investigation of dopant incorporation in silicon carbide

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Investigation of dopant incorporation in silicon carbideepilayers grown by chemical vapor deposition

Ionela Roxana Arvinte

To cite this version:Ionela Roxana Arvinte. Investigation of dopant incorporation in silicon carbide epilayers grown bychemical vapor deposition. Other [cond-mat.other]. COMUE Université Côte d’Azur (2015 - 2019),2016. English. NNT : 2016AZUR4091. tel-01466713

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UNIVERSITE COTE D´AZUR– UFR Sciences

Ecole Doctorale de Sciences Fondamentales et Appliquées

THESE

pour obtenir le titre de

Docteur en Sciences de l’UNIVERSITE Côte d’Azur

Discipline : Physique

présentée et soutenue par

Ionela-Roxana ARVINTE

Investigation of dopant incorporation in silicon

carbide epilayers grown by chemical vapor

deposition

Thèse dirigée par Marcin ZIELINSKI et Yvon CORDIER Soutenue publiquement le : 08 Novembre 2016

Jury :

D. CHAUSSENDE Directeur de recherche - Rapporteur

CNRS, U. Grenoble Alpes

Y. CORDIER Directeur de recherche - Directeur de Thèse

CNRS-CRHEA, Valbonne

G. FERRO Directeur de recherche - Examinateur

CNRS, UCB, Lyon

S. JUILLAGUET Maitre de Conférences Rapporteur

CNRS, U. Montpellier

J-F. MICHAUD Maitre de Conférences Examinateur

U. François-Rabelais, Tours

M. ZIELINSKI Docteur Ingénieur Co-directeur de Thèse

NOVASiC, Valbonne

Centre de Recherche sur l’Hétéroépitaxie et ses Applications, CNRS, Valbonne

UNIVERSITE COTE D´AZUR – UFR Sciences

Ecole Doctorale de Sciences Fondamentales et Appliquées

THESE

pour obtenir le titre de

Docteur en Sciences de l’UNIVERSITE Côte d’Azur

Discipline : Physique

présentée et soutenue par

Ionela-Roxana ARVINTE

Etude de l’incorporation des dopants N et Al

dans des films de carbure de silicium épitaxiés

en phase vapeur

Thèse dirigée par Marcin ZIELINSKI et Yvon CORDIER Soutenue publiquement le : 08 Novembre 2016

Jury :

D. CHAUSSENDE Directeur de recherche - Rapporteur

CNRS, U. Grenoble Alpes

Y. CORDIER Directeur de recherche - Directeur de Thèse

CNRS-CRHEA, Valbonne

G. FERRO Directeur de recherche - Examinateur

CNRS, UCB, Lyon

S. JUILLAGUET Maitre de Conférences Rapporteur

CNRS, U. Montpellier

J-F. MICHAUD Maitre de Conférences Examinateur

U. François-Rabelais, Tours

M. ZIELINSKI Docteur Ingénieur Co-directeur de Thèse

NOVASiC, Valbonne

Centre de Recherche sur l’Hétéroépitaxie et ses Applications, CNRS, Valbonne

Abstract

Investigation of dopant incorporation in silicon carbide epilayers grown by

chemical vapor deposition

This work is dedicated to the investigation of intentional dopant incorporation in silicon

carbide epilayers grown by chemical vapor deposition technique. The role of main process

conditions (growth temperature, dopant supply, deposition rate, growth pressure and C/Si ratio)

on both, Nitrogen and Aluminum incorporation was studied in details. Previous works have

widely explored the characteristics of dopant incorporation, especially the nitrogen incorporation

addressing a potential influence of growth equipment for the observed incorporation trends.

An exhaustive experimental study of N and Al incorporation was performed for

homoepitaxial 4H-SiC layers grown on Si- and C-faces of 4H-SiC substrates in our CVD setups

to explore such influence. It was completed by the assessment of the structural, optical and

electrical properties of the Al doped 4H-SiC films. Furthermore, the fabrication of pn diodes was

tested on the grown layers.

We have observed different experimental tendencies depending on dopant nature, crystal

orientation and chemical environment. We conclude from these observations that the mechanism

behind the experimentally obtained tendencies is widely influenced by factors such as process

conditions (i.e. growth temperature and/or pressure) and the carbon coverage at the grown

surface, especially on C-face.

Keywords: CVD epitaxy, 4H-SiC, N and Al incorporation, SIMS and C-V measurements, lattice

polarity, site competition, surface coverage, lattice constants, reciprocal space mapping (RSM),

LTPL measurements, carrier density, carrier mobility, PN diodes

Résumé

Etude de l’incorporation des dopants N et Al dans des films de carbure de silicium

épitaxiées en phase vapeur

Ce travail est consacré à l’étude de l’incorporation volontaire des dopants dans des films

de carbure de silicium épitaxiés par la technique de dépôt chimique en phase vapeur. Le rôle des

principaux paramètres de croissance (température, flux de dopant, vitesse de dépôt, pression dans

le réacteur et le rapport C/Si) sur l’incorporation d’azote et d’aluminium a été étudié en détail.

Les travaux menés jusqu’ici ont largement exploré les caractéristiques de l’incorporation de

dopants, en particulier l’incorporation d’azote et ont montré des résultats parfois très dépendants

de l’équipement de croissance utilisé. Afin d’explorer cette influence, une étude expérimentale

exhaustive sur l’incorporation de N et Al a été réalisée sur des couches homoépitaxiées 4H-SiC

sur la face carbone et sur la face silicium de substrats 4H-SiC dans nos réacteurs CVD. Cette

étude a été complétée par une analyse des propriétés structurales, optiques et électriques de

couches 4H-SiC dopé Al. Aussi, la fabrication de diodes pn a été expérimentée sur les couches

épitaxiées dans nos réacteurs.

Nous avons pu observer différentes tendances expérimentales selon la nature du dopant,

l’orientation cristalline du substrat et l’environnement chimique durant la croissance.

Nous en déduisons que le mécanisme derrière les tendances observées est largement influencé

par des facteurs comme les conditions de croissance (c'est-à-dire la température de croissance

et/ou la pression) et la couverture de carbone à la surface de la croissance, surtout sur la face C.

