layers that simulate a

native substrate. Often, it is desirable to fabricate a freestanding GaN layer – one that

has been removed from the heteroepitaxial substrate – especially in circumstances

where features such as backside electrical contact or effective heat sinking for high

power devices are necessary. In other cases when a thick heteroepitaxial layer will

suffice, it may not be necessary to undergo the difficulty in creating a freestanding

substrate. This is often the case with LED production for instance, where contacts can

be fabricated on the top surface only; but even here, using a thick doped buffer layer

could serve the dual purpose of improving the epitaxial quality while furnishing a

bottom electrical contact.

Regardless of whether a layer is freestanding or not, heteroepitaxially grown

GaN layers have high defect densities, primarily in the form of threading dislocations

originating at the heteroepitaxial interface. These defects can have damaging effects

on devices, by reducing the luminescence efficiency for LEDs and laser diodes,

increasing noise and leakage current in high frequency amplifiers, and increasing the

series resistance, leading to excessive heat generation and early device failure. It has

long been known that thicker heteroepitaxial layers can have reduced defect densities,

owing to the effects of defect entanglement and annihilation as the film is grown.

Thus, the need for a thick GaN layer is driven not just from the desire for improved

heat sinking, and backside contacts, but also for the general improvement of all sorts

of electronic and optical devices gained by using a lower-defect density substrate

layer.

3

In this dissertation, I will present a novel, robust and economical method for

producing thick high quality low defect density heteroepitaxial GaN films for use as

substrates for subsequent device layer growth. This new method combines two

previously incompatible growth techniques, hydride vapor phase epitaxy (HVPE) and

metalorganic vapor phase epitaxy (MOVPE) into a single growth system. This allows

for the first time the direct growth without interruption of thick (greater than 20 µm)

GaN layers onto a heteroepitaxial substrate by using MOVPE to provide the buffer

layer and HVPE for the rapid growth of high quality material. I will demonstrate that

the quality of the GaN grown is extremely high, with low levels of impurity

incorporation, improved dislocation density, and fine detail structure observed in the

photoluminescence spectra. In the subsequent chapters I will address aspects of this

new growth technique, from the fundamentals of the two growth methods, effects of

the growth parameters on film quality, material characterization, and further system

modifications that resulted in the first recorded HVPE growth of semi-insulating GaN.

In chapter 2, I will describe the current state of the art of GaN growth,

comparing and contrasting the chemistry and thermodynamics of HVPE and MOVPE

growth. Put briefly, until now the methods have been incompatible because they

relied on entirely different substrate heating methods; HVPE has always relied on an

external heater (hot wall system) while MOVPE has always utilized internal heating

(cold wall system). I will describe the physical and chemical reasons why this is so,

and present the hybrid HVPE system that incorporates an internal heating scheme to

allow MOVPE in the same growth chamber.

4

Chapter 3 provides details on the optimization of the hybrid VPE growth

process. To do this, I will discuss the gross film quality metrics required for a device-

quality substrate, specifically the surface morphology and its effects on film stress and

cracking. The origins of stress in heteroepitaxially grown GaN will be described, from

the effects of island coalescence, to lattice mismatch, to differences in the thermal

coefficient of expansion between the heterosubstrate and the GaN film. Growth

parameters such as temperature, growth rate, type of surface pretreatment, and ratio of

group V to group III species (V/III ratio) all have their effect on film quality, and

finding the optimum combination to maximize surface flatness, minimize stresses to

prevent cracking, while growing sufficiently thick films to have reduced dislocation

density is a complex problem. I will show that a thermal surface pre-cleaning step is

critical for obtaining high quality deposition, and that high growth rates yield low-

stress films with rough pitted morphology, while low growth rates produce smooth

films that are stressed and frequently crack. Using the two growth regimes, by first

growing a pitted high rate layer followed by a smoothing low rate layer, produces

thick smooth layers of good quality with manageable residual stress. Smoothing layer

growth over a pitted layer may be done at any time, immediately after deposition of

the pitted layer or in a second growth operation following heterosubstrate removal.

In chapter 4 the microstructure of the hybrid MOVPE-HVPE films are

characterized using X-ray diffraction (XRD) and transmission-electron microscopy

(TEM). I will show that the film evolves as it grows away from the interface in three

roughly defined zones: interfacial, transition, and bulk. In the interfacial zone the

microcrystalline islands coalesce and the disorder is greatest at their grain boundaries.

5

In the transition zone, the material rapidly becomes more ordered, from highly

oriented polycrystalline to a single crystal, leading to the onset of the bulk zone where

threading dislocations continue to entangle and annihilate. The correlation between

threading dislocation density and X-ray rocking curve linewidth will be discussed, and

I will show that there is a quantifiable improvement as film thickness increases out to

50 µm and beyond.

The optical properties of as-grown and freestanding GaN are characterized via

photoluminescence in Chapter 5. I will show that the material quality is very good

with no mid-gap yellow luminescence and the near band-edge luminescence is

dominated by donor-bound and free exciton emission profiles. The narrow spectral

lines show that the material is highly uniform and with a relatively low donor

concentration.

By adding manganese (Mn), a mid-gap dopant, to the hybrid growth system it

is possible to produce for the first time semi-insulating GaN. In Chapter 6, I will

demonstrate how I have done this, and how the growth parameters affect the

incorporation of the Mn into the GaN. From X-ray and TEM characterization results,

I will show that at low levels (less than 0.1%) the Mn is incorporated as alternating

Mn-rich and Mn-poor crystal planes in the growth direction. As the Mn concentration

increases, phase segregation occurs and MnN crystallites are observed.

Finally, I will summarize and present my conclusions in Chapter 7. I will offer

suggestions for further areas of research and development for this new hybrid growth

technique, including proposing system modifications and improvements to allow for

the production of thick alloy layers, as well as high-temperature MOVPE growth. In

6

this way it may be possible to produce a complete device structure, from nucleation

layer to thick GaN (or alloy) HVPE buffer to the final MOVPE device layers all in a

single growth system at one time.

1.2 Properties of GaN, AlN and InN

GaN occurs in two polytypes, hexagonal (wurtzite) and cubic (zincblende). Of

the two forms, hexagonal is the more stable version and is most commonly used for

semiconductor devices.1 Unlike cubic materials with a single lattice constant, the

wurtzite structure has two lattice constants: an “a” lattice constant associated with the

spacing within the (0001) basal plane, and a “c” lattice constant associated with the

unit cell spacing normal to the (0001) plane (along the [0001] crystal direction). At

room temperature, GaN has the lattice constants a = 3.189 Å and c = 5.185 Å.

Hexagonal GaN has a room temperature band gap of 3.43 eV, corresponding to

a wavelength of 361 nm, in the near ultraviolet part of the spectrum. As with the III-V

arsenide system, GaN can be alloyed with InN (bandgap 0.75 eV) and AlN (bandgap

6.2 eV) to produce direct-gap alloys with an emission spectra ranging from the

ultraviolet to infrared, although the issue of lattice mismatch between related

compounds is more severe with nitrides than for arsenides (see Table 1.1 below,

comparing the band gap and lattice constants for the group-III nitrides).

The desire for a solid-state direct-gap emitter in the blue part of the visible

spectrum arises from the fact that although red and to some extent green devices can

be produced using arsenides or arsenide-phosphide alloys of Ga, In and Al, the

bandgaps of these alloys are too small for blue emission. As the human eye perceives

7

white light as a combination of red, green, and blue wavelengths, producing a solid-

state white light source requires a blue emitting material, achievable by using InGaN-

based light emitting diodes (LEDs). An alternative approach where shorter-

wavelength emission stimulates phosphors for broad-spectrum white light also

requires a wide bandgap semiconductor source.

Table 1.1. Room temperature band gap and lattice constants for GaN, InN and AlN

room temp Eg (eV)

lattice constant “a” (Å)

lattice constant “c” (Å)

“a” mismatch relative to GaN

GaN 3.43 3.189 5.185 -

InN 0.75 3.533 5.693 10.8%

AlN 6.2 3.112 4.98 -2.4 %

The advent of nitride-based LEDs has brought forth a proliferation of novel

and exciting devices and applications. Modern GaN-based LED-based traffic lights

emit in the green and yellow wavelengths, simultaneously using less energy and

requiring replacement much less frequently, resulting in a real cost savings. Nitride-

based LEDs are everywhere in white lighting systems for automotive and consumer

electronics displays, as well as large-scale installations such as illuminated and

animated billboards and stadium-sized displays. Sony, the originator of the

“Jumbotron” stadium display, saw the advantages in the year 2000 and switched

technologies from inefficient incandescent bulbs to high-brightness nitride-based

LEDs.

Another application for GaN-based emitters is for HD DVD systems. Current

DVD systems utilize a 650 nm laser based on AlGaAs, allowing some 5 to 7 gigabytes

8

(GB) of data to be stored on a single disc. By using a shorter wavelength laser, it is

possible to directly increase the areal data density stored on the disc. The data

increase with decreasing wavelength is only limited by the absorption of the disc

material, in this case polycarbonate plastic, which strongly absorbs below 405 nm. By

utilizing a 405 nm laser and using other data compression techniques, it is expected

that it will be possible to store as much as 25 GB onto a similar-sized disc; the goal

here is to be able to fit an entire high-definition movie feature onto a single disc for

playback.

Electronic devices, both optical and non-optical, such as high power high

frequency microwave amplifiers, require aggressive cooling schemes to avoid damage

or degradation during operation. GaN’s wide bandgap affords superior protection

against thermal degradation, allowing for operation in excess of 200º C, unthinkable

for GaAs or Si-based electronic devices. By reducing or eliminating the need for

cooling systems, weight and complexity for amplifier packages can be greatly

reduced, realizing a significant cost and fuel savings for avionics and satellite-based

applications, where low mass is an important figure of merit. Even in less weight-

critical applications, such as in cellular and WiMax infrastructure system components,

the improved power handling and reduced cooling requirements are attractive.

For these reasons, as well as other niche applications, the market for GaN and

GaN-based devices has been rapidly growing throughout the first decade of the 21st

century. The demand for more and more GaN substrates requires the development of

methods for producing them more and more cheaply and with higher and higher

quality.

9

Chapter 2: A New Hybrid VPE Method for GaN

2.1 Towards a gallium nitride substrate

By combining elements from groups III and V of the periodic table, a family of

semiconductor materials of technological and scientific importance can be produced.

While these compounds may exist in an impure state in nature, the ability to create

high-purity artificially structured layers significantly leverages their natural properties.

Starting with the earliest attempts at GaAs production, these materials have been

combined in binary, ternary, quaternary, and in some cases quinary compounds to

precisely tailor their electronic properties to fit a variety of purposes, from high-speed

microwave amplifiers, to visible and IR lasers, LEDs and many others. In general,

III-V compounds have superior electronic properties compared to silicon, with higher

electron mobilities and saturation velocities, and lower intrinsic carrier concentrations,

allowing fabrication of high-speed low-noise electronic circuits. Additionally, many

III-V compounds are capable of doing something silicon cannot: efficiently emit light.

These compounds and alloys with direct electronic gaps have been the linchpin of the

optoelectronics industry, from AlGaAs lasers in CD and DVD recorders, to the

InGaAsP lasers emitting at 1.55 µm in long-distance optical fiber communications

systems..

Common to the entire class of III-V compounds and alloys is the bonding

arrangement between cations (group III) and anions (group V). When brought

together the atomic orbitals of both species undergo hybridization of the pi-bond

variety, which results in tetrahedral coordination between neighboring anions and

10

cations, as seen below in Figure 2.1. This type of bonding arrangement leads to two

different crystalline structures: a cubic form known as zincblende, one example of

which is GaAs, and a hexagonal form known as wurtzite, such as BN. The difference

between the two crystalline forms lies in the stacking order of the close-packed planes,

i.e. the (111) planes for cubic structures and the (0001) planes for hexagonal ones.

Cubic materials use an ABCABC… stacking order while hexagonal materials just use

an ABAB… sequence. While gallium nitride has been produced in both zincblende

and wurtzite forms, the wurtzite structure has a wider bandgap and is more stable at

atmospheric pressure.2

The bandgap for a particular III-V compound is directly related to the bond

strength between the cation (group III) and anion (group V). The trend is that the

lower the atomic number for either species, the closer the bonding orbitals are to the

nucleus, the stronger the bond and the larger the electronic bandgap. Thus, the

bandgap for GaAs is 1.42 eV, while GaP is 2.26 eV, AlAs is 2.16 eV and AlGaInP

Figure 2.1. (Left): Schematic diagram showing the bonding arrangement between Ga and its 4 nearest neighbor N atoms in GaN. (Right): view of the GaN unit cell along the [101�0] azimuth; the Ga (and N) sublattice follows the ABAB… stacking sequence along the [0001] direction.

11

varies from 1.91-2.52 eV depending on composition. Following this natural

progression, the nitride group of III-V compounds should be expected to have the

largest bandgaps, and it does. Gallium nitride has a room temperature bandgap of 3.4

eV, AlN is 6eV and InN is 0.75eV. In a manner analogous to the group III-arsenides

system, GaN can be alloyed with the AlN and InN to effectively allow engineering the

bandgap from the near infrared part of the spectrum through the ultraviolet. The group

III-nitride system allows for the first time the creation of high power, high efficiency

optical devices operating in the yellow, green, blue, and ultraviolet parts of the

spectrum, within certain limits. These limits are imposed by the larger lattice

mismatches between the group III nitrides compared to the group III arsenides: AlAs

has a lattice constant 0.1% larger than GaAs, while for InAs it is 7% larger; however

AlN has a lattice constant 2.4% smaller than GaN while InN has one 11% larger.

Figure 2.2 below is a plot of the bandgap and lattice constants for the group III

nitrides, with some other common materials included for reference.

Figure 2.2. Bandgap vs. lattice constant for group III-nitrides, sapphire, SiC, GaAs, and Si. The wavelength of light corresponding to each gap is given in nanometers for reference purposes.

12

Key to any method for producing nitride-based electronic devices is the

requirement for a high quality single crystal substrate. Without such a substrate, it is

not possible to epitaxially grow the thin layers used for the active regions of devices.

In the case of silicon- and arsenide-based electronics, Si and GaAs substrates produced

by a bulk growth method from the melt are available at a low cost. In the case of

silicon, a single crystal boule is pulled from a high purity silicon melt; for GaAs the

process is slightly more complex as GaAs tends to dissociate into elemental Ga liquid

and As vapor near its melting temperature. This is overcome by using a pressure

vessel operating at several tens of atmospheres of pressure and a liquid encapsulant

such as boron to maintain an appropriate overpressure of As while the boule is pulled

from the melt.

In the case of GaN the pressures are more extreme. The relative strength of the

Ga-N bond is much greater than the Ga-As bond; this increased bond strength has the

effect of raising the melting temperature of GaN to over 2420ºC. Similar to the case

of the arsenides, GaN tends to dissociate with increased temperature, but unlike the

arsenides, the dissociation products include the highly stable N2 gas, which makes the

dissociation reaction virtually irreversible except at the highest of pressures, as seen

below in the Ga-N phase diagram of Figure 2.3.

As can be seen in Figure 2.3, to approach the congruent melting temperature of

GaN requires maintaining a nitrogen overpressure on the order of 93,000 atmospheres.

While this can be done for small-sized (few mm) reaction vessels to produce single

crystals on the order of millimeters in size, the extremely high pressure requirement

13

makes this process difficult to scale up for commercially viable sized substrates, which

are a minimum of 2 inches across. Other methods have been proposed for bulk GaN

crystal growth, including flux growth through a liquid or solid medium3,4 and

somewhat reduced high pressure growth using supercritical ammonia5,6, but to date

these are experimental processes that have not yet produced wafer-scale pieces of

high quality material.

Figure 2.3. The Ga - N2 phase diagram (from Okamoto7). GaN dissociates at elevated temperatures – the equilibrium pressure between GaN and N2 for various temperatures is shown by the dotted lines on the left side of the figure. As the temperature rises, the necessary N2 overpressure increases until it reaches 93,000 atmospheres at the congruent melting temperature of 2427ºC.

The problem persists: how does one produce a high-quality single crystal GaN

substrate for device layer growth? Currently the best approach uses heteroepitaxial

growth of GaN onto other mono-crystalline substrates to produce a sufficiently thick

layer that can act as a quasi-bulk substrate. Epitaxial growth, unlike bulk growth,

14

utilizes chemical reactions on a substrate surface to produce GaN at temperatures and

pressures far below the melting point; in this way it is feasible to heteroepitaxially

produce thick “pseudosubstrates” onto dissimilar single-crystalline materials. There

are a variety of possible heterosubstrates available to use as templates for GaN

pseudosubstrate growth, and in the next section I will discuss the relative merits and

disadvantages of several of the most commonly used ones.

2.2 Substrates for GaN heteroepitaxial growth

Because there is no bulk-grown GaN substrate available, all GaN layers to be

used for device growth must themselves be grown on a heteroepitaxial substrate. The

choice of which particular substrate to use is based on compromises. Factors such as

the lattice constant (and lattice mismatch between the substrate and GaN), the

difference in thermal expansion coefficients between the materials, which results in

stress accumulation during post-growth cool down, and the cost of the substrate itself

must be weighed and considered before the optimal choice can be determined. The

growth technique used to deposit the GaN often affects substrate material choice too;

some processes use or have as by-products chemical compounds that may attack a

particular substrate while leaving other materials unaffected. In the following

sections, I will consider and discuss the relative strengths and weaknesses of various

commonly used substrates for GaN heteroepitaxial growth.

2.2.1 Silicon

By far the least expensive of the substrate choices, high quality epi-ready (111)

silicon is readily available, with a price of around $25 per 100mm wafer.8 (111)

15

silicon has a lattice mismatch of approximately 17% with respect to (0001) GaN, and a

thermal expansion coefficient less than half of that for GaN. GaN films grown on

silicon undergo tensile stresses as the temperature is lowered, potentially leading to

cracking as the substrate with GaN epitaxy is cooled. Silicon is also a reactive

surface; while an atomically clean surface can be prepared for growth, when exposed

to ammonia above 500°C, silicon has the potential to form an amorphous silicon

nitride layer several atomic layers thick. This amorphous layer makes heteroepitaxy

impossible, as crystalline registry with the lattice below, is disrupted. Special

techniques, such as lower temperature deposition, use of surface patterning for

selective overgrowth, and/or the use of one or more buffer layers are necessary for

good results.1,9 Silicon is susceptible to chemical attack by chlorine-bearing species;

this is of concern where HVPE growth is contemplated. In that particular process,

HCl is formed as a by-product of the growth reaction. HCl can react with exposed

silicon to form volatile silicon chloride, which then incorporates into the growing film.

To avoid this, it is necessary to completely passivate all exposed non-growth surfaces

prior to the commencement of HVPE growth; even small pinholes in a passivating

layer can lead to significant silicon contamination in the film. As a natively-grown

nitride layer formed by ammonia may have such pinholes, the passivation is frequently

done by depositing an additional MOVPE GaN layer onto the non-growth side of the

wafer.10

16

2.2.2 Silicon carbide

Hexagonal SiC is used as a substrate for GaN heteroepitaxial growth. While

SiC has a better lattice mismatch than silicon, 3.5%, its thermal coefficient of

expansion is almost twice that of GaN and it is extremely expensive, costing upwards

of $1000 per 4” wafer,8. SiC substrates are available in both p and n type, allowing

for the use of backside contacts for devices grown on thick GaN films. However, the

thermal coefficient mismatch puts the GaN layer under significant compressive stress

during cool down, which can lead to film cracking or breaking; thicker films

exacerbate this problem. SiC is more resistant to chemical attack than silicon, and

does not form a nitride blocking layer as readily; therefore passivation of exposed non-

growth surfaces is not necessary. SiC has a band gap of 3.0 eV, which is smaller than

GaN; optical methods for removing the GaN film, such as laser lift-off, are not

effective unless applied from the epitaxial side of the wafer, a process that might

damage or induce defects into the film. While a similar situation exists with silicon

substrates, silicon can easily be removed by selective etching, a process that is more

difficult for silicon carbide due to its highly stable nature.

2.2.3 Sapphire

Sapphire is commonly used as a substrate for GaN epitaxy, representing a

reasonable compromise between cost and performance. Although it has a 14% lattice

mismatch and the thermal expansion coefficient 33% larger than GaN, it is virtually

impervious to chemical attack in a growth reactor. Sapphire is an extremely stable

oxide compound, not susceptible to forming an amorphous nitride when exposed to

17

ammonia, nor does it melt or decompose at GaN growth temperatures, nor is it

attacked by the chlorine compounds in an HVPE reactor. At approximately $150 per

4” wafer,8 it is not expensive. Sapphire is a dielectric, however; with a bandgap of 10

eV, and so it is not possible to fabricate a device on GaN on sapphire with a backside

electrical contact. Its large bandgap makes it transparent, however, so it is possible to

use optical methods for removing the GaN film from the substrate side, minimizing

the possibility of damage to the epitaxial film during removal.

2.2.4 Lithium gallate

Lithium gallate (LiGaO2) appears to be a promising material;11 when in its

(001) orientation it has a lattice constant very close to GaN, with less than 1%

mismatch. However, this material is not currently available in a full sized wafer; 10

mm square pieces cost in excess of $100,12 making the price very high in comparison

to the other choices. Lithium gallate is also susceptible to attack by chlorine

compounds in an HVPE reactor; use of this technique requires passivation of the

exposed surface prior to growth, similar to that for silicon. Passivation of the exposed

surface also prevents the thermal decomposition of lithium gallate at the elevated

temperatures required for GaN growth.

2.3 Summary of heteroepitaxial substrate choices

While it is clear that heteroepitaxial GaN growth requires compromises when

selecting a substrate material, some choices are better than others. Table 2.1 below

summarizes the properties and characteristics of the substrates described and how they

compare to GaN itself.

18

Table 2.1. A comparison of lattice and thermal mismatch, chemical compatibility issues, and cost for GaN heteroepitaxial substrates: Si, GaAs, Al2O3 (sapphire), GaAs, and LiGaO3 (lithium gallate).

Currently, the most commonly used heteroepitaxial substrates for GaN growth

are sapphire and SiC; SiC is favored for growth of electronic devices such as high

power amplifiers primarily for its relatively small lattice mismatch and high thermal

and electrical conductivity, which allow for efficient heat-sinking and backside

electrical contacts. On the negative side, the large mismatch in coefficient of thermal

expansion (CTE) limits the thickness of GaN layers that may be grown without

cracking. Sapphire on the other hand is less expensive and virtually inert chemically,

making it most suitable for the harsh environments created by some GaN growth

processes, while having a much smaller thermal mismatch. This makes it most

suitable for the growth of thick pseudosubstrate GaN layers using the high growth rate

but chemically aggressive HVPE growth process. HVPE, as well as the other method

for vapor phase epitaxy of GaN, MOVPE, will be described in the next section.

Material Lattice Constant (Å)

Lattice Mismatch

CTE (room temp)

CTE mismatch

Compatibility Issues

Cost per wafer

Si (111) 3.83 (110) 17% 2.6 x 10-6/°C

- 53% Attacked by HCl, reacts with NH3

$258

(100 mm)

SiC a = 3.08; c = 15.12

3.5% 10.3 x 10-6/°C

+ 84% None, but Eg = 3.0eV (< GaN)

$10008

(100 mm)

Al2O3 sapphire

a = 4.758; c = 12.99

49% (14%)

7.5 x 10-6/°C

+ 34% None $1508

(100 mm)

GaAs (111)

3.26 (111) 2.4% 6.86 x 10-6/°C

+ 23% Attacked by HCl, decomposes

$12013 (100 mm)

LiGaO3 (100)

5.40 (orthorhombic)

< 1% n/a n/a Attacked by HCl, decomposes

$130-20012

(10 x 10 mm)

GaN reference

a = 3.189; c = 5.185

N/A 5.59 x 10-6/°C

N/A N/A N/A

19

2.4 Vapor-phase epitaxy of GaN

Thick layers of GaN are grown by a vapor-phase epitaxy (VPE) process, either

using metalorganic precursors (MOVPE) or chlorine-based precursors (HVPE).

While both processes have fundamental differences, they share some basic

similarities: both react a gas-phase gallium-bearing compound with a gas-phase

nitrogen-bearing compound on the surface of a substrate to form GaN. The nitrogen-

bearing compound is almost always ammonia (NH3); the differences between the two

growth processes principally lie with the gallium-bearing compounds and the chemical

reactions involved in the deposition. In this section, a brief discussion of both

MOVPE and HVPE will be presented, with some further elaboration on using the best

aspects of both techniques for optimum results.

2.4.1 MOVPE

MOVPE, metalorganic chemical vapor deposition, utilizes a group III

metalorganic compound such as trimethylgallium (TMG), trimethylaluminum (TMA),

trimethylindium (TMI), etc. A typical deposition reaction between a metalorganic

compound (in this case TMG) and ammonia is given below:

Ga(CH3)3 + NH3 → GaN + 3CH4

Thermodynamic calculations14 indicate that at a temperature of 1000°C, this is

a strongly exothermic reaction, releasing nearly 2.5x105 Joules/mole of GaN. This

deposition reaction also shows a net increase in entropy, on the order of 43

Joules/mole•K. The free energy, ∆G, available for a reaction at a given temperature T

is based on both the enthalpy, ∆H, and the entropy change, ∆S:

20

∆G = ∆H – T∆S

For MOVPE deposition, which is both exothermic and positively entropic, the

free energy change (or driving force) becomes more negative as the temperature rises.

For any reaction to occur spontaneously, the free energy change of the reaction must

be negative (i.e. the products must be at a lower energy state than the reactants). Thus,

it is notable that the driving force promoting the deposition reaction, which is just the

magnitude of the free energy change, increases as the temperature is increased. The

significance of this fact will be discussed further when comparisons are made with

HVPE.

MOVPE is a well-established technique for GaN growth, having been

developed for the production of other III-V semiconductor materials. Commercially

available deposition systems are capable of depositing smooth, high quality layers

suitable for use as active layers in device structures. While system design and

operation methods have already been optimized, these systems are mechanically

complex, involving significant capital expenditure and maintenance to preserve their

optimum growth characteristics. Additionally, these systems are usually designed

with lower growth rates in mind; typically the maximum growth rate for GaN is less

than 10 µm/hr in such a system. The production of thick layers in excess of 100 µm

requires long growth times, leading to higher maintenance and running costs. Metal-

organic compounds are expensive, long growths for thick layers consume

proportionately more of them; thick or freestanding layers can be very costly to

produce using only this method.

21

Because a variety of metal-organic compounds are available, MOVPE is well

suited for the production of alloy layers, such as GaInN for blue LEDs and lasers.

Alloys are made by introducing separate metal-organic compounds in specific ratios

into the growth reactor; this process is precise and controllable. MOVPE systems are

typically set up to allow for rapid changing of the group III precursor ratios, thus

allowing for rapid changes in film composition at nearly the monolayer level; this is

excellent for all forms of device production, from quantum wells to transistor

structures. Additionally, the metalorganic precursor compounds and reaction

byproducts are relatively non-reactive with the hetero-substrate, methane gas does not

chemically attack any of the substrates mentioned in the previous section.

Metal-organic compounds are sensitive to temperature; above approximately

500°C, these compounds undergo pyrolysis and will decompose on surfaces they

encounter. This can lead to effects such as gas-phase depletion upstream of the

deposition zone, reduced growth rate at the substrate, and wall deposition and flaking

off of particles downstream. To alleviate the problem, MOVPE systems are designed

using a heating system that heats only the substrate and its holder, while not

substantially heating the walls of the reactor (cold wall system). Typical methods for

doing this utilize RF induction heating of a graphite susceptor, or banks of high

intensity lamps impinging directly onto the substrate. In some instances the walls are

actively cooled using a forced gas stream, in other cases the walls are allowed to reach

an intermediate temperature without additional cooling. Cool or cold walls are

important for another reason: the temperature-dependent part of the free energy

equation (-T∆S) becomes more negative with increased temperature due to the

22

positive entropy change of the MOVPE deposition reaction. This increases the driving

force for the reaction on hotter surfaces; keeping the walls cooler than the substrate

reduces the chemical potential for the reaction and deposition on those surfaces.

Overall, MOVPE is an excellent method for growing less than 10 µm thick

layers of GaN and other group III nitrides. It can allow rapid switching between group

III precursors, making it ideal for growing alloyed layers of different compositions for

device structures. However, its low maximum growth rate, and its high cost of

materials, especially metal-organic precursors, makes it a less-than-ideal method for

the production of thick layers of GaN.

2.4.2 HVPE

HVPE is a deposition technique dating back to the earliest days of III-V

compound research of the 1960s.15,16,17 Metal atoms, from a reservoir of pure metal,

are transported to the substrate by first reacting upstream with HCl at 800-900°C,

forming a volatile gas-phase metal chloride:

Ga + HCl → GaCl + ½H2

Owing to factors such as reservoir chamber design, gas flow rate and residence

time, and reservoir temperature, this reaction may go forward to virtual completion as

predicted by thermodynamic models, or it may be partially incomplete, resulting in an

excess of HCl in the GaCl gas stream.

Once volatized, the metal compound is transported by carrier gas downstream

to the substrate zone, where the deposition reaction occurs:

GaCl + NH3 → GaN + HCl + ½H2

23

This reaction is also exothermic, although with an enthalpy of reaction on the

order of 6.9x104 J/mol, it is only about 30% that of the MOVPE reaction.

Interestingly, the entropy change for this reaction is negative, approximately -46

J/mol•K. Considering the reaction’s negative change in entropy, it is possible to

predict that the driving force for deposition decreases as the temperature increases. In

fact, above a temperature of approximately 1240°C the deposition reaction becomes

thermodynamically unfavorable and etching rather than deposition should result.

It is important to note that the deposition reaction results in the formation of an

HCl molecule and a hydrogen molecule; both molecules have the potential to react

with the GaN film, the substrate, and the substrate holder in the reactor. HCl, in

particular, can etch GaN to produce GaCl and nitrogen:

HCl + GaN → GaCl + ½N2 + ½H2

HCl, whether from the completion of the deposition reaction or from an

incomplete GaCl production reaction, can chemically attack the film during growth.

Often this etching is preferential; with higher etch rates in strained regions, such as the

boundaries of coalesced grains. This partial etching may result in a “textured” or

rough film morphology, often in the form of hexagonal shaped hillocks and pyramids.

The MOVPE reaction, on the other hand produces methane, which does not interact

with the reactor, substrate, substrate holder, or GaN film.

HVPE is a method best suited for high growth rate deposition; deposition rates

of 10-100 µm and higher are readily achievable.18,19 The use of HCl to transport

gallium has certain advantages. As an inherently carbon-free process, there is little

opportunity for carbon to become incorporated in the growing film. The precursor

24

chemicals HCl and especially Ga also can be obtained in purity levels that exceed

those of metalorganic materials; the purest trimethylgallium can be 99.9999% (“six

nines”) pure, while MBE-grade gallium is available at 8 nines purity at one tenth the

price. Likewise, impurities in HCl can be reduced to sub-ppb levels by passing it

through a corrosive gas purifier. The lower cost and higher purity of precursor

materials make HVPE a good choice for producing thick or free-standing layers.

While HVPE is primarily used for GaN growth, similar delivery systems

utilizing HCl gas and a metal have been devised for indium and aluminum nitrides as

well.20 While a combination of individual metal delivery systems may be utilized to

produce alloys, HVPE is not as well suited for rapid compositional changes during

growth like MOVPE. This is because the delivery of GaCl to the substrate can only

occur after HCl is passed over the gallium upstream; this leads to inherent latencies in

the process. The production of GaCl will occur only after the HCl flow is initiated,

and will continue for some brief time after it is terminated; this makes formation of

abrupt interfaces a difficult task.

Unlike the fragile metal-organic compounds, GaCl is not pyrolized by contact

with a heated reactor wall upstream of the substrate. The strength of the gallium-

chlorine bond is the reason; the average thermal energy at the growth temperature is

too low to lead to spontaneous pyrolysis, absent a catalyst such as a GaN surface in the

presence of ammonia. A “cold-wall” heating system as used for MOVPE is neither

necessary nor desirable for HVPE. At around 420°C, GaCl condenses from the gas

phase and spontaneously self-reacts to form a stable low-temperature gallium

trichloride (GaCl3) and excess gallium metal. For this reason, it is important that the

25

region between the area of GaCl production and GaN deposition must be maintained

at a temperature above this point to prevent this condensation from occurring.

The choice of substrates to use for HVPE is more constrained compared to

MOVPE. Because of the inevitable presence of HCl in the reactor environment, it is

necessary to ensure that the substrate onto which GaN is deposited is not first attacked

by HCl, or by the somewhat less reactive GaCl. Silicon, for example, readily reacts

with both compounds, forming volatile silicon chloride and gaseous hydrogen or

liquid gallium as by-products. The silicon chloride may be incorporated into the

growing film, while gaseous hydrogen diffuses away. The liquid gallium, however,

remains behind at the silicon surface, disrupting the crystalline registry between the

substrate and any subsequent epitaxial growth. Thus, films grown on silicon tend to

be heavily doped with silicon, and with a poor morphology. It is possible to mitigate

this effect by prior deposition of a passivating layer onto the silicon; this layer serves

to isolate the reactive silicon from the chloride attack. Unfortunately, passivating

layers are usually non-crystalline, which can disrupt the registry between the silicon

and the epitaxial layer, resulting in very poor morphology. Techniques have been

developed where a MOVPE layer of GaN21 or AlN22 are first grown, providing

passivation while maintaining some coherence between the silicon and the film.

Similarly, non-passivated surfaces on LGO substrates are susceptible to

chloride attack, and thermally induced dissociation. Deposition of a lower-

temperature GaN layer, often by MOVPE,11 can be useful in helping the situation.

Substrates made of more stable compounds, such as SiC and sapphire, do not require

26

surface passivation and are not attacked by the HVPE reactor environment, making

them better candidates for HVPE growth.

2.4.3 Hot and Cold Walled Reactors

When comparing both MOVPE and HVPE deposition reactions, it is useful to

consider the type and method of substrate heating systems, and how these affect the

reactor environment. As previously mentioned, the driving force (free energy change)

for MOVPE and HVPE differ significantly in magnitude and sensitivity to

temperature. Figure 2.4 above compares the variation in free energy change with

temperature for HVPE and MOVPE.

As shown in Figure 2.4, the driving force for MOVPE deposition increases

with temperature, while the much smaller driving force for HVPE decreases, crossing

Figure 2.4. Temperature variation of the thermodynamic driving force for MOVPE and HVPE. While the actual free energy change for both reactions is negative, this plot shows the magnitude of the change as a positive value for visual ease of comparison. (Thermodynamic data from Przhevalskii14)

27

zero (equilibrium) at approximately 1240°C. The fact that the driving force increases

with temperature for MOVPE indicates that a hot-walled deposition system, as

depicted below in Figure 2.5a, is a poor choice for this technique. Walls that are

approximately the same temperature as the substrate present a large upstream surface

where the reactants may react prior to reaching the substrate. This leads to wall

deposits which may flake off particles and a generally inefficient delivery of reactants

to the desired growth surface.

Figure 2.5. Schematic diagram showing the different heating schemes used for “hot-” and “cold-” walled reactors. A) A hot-walled reactor utilizes a tube furnace that heats the entire reactor to a more uniform temperature, although typically the walls are slightly hotter than the substrate/susceptor in the center of the reactor tube. This is the optimum heating scheme for HVPE. B) A cold-walled reactor utilizes a heating system that heats only the susceptor and substrate, the reactor walls are unheated and remain significantly below the substrate temperature. This is the optimum heating scheme for MOVPE.

28

Cold-wall systems, by virtue of the cooler reactor walls, present an upstream

surface with a reduced driving force for this reaction to occur, as shown in Figure

2.5b. Thus, the propensity for wall deposits and reactant losses are reduced in this

instance. It should be noted, however, that the values calculated for Figure 2.4 are the

standardized free energies for the reactions, and are calculated per mole of GaN

deposited, assuming no dilution of the reactants in the gas stream. In real-life

situations, the dilution of the TMG and NH3 into the gas stream will reduce the

magnitude of the driving force slightly, but the overall trend remains the same.

In the case of HVPE, as the driving force decreases with increasing

temperature, walls that are cooler than the substrate will have a greater tendency to

develop wall deposits. For this reason, HVPE growth is done in a hot-walled

environment, typically a furnace.

2.4.4 Near-equilibrium vs. far from equilibrium processes

It is interesting to compare the magnitudes of the driving forces for the two

reactions, as they differ so greatly. The smaller the driving force behind a reaction,

the closer to an equilibrium process it is; thus, HVPE is a near-equilibrium process,

while MOVPE is a far-from-equilibrium process. One significant aspect of this lies in

how rapidly such a process can occur, i.e. the growth rate. In deposition processes that

are far from equilibrium, an adsorbed atom has a large driving force “pushing” it to

any available suitable site on the growth surface. As the magnitude of the driving

force for the reaction increases, so does the probability of an adsorbed atom bonding

into a higher-energy defect state. Point defects such as vacancies and antisite defects

29

require additional energy to form, but since the driving force in far-from-equilibrium

process is so large, discrimination between preferred and less-preferred sites is

reduced.

In the case of a near-equilibrium process, the driving force acting to put

adsorbed atoms into a stable position is much lower. As such, the additional energy

required to form a defect is a larger fraction of the driving force, which in turn,

reduces the frequency of such defects’ formation. In a way, the reduced driving force

allows adsorbed atoms to explore multiple possible sites prior to their incorporation

into the film; this allows for the selection of the most energetically favorable site,

instead of just the first one encountered.

The driving force can be related to the growth rate by considering that the

driving forces that have been mentioned are per mole of GaN produced; the growth

rate is the product of the molar deposition rate and lattice constant. The “driving

power” for deposition is the product of the molar growth rate and the free energy

change for the particular reaction (HVPE or MOVPE). Defect generation is an

entropy-driven process; the greater the driving power, the greater the propensity for

defects to spontaneously form. At some power, the defect generation rate will be

sufficiently large to produce unacceptably high defect levels in the grown material.

Because the driving force for HVPE is just 2-4% of that for MOVPE, an HVPE

growth rate (or molar deposition rate) of 25-50 times greater would have an equivalent

driving power. These similar driving power dissipation rates could explain why

typical MOVPE growth rates are 1-5 µm/hr and HVPE’s are 10-200 µm/hr.

30

2.4.5 Low Temperature Nucleation layer

Because of the large lattice mismatch between the available heterosubstrates

and GaN, high quality epitaxial growth directly onto the substrate is difficult, and

often polycrystalline films with random orientations are the result. The usual method

for homoepitaxial growth, the Burton, Cabrera and Frank (BCF) mechanism,23

involves atoms attaching at ledges and kinks in the slightly roughened crystalline face.