Mots-clés : Épitaxie CVD, 4H-SiC, incorporation N et Al, analyse SIMS et C-V, polarité du

matériau, compétition de site, couverture de surface, paramètre de maille, cartographie de

l’espace réciproque (RSM), mesure LTPL, porteurs libres, mobilité, diodes PN

Acknowledgements

The completion of this thesis could not have been possible without the contribution and

guidance of many people that I had the opportunity to meet and work with.

First and foremost, I would like to express my deepest thank you to my advisor Dr.

Marcin Zielinski who gave me the wonderful opportunity to be his first PhD student and work

under his guidance. I really appreciate all his continuous support, time, advice and ideas that had

helped and encouraged me to widen my research work. I am also thankful for his enthusiasm and

joy towards research which was very contagious and inspirational.

Besides, I would like to send my thanks to my second advisor Dr. Yvon Cordier for all

his help regarding administrative papers and fruitful discussions about the electrical

measurements of the epilayers and diodes. I really appreciate their helpful suggestions and

corrections in order to complete the present manuscript. I will like also to express my gratitude to

Dr. Jean-Yves Duboz for allowing me to stay and finish the writing of the present thesis at

CRHEA-CNRS laboratory.

I should also thank my NOVASiC colleagues, Dr. Thierry Chassagne for his help,

support and attention to details when it came to experiments, writing and presentations. To

Thierry Bourgoin and Marc Bussel for all their help in repairing the reactors as fast as possible

(Thierry, comment fais tu?) and characterization of the epilayers when the time was not my

friend (merci Marc!). Many thanks go to all the people who have been and are at CRHEA for

creating a nice environment to work too, for their help and support during my research. Special

thanks go to Sébastien Chenot for all the help with the diodes fabrication, Virginie Brändli for

her friendship, support and help with the SEM images and Rami Khazaka with whom the

scientific discussions about SiC growth were very stimulating.

My thanks also go out to the consortium of the European project NetFISiC (young

researchers and their supervisors) who shared their views and experiences in the field during

common shared trainings events and conferences. Moreover, the wonderful environment and the

friendships created are invaluable. Some special recognition should go to Pawel Kwasnicki and

Marilena Vivona for their valuable collaboration. Also, to their supervisors, Dr. Sandrine

Juillaguet (Laboratoire Charles Coulomb) and Dr. Fabrizio Roccaforte (Insitute for

Microelectronics and Microsystems) for hosting me at their institutions where I got the

opportunity to learn and use different techniques to characterize the optical and electrical

properties of the grown films.

I should also thank Dr. Sylvie Contreras and Dr. Leszek Konczewicz for all the Hall Effect

measurements performed on the Al-doped 4H-SiC films and their strong collaboration. I must

also thank Dr. Herve Peyre for his involvement regarding the SIMS measurements.

Lastly, I would like to express my deep graditude to my family and friends for their

support and for beliving in me during all this period. And most of all to Guillaume who was

patient, supportive and encouraging specially in the last part of this PhD. I really apreciate it.

To my beloved parents.

Table of contents

General introduction ..................................................................................................................... 1

Chapter 1 Introduction to Silicon Carbide material .................................................................. 5

1.1. General remarks ..................................................................................................................... 5

1.2. History of Silicon Carbide material ...................................................................................... 8

1.3. SiC crystal structure and physical properties .................................................................... 10

1.3.1. Crystal/Polytypes/Notation/Hexagonality ....................................................................... 10

1.3.2. Band structure and related electrical properties .............................................................. 14

1.3.3. Other properties ............................................................................................................... 17

1.4. State of the art of SiC material and elaboration techniques ............................................. 18

1.4.1. Bulk growth methods ...................................................................................................... 18

1.4.1.1. Physical Vapor Transport and related growth methods............................................ 18

1.4.1.2. High Temperature Chemical Vapor Deposition method .......................................... 20

1.4.1.3. Liquid phase growth ................................................................................................. 21

1.4.2. Current status of 4H-SiC substrates ................................................................................ 21

1.4.2.1. Size evolution ........................................................................................................... 21

1.4.2.2. Defect density ........................................................................................................... 22

1.4.3. Epitaxial growth methods................................................................................................ 23

1.4.3.1 Geometry of reactors ................................................................................................. 25

1.4.3.2 Chamber heating ........................................................................................................ 25

1.4.3.3 Capacity of the process chamber ............................................................................... 27

1.4.4. Epitaxial process (CVD) ................................................................................................. 30

1.4.4.1. Chemistry/ deposition rates ...................................................................................... 30

1.4.4.2. Doping of SiC ........................................................................................................... 31

1.4.4.3. Homoepitaxial growth of -SiC ............................................................................... 33

1.4.4.4. Si-face versus C-face ................................................................................................ 36

1.4.4.5. Heteroepitaxial growth of -SiC .............................................................................. 37

1.4.5. Current status of SiC epilayers ........................................................................................ 40

1.5. State of the art of SiC devices .............................................................................................. 41

1.6. Conclusions ............................................................................................................................ 45

1.7. Bibliography .......................................................................................................................... 46

Chapter 2 Growth process of SiC epilayers and characterization techniques ....................... 55

2.1. Introduction ........................................................................................................................... 55

2.2. Hot wall Chemical vapor deposition (CVD) reactor ......................................................... 55

2.2.1. Operational principle ....................................................................................................... 55

2.2.2. Growth of SiC in a hot-wall CVD reactor....................................................................... 56

2.2.2.1. Boundary layer model .............................................................................................. 57

2.2.3 Kinetics influence ............................................................................................................. 62

2.2.3.1. Temperature influence on growth rate ...................................................................... 62

2.2.3.2 Precursor supply influence on the growth rate .......................................................... 65

2.2.3.3. C/Si ratio influence on growth rate ........................................................................... 67

2.2.3.4. Growth pressure influence on growth rate ............................................................... 69

2.2.4. Doping of the SiC layers in CVD reactor ....................................................................... 71

2.2.4.1. Trimethylaluminum (TMA) as p-type precursor gas ............................................... 71

2.2.4.2. Doping incorporation ................................................................................................ 74

2.3. Characterization techniques ................................................................................................ 76

2.3.1. Surface inspection techniques: ........................................................................................ 76

2.3.1.1. Optical Microscopy (OM) ........................................................................................ 76