In this manner the film grows in a layer-by-layer method often referred to as

2-dimensional (or 2-D) growth. 2-D growth yields the highest quality epitaxy, but it

only works for similar crystalline structures, such as homoepitaxial growth or very

small mismatch heteroepitaxy (such as AlAs on GaAs).

In the case of the large mismatches between GaN and the available

heterosubstrates, 2-D growth cannot begin directly onto the substrate. Instead, an

alternate mechanism is proposed for the initiation and growth of GaN, involving the

nucleation of small strained 3-dimensional (3-D) islands.24,25 These 3-D islands are

each coherently aligned to the crystalline structure of the heterosubstrate, although the

two lattices may be rotated with respect to each other (as in the case of GaN and

sapphire, where there is a 30º rotation). Once the islands have nucleated, the epitaxial

growth commences as adsorbed reactants migrate to the energetically favorable sites

on the edges of the islands, making them grow in diameter, and height. The

translational periodicity of the rotated sapphire lattice is 14% smaller than the lattice

constant of GaN so it is expected that the GaN islands will be subject to biaxial

compression. This compressive stress can be partially relieved by threading

dislocations along the interface.25 Eventually, neighboring islands encroach upon each

31

other and the film coalesces into a mosaic-like continuous crystalline layer, with low

angle grain boundaries formed by arrays of threading dislocations.25,26

The quality of the resulting epitaxy depends on the density of islands initially

nucleating onto a bare heterosubstrate. If the density is too low, there is a likelihood

that adsorbed reactants may spontaneously form a misaligned or polycrystalline

nucleus between islands, resulting in non-epitaxial growth and poor quality. The

island density may be increased by the deposition of a nucleation enhancement or

buffer layer prior to the epitaxial growth. The purpose of this layer is to provide an

intermediate region with good crystalline registry to the substrate as well as to GaN.

While thin layers (<200 Å) of ZnO have been effectively used for this purpose,27,28 it

may desorb or contaminate the growing film at typical growth temperatures in excess

of 1000˚C. Alternatively, a thin 100-400Å layer of GaN29,30 or AlN31,32 has been

shown to be very effective without this limitation. The growth conditions for this GaN

nucleation enhancement layer require careful control. If the growth rate is too rapid or

the surface reaction kinetics too energetic, spontaneous nucleation of randomly

oriented islands will result, and the epitaxial film will be polycrystalline.

Nakamura29 discovered that the optimum condition for producing GaN

nucleation layers is to use a low temperature MOVPE process, typically 500ºC. At

this temperature the driving force for deposition is large, but the surface reaction

kinetics are sluggish, and the result is the slow growth of smooth but mostly

amorphous GaN film. As the substrate is heated to the growth temperature the

disordered film undergoes a solid-state crystallization process and self-organizes into

coherent 3-dimensional islands, as shown by the atomic force microscope (AFM)

32

images below in Figure 2.6. As deposited, the nucleation layer is flat and featureless.

After annealing in an ammonia atmosphere at 1025°C for 30 minutes the film self-

assembles into discrete islands ~50 nm in diameter and 2 nm high.

MOVPE is the preferred method for making the low temperature nucleation

layer because MOVPE is more readily adaptable to the low growth rate and low

temperature regime than HVPE. In MOVPE metalorganic compounds are piped into a

growth system from an external source held at or near 0ºC, whereas HVPE requires a

reaction between HCl and Ga metal at 850ºC, and this reaction must occur in

reasonably close proximity to the substrate in order prevent decomposition of the

GaCl. To deposit low temperature HVPE would require inverting the temperature

profile in the reactor, with the hottest zone upstream of the substrate. This leads to an

unstable situation where the growth rate and substrate temperature are difficult to

control precisely. For this reason, MOVPE nucleation layers are typically the standard

used for heteroepitaxial GaN growth. Even for HVPE growth the best films are

Figure 2.6. Left: atomic force microscope (AFM) image of a 250 Å GaN MOVPE nucleation layer after deposition; the film appears featureless. Right: image of the same film after 30 minute anneal at 1025ºC in NH3 ambient; 50 nm diameter GaN islands have formed with an approximate density of ~1010/cm2.

33

produced on substrates with prior-deposited MOVPE nucleation layers.1 This then

requires access to an MOVPE apparatus in order to provide optimal nucleation layers

for HVPE GaN.

2.4.6 The Hybrid VPE system

One approach to best utilize the capabilities of both an HVPE and MOVPE

system is to physically combine the two into a single hybrid reactor chamber, operable

in two modes, hot-walled and cold-walled, as shown in Figure 2.7, below. In MOVPE

mode the system uses a cold-wall heating system, such as an internal substrate heater,

and reacts metal-organic compounds with ammonia. When the HVPE mode is

desired, a hot-wall heating system (such as a tube furnace) is employed, and HCl is fed

over a gallium source to provide GaCl. For both modes, ammonia is fed into the

reactor separately, and the growth reaction occurs at the substrate.

There are some significant advantages to a hybrid system capable of switching

growth modes in situ. In such a system, for example, MOVPE can be used to provide

a nucleation layer on the hetero-substrate. Then, HVPE can be used to grow a thick

GaN layer utilizing the nucleation layer as a template and providing the bulk of the

GaN at a high growth rate with lower cost. After the HVPE, it may be possible to

switch back to MOVPE mode in situ, to grow the desired layers for device structures.

Because these growths can all be done in the same reactor, there is no need to cool

down and unload the substrate between steps, and there is no need for separate HVPE

and MOVPE reactors. This reduces the number of reactors necessary for the entire

process, increases the throughput of the reactor by eliminating transfers between

34

systems and waiting for substrates to heat up and cool down, and reduces the potential

for surface contamination that may occur while a substrate is moved between reactors.

Designing a hybrid system is more complicated than designing a standard

HVPE system. A hybrid system may be constructed by modifying an existing HVPE

system, adding additional input ports for metal-organic sources, for instance. A cold-

wall heating system compatible with the hot-wall heating system must be devised as

well; one such solution is the use of a resistive heater within the substrate holder or

susceptor. While these modifications are not trivial and require careful engineering

solutions, the cost savings realized by retrofitting an existing HVPE system instead of

acquiring a separate MOVPE system can justify the effort.

Figure 2.7. Model of a novel hybrid MOVPE/HVPE deposition system featuring hot-walled and cold-walled heating systems. In hot-walled mode, the tube furnace is used to heat the entire reactor for HVPE deposition. In cold-walled mode, the internal heaters in the susceptor provide the heating for only the substrate for MOVPE deposition

35

Chapter 3: Hybrid MOVPE/HVPE GaN process optimization

In this chapter, I will discuss the optimization of the hybrid MOVPE/HVPE

GaN growth process as a method to produce pseudosubstrates for subsequent device-

layer growth. To explore this in detail, it is necessary to discuss the metrics for

determining the quality of the material, from a morphological, microstructural, and

electro-optical perspective. Without the ability to consistently produce a smooth and

flat surface, further discussion of using GaN as a pseudosubstrate is moot. Only after

these metrics are achieved can the microstructural, electronic, and optical properties of

the epitaxial GaN be optimized for subsequent device overgrowth.

One of the fundamental issues related to the heteroepitaxial growth of GaN is

film stress induced by the interface between the heterosubstrate and the epitaxial film;

if this stress is not properly managed the epitaxial film can crack, buckle, and in

extreme circumstances peel away from the substrate. In the first part of this chapter I

will describe three origins of stress: coalescence-induced, lattice mismatch, and

thermal mismatch. Subsequently, I will discuss the effects of process parameters on

the evolution of these stresses and the morphological changes they bring about at a

microscopic and larger scale. At the end of the chapter a method for controlling

stresses to produce thick smooth and uncracked films will be presented.

3.1 Stress in heteroepitaxial GaN

One of the most important challenges in making thick or freestanding GaN

layers is the management of stress in the film. 33 As stress accumulates, it can exceed

the fracture limit of the film and substrate, leading to cracking and outright breakage.

36

Even at lower levels, stress can cause wafer bowing or warping, limiting the utility of

such a film for subsequent device growth or processing

Stresses arise during the initial nucleation and growth phases, they increase as

the film grows thicker, and they are affected by the thermal expansion mismatch

between the GaN and substrate as the wafer cools to room temperature. These effects

must be managed in order to reliably produce thick or freestanding layers. In the

following sections I will discuss the origins of these stresses and how they manifest

themselves.

3.1.1 Lattice mismatch stress

Much of the stress present in heteroepitaxially grown GaN is caused by the

lattice mismatch between GaN and the substrate. The room temperature in-plane

lattice constant of (0001) GaN is 3.189Å. Of the substrate choices mentioned in

Chapter 2, only lithium gallate, an orthorhombic material, has a similar lattice spacing

along its <100> direction. Silicon carbide, with the next closest spacing has a

mismatch of approximately 3.5%, which is comparable to the 4% mismatch between

silicon and germanium. For mismatch levels in this range, extensive networks of

dislocations can form at the interface to help partially relax the strain. As a result,

such films are mostly coherent, with periodically spaced regions of dislocations.

Sapphire has a lattice constant of 4.758Å in the (0001) plane, much larger than

the 3.189Å value for GaN. While this would appear to yield a mismatch of 49%, the

GaN lattice is rotated by 30° with respect to the sapphire, as seen below in Figure 3.1.

In this particular orientation, the densely packed <11�00> direction in sapphire, with a

37

translational period of 2.759Å, is parallel to the “a” <12�10> direction of GaN, and

vice versa. This makes the actual mismatch 13.9%, which while reduced, is still large.

With such a mismatch, initial GaN growth on sapphire occurs by way of 3-

dimensional island nucleation, with each island coherently aligned to the substrate. As

growth continues, the islands grow vertically and laterally, at roughly the same rate,

and eventually coalesce into a continuous crystalline layer.

Because the translational periodicity of the rotated sapphire lattice is smaller

than the lattice constant of GaN, it is expected that the GaN islands will be subject to

biaxial compression with respect to the sapphire. This compressive stress is partially

relieved by threading dislocations along the interface, where the islands meet to form

low angle grain boundaries. The stress effects of these islands meeting and joining

together will be discussed in the following section.

Figure 3.1. (Left) A schematic drawing of a GaN (0001) unit cell overlaid onto the (0001) sapphire unit cell. In this orientation, the [112�0] directions are coincident for both sapphire and GaN, and the resulting mismatch is large, approximately 33%. (Right) Rotating the GaN lattice 30º with respect to sapphire, so that the GaN [112�0] direction lies parallel to the [101�0] direction, the mismatch is reduced to a smaller value of 13.9%

38

3.1.2 Coalescence stress in GaN on sapphire

Owing to the large lattice mismatch between sapphire and GaN, initial

heteroepitaxial growth of GaN begins as coherent 3-dimensional islands nucleate

randomly onto the sapphire surface. The density of these nuclei is dependent on

factors such as the thickness of a nucleation layer, if any, the growth rate, substrate

temperature, and V/III ratio. Typically, the inter-island spacing is on the order of 500-

5000 Å.

These islands may be slightly misaligned, tilted and/or twisted, with respect to

the substrate. This misalignment is very slight, on the order of 0.01-0.1 degree, but it

gives rise to a tensile stress as the grains coalesce and the growth mode changes from

a 3-dimensional to a 2-dimensional mode. As adjacent grains grow together, they

create a very low angle grain boundary. The formation of such a grain boundary is

energetically favorable at the point of coalescence, as the energy of having two free

solid-gas surfaces is much higher than the energy of one low angle grain boundary.25,26

The effect of the surface energy reduction is most significant where the islands have a

high surface area to volume ratio; such is the case of GaN nucleation on sapphire. In

this situation, as the islands’ neighboring free surfaces begin to approach, the driving

force of the free surface area reduction can cause the islands to stretch slightly at the

interface to close the minute gaps between the grains. The stretching locally distorts

the lattice, putting it under tension; the energy required to cause this distortion is more

than offset by the reduction in free surface energy that occurs when the grains form a

low angle boundary. Thus, while the islands themselves may nucleate under

compression, as they coalesce, the boundary regions experience tension.26,33

39

Depending on the density of nuclei in the buffer layer, growth rate, V/III ratio,

temperature, and other factors, the thickness of the film at coalescence will vary from

0.1 – 1.0 µm. Once the film is fully coalesced, the film will begin to grow as a mosaic

structure of grains in 2-dimensional step-flow mode. As the film thickness continues

past the coalescence thickness, the tensile stresses at the grain interfaces become

distributed into a more uniform tensile stress over the entire film surface. The

stretching of the grains at the boundaries becomes spread throughout the entire film, as

a uniformly distributed tensile distortion. Subsequent growth on top of this tensile-

stressed layer does not allow for relaxation to occur, as each layer’s growth uses the

previous layer as a template; under these circumstances absent a relief mechanism

strained templates lead to strained films. Each atomic layer that grows is

approximately as stressed as the layer beneath it, and the overall stress within the film

increases with the thickness.33,34 This process is shown schematically below in Figure

3.2.

At some point, the accumulated strain energy becomes so great that cracks can

spontaneously form to relieve the stress. As GaN does not exhibit an efficient

dislocation glide mechanism,35 cracking is the primary method of stress relief. These

cracks form in-situ, at the growth temperature,33 and are not a result of thermal

mismatch stress. We have observed that these cracks occur at the growth temperature,

typically for films grown in excess of 2-5 µm. The actual maximum thickness that can

be grown before cracking occurs is dependent on such factors such as the growth rate,

grain size and spacing, growth temperature, and V/III ratio.

40

Figure 3.2. A schematic representation of the evolution of coalescence stress in heteroepitaxial GaN on sapphire. Nucleation of GaN forms compressively strained islands. As the islands grow together surface energy reduction induces a tensile strain at the interface, which propagates upward through subsequent layers.

3.1.3 Thermal mismatch stress

Another source of stress in thick heteroepitaxially grown GaN films arises

from differences between the coefficients of thermal expansion of GaN and sapphire.

At room temperature, GaN has an approximate linear coefficient of thermal expansion

of 5.6 x 10-6/°C, while sapphire’s is 33% greater, at 7.5 x 10-6/°C. Although the

coefficient of thermal expansion is not a constant and does vary with temperature,

41

throughout the entire range up to 1100°C sapphire will always expand and contract

more than the GaN layer. As the temperature drops, the sapphire contracts at a greater

rate than the GaN, and as a result the GaN becomes compressively stressed while the

sapphire develops a tensile stress.

Microscopic cracks in brittle materials such as GaN and sapphire are more

likely to propagate under tensile than compressive stress.36 As the cooled GaN film is

under biaxial compression, catastrophic cracking is inhibited to a degree. However, if

the GaN layer is sufficiently thick (typically greater than 100 µm), the tensile stresses

imposed on the sapphire substrate during cool down can cause it to crack and break

under tension. Once this substrate cracks, the crack usually propagates upward

through the epitaxial film resulting in total wafer breakage. 37

Figure 3.3. Illustration of the effects of the 33% mismatch in thermal expansion coefficient between sapphire and GaN. GaN is grown in a state of biaxial tension induced by the coalescence stress due to the lattice mismatch between GaN and sapphire. As the wafer cools, the sapphire contracts at a greater rate than GaN, resulting in convex bowing. If the stresses are too great, the GaN film can develop cracks, buckle, and delaminate from the sapphire.

It is interesting to note, however, that the coalescence and thermal stresses are

in opposite directions. While coalescence leads to tensile stress development in the

42

film, as it cools the sapphire substrate contracts more and imparts a net compressive

stress state onto the GaN. Empirically, it has been observed that these two stresses

appear to balance each other at approximately 600°C;38 at this temperature the

substrate lays almost flat, exhibiting minimal warping or curvature. The effects of the

thermal mismatch stress causing a reversal in the GaN stress state is graphically

illustrated above in Figure 3.3.

3.2 Effects of stress

As sapphire has a greater thermal coefficient of expansion than GaN, a

coherently-grown GaN film will be subject to biaxial compressive stress imposed as

the substrate and film cool down to room temperature. Conversely, the sapphire

substrate will be subject to a biaxial tensile stress imposed by the GaN film. To

minimize the overall strain energy under these imposed stresses, the wafer becomes

bowed, distorted such that the top surface of the epitaxial layer is bent to a convex

shape (as seen in Figure 3.3) When the film thickness is in excess of 50 µm, a

coherently strained 2” sapphire wafer may exhibit significant bowing; the center of the

wafer may rise 700 µm or more above the edges of the wafer. Often, with bowing this

severe we have observed that the wafer can crack spontaneously during storage.

Wafers that do not lie flat are difficult to process into devices; lithographic

techniques require a relatively flat substrate for accurate pattern transfer. Attempts to

hold flat or artificially flatten a bowed wafer for lithography are not likely to be

successful, because of the likelihood of wafer breakage under the applied force. With

43

this in mind, it is desirable to minimize the bowing of wafers with thick GaN layers, as

these have the greatest tendency to become distorted.

3.2.1 Cracking

In cases where the stresses exceed the yield-strength threshold, cracking will

occur. The nitrogen-gallium bond in GaN is significantly stronger than the gallium-

arsenic bond in GaAs. Comparing the Pauling electronegativities of the elements

involved, gallium has an electronegativity of 1.81, arsenic has 2.18, and nitrogen has

3.04. Many properties of a compound can be qualitatively inferred by considering the

nature of the bonding. When the electronegativity difference is less than 0.5, for

instance, the nature of the bonding is considered covalent, but non-polar. As the

electronegativity difference rises above 0.5, the nature of the bond becomes more and

more polar in nature. This means one element has a significantly stronger affinity for

the bonding electron than the other, so the electron tends to become more localized

around one atom. Comparing GaAs to GaN, the Pauling electronegativity differences

are 0.37 and 1.23 respectively. Under these criteria GaAs is non-polar covalent and

GaN is a polar material.

Practically speaking, this is significant when considering the ability of

dislocations to glide through a material to relieve an imposed stress.35 The glide

mechanism involves the sequential breaking and remaking of bonds around the core of

the dislocation. In order for a dislocation to slip through a material, the bonds around

the dislocation core must be broken, while the half-plane of atoms on one side of the

44

dislocation moves with respect to the other side. The dislocation sweeps through the

material, breaking and remaking bonds on either side of the core.

Breaking bonds to allow dislocation glide is an energetic process. There must

be a driving force to “push” the dislocation; this is in the form of the resolved shear

stresses originating from the biaxial stresses described previously. Resisting this

driving force is the energy barrier presented by the breaking/remaking mechanism; the

stronger the binding between the atoms, the greater the energy barrier to the making

and breaking of bonds. In this way, the physical process of dislocation slip can be

modeled as an activated process, pitting the shear driving force against the binding

energy barrier. Qualitatively speaking, the more polar the bonding in a material, the

less ability there is for a dislocation to glide across the bonds. Thus it can be

explained that in GaAs, dislocations may slip and allow stress relief, while this

mechanism is less favorable in a more polar material like GaN.35

While dislocation glide can be an effective way to allow material to plastically

deform, it is not the only way that stresses can be relieved. Another mechanism is

crack formation. The Griffith crack propagation model for brittle materials26,36 can be

a useful way to understand the mechanism. In its simplest and most basic concept, the

Griffith model states that a crack will propagate to relieve stress, and thus release

energy, only if the energy released is greater than the energy required to form the new

free surfaces on either side of the crack. The energy “cost” to produce the free

surfaces must be less than the energy “gain” given by the release of strain energy, or

the crack will not spread.