2.3.1.2. Scanning Electron Microscopy (SEM) ..................................................................... 77

2.3.1.3. Atomic Force Microscopy (AFM) ............................................................................ 78

2.3.2. Optical characterization techniques: ............................................................................... 81

2.3.2.1. Fourier Transform Infrared Spectroscopy (FTIR) .................................................... 81

2.3.2.2. Low Temperature PhotoLuminescence (LTPL) ....................................................... 84

2.3.2.3. Raman scattering spectroscopy ................................................................................ 89

2.3.3 Structural characterization: X-ray Diffraction ................................................................. 92

2.3.4. Chemical characterization: Secondary Ion Mass Spectrometry (SIMS) ......................... 95

2.3.5 Electrical characterization ................................................................................................ 98

2.3.5.1 Capacitance-Voltage technique (C-V) ....................................................................... 99

2.3.5.2 Hall Effect measurements ........................................................................................ 104

2.4. Conclusions .......................................................................................................................... 108

2.5. Bibliography ........................................................................................................................ 108

Chapter 3 Dopant incorporation in SiC epitaxial layers: the case of nitrogen and aluminum113

3.1. Introduction ......................................................................................................................... 113

3.2. Evolution of dopant incorporation .................................................................................... 114

3.2.1 Influence of polytype and off-angle ............................................................................... 114

3.2.1.1 Nitrogen case ........................................................................................................... 115

3.2.1.2 Aluminum case ........................................................................................................ 118

3.2.2 Influence of crystal orientation ...................................................................................... 121

3.2.2.1 Nitrogen case ........................................................................................................... 122

3.2.2.2 Aluminum case ........................................................................................................ 123

3.3 Dopant incorporation trends .............................................................................................. 125

3.3.1 Growth temperature dependence .................................................................................... 125

3.3.1.1 Nitrogen case ........................................................................................................... 125

3.3.1.2 Aluminum case ........................................................................................................ 127

3.3.1.3 Discussions .............................................................................................................. 129

3.3.2 Dopant supply dependence ............................................................................................ 131

3.3.2.1 Nitrogen flow rate .................................................................................................... 131

3.3.2.2 TMA flow rate ......................................................................................................... 132

3.3.2.3 Discussions .............................................................................................................. 133

3.3.3 Growth rate dependence ................................................................................................. 135

3.3.3.1 Nitrogen case ........................................................................................................... 135

3.3.3.2 Aluminum case ........................................................................................................ 136

3.3.3.3 Discussions .............................................................................................................. 137

3.3.4 Reactor pressure dependence ......................................................................................... 138

3.3.4.1 Nitrogen case ........................................................................................................... 138

3.3.4.2 Aluminum case ........................................................................................................ 139

3.3.4.3 Discussions .............................................................................................................. 140

3.3.5 C/Si ratio dependence..................................................................................................... 141

3.3.5.1 Nitrogen case ........................................................................................................... 142

3.3.5.2 Aluminum case ........................................................................................................ 144

3.3.5.3 Discussions .............................................................................................................. 147

3.4 Conclusions ........................................................................................................................... 150

3.5 Bibliography ......................................................................................................................... 152

Chapter 4 Material properties .................................................................................................. 156

4.1 Introduction .......................................................................................................................... 156

4.2. Samples preparation ........................................................................................................... 157

4.2.1 Doping estimation .......................................................................................................... 159

4.3. Surface morphology ............................................................................................................ 160

4.4. Structural characterization ................................................................................................ 164

4.4.1. X-ray rocking curve measurements ............................................................................... 164

4.4.2. Determination of lattice constants ................................................................................. 165

4.4.3 Reciprocal space mapping (RSM).................................................................................. 169

4.4.4 Polytype analysis ............................................................................................................ 172

4.4.4.1 Discussions: 3C-SiC polytype ................................................................................. 174

4.5 Physical properties of Al doped epilayers .......................................................................... 179

4.5.1 Optical properties ........................................................................................................... 179

4.5.1.1. Low Temperature Photoluminescence (LTPL) measurements .............................. 179

4.5.2 Electrical properties........................................................................................................ 182

4.5.2.1 Hall Effect measurements ........................................................................................ 182

4.6 Conclusions ........................................................................................................................... 190

4.7. Bibliography ........................................................................................................................ 191

Chapter 5 Fabrication of PN junction power diodes .............................................................. 195

5.1. Introduction ......................................................................................................................... 195

5.2. Fundamentals of PN junction ............................................................................................ 196

5.2.1. PN junction at thermodynamic equilibrium .................................................................. 197

5.2.2. P-N junction under forward bias ................................................................................... 198

5.2.3. P-N junction under reverse bias .................................................................................... 199

5.2.4 Current-Voltage (I-V) characteristics of P-N junction ................................................... 200

5.3. Fabrication of SiC P+-N junction diodes ........................................................................... 203

5.3.1. Epitaxial process ........................................................................................................... 203

5.3.2 Sample cleaning ............................................................................................................. 204

5.3.3. Diode fabrication procedure .......................................................................................... 205

5.3.3.1 Photolithography ..................................................................................................... 205

5.3.3.2 Etching ..................................................................................................................... 206

5.3.3.3 Ohmic contact deposition ........................................................................................ 208

5.4 Electrical analysis of the diodes .......................................................................................... 211

5.4.1 Linear TLM structures ................................................................................................... 214

5.4.2. I-V measurements of the P+-N and PiN junction diodes ............................................... 216

5.5 Conclusions ........................................................................................................................... 219

5.6. Bibliography ........................................................................................................................ 220

General conclusions and perspectives ...................................................................................... 224

Annex I Growth of 3C-SiC epilayers on -SiC substrates ..................................................... 227

Annex II Discussions: growth temperature dependence of Al and N incorporation ........... 236

Annex III Surface coverage....................................................................................................... 242

Introduction générale ................................................................................................................ 244

Résume détaillées ....................................................................................................................... 247

Conclusions générales et perspectives ...................................................................................... 254

General introduction

1


Nowadays, we are leaving in the “technology age” where the electronic devices are

surrounding our life. With the technology progress more energy is consumed which brings us to

the important problem of this century: power consumption. In order to efficiently reduce the

power consumption and dissipation we have to look towards the performances of semiconductors

devices. The electronics market for decades now, uses the Silicon (Si) semiconductor as the

principal semiconductor material for power devices fabrication due to the relative mature

technology. However, the industry demands are increasing towards new applications such as

photovoltaics and wind turbines, motor drivers, railcar traction and electrical vehicles, and Si

seems to achieve its material limitation which restricts the overall applications.

To continue the development of electronic devices, innovative solutions have been

considered and Wide BandGap Semiconductors (WBG) have made their entry into the power

electronics industry. Silicon carbide (SiC) is one of the WBG semiconductors which possess

intriguing physical properties which can lead to higher power density, higher junction

temperature and higher voltage capabilities.

SiC has been extensively studied for decades in academic institutions and at industrial

level, but recent advancements in SiC growth technology (commercially available 6-inch 4H-SiC

wafers with reduced defect density and zero micropipes) demonstrate the material benefits for

power electronics regarding power energy, efficiency, performance and a reduction of device

costs. All this progress in material technology attracted more attention from industrial level as

more and more SiC power devices are reaching the market.

Although technological breakthroughs are evident, the SiC epitaxial growth still confronts

with several important challenges that are limiting the device performances. To reduce such

limitations, we need to continue the understanding of growth mechanisms and impurity

incorporation. Doping of SiC is performed usually during in-situ epitaxial growth in order to

adjust the properties of the grown material to the device requirements. Therefore, process

optimization and comprehension of doping mechanism in the epitaxial layers are necessary.

This thesis aims to bring new insights into the doping mechanism in the field of epitaxial

growth of SiC by chemical vapor deposition technique. Most of the efforts have been oriented

towards the study of the influence of growth parameters on Nitrogen and Aluminum doping of


2

4H-SiC films. Moreover, the work presented in this manuscript has been carried in close

collaboration with two partners from the NetFISiC (Network on Functional Interfaces for Silicon

Carbide) European project: Laboratoire Charles Coulomb (L2C) in Montpellier and the Institute

for Microelectronics and Microsystem (CNR-IMM) in Catania, since most of the work was

performed in the framework of this project.

The manuscript is structured in five chapters, as follows:

In the first chapter an overview of the SiC material is presented. It includes details about

the chemical, physical and electrical properties of the most used SiC polytypes in

microelectronics. The bulk and epitaxial methods for SiC growth have been introduced, focusing

on the significant progress and challenges regarding the growth of both 4H-SiC and 3C-SiC

polytype during the years. Finally, the current status of the SiC epilayers and devices was also

presented.

The first part of the second chapter was attributed to the description of the hot-wall

chemical vapor deposition (CVD) reactor and the growth process of epitaxial layers. A

qualitative model was applied to understand the SiC growth kinetics inside the growth chamber,

with a particular attention to the influence of the growth kinetics on deposition rate. Moreover, an

attention was given to the impurity sources used for n- and p-type doping of the layers. The

second part of the chapter presents the different characterization techniques used to analyze

various features of the grown epilayers.

The third chapter was assigned to study the trends of voluntary incorporation of Nitrogen

(N) and Aluminum (Al) in 4H-SiC layers grown on both Si- and C-face 4H-SiC substrates. The

beginning of the chapter presents the influence of polytype, substrate off-angle and crystal

orientation on dopant incorporation. Further on, the influence of various process parameters such

as: growth temperature, dopant flow rate, growth rate, reactor pressure and C/Si ratio on N and Al

incorporation is described and discussed. This study helps to have a more complete vision of the

dopant incorporation mechanism during CVD growth, as efforts were made to explain the

experimentally obtained tendencies based on the mechanisms of adsorption/desorption at the

grown surface.

Chapter four describes the structural, optical and electrical properties of the Al doped 4H-

SiC epilayers. The impact of Al incorporation on the lattice constants and surface morphology of


3

the grown layers has been investigated. Besides, the optical and electrical properties of Al doped

4H-SiC layers have been also explored and displayed herein.

The fifth chapter, which is the last chapter of this manuscript, includes the fabrication and

electrical analysis of p-n junction diodes prepared on various substrates.

The manuscript is closed by some general conclusions of the main results obtained during this

study and perspectives of further studies.

Chapter 1 Introduction to Silicon Carbide

4

Contents

Chapter 1 Introduction to Silicon Carbide material ................................................................................ 5

1.1. General remarks ................................................................................................................................... 5

1.2. History of Silicon Carbide material .................................................................................................... 8

1.3. SiC crystal structure and physical properties.................................................................................. 10

1.3.1. Crystal/Polytypes/Notation/Hexagonality ..................................................................................... 10

1.3.2. Band structure and related electrical properties ............................................................................ 14

1.3.3. Other properties ............................................................................................................................. 17

1.4. State of the art of SiC material and elaboration techniques .......................................................... 18

1.4.1. Bulk growth methods .................................................................................................................... 18

1.4.1.1. Physical Vapor Transport and related growth methods .......................................................... 18

1.4.1.2. High Temperature Chemical Vapor Deposition method ........................................................ 20

1.4.1.3. Liquid phase growth ............................................................................................................... 21

1.4.2. Current status of 4H-SiC substrates .............................................................................................. 21

1.4.2.1. Size evolution ......................................................................................................................... 21

1.4.2.2. Defect density ......................................................................................................................... 22

1.4.3. Epitaxial growth methods .............................................................................................................. 23

1.4.3.1 Geometry of reactors ............................................................................................................... 25

1.4.3.2 Chamber heating ...................................................................................................................... 25

1.4.3.3 Capacity of the process chamber ............................................................................................. 27

1.4.4. Epitaxial process (CVD) ............................................................................................................... 30

1.4.4.1. Chemistry/ deposition rates .................................................................................................... 30

1.4.4.2. Doping of SiC ......................................................................................................................... 31

1.4.4.3. Homoepitaxial growth of -SiC ............................................................................................. 33

1.4.4.4. Si-face versus C-face .............................................................................................................. 36

1.4.4.5. Heteroepitaxial growth of -SiC ............................................................................................ 37

1.4.5. Current status of SiC epilayers ...................................................................................................... 40

1.5. State of the art of SiC devices ............................................................................................................ 41

1.6. Conclusions ......................................................................................................................................... 45

1.7. Bibliography .......................................................................................................................... 46


5

Chapter 1 Introduction to Silicon Carbide material

1.1. General remarks

Nowadays, we found ourselves being described as a “Connected” generation where the

technology is expanding, and the development and use of new devices bring higher energy

consumption. Accordingly, the most critical question of this century is how to lower the energy

consumption, to conserve the fossil fuels and reduced environmental pollution. The answer to

such a substantial question can be given by the materials namely, the wide bandgap

semiconductors (e.g. silicon carbide, gallium nitride, zinc oxide, and diamond) which offer strong

capabilities to revolutionize the modern power electronics and get clean energy innovations [1].