45

In a material with non-polar covalent bonding, such as GaAs, the dislocation

slip/glide mechanism has a sufficiently low threshold such that stresses may be

relieved in this way before they accumulate to the Griffith cracking threshold. In a

more polar-bonded material such as GaN, the dislocation slip mechanism has a much

higher stress threshold; in fact, this slip threshold exceeds the Griffith cracking

threshold, and thus cracks, rather than dislocation slipping, are more likely to form in

GaN to reduce the stress.

Figure 3.4. A comparison of tensile and compressive stress cracking mechanisms. a) At the onset of tensile stress cracking, a tiny crack forms at the GaN/Sapphire interface. b) As the crack spreads, two free surfaces are formed by the release of the localized elastic strain energy. c) Under compressive stress, the crack edges are pushed together. d) As the compressive crack propagates, the film must buckle away from the substrate, creating four free surfaces.

The Griffith model also explains why cracking more frequently occurs in films

under biaxial tension, and less frequently in films under biaxial compression. In a film

under tension, a crack allows the two separate sections to move apart, this movement

under applied stress translates into work (stress multiplied by the crack’s expanding

volume); this work acts to form two new free surfaces, as depicted in Figures 3.4a and

46

3.4b. Under biaxial compression, the stress acts to push the leading edge of the crack

together, rather than pull it apart. In order for the crack to propagate, the film must

buckle, breaking away from the substrate, creating four free surfaces in the process

(Figure 3.4d, and Figure 3.5, below). Energy is released by the buckling action, which

relieves the elastic compressive stress in the vicinity of the crack. Because

compressive cracking involves the formation of more free surface area than tensile

cracking, and there is always additional energy associated with the formation of free

surfaces, compressive cracks require a higher stress state before the onset of crack

propagation.

Figure 3.5. Plan-view micrograph of a 10 µm thick GaN film that has cracked and buckled during the cool down process

When the GaN film is under compressive stress, the sapphire substrate is held

under tensile stress. While a compressively strained GaN film may or may not

47

experience stress related buckling, the sapphire can, and frequently does crack under

these circumstances. Often, we have observed the sapphire substrate develop a dense

network of cracks along multiple directions, resulting in its near disintegration into

small millimeter-sized irregular pieces. In such circumstances, it is common for the

cracks, originating in the sapphire, to propagate upward through the epitaxial layer,

resulting in a broken substrate with a broken film. Thus, even under compressive

biaxial stress, cracking can have a detrimental effect on the growth of thick and

freestanding GaN layers on sapphire, albeit after growth during cool-down.

3.2.2 Peeling and delamination

The room-temperature interface between GaN and sapphire is the region of the

highest stress concentration and stress gradient. There is compressive stress in the

film above the interface, and tensile stress in the substrate below the interface, with the

maximum shear stress occurring directly at the interface. Under these conditions, the

interface between sapphire and GaN can fail, causing portions of the film to become

separated from the substrate, with parts of the film buckling upwards (shown

schematically in Figure 3.4d and the micrograph of Figure 3.5). This is a form of the

Griffith cracking discussed previously, with crack direction parallel to the substrate.

As before, the reduction in the stress-volume product of the newly formed crack

releases energy, some of which is expended creating four free surfaces: two are

formed parallel to the growth plane (between GaN and sapphire), and two are parallel

to the growth direction. Because more surface energy is required for this type of crack

48

formation, this process requires a greater driving force, in the form of higher stresses

before it can occur.

In practice, the problems of film delamination, peeling, and buckling can be

severe. Films grown at high temperature are coherent at the growth temperature, but

upon cooling down, the greater contraction of the sapphire imparts a compressive

stress onto the GaN film. In some instances the stresses are sufficient to make the

GaN film spontaneously self-separate.39 In our experience we have observed film

pieces up to 10 mm across to spontaneously spring free from the sapphire substrate.

While it might seem that this would present an opportunity for the formation of

freestanding layers, these delaminated pieces typically contain a high density of cracks

throughout, rendering them very fragile and difficult to utilize for device processing.

Additionally, it is difficult to reproducibly induce this sort of peeling over the entire

surface of a wafer, or even a part of it. As such, the pieces which spontaneously

spring free of the sapphire are of limited value; maintaining control of the sapphire-

GaN interface to prevent the cracking and peeling from occurring is a more favorable

approach in terms of process uniformity.

3.3 The surface morphology of HVPE GaN films

The biaxial stress imposed onto the GaN film at the sapphire interface can

determine whether or not the film is of suitable quality to be used as a substrate for

device layer growth (cracking or bowing will destroy its utility, even if other material

qualities are excellent). However, other figures of merit are necessary for comparing

the film quality to that produced by other methods, such as MOVPE. For instance, a

49

second important figure of merit for an epitaxial film destined to be used as a substrate

relates to the surface morphology. Specifically, if the surface is smooth and uniform,

it will provide a better platform for device layer growth than compared to a rougher

film with many topological features such as hillocks, microfacets, roughness, etc.

The as-grown surface of an HVPE film exhibits a characteristic roughness not

seen with thinner MOVPE films. While this is partially a result of the greater film

thickness, the size, scale, and amplitude of this roughness can also vary depending on

such parameters as growth rate, temperature, and V/III ratio. Surface features such as

hillocks, pits, oversized hillocks (“pimples”), and defects such as polycrystalline

inclusions can all affect the surface roughness and its suitability for use as a template

layer for device layer growth. In this section, the various manifestations of surface

roughness will be presented, and methods employed to minimize their effects will be

discussed.

3.3.1 Hillocks

The most characteristic features of HVPE-grown GaN films are hexagonal

pyramid-shaped hillocks,19,37,40,43 one example of which is shown below in Figure 3.6.

These hillocks may occur in various sizes, from approximately 20 µm across to

several hundred microns. Regardless of size, the hillocks all have some common

characteristics: the sloping sides of the hillocks usually appear smooth and featureless

when viewed under an optical microscope; the hillocks are higher in the center than at

the edges, with the thickness of the hillock at its center typically 1-3 µm greater than

50

the thickness at the edge; and hillock size distribution tends to be tightly centered

around an average size, which in turn depends on the growth conditions.

Figure 3.6. 100x optical micrograph of a 2 µm film grown at 1050ºC, showing presence of numerous hexagonal hillock-shaped prominences.

To the unaided eye, these hillocks are visible as a surface texture, similar to

that of an orange peel, on an otherwise smooth film. This is due to the scattering of

light off of the very low-angle hillock sides. Under the microscope, these hillocks are

clearly visible, although at higher magnifications (500x and greater) the hillocks are

frequently larger than the viewing area, so the surface morphology appears mirror-

smooth. While some hillocks appear to have hexagonal spiral steps on their sides,

most hillocks are smooth-sided, implying that any growth steps that are present are

51

much smaller than the wavelength of light. The sides of these hillocks are not true

crystal facets, however, as the angle of the sidewalls is very low, on the order of 1-5

degrees, and variable depending on growth conditions. From their hexagonal shape

however, it is clear that the sidewalls grow in a crystalline direction; the sidewalls are

a combination of (0001) surfaces and (101�0) steps,41,42 based on the hillocks’

orientation.

The hillocks are caused by a localized region of higher growth rate in the

center of the hillock, as compared to the edges.19,43 The cause of this higher growth

rate region is believed to be from the effect of threading dislocations penetrating the

growth (0001) surface, as elaborated by the Burton, Cabrera and Frank (BCF)

theory.23,24,44 Under typical 2-dimensional growth conditions, adsorbed atoms either

attach themselves to a nearby atomic step, or desorb from the surface. The “sticking

coefficient” of adsorbed species is thus partially dependent on the relative density of

nearby atomic steps; the greater the density of steps, the less likely an adatom will

desorb, and the more likely it will become incorporated into the film. Threading

dislocations penetrating the top surface of the growing film present a localized excess

of atomic steps, which in turn provide an excess of sites available for adatoms to

attach themselves. The steps surrounding a threading dislocation are in the form of a

spiral around its core, which promotes the characteristic spiral-shaped growth

morphology observed in such hillocks.41

According to the BCF model, higher temperature growth enhances hillock

formation by promoting desorption of adsorbed gallium atoms; the residence time an

adatom has on the surface is reduced as the temperature is raised. Adatoms that find

52

their way to a region with step sites will be readily incorporated, but those that do not

will desorb more rapidly. Thus, raising the substrate temperature during growth can

increase the overall texture of the growing film. Additionally, the BCF model predicts

that dislocation-driven growth is most significant for lower supersaturation levels in

the gas phase. Thus, low growth rates and low V/III ratios reduce the overall

supersaturation of GaN species in the gas phase, leading to preferential growth

wherever dislocations present an excess of surface steps. Therefore, to inhibit hillock

formation and growth, it is advisable to use a high V/III ratio, a higher growth rate,

and a lower substrate temperature.43,42

Hexagonal hillocks can occasionally occur in extra-large sizes, to the extent

that they are individually visible to the unaided eye. Hillocks such as these are

hundreds of µm across, and more than 20 µm high. When viewed with an optical

microscope, spiral growth steps are usually visible; at the very center of the spiral is a

small (1-5 µm) hexagonal depression. This depression is the location where multiple

threading dislocations penetrate the surface, each contributing an extra growth-

promoting step, which when combined leads to the greatly enhanced localized growth

rate at the center of the hillock.19

3.3.2 Pits

Naturally occurring pit-shaped defects appear in GaN grown on sapphire in

two major morphological types: hexagonal and irregular. These pits or localized

depressions are different from the pits that are revealed when the surface is chemically

etched, for instance, in a hot KOH solution.45 Etch pits are correlated to threading

53

dislocations penetrating the surface46, as the strain field surrounding the dislocation

core locally weakens the crystal structure,47 making it more susceptible to chemical

attack. On the other hand, these naturally occurring pits spontaneously form on the

surface of the film during growth, and are characteristic of the fine-scale morphology.

3.3.3 Hexagonal pits

By far the most important and common pit type, hexagonal pits are sometimes

referred to as V-pits or inverted hexagonal pyramids,48 because of their cross-sectional

shape, as seen below in Figure 3.7. These pits are hexagonal, with clearly defined

faceted sidewalls, and a width-to-depth ratio of approximately 1.4:1.

Figure 3.7. Cross-section optical micrograph of a hexagonal-shaped pit in a 24 µm HVPE GaN film. The faceted sidewalls have angles consistent with the {101�1} family of planes.

54

The origins of these pits are inversion domains, which form as the low

temperature nucleation layer undergoes a solid-state recrystallization prior to the onset

of HVPE growth.49 During this recrystallization, the mostly amorphous low

temperature layer self-organizes into small 3-dimensional coherently aligned nuclei on

the sapphire surface. Some of these nuclei, however, are inverted: instead of

presenting a Ga-polar (0001) GaN top surface, they instead present an N-polar (0001�)

surface. Because GaN is a polar material, these two surfaces have different atomic

bond coordinations with adjacent layers in the [0001] direction. These different bond

geometries affect the relative growth rates for the two directions; to understand why

this happens it will be helpful to describe the bonding geometry of GaN in some

further detail.

Figure 3.8. The basic tetrahedral bonding arrangement between Ga and N atoms in GaN. In this orientation, the [0001] arrow is indicating a Ga-polar direction, as there is one Ga bond pointing “up” with three bonds with “downward” components into the basal plane.

55

The symmetry of the bonding between the gallium and nitrogen atoms in

hexagonal GaN is tetrahedral; that is each gallium atom is bonded to 4 nitrogen atoms,

and vice versa. If, for instance, a gallium (or nitrogen) atom was placed at the center

of a tetrahedron, as above in Figure 3.8, the four nitrogen (or gallium) atoms

surrounding it would define its exterior vertices.

Figure 3.9. The GaN unit cell as viewed along the [101�0] azimuth. The bonding arrangement is such that along the [0001] direction there is a single parallel Ga bond and 3 N bonds projecting at 19.5° above the (0001) plane. This orientation is referred to as Ga-polar. In the opposite direction the situation is reversed: there is a single N bond parallel to [0001�] and 3 Ga bonds projecting at 19.5° above the (0001�) plane, for the N-polar surface. (The geometry of the [101�0] projection distorts the appearance of two of the three projecting bonds, making them appear to be at a larger angle than 19.5°).

If one considers the tetrahedron of a gallium atom surrounded by 4 nitrogen

atoms, the base of this tetrahedron, an equilateral triangle of 3 N atoms, would lie in

the (0001) plane. The vertical direction of this tetrahedron is the [0001] direction, the

56

growth direction for hexagonal GaN. When viewed along the [0001] direction from

the “top” of the tetrahedron, the Ga atom has one bond, pointing “up” (parallel to

[0001]), and 3 bonds with “downward” components at the tetrahedral angle (109.5°

from [0001]). When growth occurs in the [0001] direction, the Ga atom presents a

single perpendicular bond pointing out of the growing film’s surface, an orientation

referred to as Ga-polar. If this tetrahedron is inverted such that growth would occur in

the [0001�] direction, the nitrogen atom would present its perpendicular bond along this

[0001�] direction, an orientation referred to as N-polar. Figure 3.9, above, shows the

bonding arrangement for a full unit cell of GaN, as viewed along the [101�0] azimuth,

oriented with the Ga-polar direction pointing upward.

Like most crystal growth techniques, HVPE is done under group-V-rich

conditions. In this situation, growth on an N-polar surface proceeds more slowly than

on a Ga-polar surface. Whenever an excess of reactive nitrogen is present, the growth

rate depends on the attachment rate of Ga atoms onto available surface bonding sites.

If the surface is N-polar, the density of available sites for Ga attachment is just one

bond per N atom. However, when the surface is Ga-polar, the N atoms on the surface

present 3 tetrahedrally coordinated bonds, each capable of forming a shared bond with

a Ga atom. The probability of Ga attachment will be greater in the case of a Ga-polar

surface, as each adsorbed gallium atom will form 3 bonds with 3 nitrogen atoms on

the surface. This description is not complete, as it ignores effects such as surface

reconstruction during growth, but qualitatively it can explain the significant observed

growth rate difference between the two orientations.

57

As the HVPE growth commences, the adjacent Ga-polar (0001) faces grow

more rapidly than the inverted (0001�) N-polar faces. The proximity of fast and slow

growing domains on the face of the film does not necessarily lead to hexagonal pits,

however. Pits only form under certain growth conditions, specifically those favoring

vertical growth over lateral growth,41,42 such as a high growth rate and low growth

temperature. In this growth regime, the Ga-polar faces will grow rapidly upward, but

more slowly sideways. At the same time, the N-polar faces will grow upward more

slowly, resulting in a morphology consisting of high spots (thicker areas of Ga-polar

GaN) adjacent to these thinner N-polar areas. The sidewall facets of the pits are of the

{11�01} family, the permutations of which define the six separate facets (see Figure

3.7). This facet has fewer out-of-plane bonds available for growth, thus resulting in a

reduced relative growth rate compared to the (0001) direction.

To the unaided eye, a thick pitted layer looks anything but crystalline. The

dense network of pits scatters light, giving the surface a hazy sheen, with a color that

is beige, yellow, or brown. Under the microscope, things hardly look better; the stark

relief of the flat regions adjacent to the pits, compared to the depths of the valleys

inside the pits, indicates a truly rough surface not appearing suitable for further device

growth.

However, these pitted layers demonstrate a vastly reduced stress state. A thick

pitted film does not crack, or break, or peel, and the wafer bowing is greatly reduced

as well. While it can be challenging to grow a smooth layer greater than 5 µm thick,

pitted layers in excess of 100 µm are entirely feasible. This has enormous significance

when contemplating the fabrication of very thick, or freestanding, GaN layers to use as

58

substrates for subsequent device fabrication. The surface morphology must eventually

be changed for device-quality films, however.

These pitted layers exhibit significantly less residual stress based on the

geometry of the pits.48 The top surface of the GaN film, having a network of pits, is

an “open structure”; the film is not continuous and the shape and size of the pits can be

easily distorted under tensile or compressive stress. Thus, as the substrate contracts

more rapidly during the cool down from growth temperature, the epitaxial GaN film

can adjust to accommodate the shape of the sapphire substrate by distorting the

topmost surface, as depicted in Figure 3.10. This distortion slightly changes the

sidewall angle of the pit, rather than imposing a bulk biaxial compression to a

continuous film. The open structure is more compliant, strain imposed by a thermal

mismatch will induce lower stresses than would be the case for a continuous film; the

strain, while the same in both cases, has a correspondingly lower stress state in the

more compliant pitted film.

Figure 3.10. Schematic diagram of the distortion of hexagonal pits as a mechanism for strain relief. (Left) at the growth temperature (left) the pit angle α corresponds to the facet angle. (Right) as the wafer cools, the thermal mismatch strain is accommodated by a slight distortion of the pit angle.

To grow a pitted layer, it is necessary to promote a growth regime wherein the

lateral {11�01} growth rate is lower than the vertical (0001) growth rate.42 This can be

59

readily accomplished by any combination of lowering the growth temperature,

increasing the growth rate, or reducing the V/III ratio. All of these changes have the

effect of increasing the density and size of pits on the top surface. Typically, we have

grown pitted layers with a high growth rate, often greater than 50 µm/hr.

3.3.4 Irregular pits

On certain HVPE growths, the top surface of the film was “decorated” with a

network of very shallow, irregular shaped pits. This occurred when the growth rate

was low and the substrate temperature was high. Under these conditions the lateral

growth rate was high compared to the vertical growth rate, e.g. low growth rate and

low V/III ratio.

The morphology of these pits is irregular; instead of having a hexagonal

characteristic when seen from a plan-view, the edges of the pits meander in various

random directions. The width of the pits varies, from several microns to over 100 µm.

Interestingly, the depth of the pits is not related to the size; instead the pits are very

shallow, in cross section the typical depth was 1000-2500 Å.

To the unaided eye, such films with irregular pits appear hazy and slightly

brown due to the scattering of light from the edges of the pits. While the pits are

shallow, they are effective at relieving the stress, as no cracks were observed in such

films, even when grown to a thickness in excess of 50 µm.

The presence and density of these irregular pits was inconsistent from growth

run to run, leading to postulation that these pits might be related to the effect of some

sort of contamination during the processing or pre-growth sequence. To test whether

60

this is the case, we investigated new in-situ pretreatments. Identical growth runs were

made in rapid succession, on identically prepared sapphire substrates, using two

different in-situ surface pre-treatments as a variable. Prior to the deposition of the

nucleation layer, the sapphire was heated to a point above the HVPE growth

temperature where it was exposed to either inert nitrogen gas or ammonia.

Comparison of the post-growth film morphologies (Figure 3.11, below) shows that

preheating the sapphire in an inert nitrogen flow resulted in the complete elimination

of irregular pits, with a concomitant increase in film stress as exhibited by the onset of

cracking. However, preheating in an ammonia atmosphere resulted in inverted,

(0001�) N-polar films;50,51 their resulting morphology is full of truncated hexagonal

pyramids of various sizes and heights, with an overall rough surface. Molecular

nitrogen has a binding energy of 9.763 eV and is essentially chemically inert under the

conditions encountered in our growth system (atmospheric pressure, 1100°C

maximum). Because nitrogen treatment alone did not show this growth morphology,

we assume that the ammonia chemically altered the sapphire surface via a nitridization

process, changing the subsequent GaN epitaxial growth mode from Ga- to N-polar.