A rational use of electrical energy can be attained with improvement in power electronics

efficiency. Among the possible wide-bandgap semiconductor candidates we have in hand Silicon

Carbide (SiC) and Gallium Nitride (GaN) which attracted increasing attention in the last decades

due to its impressive intrinsic properties as shown in fig. 1.1. Currently, Silicon (Si) dominates

the power electronics field due to a mature and well established technology. However, Si device

technology approached its physical limitation and the adoption of SiC semiconductor can enable

alternative power device designs for faster and more energy efficient applications [2].

Fig. 1.1 Comparison of material properties among Si, GaN and SiC [3]


6

Of course from the theoretical characteristics display by fig.1.1, GaN seems to offer better

high-frequency and high-voltage performance. Although due to the low quality of bulk

substrates, the high-voltage applications were granted to the SiC technology. Even though the

industrial interest of both GaN and SiC technology is increasing daily, the SiC remains the

material of choice for power devices in high-voltage and high-temperature applications.

Figure 1.2 shows the application range of both GaN and SiC power devices nowadays. As

can be observed, GaN power devices found to be suitable for very high-frequency and low-

voltage operations. Therefore, GaN can be an alternative but cannot compete yet with SiC.

Fig. 1.2 Major applications areas of GaN and SiC power devices [4]

SiC power devices capabilities are proven day by day more advantageous and efficient,

showing their ability when employed for different applications such as reduction of resistive

losses, more efficient switching, higher operation temperatures, raise of the conversion efficiency

from the current and a boost on overall energy reduction.

The SiC power devices found use in applications such as: power supplies, motor control,

heating, robotics, electric/hybrid vehicles, electric power transmission, and traction, home

appliances, biomedical, electric power for sensors, communication, high precision navigation,


7

motor drivers and converters for missile systems, ground and air military vehicles, motor drivers

for space exploration and others [5]–[7].

Unfortunately, the SiC power semiconductor devices display a higher cost compared with their

silicon counterparts which reflects the difficulties in the material manufacturing process.

As can be observed from fig. 1.3 in which the material manufacturing process is represented until

the final product is obtained, five main process steps are required. Thus, the cost per wafer

accounts the substrate material-epitaxial layer-processing and packaging of SiC for power

electronics. A cost reduction can be achieved by an improvement of material quality and high

volume processing of wafers.

Fig. 1.3 Fabrication chain of SiC power devices

Although SiC substrates and epilayers are more expansive than Si, they still can compete

with Si devices regarding the cost/area, considering a higher performance and efficiency of SiC

devices.

My PhD thesis was elaborated in the framework of NetFiSiC project (ITN Marie Curie

action) which aimed to improve the functional interfaces of SiC material (of various polytypes)

for obtaining a step forward in electronic devices performance. My work was part of the work-

package number 1 (WP1) dedicated to material growth and related aspects. The objective of my

PhD thesis included the improvement of SiC epitaxial layers grown by Chemical Vapor

Deposition (CVD) technique for electronic applications. Hence, for attaining such an objective,

my work involved an elaborated study of n-type (Nitrogen) and p-type (Aluminum) doping of

4H-SiC and 3C-SiC epilayers to understand the dopant incorporation dependence on growth

parameters. Another part of my study was directed towards the exploration of 3C-SiC epilayers

grown on well-oriented hexagonal ( -SiC) substrates for an enlargement of the 3C-SiC domains.


8

However, the time attributed for this study was limited from where smaller amount of results

were acquired.

The motivation behind the topics studied during my PhD work is associated with the

requirements of the power industry to improve the electronic devices performance, in particular

the needs for high quality epilayers with specific doping density and thickness. Thus, both I and

NOVASiC set our interest in understanding the doping mechanism on both surface polarities of

SiC epilayers for obtaining the uniformity and dopant control demanded for electronic

applications. The cubic material attracted our attention due to its potential in electronic

applications, and the demand for perfecting the material which is inclined to high amount of

surface defects. That is why our target we set on obtaining single domain 3C-SiC layers grown on

6H- and respectively, 4H-SiC on-axis substrates.

Moreover, for NOVASiC it was an opportunity to gain a deeper understanding of p-type

(Aluminum) doping process and its memory effect. This made easier the transfer of Aluminum

doping process from a smaller reactor R1 (2 inch capacity) towards a bigger one R2 (4 inch

capacity). For me it was an opportunity to enlarge the domain of knowledge, to acquire more

information and understanding of everything that involved growing a high-quality SiC epilayer

(growth, doping, characterization and processing) that can boost the material future.

1.2. History of Silicon Carbide material

Silicon carbide (SiC) is a material that can be found in natural state, but only in minute

quantities in exceptional geological places (diamond containing chimneys, kimberlitic volcanic

openings or certain types of meteorite).

Originally, the existence of SiC was suggested in Sweden in 1824 by Jöns Jacob Berzelius who

proposed the idea that a Si-C chemical bond might exist after his experiment on diamond

synthesis [8]. The SiC material was first observed in laboratory experiments in 1849 by C.M.

Despretz, and only in 1891 it was elaborated at industrial scale by Eugene G. Acheson. While

attempting to synthesize diamond at high temperature (T>2000°C) from carbon and aluminum

silicate reaction, blue crystals were obtained and described under the general term

“carborundum” [9]. The term came as Acheson thought his synthesis compound was formed

from carbon and aluminum oxide (Al2O3) also called “corundum”. Since carborundum holds an


9

exceptional hardness (9–9.5 on the Mohs scale, compared to 10 for diamond), chemical inertness

and resistance to high temperatures, it become quickly an abrasive material commercialized by

the Carborundum Company.

In 1905, hexagonal grains of SiC were found by Henri Moissan in the Diablo Canyon

meteorite in Arizona and hence, the name Moissanite was given to natural SiC [10]. Moissanite is

found only in few locations on Earth and in small sizes, making its use difficult. Though, Charles

& Colvard® [11] created and patented a process to obtain Moissanite gemstones used in jewelry

industry. In Europe different restrictions depending on the country are given for Moissanite

jewelry sales due to the resemblance with diamond. For example, Moissanite jewelry distributed

by Charles & Colvard® can be purchased from UK.

In 1907, H.J. Round have observed the electroluminescence properties of the SiC,

allowing the SiC to be known as one of the oldest semiconductor materials [12].

For a long time the SiC was limited as an abrasive material due to the difficulties in obtaining

quality crystalline materials. The SiC research was stimulated after Lely discovery of the

sublimation process for production of high quality crystals in 1955 [13]. In the 1960s, different

studies to determine the optical properties of SiC polytypes were done by Choyke and coworkers

[14]. In 1978, the seeded sublimation method (a modified Lely method) for SiC growth was

developed by Tairov and Tsvetkov [15] bringing the attention back to SiC, as larger SiC

diameters wafers could be obtained. This impacted the development of CVD systems for

epitaxial SiC. In 1980s, Nishino et al. [16] demonstrated the heteroepitaxial growth of 3C-SiC on

Si by introducing the carbonization process. Another breakthrough for SiC was the step-

controlled epitaxy developed by Kuroda et al. [17] that allowed the homoepitaxial growth. In the

1990s, SiC was commercial release leading to an increase of the research and development

activity regarding the material and devices which lead to rapid improvements (e.g. site

competition, Schottky diodes). Infineon had started to commercialize the SiC Schottky Barrier

Diodes (SBDs) in 2001 [18] which resulted to a switching in interest from academic to industrial

domain and, from 2010 the SiC started his industrial period. This is confirmed from the presence

of industry at the ICSCRM (International Conference on Silicon Carbide and Related Materials

that debuted in 1987 in Washington D.C, USA) and the ECSCRM (European Conference on

Silicon Carbide and Related Materials) conferences. These companies have been active involved

in the development of the SiC market. Here are some examples regarding the SiC development:


10

Cree showed an increase of 4H-SiC substrates from 2 inch up to 150mm and epitaxial layers with

very low basal plane dislocations (density <1cm-2

). II‐VI Advanced Materials have introduced

since July 2015 the development of 200mm SiC substrates, which will reflect on the final cost of

power devices. Infineon showed an increase in efficiency and performance of SBDs and of the

transistors such as Metal-Oxide Semiconductor Field Effect Transistor (Cool SiC TM

MOSFETs).

1.3. SiC crystal structure and physical properties

1.3.1. Crystal/Polytypes/Notation/Hexagonality

Silicon carbide (SiC) is an interesting material since the crystal structure gives rise to a

special case of polymorphism, called polytypism, which allows a multitude of SiC crystal lattices

(polytypes) to form. The polymorphism represents the ability of a material compound to occur

under different structures maintaining the same chemical composition. However, when we talk

about polytypism, the actual difference between the polytypes is given by the stacking sequence

along one direction. When is coming to SiC more than 250 polytypes can be formed [19].

The SiC crystal is created in equal proportions by silicon and carbon atoms arranged in a

tetrahedron. Each silicon atom (or carbon atom) is bounded to four carbon atoms (or silicon

atoms) arranged in tetrahedral configuration (valence band filled with 8 electrons) as represented

in fig 1.4. The lattice constant “ ” of the Si-C tetrahedral (i.e. distance between Si-Si or C-C

bond) is about 3.08 while the distance of Si-C bound is almost 1.89 [20]. Further, the bonding

between Si and C atoms is considered to be 88% covalent and only 12% ionic [21].

SiC crystal structure with 0001 planes is actually a polar material in which the surfaces

can be terminated by either silicon (Si) or carbon (C) atoms. The developed polar surfaces are

usually marked as Si-face (or indicated (0001) SiC) or C-face (or indicated (000 ) SiC) as

represented in fig 1.4(b-c).


11

Fig.1.4 (a) Basic building block of SiC crystal structure with a tetrahedron having a Si-atom in

the middle bounded to four C-atoms; (b) Si-face; (c) C-face of a SiC crystal

During the stacking of the tetrahedral Si-C bonds along the c-axis (i.e. perpendicular to

the compact plan [11 0]) two different orientations are possible: either identical or rotated by 60°

with respect to each other are represented in fig. 1.5. The rotation of the SiC tetrahedron gives

rise to a twin structure (fig. 1.5b) that mirrors the normal one, and if we stack the same type of

tetrahedron, the cubic structure (3C-SiC polytype) can be obtained.

Fig. 1.5 (a) SiC tetrahedron with the Si-C bond along the c-axis and (b) twined SiC tetrahedron

rotated by 60° around the c-axis

The stacking sequence of Si-C bilayers along the c-axis determines the material polytype.

Thus, the Si-C bilayers are denoted by letters such as A, B or C which identifies their stacking

position in the (0001) plane. The SiC crystal is thought to be composed at the base by

hexagonally closed packed layers stacked one on top of the other. If the A represents the position

of the first Si-C bilayer, the next bilayer can get the B or C position which directs the bottom


12

atoms on top of hexagonal hallow of the upper atoms of the reference layer [22]. For example if a

stacking of the bilayers started with AB than the third bilayer can grow positioned A or C.

A notation of the different SiC polytypes was proposed by Lewis S. Ramsdell [23]. It consists of

a number corresponding to the bilayers existing in one repeating unit and a letter describing the

crystal structure (C for cubic, H for hexagonal and R for rhombohedral). The common polytypes

are the 3C-SiC named also -SiC (ABC….. stacking), 4H-SiC (ABAC ….. stacking), 6H-SiC

(ABCACB ….. stacking), 2H (AB… stacking) and 15R-SiC (ABCACBCABACABCB…

stacking) that are described in fig. 1.6 and table 1.1.

A difference between -SiC (describing the cubic polytype) and -SiC (describing the hexagonal

and rhombohedral polytypes) [22] lies also in the crystallographic structure (polytype dependent).