The discovery that irregular pits are eliminated by a thermal cleaning

procedure improved the run-to-run repeatability, yielding consistently mirror-smooth

films. Unfortunately these mirror-smooth films are highly stressed; cracking in GaN

grown on pre-baked sapphire occurs in films as thin as 2 µm when growth conditions

are such that lateral {11�01} growth predominates. Further investigation into the

precise cause of irregular pits and the nature of the stress relaxation mechanism could

yield valuable insight into the production of mirror-smooth thick crack-free layers.

61

Figure 3.11. The effect of surface thermal pretreatment prior to deposition of the low temperature MOVPE layer. a) Control, no thermal pretreatment used, irregular pits are present. b) Thermal pretreatment in nitrogen ambient. Surface has no pits and is much smoother, but increased stress has caused extensive cracking. c) Thermal pretreatment with ammonia ambient. Truncated hexagonal pyramids are characteristic of N-polar GaN.50 All samples have 5 µm HVPE GaN grown under identical conditions after surface pretreatment.

3.3.5 Quantifiable roughness measurement

The observation of roughness and thickness variation over different length

scales on the as-grown HVPE surface leads to the need for quantifiable methods for its

characterization. The measurement of roughness, or height variation, can be done

over a fine scale of less than 1 µm with atomic-step resolution using an AFM, or it

may be done over a larger scale of hundreds of microns using an optical microscope

viewing in cross-section. In either case, different values may be obtained for the same

sample.

For example, a pitted sample may have large regions of flat material between

pits. In some instances, the spacing between pits may be tens of microns, but more

typically the spacing is less than 10 µm. An AFM scan of such a flat region will not

detect that the surface is roughened by pits, although it may detect growth steps on the

planar regions between pits. Similarly, an AFM scan of a sample with hillocks may

62

detect the presence of atomic steps on the surface, but will not give much useful

information about the overall height and variation of the hillocks.

The longer-scale height variation or roughness takes on importance when a

thick or freestanding GaN layer is desired as a substrate for device growth. A surface

that has a large thickness variation due to hillocks will present an uneven template for

further layers, possibly degrading device performance. Additionally, height variations

in the film will degrade the ability to perform precision lithographic and other

fabrication processes. Similarly, a film with pits has a discontinuous surface, leading

to discontinuities in any device structure. As such, the quantification and ultimate

reduction of longer-scale (microns to millimeters) roughness is important for the

growth of thick and freestanding layers. A simple technique for this kind of

characterization uses cross-sectional optical micrographs, with thickness

measurements taken at regular intervals across the photograph.

The optical microscope we used for cross-sectional analysis has two

magnifications, 100X for analyzing thicker films, and 500X for analyzing thinner

films or films with higher spatial frequency features. A 1 megapixel CCD camera is

attached to the microscope photo mount for digital image acquisition. The micrograph

is then overlaid with a comb-shaped template; each “tooth” of the comb indicates one

of the 16 positions on the photograph where thickness measurements are made, as

shown in the examples below in Figure 3.12. The spacing between measurement

positions on the comb is uniform; for 100X photographs the spacing is 40 µm, for

500X it is 8 µm. In each photograph, 16 measurements are taken, and statistical

63

analysis of the measurements yields average thickness and standard deviation (as RMS

roughness).

Figure 3.12. (Top) 500x plan-view micrographs of pitted (left) and smooth (right) HVPE GaN films. (Bottom) respective cross-section micrographs with comb overlay and measured thickness values for each “tooth”. The rougher-appearing pitted film has a thickness variation at least 7 times greater than the smooth, cracked film at this lateral scale level.

Analysis of pitted samples indicates that the typical RMS roughness of a

thicker (> 40 µm) film is on the order of 5%, meaning that the measured thickness of

the film at any given location may vary by as much as ±5% of the average. For a

spatial sampling frequency of 8, or even 40 µm, such a variation is unacceptable for a

64

device quality substrate; additional surface treatment (i.e. polishing) would be

necessary to planarize the film. In contrast, a very smooth film with gently sloping

hillocks has an RMS roughness less than 1% of the total thickness, while some films

have no detectable variation.

With a quantifiable method for determining the film roughness or thickness

variation of a given sample, it is possible to measure the relative success of methods

for producing smoother films. With this in mind, the following section will discuss

selected methods and their relative success at producing smoother, flatter films.

3.4 Effects of substrate temperature, growth rate and V/III ratio

Growth conditions that change the mobility and residence time of Ga adatoms

on the surface affect the resulting film morphology. The longer a Ga adatom has to

diffuse on the surface before being fixed in place by a nitrogen atom, the greater the

probability that it will find a lower-energy position with multiple bonds, such as an

atomic ledge or kink site. Under these conditions islands in the nucleation layer grow

together rapidly and any slow-growing grains are overgrown. Conditions that reduce

the residence time for unattached Ga adatoms lead to shorter surface diffusion lengths

and a more columnar structure, and result in a high density of smaller grains.

Lattice vibrations from the substrate, in the form of heat, provide the energy to

move adsorbed Ga atoms around and away from the surface. The effect of the

substrate temperature in a 2 µm film is shown below in Figure 3.13. At the low

temperature of 1000ºC, adatom mobility is reduced and the hillocks that form are

smaller (less than 10 µm) and numerous. At the higher temperature of 1050ºC,

65

significantly greater Ga adatom surface diffusion length reduces the number of

hillocks and greatly increases their size to well over 100 µm across.

Figure 3.13. The effect of substrate temperature on resulting surface morphology on 2 µm HVPE films. At 1000ºC Ga adatom mobility is low, leading to the formation of numerous small hillocks. At 1050ºC the Ga adatoms have higher surface diffusivity leading to fewer but larger hillocks.

Figure 3.14. The effect of growth rate on 2 µm thick HVPE GaN films grown at 1025ºC. At the low growth rate of 5 µm/hr the surface is smooth but under stress, leading to cracking. The high growth rate film is rough with small hexagonal pits, but the overall stress state is far lower as there are no cracks evident.

66

The rate at which the Ga atoms arrive at the substrate surface (i.e. the growth

rate) also affects the morphology. Low growth rates show no preference for vertical

(0001) growth over lateral {11�01} growth; there is no appreciable facet-dependent

growth rate variation. High growth rates exhibit greater vertical (columnar) growth,

leading to enhanced faceting and the formation of hexagonal pits. This effect is shown

in the micrographs of Figure 3.14 where 2 µm thick films show markedly different

surface morphologies when the growth rate is raised from 5 µm/hr to 50 µm/hr. The

low growth rate film is smooth but cracked, indicating a highly stressed condition,

while the pitted high growth rate film is uncracked.

Figure 3.15. The effect of V/III ratio on the surface morphology of 2 µm HVPE GaN films grown at 1025ºC. The lower ratio film exhibits larger hillocks as a result of higher Ga adatom surface mobility. Under high V/III ratio the adatoms have shorter diffusion lengths and the hillocks are correspondingly smaller.

For HVPE growth the V/III ratio is determined by the flow rates of ammonia

and HCl, presuming that a well-designed GaCl production cell will have near-unity

conversion efficiency of HCl to GaCl. Figure 3.15 above shows the effect of raising

the V/III ratio for 2 µm films grown at 1025°C and a moderate growth rate of 12

67

µm/hr. At the lower ratio of 80:1 Ga adatoms have more mobility and migrate across

the surface to the highest growth rate hillocks, resulting in a larger grain structure.

Raising the ratio to 300:1 inhibits Ga adatom migration and results in smaller tightly

packed grains.

To summarize, the effects of the growth conditions on Ga adatom diffusion

determine to a large part the resulting surface morphology of the film. Conditions that

increase the diffusion length of Ga adatoms (increased substrate temperature, reduced

growth rate, and reduced V/III ratio) enhance the 2-dimensional lateral growth mode,

leading to rapid grain coalescence and smoother surfaces with larger grains.

Unfortunately, these conditions also maximize the stress state of the film, and cracking

is a frequent result in otherwise high quality films. Conditions that reduce the Ga

adatom diffusion (decreased substrate temperature, high growth rate, and high V/III

ratio) have the effect of enhancing the columnar growth mode, producing films that

are rougher with hexagonal pits and smaller grains. This roughness allows for some

stress relaxation however, and these films can be grown far thicker than smoother ones

without cracking.

3.5 smoothing layer growth

By its nature, a rough pitted layer is not useful as a substrate for device growth.

The large variation in surface height and discontinuous drop-offs adjacent to the pits

makes the growth of a smooth device layer directly on this surface difficult. However,

under conditions favorable for enhanced 2-dimensional lateral HVPE growth of GaN,

the sidewalls of the pits will grow together, in effect filling in the pits from the sides

68

with Ga-polar material.52 The N-polar inversion domain at the bottom of the pit is

buried under the sidewall growth of the adjacent Ga-polar material. Thus, it is

possible to fill in a pitted surface by changing the growth conditions from vertical

(columnar) to lateral (2-dimensional). To do so we use growth conditions optimized

for 2-dimensional HVPE growth: by using a combination of higher substrate

temperature, lower growth rate and lower V/III ratio.

Figure 3.16. The effect of smoothing layer regrowth on a pitted film. Top row: plan-view 100x, 500x, and cross-section of the as-grown pitted film. Middle row: after 5 µm of smoothing layer growth, the pits are largely filled in but the surface has significant roughness remaining. Bottom row: 11 µm of smoothing layer regrowth has filled in the pits and largely smoothed the surface, although some growth hillocks remain.

69

Following a pitted layer growth, a smoothing layer can be formed by lowering

the growth rate, increasing the growth temperature and/or decreasing the V/III ratio.

This can be done in-situ, immediately after the growth of the rough pitted layer, or it

can be done later, as a regrowth after some intermediate processing, such as substrate

removal for instance. In either circumstance, the regrowth proceeds as laterally

enhanced growth, which preferentially fills in the hexagonal pits on the surface.

The actual reduction in the surface roughness has been quantified by direct

observation of a pitted film, followed by successive smoothing layer regrowth steps,

as shown above in Figure 3.16. While the Figure shows that regrowth significantly

improves the surface morphology and fills in pits, it does not result in a perfectly flat

surface. Other features remain, or become dominant in their absence, especially

localized pyramidal hillocks. We refer to our novel process as the 2-step growth

method.53

By filling in the pits at the top surface, the smoothing layer reduces the

compliance of the pitted layer, increasing the biaxial stress levels in the film as the

wafer cools down. Thus, a thick pitted layer that would not crack during cool-down

can crack when a smoothing layer is applied. This is a direct result of the smoothing

layer, as cracking has been observed in uncracked pitted films that have had

smoothing layers regrown on subsequent growth runs. The overall stress state is still

apparently lower than for similarly thick layers grown in a single-step process, as we

were unable to produce uncracked thick films without using the 2-step growth process.

This increase in stress is not observed when the regrowth is done on freestanding

70

pitted films, as there is no thermal mismatch when GaN is grown on GaN. This

provides a possible avenue for the formation of thick freestanding GaN layers.

3.6 The 2-step growth process

To produce GaN films with thickness in excess of 15 µm we have developed a

technique called the two-step growth process. In the first step a pitted layer of good

crystalline quality is made. Because of the extra lateral surface area in the rough film

it can be grown thick without cracking. To complete the growth a thin, smoothing

layer is added to planarize the film. In Figure 3.17, below, a schematic of the process

is shown along with scanning electron microscope (SEM) images of the rough layer

and the finished two-step growth with the high-quality smoothing layer.

Figure 3.17. The two-step growth process to achieve GaN film thickness in excess of 15 µm. (a) A schematic of the two-step process in which a rough, but high quality GaN layer is made, after which a smoothing layer is deposited. (b) Scanning electron microscope (SEM) image of the rough layer, here grown to slightly less than 40 µm thick. (c) SEM image after the smoothing layer has been applied, where the total film thickness is now approximately 40 µm.

71

Single-step planar growth in excess of 15 µm will nearly always result in a

high density of cracks, while under carefully controlled conditions 40 µm planar and

crack-free samples can be produced, as seen below in Figure 3.18. While the regrowth

on pitted layers is effective at filling in the pits, several challenges arise from the

process. First, the pit filling-in does not perfectly planarize the film’s surface. As the

pits fill in, growth hillocks become predominant, leading to an imperfect smoothing

situation. Additionally, the elimination of the stress-relieving pits on the surface leads

to increased stress levels within the film, limiting the crack-free film thickness.

Figure 3.18. 500x cross-sectional optical micrograph of a 40 µm film grown by the 2-step method, with comb overlay showing thickness variation measurements. The data for 3 different readings in different locations across the wafer are presented, showing that the overall thickness variation is held to the order of a single percent of total thickness.

One possible solution to the stress cracking problem is to remove the pitted

film from the sapphire substrate before growing the smoothing layer. Without the

added strain from thermal coefficient mismatch, the driving force for cracking during

72

cool-down is eliminated, and smoothing layer growth will not lead to breakage.

Removing a low stress pitted layer from the sapphire substrate is also significantly

easier than attempting to remove a higher stress smoothed layer. During the removal

process, the stress state between film and substrate undergoes continuous change; to

reduce the risk of breakage, it is best to have a minimal stress condition at the outset.

3.7 Summary of GaN deposition process optimization techniques

The surface morphology of HVPE grown GaN is highly dependent on the

conditions at the commencement of growth on the recrystallized nucleation layer.

Lower temperature growth with higher growth rate enhances the columnar growth

habit, resulting in films with a hexagonal-pitted morphology. These films are rough,

but partially relaxed and do not exhibit cracking. At lower growth rates and higher

temperatures, 2-dimensional lateral growth is enhanced and the films tend to have

smoother surfaces with larger sized hillocks and grains. This comes at the cost of a

much higher stress state with frequent cracking at film thicknesses greater than 2 µm.

The two different regimes can be switched at will, and the growth mechanism

will change as a result of changing growth conditions. In this way a rough pitted film

can be made smooth by switching to a lateral-enhanced mode, and this is the basis for

the 2-step growth technique that has allowed us to routinely achieve smooth uncracked

films as thick as 40-50 µm, and in some cases in excess of 100 µm. The smoothing

growth does not have to occur immediately after growing the rough layer; layers

deposited on previously-grown rough films are as smooth as films produced in a single

growth run. This opens up the possibility of producing thick freestanding GaN

73

substrates by first growing pitted thick layers, removal of the substrate from the film in

an external process, and finishing the rough freestanding films with a subsequent

smoothing layer growth.

74

Chapter 4: Microstructural Characterization of VPE-grown GaN

4.1 Structure of thin GaN layers

Having demonstrated how to produce smooth uncracked GaN layers in the

hybrid VPE system, the question of the material quality at a microstructural level

comes to the fore. Defects arising from the heteroepitaxial growth of GaN, especially

threading dislocations, can propagate upward into subsequently overgrown

homoepitaxial device layers. Threading dislocations are known non-recombination

centers in GaN54,55 and act like charged scattering centers reducing carrier mobility56

and increasing device leakage current.57 In order to evaluate the crystallinity and

suitability of VPE-grown GaN for use as a device growth substrate, it is important to

use characterization methods such as X-ray diffraction58 and transmission electron

micrography (TEM)59,60 to peer into the microstructure. In this chapter I will

demonstrate that the material grown via the hybrid VPE process produces material that

is of high crystalline quality, directly on sapphire. I will also present a way to

quantitatively estimate the threading dislocation density, and will demonstrate that this

density decreases with increasing film thickness.

To begin, an X-ray diffraction (XRD) ω-scan (“rocking curve”) of the (0002)

crystallographic reflection on a 1.5 µm thick GaN layer is shown below in Figure 4.1.

Rocking curve scans are done by changing the incident beam angle around a particular

diffraction condition, while keeping the detector stationary and fully open. In this

way, small variations in the plane spacing or alignment due to tilt or strain can be

detected.58

75

Figure 4.1. High-resolution X–ray diffraction (XRD) ω-scan of a 1.5 µm thick GaN film grown in the hybrid VPE reactor. Using a linear scale on the left, the full-width, half maximum (FWHM) of the XRD peak is 317 arcseconds. When plotted using a log scale on the right, it is apparent that this is not a simple Gaussian peak. When fitted to a single Gaussian, the best fit has a FWHM of 105 arcsec.

The full-width at half-maximum (FWHM) of this reflection is 317 seconds of

arc (arcsec) according to the raw data, the linear scale scan on the left of the Figure.

This linewidth compares favorably to similar-thickness material grown by MOVPE43

and HVPE.61 The dominant broadening mechanism in these peak reflections is due to

inhomogeneous variations in structural properties such as the plane spacing, strain

fields and mosaic-like grain misorientation. Thus, the shape of this XRD reflection

should be expected to be Gaussian in nature. However, when plotted using a

logarithmic scale as on the right side of Figure 4.1, it is clear the data does not fit a

single Gaussian curve, as there are wide shoulders indicating the superposition of

multiple peaks of different widths. When fitted to a single Gaussian peak, the largest

76

(0002) XRD reflection has a linewidth of only 105 arcseconds, indicating this material

contains a very high quality component, presumably that furthest from the sapphire

interface.

Figure 4.2. High resolution (0002) XRD Gaussian curve fits to the data shown in Figure 4.1. The 2-Gaussian model on the left fits the center peak well but displays a wide shoulder that is accounted for on the right by a 3rd low intensity (0.03%) Gaussian with much larger linewidth.

While the majority of the XRD peak reflection can be fitted to a single

Gaussian, there are other components to the (0002) reflection. If these components are

distinct, we can fit them with additional Gaussians. This is done above in Figure 4.2;

on the left side of the Figure two Gaussians are used to fit the XRD data. An

improvement can be seen when comparing to the single Gaussian peak fit in Figure

4.1. There remains a small background that is not contained by the 2-Gaussian fit,

however. We can include this in the fitting process by adding an additional, 3rd,

Gaussian. The 3-Gaussian fit is shown on the right side of Figure 4.2 and the match

with the X-ray scan is very good.

77

Since the film thickness is small enough to neglect absorption losses, we can

roughly estimate the thickness of material contributing to each Gaussian component

by comparing the relative amplitude of the 3 fitted peaks. The primary peak on the

right side of Figure 4.2 comprises 85% of the total amplitude with a linewidth of 103

arcseconds; indicating that the majority of this film is of high quality. The secondary

peak has a linewidth of 228 arcseconds with 15% of the relative amplitude, showing

the presence of greater disorder in a modest fraction of the film. The third fitted peak

has a linewidth of 659 arcseconds, but with a very weak relative intensity of 0.03% it

is a very minor component of the total.

4.2 A 3-zone layered growth model

From the analysis associated with Figures 4.1 and 4.2, the XRD pattern can be

accurately modeled with a 3-Gaussian curve fit. The data does not indicate whether

these three components are mixed in a single uniform phase, or if they are layered

vertically. While the uniform mixed model may be correct, there is more evidence for

a three-layered explanation.49 Because there is a low-temperature nucleation layer of

GaN that crystallizes at higher temperature, it makes logical sense that there is a very

thin rough layer where the GaN meets the sapphire. The small background has a

relative amplitude ratio of only 0.03% so it contributes an equivalent thickness of only

4.5 Å for a 1.5 µm thick film. Thus, the layer is less than a unit cell of GaN and

would constitute the interface region between the GaN and sapphire. The next region

contains the original crystallites and the region immediately above where the film

coalesces and the tilt and twist variations of the crystallites are accommodated. It

contains about 15% of the amplitude or roughly 0.