The cubic structure (zincblende) denotes the 3C, the hexagonal structure (wurtzite) denotes the

2H polytype structure while all the rest of polytypes are a mix of cubic and hexagonal structure

[24], respectively.

Fig 1.6 Stacking sequences for the three most common SiC polytypes represented in the [11 0]

plane [24].

Due to the stacking possibilities of Si-C bilayers on top of each other, there exist in the

SiC unit cell two kinds of available sites which are not equivalent. These kind of available lattice

sites are differentiated by their second-nearest neighbor. In other words if the bilayer stacking is

made having an antisymmetric bond with the nearest neighbor, the atom resides on a cubic site


13

(denoted k), while if the bonding is symmetrical the atom resides on a hexagonal site (denoted h)

[25] as illustrated in fig 1.7. The 3C-SiC polytype contains only cubic (k) sites, the 4H-SiC

polytype contains a cubic (k) and a hexagonal (h) site and the 6H-SiC holds two cubic sites

(k1,k2) and an hexagonal site (h). The relative proportion between the cubic and hexagonal sites

along the c-axis allows the hexagonality rate H of a certain polytype to be determined as follows:

(1.1)

where h and k are the number of cubic and hexagonal sites, respectively existing in the unit cell.

Fig 1.7 Illustration of the Si-C double atomic layer arrangement with their cubic and hexagonal

crystal symmetry along the c-axis direction for the three most common polytypes [26]

From all known SiC polytypes, only three of them (i.e. 4H-, 6H- and 3C-SiC) are used

today for technological purposes in electronic device applications. At industrial level, the most

widely used polytypes are the 6H-SiC and 4H-SiC. However, since the availability of 4H-SiC

increases, it has replaced almost totally the use of 6H-SiC in electronic device applications. Yet,

the 6H-SiC found its place as a substrate used for the growth of different wide-bandgap

semiconductors such as Gallium or Aluminum Nitride. Regarding the 3C-SiC polytype, until now

from my knowledge no real bulk crystal was grown due to lack of 3C-SiC seeds. Although, the


14

3C-SiC polytype can be found as free-standing wafers grown on Si substrates or only as small

crystals (few mm²).

Table 1.1 contains information regarding the crystallographic characteristics of different

SiC polytypes.

Polytype Unit cell

sequence

Lattice parameters

(nm)

Inequivalent

sites

Structure Hexagonality

rate (%)

2H-SiC AB a=0.3081 c=0.5048 1 hexagonal

(wurtzite)

100

3C-SiC ABC a=0.4349

a=0.3081c=0.75

1 cubic(zincblende)

hexagonal (hcp)

0

4H-SiC ABAC a=0.3081 c=1.08 2 Hexagonal closed

packed (hcp)

50

6H-SiC ABCACB a=0.3081c=1.501 3 hexagonal (hcp) 33

15R-SiC ABCACBCABACABCB a=0.3081 c=3.77 5 rhombohedral 40

Table 1.1 Crystallographic characteristics of the SiC polytypes

1.3.2. Band structure and related electrical properties

In the recent years the Si-based power devices started to show material limitations and

most of the attention returned to Wide Bandgap (WBG) semiconductors such as Silicon Carbide.

The SiC captured an increase interest in electronic industry due to its unique physical and

chemical characteristics which makes it suitable and advantageous for power device fabrication.

The extraordinary electrical properties of SiC include wide bandgap, high breakdown field

strength, high thermal conductivity and high saturated electron drift velocity compared with their

Si counterparts.

If we focus on the electrical properties of the SiC semiconductor, a variation of these

properties is observed from one polytype to the other. The indirect band gap varies in the range

2.3 to 3.3eV when passes from cubic structure of 3C-SiC to the wurtzite structure of 2H-SiC

polytype. The SiC wide bandgap brings extremely low intrinsic carrier density (for 4H-SiC

=5x10-9

cm-3

at 300K) compared with the value of Si with =1x1010

cm-3

. Consequently, SiC is


15

an attractive candidate for high-temperature applications (in theory even above 1000K) without

suffering from intrinsic conduction effects [5].

SiC also exhibits ~8-10 times higher breakdown electrical field strengths ( )

compared with Si, which reflects that SiC can handle higher voltages before undergoes an

avalanche breakdown. Therefore, the thickness of the SiC power devices can be reduced to

approximately one tenth than the one needed in Si devices for sustaining the same voltage. For

similar blocking voltage, the doping density on SiC layer can be two orders of magnitude higher

than on Si layer. This offers an advantage in manufacturing high-voltage devices with lower

specific on-resistance ( or drift-layer resistance) compared with Si devices. Lower drift

resistance of SiC devices leads to a reduction of the on-state losses ( ) and higher converting

efficiency. This enables the fabrication of SiC unipolar devices such as SBDs, Metal-Oxide

Semiconductor Field-Effect Transistors (MOSFETs) and Junction Gate Field-Effect Transistor

(JFETs). They are suitable for medium-voltage applications due to their higher breakdown

voltages compared to Si unipolar devices as shown in fig 1.8. For medium and high-voltage,

bipolar device structures such as PiN diodes, Bipolar Junction Transistors (BJTs) and even

thyristors where enabled (see fig.1.8).

Fig. 1.8 Major domains of the individual unipolar and bipolar power devices for Si and SiC in

terms of rated blocking voltage [5]

The higher electron drift velocity is two times higher than in silicon, allowing fast

switching frequencies and operations. Due to thinner voltage blocking layers the minority carrier


16

storage is reduced, which bring faster switching and lower reverse recovery losses. Figure 1.8

presents the working blocking voltage domains for the SiC and Si unipolar and bipolar devices.

As can be observed, the SiC devices can withstand higher voltages (>20kV) compared with Si

devices making them more attractive for ultrahigh voltage applications. It is expected that SiC

unipolar devices will replace Si bipolar devices in the blocking voltage range from 300V to

~6500V. The bipolar devices predicted to increase towards >10kV blocking voltages in the near

future [5].

The main physical properties of the most stable SiC polytypes in electronic industry and

their semiconductor counterparts such as Si and GaN are represented in table 1.2.