The uppermost layer has the highest quality and constitutes the thick

film, 1.3 µm. The model assumes that as more material is deposited the material

quality improves. This follows up to a point when the lattice and thermal mismatch

stresses begin to degrade the film. For our growth this occurs

on growth conditions.

the zones appear distinct in Figure 4.3, they are only depicted this way for descriptive

purposes, and if the model is correct, there are likely transition regions in between

each zone.

Figure 4.3. Schematic growth on a sapphire substrate. The loweand sapphire has a disordered crystal structure.low angle grain boundaries with dislocations at their interfaces, shown by the dense network of dark dislocations in the middledislocations merge and annihilate forming the high quality layer with comparatively few defects.

While we consider these films thin, because our hybrid growth process is tuned

to high growth rates and thick epitax

growth.29 These films are high quality by any standard, even though the growth rate is

in excess of 10 µm/hr.

crystal quality can be maintained or improved with increas

78

5% of the amplitude or roughly 0.25 µm of the 1.5 µm thick

layer has the highest quality and constitutes the thick

µm. The model assumes that as more material is deposited the material

quality improves. This follows up to a point when the lattice and thermal mismatch

grade the film. For our growth this occurs above 2 µm,

on growth conditions. Figure 4.3, below, summarizes the three-zone model. While

the zones appear distinct in Figure 4.3, they are only depicted this way for descriptive

e model is correct, there are likely transition regions in between

Schematic (left) and TEM bright field image (right) of the three-zone GaN sapphire substrate. The lowest zone at the interface between the GaN has a disordered crystal structure. As the crystallites coalesce they form

low angle grain boundaries with dislocations at their interfaces, shown by the dense network of dark dislocations in the middle zone on the right. As growth continues, dislocations merge and annihilate forming the high quality layer with comparatively

While we consider these films thin, because our hybrid growth process is tuned

to high growth rates and thick epitaxy, the films are of typical thickness for

films are high quality by any standard, even though the growth rate is

µm/hr. In the next section, we will address the question of whether the

crystal quality can be maintained or improved with increasing layer thickness.

25 µm of the 1.5 µm thick film.

layer has the highest quality and constitutes the thickest part of the

µm. The model assumes that as more material is deposited the material

quality improves. This follows up to a point when the lattice and thermal mismatch

2 µm, depending

zone model. While

the zones appear distinct in Figure 4.3, they are only depicted this way for descriptive

e model is correct, there are likely transition regions in between

zone GaN zone at the interface between the GaN

As the crystallites coalesce they form low angle grain boundaries with dislocations at their interfaces, shown by the dense

zone on the right. As growth continues, dislocations merge and annihilate forming the high quality layer with comparatively

While we consider these films thin, because our hybrid growth process is tuned

y, the films are of typical thickness for MOVPE

films are high quality by any standard, even though the growth rate is

In the next section, we will address the question of whether the

layer thickness.

79

4.3 Structural improvement with increasing thickness

Figure 4.4. Mosaic of images of a series of TEM micrographs taken of a 12 µm thick hybrid VPE-grown GaN film. The cross-sectional bright-field images are recorded with g = [112�0] detecting dislocations with Burgers vector b = a. Dislocations appear as thread-like lines which are increasingly entangled and/or bent over into the basal plane as the film grows away from the interface with the sapphire substrate.

Figure 4.4 above shows a series of successive cross-sectional bright-field TEM

images of a 12 µm film. This particular film was prepared in a single HVPE growth

step on sapphire with a recrystallized MOVPE buffer layer, using a moderate growth

rate of 25 µm/hr. The film was mirror-smooth and uncracked. Starting at the sapphire

interface, there is a dense tangle of dislocations, denoted by the dark thread-like lines.

As the film grows away from the interface, the dislocations follow curved paths and

80

some are bent over into the basal plane, parallel to the growth direction – effectively

ending their propagation upward. The interaction, entanglement, and annihilation of

threading dislocations happens efficiently as growth proceeds; within approximately 1

µm the dense tangle of dislocations is dramatically diminished. This effect continues

with increasing thickness; the counted dislocation density drops into the range of

107/cm2 near the surface.

While TEM analysis can be used to directly count threading dislocation

density, the difficulties of sample preparation make this impractical. An alternative

method can be devised using a series of symmetric (0002) and asymmetric (202�1)

high resolution XRD (HXRD) ω-scans, which will be discussed subsequently.

4.4 X-ray methods for determining approximate dislocation density

The linewidth of a high resolution ω-scan is representative of the amount of

crystalline disorder present over the sampling depth of the impinging x-ray beam.

Another effect that contributes to the broadening of the linewidth is residual strain,

which increases with the film thickness. The competing residual stresses between the

compressively strained GaN and the sapphire under tension induces a physical

warping (bowing) of the as-grown wafer; this bowing-induced strain elastically varies

the lattice constant. To extract useful information about the disorder caused by

threading dislocations, it is necessary to decouple that effect from the broadening

caused by wafer bowing.

81

4.41 Accounting for the effect of bowing on XRD linewidth

The linewidth broadening effects of an x-ray scan of a bowed wafer surface are

geometrical in nature, and depend on the spot size of the beam on the surface.

Consider figure 4.5, below, which schematically shows an x-ray diffraction situation

for a bowed sample:

Figure 4.5. Schematic diagram of the X-ray diffraction geometry for a bowed substrate with bending radius r. The X-ray beam has a non-zero slit width angle α, which, when projected at angle ω onto the surface illuminates an area with lateral dimension l. The wafer curvature over beam width l changes the local orientation of crystalline planes by an angle up to ±∆ω/2 (∆ω << ω, the diagram exaggerates for visual clarity). This expands the available ω values that allow the diffracted beam to reach the detector, resulting in a broadening of the XRD peak.

Because α is usually 0.5º or less, we can use the small angle approximation to

extract l, the X-ray spot size as projected onto the surface at an angle ω:

l = �����

α

�

���(ω) ≈

��(α

�)

���(ω) ≈

�α

���(ω)

82

This larger-than-infinitesimal sampling length of a curved surface leads to a

natural broadening mechanism where bent crystal planes will diffract around the ideal

diffraction condition at ω, offset by a much smaller angle ∆ω. Using the small angle

approximation again, spot length l can be similarly expressed in terms of the wafer’s

bowing radius r, and the broadening ∆ω:

l = r sin(∆ω) ≈ r∆ω

Using these approximations, there is a linear relationship between slit angle α

and curvature-induced broadening ∆ω:

r∆ω ≈ l ≈ �α

���(ω)

∆ω ≈ �α

����(ω)

This is significant because a series of x-ray scans can be done with

successively smaller slit angles, and a plot of FWHM vs. slit angle will show a zero-

intercept value that should correspond to the intrinsic broadening due to other forms of

crystalline disorder (such as dislocations).

An example is shown for a 12 µm thick GaN layer grown on sapphire in

Figure 4.6, below. In this Figure, the variation in linewidth (actually FWHM) of scans

from two different diffracting planes are plotted with respect to decreasing slit angle,

with extrapolation to the zero-angle slit (i.e. intrinsic) condition. The FWHM

decreases with slit angle for both the symmetric (0002) and the asymmetric (202�1)

scans. Because the (0002) plane is the basal plane, intrinsic broadening is caused by

variation in the value of the c-axis lattice constant, which is correlated with tilted grain

boundaries and the pure screw dislocations that comprise them, with Burgers vector b

parallel to the [0001] direction.62 The asymmetric (202�1) diffraction plane is tilted 62˚

83

with respect to (0002), and it contains components of tilt and twist in the low angle

grain boundaries, with twist correlated to edge dislocations with Burgers vector b

along the [112�0] direction, entirely within the basal plane. The linewidth broadening

effect due to slit width variation is noticeably stronger with the basal (0002) scans than

it is with the higher-angle (202�1) scans, as would be expected for diffraction on planes

parallel to the imposed bending moment instead of at a 62º angle to it.

Figure 4.6. The variation in full-width half maximum (FWHM) X-ray diffraction linewidth with X-ray source slit width for a 12 µm thick GaN layer on sapphire. Symmetric (0002) and asymmetric (202�1) diffraction scans have been extrapolated to zero slit width, indicating the bowing-free linewidth.

One point should be made about the fitting of the X-ray diffraction data.

Gaussian fits were used in all cases, since dislocations, grain boundaries, etc. are

expected to contribute to inhomogeneous broadening. The full-width half maximum

of the Gaussian is not a fundamental property but can be easily derived. The Gaussian

formula is:

84

���� = �

√��(���)�

��� .

Here, µ is the mean and σ the variance from the mean. The amplitude �

� √�

normalizes the integration to 1. To solve for the FWHM we set P(x) to 0.5 times the

normalization, µ to zero for convenience, and solving for 2(x − μ)� yields:

FWHM = 2x = 2σ �−2 ���� = 2σ √22

The approximate value of√2ln2 is 1.1774, so the difference between FWHM

and 2σ is a factor of about 18%, an amount small enough to neglect when comparing

orders of magnitude differences in dislocation density.

4.4.2 The relationship between FWHM and dislocation density

While the lattice disorder caused by the presence of dislocations can induce

broadening in X-ray diffraction scans, it is not the only reason for broadening. In his

study, Ayers63 discloses that the total XRD linewidth Δω����� is the root of the sum of

the squares of individual components from various sources:

�ω������ = �ω������

+�ω����� +�ω���� +�ω�����

In the above relation, Δω����� represents the intrinsic linewidth broadening for

the particular diffractometer and crystal setup. Typically this is very small, fewer than

10 arcseconds, which is at least one order of magnitude smaller than the total

broadening. The broadening due to dislocations, Δω����, is induced by lattice

distortion tilting and twisting locally around the dislocation core, as well as the

extended strain field further away through the lattice. The magnitude of this effect is

85

inversely proportional to the average spacing between dislocations, which is related to

the dislocation density D as follows:

�ω����~ 1{� − �������}���������������������� = 1 �

√�� = √�

Of the remaining two factors contributing to XRD linewidth broadening,

�ω��� is the intrinsic broadening effect caused by the crystal size, i.e. film thickness.

Its magnitude is inversely proportional to the film thickness, and amounts to fewer

than 20 arcseconds for a 1.5 µm film so it is not a significant source of broadening and

can be usually neglected, especially for thicker films. The effect of wafer curvature,

�ω����, is significant – especially with the more highly stressed thicker films, but this

component can be removed from the equation by extrapolating to the zero-width

FWHM value as described in Section 4.4.1 above.

Using these approximations we can simplify the previous root sum equation:

�ω� ������ = ������ + �ω������

����� represents a synthetic proportionality constant linking the dislocation

density and dislocation-induced broadening. If the effect of Δω!"#$% is discounted,

K&$'� may be may be estimated by correlating a particular X-ray linewidth to a counted

dislocation density as determined by using another method such as TEM analysis64,65

or etch pit density.47,66 This gives us a qualitative guide to determining the dislocation

density based on square of the magnitude of the X-ray linewidth.

4.43 Reduction in the dislocation density with increased thickness

By plotting the variation of the square of the magnitude of the zero-slit

extrapolation of the FWHM for different film thicknesses, we can estimate the

86

reduction in dislocation density with increased layer thickness. This is shown below

in Figure 4.7, using two different XRD reflections.

Figure 4.7. The variation of the square of the magnitude of the zero-slit extrapolation FWHM (i.e. dislocation density) with GaN layer thickness. There is a clear reduction as film thickness increases, although the data for the asymmetric (202�1) scan is noisy. As the slope of the asymmetric (202�1) scan is greater than that for the (0002) scans, it is apparent that tilt-inducing screw dislocation density is reduced at a lower rate than for the twist-inducing edge dislocations. The scans labeled “(0002) 0°” and “(0002) 90°” are for the same reflection with the sample rotated 90º in the sample mounting apparatus, the offset may be an artifact caused by the diffraction equipment.

As discussed in Section 4.41, the (0002) reflection yields information about

screw dislocations causing tilt while the (202�1) reflection contains a combination of

effects from screw and edge dislocations. An additional scan of the (0002) reflection

was also done with the sample mounted at a 90° rotation parallel to the sample stage;

for an unknown reason the zero-slit extrapolation was notably larger for this scan

87

when compared to the non-rotated version, perhaps due to some effect caused by

changing the sample mounting position.

While the magnitudes of the two (0002) FWHM reflections differ, they share a

similar slope of approximately -2750 arcsec²/µm. The zero-slit FWHM extrapolation

for the (202�1) scan shows an anomalously high value for the 28 µm sample but the

best-fit slope is almost 80% larger, at -4860 arcsec²/µm.

The broadening effect of dislocations depends on the angular projection of the

Burgers vector onto the diffraction plane.67 The (0002) plane is parallel to the growth

surface; twist dislocations having in-plane displacements will not broaden the X-ray

scan; only tilt-inducing screw dislocations will have an effect. The (202�1) plane is

inclined at a 62° angle with respect to (0002); using the (0002) value as the measure of

broadening due to tilt only, is possible to use trigonometric relationships to isolate the

effect of edge twist dislocations.

�()))�*� =������ ����0° +��+���� ���0° =������ ;

�(�)�,�)� =������ ����62° +��+���� ���62° ≈ 0.22������ + 0.78��+����

Where ������ and ��+���� are the broadening effects due to tilt (screw dislocation

density) and twist (edge dislocation density). Differentiating the above equations with

respect to film thickness, h, yields the slopes of the lines in Figure 4.7:

� �ℎ� [�()))�*� ] = � �ℎ� ������� � = -2750 arcsec²/µm;

� �ℎ� [�(�)�,�*� ] = -4860 arcsec²/µm

≈ 0.22 � �ℎ� [������ ] + 0.78� �ℎ� [��+���� ]; � �ℎ� [��+���� ] ≈ −5380 arcsec²/µm = 1.96 � �ℎ� [������ ]

88

As the dislocation density tracks the square of the variance, it appears that edge

dislocations are eliminated twice as quickly as tilt-inducing screw dislocations as the

film thickness increases. Using the directly measured cross-sectional transmission-

electron microscope dislocation density (counting both types) of 5x107/cm2 for the 12

µm thick sample, we can use Figure 4.7 to infer that the 55 µm thick sample has

roughly 2.4x107/cm2. Extrapolating the (0002) lines out to greater thicknesses, we

would expect a reduction in dislocation density to fewer than 107/cm2 between 100

and 120 µm, mostly consisting of tilt-inducing screw dislocations.

4.5 Summary of microstructural characterization

We have characterized the microstructure of GaN layers deposited with the

hybrid VPE deposition system. X-ray rocking curve analysis of a 1.5 µm film shows

three distinct Gaussian peaks corresponding to a 3 zone growth model: a very thin

highly disordered layer at the sapphire interface, an intermediate zone some 0.25 µm

thick where nucleated grains coalesce and misfit dislocations entangle and annihilate,

and a 1.25 µm high quality region where the bulk of the material grows. The rocking

curve linewidth for the topmost region is only 105 arcseconds, an indication that the

crystallinity is excellent for such a thin layer.

TEM imaging of a 12 µm film shows that the material quality improves and

the dislocation density decreases with increased thickness. A method to correlate the

ω-scan linewidth to the measured dislocation density has been developed; once the

effects of wafer bowing due to residual stress have been accounted for, the square of

the value of the rocking curve FWHM is directly related to the crystalline disorder

89

caused by dislocations. In this way it is possible to predict the reduction in threading

dislocations as further material is grown. Using symmetric (0002) and asymmetric

(202�1) scans we can separate the disorder-inducing effects of tilt (screw) dislocations

from twist (edge) dislocations, and we see that twist is more rapidly reduced than tilt

with thicker growth. Using the method described, we have produced a 55 µm thick

layer with a dislocation density of 2.4x107/cm2 and predict that a reduction to densities

below 1x107/cm2 can be achieved with thicknesses of approximately 100-120 µm.

90

Chapter 5: Photoluminescence characterization

5.1 Photoluminescence characterization of GaN

X-ray analysis of GaN films can reveal much about the nature and extent of

disorder within the crystal lattice, as well as giving information about the dislocation

density and its decrease with increasing thickness. However it cannot reveal much

about the electronic structure of the material, nor can it detect the presence of low-

level (doping) impurities. Photoluminescence, on the other hand, can reveal much

information about the electronic band structure68, the nature of the impurities that are

present,69,70,71 and the extent of residual substrate-induced strain.72

5.1 10K Photoluminescence

As the ionization potential for a donor in GaN is 30-32meV,73

photoluminescence at low temperatures (typically below 15K68,74) exhibits emission

from excitons. An exciton is an electron-hole pair bound by a Coulombic electrostatic

attraction with a binding energy dependent on the carrier effective masses,

approximately 20-21meV in GaN.75,76 When excitons collapse they spontaneously

emit a photon, the energy of which is determined by the band gap, reduced by the

exciton self-binding energy and exciton-impurity binding energy (if the exciton is

bound to an impurity).

Excitons may behave as hydrogen-like free particles moving within the lattice,

in which case they are appropriately called free excitons. The dielectric nature of GaN

induces a screening effect increasing the effective radius of the exciton, decreasing the

91

wave function overlap and so lengthening the lifetime for free excitons, creating a

bottleneck for emission. Excitons may also become weakly bonded to donor

impurities in the lattice, even non-ionized donors (which at these low temperatures are

the only type available). This binding energy is weak, on the order of 3-4 meV,70,77

but it is sufficient to cause the exciton to become more spatially localized in its

vicinity, increasing the decay rate. For that reason the dominant emission from GaN

at low temperatures is from neutral donor-bound excitons.

Figure 5.1. The calculated band structure around the Γ point in wurtzite GaN. At k=0 the valence band is split by crystal field and spin orbit coupling into the A (Γ9), B (Γ7), and C (Γ7) states, with separations of 6 meV between A-B and 37 meV between B-C. Binding energies E�

�, E��, and E

� for excitons XA, XB and XC are shown directly under the conduction band. (Source: Chen.75)

Wurtzitic GaN is a direct bandgap semiconductor, and as is typical with other

III-V direct bandgap crystals, it has a bandgap minimum at the Γ point, but with C6V

symmetry. Figure 5.1 shows the calculated band structure with the exciton binding

energies appearing below the conduction band.

92

There is no degeneracy in the upper valence bands at Γ. Γ-. (A) is the highest

energy valence band, followed by Γ/. (B) and finally the other Γ/

. band (C). The

splitting between the three bands is related to the spin-orbit interaction and the crystal

field splitting, and in the case of GaN the crystal field splitting is approximately 22

meV and about twice as large as the spin-orbit term.68 In practice there is an effective

6meV energy split between the higher-energy heavy hole band (A) and the next-

lowest lighter hole band (B), with an additional 37meV split to the third band (C).75,78

At low temperatures, we expect to see exciton transitions from conduction to valence

bands A, B, and C, each with an increasing energy and denominated as XA, XB, and

XC respectively. In practice, biaxial compression commonly found in heteroepitaxial

films grown on sapphire decreases the oscillator strength of the XC transition and

increases the strength of the XB transition, so often the XC transition is not seen in

photoluminescence scans.79

5.1.2 Photoluminescence setup

Photoluminescence characterization was done at the Paul Drude Institute in

Berlin, Germany. A He-Cd laser operating at 325 nm with confocal optics allowed for

spot sizes as small as 2 µm, with excitation densities up to 200 W/cm2. Samples were

maintained at 10K by a liquid helium cryostat. The 0.85 m spectrometer used a liquid

nitrogen cooled CCD array for detection, and has a spectral resolution of 0.03nm

(approximately 0.3 meV at E = 3.5 eV).

93

5.2 Photoluminescence characterization of a 12 µm sample

Figure 5.2 below shows a 10K photoluminescence spectra from a 12 µm thick

GaN layer grown using the hybrid VPE 2 step growth technique.

Figure 5.2. 10K PL spectra of a 12 µm thick HVPE-grown GaN film. In the wider log scale scan the primary emission is from neutral donor-bound excitons (D0X) and free excitons (FE). The low energy shoulder indicates presence of an acceptor-bound exciton (A0X), with free exciton-optical phonon replicas FE 1LO and FE 2LO at 92 and 184 meV below the free exciton energy. Inset: the linear scale expansion of band-edge luminescence shows that neutral donor-bound exciton emission consists of two peaks, D

�X and D��X, offset by approximately 0.8 meV. Free exciton emission to

valence band A (FEA) and valence band B (FEB) is also present.

The wide-span logarithmic scale scan shows strong emission from neutral

donor-bound excitons with a high-energy shoulder from free exciton emission. On the

low energy side of the emission peak there is evidence of a neutral acceptor-bound

exciton (A0X) emission some 17meV below the neutral donor-bound emission,

consistent with an acceptor impurity such as Mg or C.71,80 Free exciton annihilation

can also result in the emission of one or more lattice-optical phonons each having an

94

energy of 92 meV.74 These so-called first and second “phonon replica” peaks can be

seen at energies 92 and 184 meV below the free exciton emission peak. Notably

absent from this scan is any evidence of so-called yellow luminescence associated

with a deep-level carbon-gallium vacancy complex.71,81

The inset portion of Figure 5.2 has been expanded around the band-edge

emission portion. The neutral donor-bound exciton emission is comprised of two

discernable peaks, D�)X at 3.4764 eV and D�

)X at 3.4773 eV. These peaks have been

associated with impurities in GaN: donor D2 has been identified as an oxygen atom on

a nitrogen site (ON) and donor D1 is a silicon atom on a gallium site (SiGa).69,70 The

~0.9 meV difference in the emission is directly related to the difference in the exciton

binding energies for silicon and oxygen. These impurities are presumed to have come

from reactions between the liquid gallium and the fused quartz (SiO2) reservoir in the

hybrid HVPE growth system.82 Free exciton emission lines FEA and FEB are visible,

although the emission efficiency at low temperature is much reduced compared to the

donor-bound emission, due to the localization effect described previously.

The effect of biaxial strain induced by the sapphire substrate is evidenced by

the shift of D0X lines from the reported unstrained values of D�)X = 3.471 eV and D�

)X

= 3.472 eV.69,70,83 Compressive strain in the GaN film increases the transition

energies at a rate of 55 meV per percent of strain,72 indicating that the film still has a

residual compressive strain of 0.09%. Deconvolving the two D0X emission lines

shows a FWHM linewidth of 3.5 meV, indicative of good material quality.72

95

5.3 Photoluminescence of 28 micron layer

The 10K photoluminescence spectrum for a 28 µm thick GaN layer is shown

below in Figure 5.3. Similar to the spectrum for the 12 µm layer, there are D�)X and

D�)X lines, free exciton transitions to A and B valence bands, as well as the weak

presence of a neutral acceptor-bound exciton A0X on the low-energy shoulder of the

emission. The positions of D�)X and D�

)X are approximately 8 meV above the

unstrained reference values, indicating a residual compressive strain of 0.14%, greater

than for the 12 µm layer, but consistent with the observed greater bowing of this

sample. Deconvolving the two D0X peaks yields a narrower FWHM of 2.8 meV

compared to 12 µm, indicating further material quality improvement as the thickness

increases.

Figure 5.3. 10K PL spectrum of a 28 µm thick HVPE GaN layer. As with the 12 µm layer, the D0X emission is separable into distinct D

�X and D��X lines, with narrower

FWHM compared to the thinner layer. There is strong free exciton emission into the A and B valence bands, as well as some evidence of residual acceptor-bound exciton emission A0X. The offset position of the D

�X and D��X lines imply a 0.14%

compressive strain state.

96

5.4 Emission from Freestanding 60 µm GaN

During the post-growth cooling down process, the thermal strain mismatch

between substrate and epilayer can become so great that film delamination can

sometimes spontaneously occur. In such a case, millimeter-to-centimeter-sized pieces

of GaN can break free of the substrate and become unstrained freestanding GaN

layers. In Figure 5.4 below, photoluminescence spectra were taken from such a piece

of a 60 µm thick layer that sprang free of the substrate.

Figure 5.4. 10K PL spectra from a piece of 60 µm freestanding GaN. (Left) linear scan showing strong D0X emission with higher-energy side peaks. (Right) expansion of the near band-edge emission shows D

�X�and D��X� lines at their expected strain-

free energies and a FWHM of 0.40 meV. Free exciton peaks A and B are present, as well as a third emission line approximately 6 meV above the D0XA lines which is attributed to D0XB. Two-electron transitions involving excitons exciting a neutral donor’s electron are also visible some 21 meV below the D0XA and D0XB emission lines.

97

The linear graph on the left side of Figure 5.4 shows strong primary emission

from the D0X lines, but there are two satellite peaks on the high energy side.

Expanding the emission spectra on the right side of the Figure shows the structure

more clearly. The D�)X and D�

)X lines are present at the expected separation of 0.9

meV, and their values correspond well with the unstrained calculated values as

reported elsewhere.69,70 Deconvolution of the FWHM of these two peaks yields a

FWHM of approximately 400 µeV, close to the resolution limit of 300 µeV and

indicative of very high quality material.

On the higher-energy shoulder of the D0X lines are three distinct peaks: FEA

and FEB are present, as expected, but there is a third peak that is 6 meV higher than

D0X, which is ascribed to donor-bound excitons in the B valence band, or D0XB .

Thus the lines at 3.4710 and 3.4720 are attributed to D�)X0 and D�

)X0 respectively. On

the lower energy side are two additional peaks with energies 21 meV lower than the

D)X0 lines, and these are presumed to be from two-electron transitions, which shall be

discussed below.

5.5 Two-electron transitions

Luminescence from a two-electron transition (TET) occurs as a neutral donor-

bound exciton collapses and in the process excites the non-ionized donor. The neutral

donor typically sits on a substitutional site, and has an excess electron compared to the

basis sublattice atom. The non-excess electrons hybridize with the lattice much like

other sublattice electrons, while the excess electron is left weakly bound to the donor

atom. The neutral donor is well described by a positive core (of electrons and

98

nucleus) and a single electron as a hydrogenic-like state. The TETs occur when the

decay of a bound exciton leads to the excitation of the hydrogenic donor electron to an

excited state and emission a photon that is red-shifted from the bound-exciton

emission energy.84 Thus, the energy difference between the D0X and TET photons

should contain (among other transitions) a transition representing the energy

difference between the n=2 and n=1 donor-bound exciton states.

In figure 5.4, the presence of two peaks approximately 21 meV below the

D)X0 and D

)X1 lines strongly indicates that these are of TETs. Experiment and

theory78 show that exciting an unionized donor from the n=1 to n=2 state requires

about 75% of the ionization energy, which in the case of neutral donors in GaN is on

the order of 28-32 meV.73,78 The presence of two TET peaks offset by the A-to-B

valence band separation of 6 meV also lends credence to the interpretation of the two

peaks as TET: D)X0 and TET: D

)X1. TETs are not seen except in material with

extremely low background carrier concentrations or other significant recombination

centers,84 and their presence in these scans indicates that the material quality is very

high.

5.6 Summary of photoluminescence results

Low-temperature photoluminescence characterization was done on GaN films

with thicknesses of 12 µm, 28 µm and a 60 µm piece of freestanding material that

spontaneously sprang free of the sapphire substrate. None of the samples show

evidence of mid-band gap yellow luminescence associated with deep-level carbon

impurities; this is a clear advantage of the carbon-free chemistry of the HVPE

99

deposition reaction. The detailed structure of the near band-edge luminescence

consistently shows two distinct neutral donor-bound exciton transitions separated by

approximately 0.9 meV, with the lower energy transition (D�)X) attributed to an

oxygen atom on a nitrogen lattice site (ON) and the higher energy transition (D�)X) as

silicon on a gallium site (SiGa). It is possible that their presence is a result of the

reaction of gallium metal with the fused silica walls of the hybrid growth system. The

thinner samples (12 and 28 µm) show evidence of a neutral acceptor-bound exciton

transition, which may be due to carbon or magnesium impurities, while the thicker 60

µm sample does not. The source of the acceptor impurity could be carbon from the

initial MOVPE buffer layer deposition step or it could be contamination from the

sapphire substrate, but in either case the effect decreases with layer thickness and does

not appear to be inherent to the HVPE deposition process.

The FWHM linewidth of the band-edge luminescence decreases with

increasing film thickness, indicating that the material quality improves with further

distance from the sapphire interface. At 60 µm thickness, the 0.4 meV linewidth is

slightly larger than the spectral resolution of the spectrometer setup (0.3 meV).

Residual biaxial compressive strain is evident in the two thinner samples as seen by

the higher-energy shift in the emission lines; the 12 µm film has 0.09% compressive

strain and the 28 µm sample has 0.14%. The 60 µm freestanding piece is unstrained

and shows further fine excitonic structure including transitions to the A and B valence

bands as well as two-electron transitions, both of which indicate high material quality.

100

Chapter 6: A semi-insulating GaN Alloy: GaMnN

6.1 Semi-insulating GaN for use in microwave amplifiers

In addition to its use as a light-emitting semiconductor, GaN has other material

properties that make it an attractive choice for high-frequency, high power microwave

amplifiers. Its wide bandgap allows for higher temperature operation, reducing the

demands on any thermal management system, corresponding to a savings in cost and

weight. With a theoretical breakdown electric field of approximately 3.3 MV/cm

(50% greater than SiC and 5.5 times that for GaAs),85,86 it is possible to fabricate

devices with smaller active regions, reducing serial resistance and power consumption.

Recently AlInN/AlN/GaN heterostructure high-electron-mobility transistors (HEMTs)

have been fabricated with output power densities in excess of 10 Watts/mm of gate

length and power-added efficiencies of 51% at 10 GHz;87 others have fabricated

AlGaN/GaN HEMT structures with fmax of 300 GHZ.88

SiC is the most commonly used substrate for high power GaN electronics, as

SiC has a thermal conductivity of 4.5 W/cm/K which is almost 13 times greater than

sapphire (0.35 W/cm/K).89 GaN HEMT devices fabricated on sapphire substrates

have severe thermal management issues that limit overall device performance and

limit the minimum feature size.90

Semi-insulating substrates are necessary to reduce capacitance losses for high

frequency devices however, and it is currently difficult to make semi-insulating SiC.91

To produce higher quality material (and devices), an initial semi-insulating GaN

buffer layer must be added to the SiC substrate before the active transistor layer can be

101

made. However, a semi-insulating GaN (or the appropriate group III-N) substrate

would be an important development, as the 4% lattice mismatch with respect to SiC

puts severe limits on device design and growth compared to homoepitaxy.

While MOVPE GaN is typically n-type in the mid 1016/cm3 range,92 thick

HVPE GaN can have carrier concentrations as low as 1014/cm3.37 The thermal

conductivity of GaN is sufficient for the thermal dissipation requirements necessary

for high-power applications, provided the material can be reliably made in semi-

insulating form. In this chapter I will discuss the development of a new semi-

insulating GaN technology. The technology developed is similar to the methods often

employed by SiC manufactures to produce highly semi-insulating SiC by adding deep

level dopants to pin the Fermi level at midgap.91 Similarly, we add alloy and dopant

amounts to make the GaN semi-insulating. The approach is slightly different because

the amounts added are different and the resulting material has a slightly different

crystal structure. Our semi-insulating GaN layers can be so depleted of carriers that it

is difficult to make proper resistance measurements. That the material quality is

maintained is evident from high resolution X-ray diffraction and transmission-electron

microscopy (TEM). We can produce very thick layers of this material, in excess of 60

µm, an essential requirement for providing proper thermal dissipation during device

operation.

6.2 Semi-insulating GaN through the incorporation of Mn using HVPE

Mn doping in GaAs (not GaN) has been investigated for several years by Ohno

to develop dilute magnetic semiconductors for Spintronics applications.93,94,95 For

102

many of the early years of Ohno’s initial work, his attempts to dope GaAs p-type

using Mn were disappointing because the GaAs was semi-insulating.96 Under higher

doping concentrations of Mn, the material turned from semi-insulating to metallic.

While Ohno spent many years resolving this problem using kinetically limited

molecular-beam epitaxy, we saw his early results as an opportunity to produce semi-

insulating GaN by doping with Mn.

Figure 6.1. Schematic diagram of Ga(Mn)N deposition system, a modified version of the hybrid HVPE system used for all growths in this dissertation. Pieces of pure metallic Mn are placed in the annular space between the Ga nozzle and the nozzle sheath. During GaMnN deposition, HCl flows over it to produce MnCl2.

Our modified HVPE system is schematically shown above in Figure 6.1. In

the HVPE process hydrogen chloride (HCl) gas buffered with carrier gas flows over

liquid gallium (Ga) to form gallium chloride (GaCl). GaCl is transported in the N2

carrier gas along with H2 and any excess HCl to the growth surface. Here, GaCl reacts

with ammonia (NH3) to form GaN. Mn is added to the epitaxy in a similar way: we

inject HCl gas over Mn metal into a nozzle arrangement separate from the Ga source.

This HCl reacts with solid Mn and MnCl2 is the product.

103

The reaction is slightly different from the Ga reaction with HCl. At the GaCl

source temperature (850°C), Mn is a solid, and the reaction between Mn and HCl

produces a liquid with a vapor pressure of 1400 Pa (10.45 Torr)97:

2HCl + Mn(s) → MnCl2(l,v) + H2

Figure 6.2. The calculated MnCl2 vapor pressure as a function of temperature, using thermodynamic data from Kritskii.97 MnCl2 has a vapor pressure of 10.45 Torr at the typical GaCl cell operating temperature of 850°C and boils at 1230°C.

As can be seen in the vapor pressure curve in Figure 6.2 above, MnCl2 is a

volatile liquid well below its boiling point at 1230°C. While the vapor pressure of

MnCl2 is significantly below atmospheric level, there is significant vapor transport to

the GaN surface even at the GaCl cell temperature of 850°C.

If we assume the Mn is incorporated into the GaN in an independent manner,

the possible reactions to incorporate Mn as a dopant are:

MnCl2 + NH3 → MnN + 2HCl +�

�H2 (MnN reaction); or

3MnCl2 + 4NH3 → Mn3N4 + 6HCl + 3H2 (Mn3N4 reaction); or

MnCl2 + H2 → 2HCl + Mn (metallic Mn growth)

104

We are not yet sure which of these reactions occurs to incorporate Mn into the

GaN.

6.3 Initial Characterization of GaMnN

To show good control of our MnCl2 delivery system and that we can

controllably incorporate Mn into the GaN films, we first demonstrate that there is a

linear relationship between the HCl flow delivered through the Mn injector (denoted

as HClMn), and the amount of Mn incorporated into our GaN. This is shown in Figure

6.3 below; when the ratio of HClMn to HClGa is kept below 0.5 there is a linear

relationship between the HCl flow rate and the atomic Mn incorporation. We

therefore can conclude that the Mn is incorporating uniformly.

Figure 6.3. The linear relationship between the ratio of HCl flow over Mn to HCl flow over Ga and the resulting GaMnN film’s Mn content, as determined by electron microprobe analysis. This relationship holds for flow ratios up to 0.5.

105

The data in Figure 6.3 should not necessarily be interpreted as indicating there

is a uniform alloy of MnGaN in these initial layers. Indeed, for lower Mn

compositions the alloy appears uniform from electron-microprobe measurements, but

this method cannot distinguish micron-scale or smaller composition variations due to

the extent of the excited volume being sampled. In the higher-Mn concentration

samples the microprobe indicates the alloy is not uniform and segregation is present.

Notes on the electron-microprobe analysis can be found in Table 6.1, below. There

are apparently two phases present in the GaMnN sample with the flow ratio of 3.43,

while there is a large range of compositions (varying by 50%) in the sample with flow

ratio of 1.99. We can speculate that regions of Mn-rich inclusions could be present in

these layers, but further analysis of the films using transmission-electron microscopy

(TEM) will clarify that later in this chapter.

Table 6.1. Characterization summary of 5 µm thick GaMnN layers grown with various HCl flow ratios. Mn content was measured using electron microprobe, Hall effect measurements were attempted on samples 1501-1504; all were semi-insulating.

Growth reference #

Ratio: �(��)

�(��)

Mn content (atomic %)

Notes

1498 6.46 - Not single crystal

1499 3.43 - Two phases present: GaN with ~4% Mn, and MnNx

1500 1.99 6.30% Mn content varies from 6-9% Mn over sample

1501 0.385 3.10% Rough surface morphology

1502 0.223 2.00% Rough surface morphology

1503 0.160 1.14% Hall Measurement indicates semi-insulating

1504 0.065 0.67% Semi-insulating

Hall measurements were conducted on approximately 15 samples at room

temperature. In all GaMnN samples, the alloyed In contacts would not produce

106

Ohmic contacts and thus Hall measurements could not be accurately made, unlike the

case of pure HVPE-grown GaN. The contacts were alloyed in forming gas

(95%N2+5%H2) in a rapid thermal annealing (RTA) furnace up to 750˚C for 3

minutes. This inability to create Ohmic contacts is an indicator that the Mn-doped

material is semi-insulating.

The morphology of the GaMnN films was affected by the Mn concentration, as

shown in the optical micrographs below in Figure 6.4. At 500x magnification the

surface morphology is typically smooth and nearly featureless for pure HVPE GaN;

the addition of 1.1% Mn increases the mosaicity of slightly mismatched hexagonal

grains, but otherwise the surface is smooth. Increasing the Mn concentration to 6.3%

shows evidence of a second phase (presumably manganese nitride) nucleating on the

surface as well as greater surface distortion along 120° crystallographic orientations.

Figure 6.4. Plan-view 500x optical micrographs of 5 µm thick GaN and GaMnN alloys. The surface feature on the pure GaN was used to focus on an otherwise featureless surface. At 1.1% Mn concentration the surface shows increased hexagonal mosaicity; at 6.3% there is a second Mn-rich phase present on the surface.

High-resolution X-ray diffraction (HXRD) was used to determine the crystal

quality of the GaMnN layers. Figure 6.5 shows the relationship between the FWHM

of the ω-scan rocking curve and the manganese content of the 5 µm thick films

107

described in Table 6.1. The presence of Mn in the GaN lattice causes disorder as

evidenced by the increase in FWHM with manganese concentration.

Figure 6.5. The relationship between ω-scan rocking curve FWHM and Mn concentration in 5 µm thick GaMnN HVPE films. Adding Mn to the GaN lattice increases disorder, broadening the linewidth of the ω-scans.

6.4 TEM analysis: a second phase and a new crystal structure

Transmission-electron microscopy (TEM) uses the diffraction pattern from

coherently scattered electrons to discern the crystal structure and hence any crystalline

imperfections. It is a particularly useful way to characterize new materials such as our

GaMnN. The TEM image shows two significant results: first, there is a second phase

of what is probably MnN present in the film under some situations, but in small

quantities. Second, the crystal structure has changed in a slight but important way.

There is now sublattice ordering in the usual GaN hexagonal structure containing

atomic layers of higher-Mn and lower-Mn concentrations.

108

6.4.1 Mn-rich second phase in GaMnN

TEM images of a semi-insulating GaMnN film grown with a flow ratio of 0.02

(approximate Mn concentration of 0.16% estimated using Figure 6.3) are shown below

in Figure 6.6. The growth direction is vertical. In the bottom portion of the image the

low temperature buffer layer region can be seen, with very poor crystal structure. The

wavy almost vertical lines about 400 nm long are dislocations.

Figure 6.6. Transmission-electron microscopy (TEM) images of a GaMnN sample with estimated 0.16% Mn. (Left) Wide-field image showing the sapphire substrate and GaN buffer layer; the strain field around inclusions can be seen. (Right) Close-up images of selected regions from the micrograph on the left, showing the inclusions (circled in red).

From the slices enlarged from the image (on the right side of Figure 6.6), the

strain fields from second phase inclusions can be seen. The second-phase inclusions

are sparse, on the order of 2x1014cm-3, for an assumed 50 nm TEM sample thickness.

We will show below that the inclusions are most likely some phase of MnN. Since

MnN is stable at the 1025°C growth temperature, it is likely that the Mn locked into

the second phase is not likely to diffuse into the device active region, and these

inclusions are not expected to harm the long-term device performance.

109

The composition of the inclusions can be determined by energy-dispersion X-

ray (EDX) analysis. EDX was conducted in two regions of the GaMnN layer: a region

with second phase inclusions and a region without second phase inclusions. This is

shown in Figure 6.7, below.

Figure 6.7. Energy Dispersive X-ray (EDX) analysis on a TEM sample of GaMnN (estimated 0.16% Mn). EDX of a background inclusion-free region shows no detectible Mn Kα or Kβ radiation, whereas the region surrounding the inclusion does.

The EDX analysis in an inclusion-rich region shows the Mn Kα and Mn Kβ

emission lines, indicating the presence of Mn in this region. The EDX analysis in a

region free of inclusions shows no such Mn lines indicating the Mn composition here

is below the detectible limit of the EDX, approximately 0.5%. Thus, from this

analysis we know the inclusions are enriched in Mn. We cannot say with absolute

certainty they are MnN as opposed to GaMnN, but if they were GaMnN we might see

110

a more gradual compositional gradient, while the inclusions in Figures 6.6 and 6.7 are

quite distinct.

6.4.2 Sublattice ordering – a new crystal structure

The crystal structure of the GaMnN regions without inclusions is also of

interest. X-ray rocking curve data (Figure 6.5) showed the material has more random

ordering with increasing Mn content due to the alloying of GaN with Mn. High-

resolution TEM gives a more detailed picture of the GaMnN microstructure as shown

in Figure 6.8, below. In the sample imaged, the GaMnN layer is deposited on top of a

HVPE GaN buffer layer. The Mn concentration was measured by EDX in both

regions, and the Mn Kα,β emission data is shown on the right; it is apparent that the

upper region contains Mn and the lower region does not.

An expanded view of the image intensity contrast is shown on the left side of

the image. For the GaN region (lower left of Figure 6.8), the hexagonal wurtzite

structure is evident, evenly-spaced bands of a period of approximately 0.5 nm (5 Å)

show good agreement with c = 5.189 Å of the GaN unit cell. For the GaMnN region

the intensity contrast is shown on the upper left. Here, the intensity contrast is

different; the peak intensity spacing is half that of the GaN region, and there appear to

be two alternating peak intensities. This indicates that the crystal structure of the

GaMnN region has been altered from that of normal GaN, there is now additional

ordering on the scale of 2

� in the wurtzite lattice.

111

Figure 6.8. (Center) High resolution TEM image of GaMnN deposited on GaN on sapphire. The GaN buffer shows the normal wurtzite structure with intensity variations (lower left) with a period of c = 5.189 Å. EDX of this region (lower right) shows no characteristic radiation for Mn. The GaMnN layer shows intensity variations (upper left) at twice the spatial frequency (c/2 = 2.595 Å) with alternating maximum peak intensity, indicative of sublattice ordering. EDX of the GaMnN (upper right) verifies the presence of Mn.

It is conceivable that this sublattice ordering contains alternating layers of Mn-

rich and Mn-poor composition, and that an excess of Mn above some level will lead to

the nucleation of MnN inclusions. However, more careful studies are necessary to

verify this, and it may not matter to the overall goal of semi-insulating GaN. The

formation of MnN inclusions will probably serve to stabilize the Mn against diffusion

into the upper device region. Similarly, the formation of sublattice ordering indicates

112

there is a low-energy configuration for Mn in GaN, and this too may make the Mn

stable against vertical diffusion.

6.5 Summary: semi-insulating GaMnN

In this chapter I have shown that GaN alloyed or doped with Mn is both a

feasible and promising approach to provide a thermally dissipative, high frequency

isolation platform for high power, high frequency devices. We believe this system

may offer superior overall performance compared to SiC systems.

We can control the amount of Mn in the films by controlling the flow rate of

HCl over metallic Mn. It is a very simple, low cost process. Hall measurements show

the films are all semi-insulating and in all likelihood the films will remain semi-

insulating with much lower Mn content. Transmission-electron microscope (TEM)

analysis shows there are inclusions of MnN in the layers as well as alloying with Mn

in the GaN regions. There is also evidence that there are alternating layers of Mn-rich

and Mn-poor composition along the (0001) direction. Because the Mn is

accommodated in a new crystal structure or in the MnN inclusion and not interstitially,

we believe the Mn will be stable in the layer and will not diffuse into the upper device

layer. Taken together, the characterization shows that the GaMnN is a promising

semi-insulating material for high-frequency, high-power transistor applications.

The challenge that lies ahead, to allow for the commercialization of this

technology, is to develop a better and more robust method to lightly dope GaN with

Mn. The current method, in which Ga and Mn are located in the same section of the

growth reactor, is not optimal. Currently, the Ga and Mn sources must be held at the

113

same temperature, and a compromise must be made since the vapor pressures of the

two reactive species, GaCl and MnCl2 are dissimilar. If the temperature is optimized

for GaCl production, the vapor pressure of the MnCl2 is too low. If the temperature is

raised to better evaporate the MnCl2, excess Ga metal starts to evaporate with the

GaCl, which leads to metallic inclusions and defects in the growing film.

114

Chapter 7: Conclusions and future directions

7.1 Conclusions

In this dissertation I have presented a novel hybrid MOVPE-HVPE growth

system that can produce high quality epitaxial GaN on sapphire. This combined

method presents a unique and scalable means for the production of GaN-based

epilayers and devices without the need for a separate reactor for each deposition

process.

Nucleated heteroepitaxial GaN grains coalesce in a state of biaxial tension on

sapphire; as the film thickness increases so does the stress, which can cause cracking

and breakage of the wafer during growth, or cooling down. I have investigated the

effect of growth parameters on the surface morphology and shown that conditions

favoring vertical over lateral growth (such as high growth rate, low temperature, and

high V/III ratio) produce a thick hexagonally pitted film with {101�1} sidewall facets

that serve to relieve stress and allow for thicker layer deposition. Regrowth on a pitted

film under conditions favoring lateral growth (low growth rate, higher temperature,

and lower V/III ratio) fills in the hexagonal pits and smoothes the overall morphology.

Using this newly developed 2-step growth technique I have demonstrated that it is

possible to produce moderately thick GaN layers that are smooth enough to be near

device-quality.

X-ray diffraction and transmission electron microscope analysis of the

microstructure of a thin film shows that there is a thin region of highly disordered

material at the sapphire interface, followed by a less disordered layer of approximately

115

0.25 µm where many of the dislocations induced by coalescence become entangled,

leading to improved crystal quality. The linewidth of the X-ray diffraction rocking

curve scans narrows with increased thickness, once the effects of wafer bowing have

been compensated for by extrapolating to the zero slit width condition. Transmission

electron microscopy imaging was used to count the dislocation density, and a method

to correlate symmetric and asymmetric rocking curve linewidth to edge and screw

dislocations was developed. Using that methodology, I showed that the dislocation

density for a 55 µm film was 2.4x107/cm2 and could be reduced to levels below

107/cm2 for thicknesses around 100 µm.

Low temperature photoluminescence spectra show high material quality with

no detectable mid-gap yellow luminescence. Optical transitions based on neutral

donor-bound excitons indicate that the primary donors are silicon and oxygen,

presumably impurities from the fused quartz system walls. The emission linewidth

becomes narrower as film thickness increases, indicating that the material continues to

improve further from the sapphire interface. The presence of a neutral acceptor-bound

exciton was seen in the 12 µm film and weakly in the 28 µm layer, but was absent in

the freestanding 60 µm piece, indicating that the acceptor may be an artifact of the

MOVPE nucleation layer deposition process or possibly contamination due to the

sapphire substrate itself. By correlating the energy shift of the exciton transitions from

their predicted values, I demonstrated that the biaxial compressive strain state

increases with film thickness but disappears when the layer becomes freestanding

from the sapphire substrate.

116

Finally I showed how a modification to the HVPE system to include a metallic

manganese source can be used to grow semi-insulating Mn-doped GaN layers for use

as a possible substrate for high frequency electronics. For low concentrations, the

manganese incorporation rate in GaN is linear with respect to the HCl flow rate over

the manganese pieces. Transmission electron microscopy shows that a second Mn-

rich phase can develop within the GaN lattice, while at lower concentrations it appears

that there is a sub-lattice ordering effect wherein there are alternating layers of Mn-

rich and Mn-poor atomic planes.

7.2 Future directions and research

While this hybrid MOVPE-HVPE reactor has been shown to be a robust and

effective method for the direct deposition of GaN onto sapphire in a single growth

system, there are still some challenges and areas for further research. First, while it is

possible to use MOVPE to deposit low temperature nucleation layers, the system is

not capable of producing device layers on top of HVPE GaN. A method to integrate a

cold-wall heating method for high temperature MOVPE deposition would allow for

the direct growth of a nucleation layer, HVPE buffer, and device layers all in a single

system without necessity of unloading or interruption.

Furthermore, while the relationship between cracking, stress, and surface

morphology has been characterized, further research into the mechanisms at work and

a better understanding of how to control these stresses would be desirable to allow

production of very thick layers in excess of 100 µm. Currently it is possible to

produce these thick layers, but the imperfect understanding of the forces at work

117

results in a low process yield. It is known that pitted layers are less susceptible to

cracking, but the application of a smoothing layer increases the stresses in ways that

are not perfectly predictable. The smoothing layer itself is also an area that could be

further optimized by finding the right conditions for growth that give truly effective

planarization.

Further research into the Mn-doping of GaN by HVPE can also help optimize

the process. The crude modifications to the hybrid growth system showed that

manganese is an effective dopant to produce semi-insulating GaN, but the method of

delivery has hysteresis and lag effects that make commercialization of this technique a

work for the future.

Finally, the presence of silicon and oxygen donors in the photoluminescence

spectra indicates that there may be a reaction between hot liquid gallium and the fused

quartz gallium containment system. Further research and development of a gallium

nozzle consisting of materials that are inert under these conditions (such as PBN)

might well eliminate their presence, resulting in even higher material quality.

118

Chapter 8: List of References

1 L. Liu and J. H. Edgar, Materials Science and Engineering R 37, 61 (2002) 2 S. Li and C. Ouyang, Physics Letters A 336, 145 (2005) 3 M. Aoki et al., Journal of Crystal Growth 218, 7 (2000) 4 M. Aoki et al., Journal of Crystal Growth 242, 70 (2002) 5 T. Hashimoto et al., Nature Materials 6, 568 (2007) 6 D. R. Ketchum and J. W. Kolis, Journal of Crystal Growth 222, 431 (2001) 7 H. Okamoto, Journal of Phase Equilibria and Diffusion 27, 545 (2006) 8 J. Perkins, “LED Manufacturing Technologies & Costs”, presentation at DOE workshop, available at

<http://apps1.eere.energy.gov/buildings/publications/pdfs/ssl/perkins_fairfax09.pdf>, (2009) 9 H. Ishikawa et al., Japanese J. Appl. Phys. 38, L492 (1999) 10 Y. Honda et al., Physica Status Solidi (C) 2, 2125 (2005) 11 O. Kryliouk et al., Materials Science and Engineering B59, 6 (1999) 12 MTI Corporation offered price, <http://www.mtixtl.com/ligao2_1.aspx> , retrieved April 2011 13 University Wafer Company offered price, <http://www.universitywafer.com>, retrieved April 2011 14 I. N. Przhevalskii et al., MRS Internet J. Nitride Semicond. Res. 3, 30 (1998) 15 D. W. Shaw, J. Electrochem. Soc. 115, 777 (1968) 16 J. Maruyama et al., J. Electrochem. Soc. 116, 413 (1969) 17 H. Maruska and J. J. Tietjen, Applied Physics Letters 15, 327 (1969) 18 W. Götz et al., Applied Physics Letters 69, 242 (1996) 19 E. Valcheva et al., Applied Physics Letters 80, 1550 (2002) 20 A. Usikov et al., Physica Status Solidi (C) 5, 1825 (2008) 21 K. D. Soo et al., Proceedings of SPIE 6894, 689406 (2008) 22 H. J. Lee et al., Journal of the Korean Physical Society 45, S813 (2004) 23 W. K. Burton et al., Phil. Trans. R. Soc. Lond. A 243, 299 (1951) 24 J. A. Venables et al., Rep. Prog. Phys. 47, 399 (1984) 25 C. V. Thompson, “Stress Evolution During Volmer-Weber Growth of Thin Films”,

<http://www.icf11.com/proceeding/EXTENDED/5678.pdf> (2005) 26 W.D. Nix and B.M. Clemens, Journal of Materials Research 14, 3467 (1999) 27 T. F. Huang et al., Compound Semiconductors 156, 11 (1997) 28 D. J. Rogers, et al., Proceedings of SPIE 5732, 412 (2005) 29 S. Nakamura, Japanese J. Appl. Phys. 30, L1705 (1991) 30 D. Kapolnek et al., Applied Physics Letters 67, 1541 (1995) 31 K. Hiramatsu et al., Journal of Crystal Growth 115, 628 (1991) 32 H. Amano et al., Applied Physics Letters 48, 353 (1986) 33 S. Hearne et al., Applied Physics Letters 74, 356 (1999) 34 J. A. Floro, Journal of Applied Physics 89, 4886 (2001) 35 L. Sugiura, Journal of Applied Physics 81, 1633 (1997) 36 B. R. Lawn et al., Journal of Materials Science 15, 1207 (1980) 37 P. R. Hageman et al., Journal of Crystal Growth 255 241 (2003) 38 J. Keckes, Applied Physics Letters 79, 4307 (2001) 39 A. Usikov et al., Physica Status Solidi (C) 0, 2580 (2003) 40 D. Gogova et al., Physica Status Solidi (A) 200, 13 (2003) 41 T. Sasaki and S. Zumbutsu, Journal of Applied Physics 61, 2533 (1987) 42 T. Sasaki, Journal of Crystal Growth 129, 81 (1993) 43 H. X. Wang et al., Journal of Crystal Growth 233, 681 (2001) 44 H. Holloway and L. C. Bobb, Journal of Applied Physics 38, 2893 (1967) 45 J. L. Rouvière et al., Applied Physics Letters 73, 668 (1998) 46 J. L. Weyher et al, Journal of Crystal Growth 210, 151 (2000)

119

47 C. Youtsey et al., Applied Physics Letters 74, 3537 (1999) 48 T. B. Wei et al., Materials Letters 61, 3882 (2007) 49 X. H. Wu et al., Journal of Applied Physics 80, 3228 (1996) 50 M. Sumiya et al., Journal of Applied Physics 88, 1158 (2000) 51 T. Matsuoka et al., Physica Status Solidi (B) 243, 1446 (2006) 52 P. Y. Lin and Y. C. S. Wu, Materials Chemistry and Physics 80, 397 (2003) 53 G. S. Solomon et al., US Patent #6,673,149, “Production of low defect, crack-free epitaxial films on a

thermally and/or lattice mismatched substrate” 54 S. J. Rosner, et al., Applied Physics Letters 70, 420 (1997) 55 T. Sugahara et al., Japanese J. Appl. Phys. 37, L1195 (1998) 56 N. G. Weimann et al., Journal of Applied Physics 83, 3656 (1998) 57 P. Kozodoy et al., Applied Physics Letters 73, 975 (1998) 58 M. E. Vickers et al., J. Phys. D: Appl. Phys. 38, A99 (2005) 59 N. Kuwano et al., Journal of Crystal Growth 115, 381 (1991) 60 W. Sarney et al., MRS Internet J.Nitride Semicond.Res. 2000, W3.47 (2000) 61 W. Zhang et al., Physica Status Solidi (A) 188, 425 (2001) 62 H. Heinke et al., Applied Physics Letters 77, 2145 (2000) 63 J. E. Ayers, Journal of Crystal Growth 135, 71 (1994) 64 X. Chen and T. Uesugi, Applied Physics Letters 88, 031916 (2006) 65 A. Sakai et al., Applied Physics Letters 71, 2259 (1997) 66 S. Nakamura et al., Applied Physics Letters 72, 211 (1998) 67 S. K. Mathis et al., Physica Status Solidi (A) 179, 125 (2000) 68 R. Dingle et al., Physical Review B 4, 1211 (1971) 69 A. Wysmolek et al., Physical Review B 66, 245317 (2002) 70 J. A. Freitas et al., Physical Review B 66, 233311 (2002) 71 T. Ogino and M. Aoki, Japanese J. Appl. Phys. 19, 2395 (1980) 72 D. Volm et al., Physical Review B 53, 16543 (1996) 73 H. Wang and A. B. Chen, Journal of Applied Physics 87, 7859 (2000) 74 W. Liu et al., Semicond. Sci. Technol. 13, 769 (1998) 75 G. D. Chen et al., Applied Physics Letters 68, 2784 (1996) 76 J. F. Muth et al., Applied Physics Letters 71, 2572 (1997) 77 G. Neu et al., Applied Physics Letters 77, 1348 (2000) 78 B. Monemar et al., Physica Status Solidi (B) 245, 1723 (2008) 79 B. Gil et al., C. R. Acad. Sci. (Paris) 1, 51 (2000) 80 G. D. Chen et al., Applied Physics Letters 68, 2784 (1996) 81 J. L. Lyons et al., Applied Physics Letters 97, 152108 (2010) 82 M. E. Weiner, J. Electrochem. Soc. 119, 496 (1972) 83 W. Rieger et al., Applied Physics Letters 68, 970 (1996) 84 P. J. Dean, Physical Review Letters 18, 122 (1967) 85 J. C. Zolper, Solid-State Electronics 42, 2153 (1998) 86 Y. Yoshizumi et al., Journal of Crystal Growth 298, 875 (2007) 87 N. Sarazin et al., IEEE Electron Device Letters 31, 11 (2010) 88 J. W. Chung et al., IEEE Electron Device Letters 31, 195 (2010) 89 U. K. Mishra et al., Proceedings of the IEEE 96, 287 (2008) 90 E. A. Douglas et al., J. Vac. Sci. Technol. B 29, 020603-1 (2011) 91 J. Hao et al., Journal of Electronic Materials 39, 530 (2010) 92 G. Y. Zhang et al., Applied Physics Letters 71, 3376 (1997) 93 H. Ohno et al., Journal of Applied Physics 69, 6103 (1991) 94 H. Ohno et al., Applied Physics Letters 69, 363 (1996) 95 A. Shen et al., Journal of Crystal Growth 175/176, 1069 (1997) 96 H. Ohno et al., Physical Review Letters 68, 2664 (1992) 97 V. E. Kritskii, Russian Journal of Applied Chemistry 79 190 (2006)