Property Si 4H-SiC 6H-SiC 3C-SiC GaN

Bandgap Eg (eV) 1.12 3.2 3.0 2.3 3.4

Relative dielectric constant εs 11.9 9.7 9.7 9.7 9.5

Breakdown field EB

ND=1017

cm-3

(MVcm-1

)

0.6 //c-axis: 3.0

c-axis:2.5

//c-axis: 3.2

c-axis: >1

1.8 2-3

Thermal conductivity

(W/cm-K)

1.5 3-5 3-5 3-5 1.3

Electron mobility at ND=1016

cm-3

(cm2V

-1s

-1)

1200 800 //c-axis: 600

c-axis:400

800 900

Hole mobility at NA=1016

cm-3

(cm2V

-1s

-1)

420 115 90 40 200

Saturated electron velocity

(107 cm s

-1)

1.0 2 2 2.5 2.5

Thermal expansion coefficient

(x10-6

K-1

)

2.59 //c-axis: 3.2

c-axis:3.3

//c-axis: 4.7

c-axis:4.3

3.8 //c-axis: 3.17

c-axis: 5.59

Melting /Sublimation Temperature

(°C)

1414 2830 2830 2830 2500

Table 1.2 Main physical properties of the most common SiC polytypes (4H, 6H and 3C-

SiC) and their counterparts Si and GaN at room temperature [27]

Even though, the 3C-SiC quality is not as good as the one of hexagonal polytypes it still

presents a wide interest due to low temperature stability and its superior saturated electron

velocity compared with their -SiC counterparts, as shown in table 1.2. The 3C-SiC layers on Si

substrates were grown in order to take advantages of the promising properties at a lower cost.

Except that the low crystalline quality hampered the use for device fabrication, although they are

used for MicroElectroMechanical Systems (MEMS) applications. Moreover, 3C-SiC layers


17

grown on -SiC substrates found their use also in fabrication of MOSFETs devices for power

switching applications [28]. They seem to be a good candidate due to smaller ionization energy

and lower state density at the 3C-SiC/SiO2 interface compared with 4H-SiC.

The advantages of the three most stable SiC polytypes brought a boost and stronger

interest for electrical devices in high-temperature, high-power and high-voltage applications such

as: computer and cell-phones industry, motor and traction control, smart power grid, electrical

vehicles, power converters for renewable energies, energy storage, naval power system, solar

cells, medical and biotechnology, harsh environment sensors [29].

1.3.3. Other properties

The common physical properties such as mechanical, chemical and thermal ones of all

SiC polytypes are given by the specific nature of the Si-C chemical bond. The SiC is a hard

material provided by the short bond length between Si and C atoms (of 1.89 ), which makes it a

perfect candidate for cutting tools or abrasive/protective material. The high thermal conductivity

presented by SiC material promotes the fabrication of electronic devices that sustain high-

temperatures (>250°C). Theoretically, the operating temperature for SiC devices can go up to

700°C although, the performance decreases due to a lack in packaging technology. It is a very

important characteristic which attracts the use of SiC power devices for high-temperature and

high-power applications (e.g. car engines). Moreover, leads to a decrease of the system heat,

impacting the overall system volume and cost due to smaller/lighter cooling systems. Considering

this property, the SiC becomes an important material in electronic components for engine and

cooling cycles in hybrid vehicle system and railcars (e.g. Tokyo Metro's Ginza Line subway) as

reported by Toyota and Mitsubishi [30], [31].

SiC material displays also a chemical inertia, which conducts impossible the reactions at

room temperature with the known chemical substances, making it a good corrosion resistant

material. However, an etching of SiC at higher temperatures of around 600°C with Molten

Potassium Hydroxide (KOH) was found to be feasible. This property makes it a perfect candidate

for harsh environments applications such as: combustion environments, space exploration, energy

and chemical production, aviation and others [20], [29]. The mechanical properties of SiC are

attractive for microelectronics due to its high Young’s modulus that ranges between 330 and

700GPa [32] depending on the polytype, which makes it an intriguing material for Micro-Electro-


18

Mechanical-System (MEMS) applications. In particular, the 3C-SiC grown on Si substrates found

out to be a perfect candidate for MEMS applications, especially for fabrication of Atomic Force

Microscopy (AFM) cantilevers [33]. SiC material is also known as a biocompatible material that

can be used for biochemical or medical sensors. It has been reported the use of SiC material as

hard coating for numerous biomedical applications: non-fouling coatings on coronary heart stents

and/or a passivation layer for prosthetic bone implants [34]. Moreover, SiC material found its use

as nuclear fuel particle coatings in Nuclear Power Plants.

1.4. State of the art of SiC material and elaboration techniques

Modern semiconductor technology relies on the crystal growth capacities to achieve high

quality wafers that can boost the performances of the SiC devices. Although, to reach such

performances we need first an established growth process that produces high quality crystals to

be further used as substrates. However, SiC crystal growth turns out to be extremely difficult due

to physical-chemical nature of the material, hampering the development of the SiC

semiconductor devices and their applications in last century.

1.4.1. Bulk growth methods

With time, alternative growth methods for obtaining single crystalline SiC bulk material

were found, that includes solution phase growth from non-stoichiometric solutions and vapor

phase growth, respectively [20].

1.4.1.1. Physical Vapor Transport and related growth methods

Nowadays, the process of choice especially in industry for obtaining volume production

of SiC semiconductor wafers is the Physical Vapor Transport (PVT) method. The PVT method

also referred to as seeded sublimation growth (Modified Lely Method) is the most common

technique, finding its fundamentals in the works of Lely [13] and Tairov and Tsvetkov [15]. The

PVT technique [35] is based on the chemical process of sublimation where the main challenges in

process optimization are related to temperature and gas species concentration control. The PVT

process is carried out in a closed inductively heated graphite crucible surrounded by graphite

foam isolation, where the growth temperature exceeds 2000°C. An argon atmosphere and


19

pressures in the range 15-50 mbar are used. Inside the crucible, the SiC powder presented at the

bottom is heated and begins to decompose to gaseous species (e.g. Si, Si2C and SiC2) that will

condense on the crystalline seed fixed onto top part as represented in fig. 1.9. The process driving

force is actually the temperature gradient between the heated polycrystalline SiC powder and the

somewhat cooler SiC seed.

The advantages of this bulk process method are the crystal growth rate that typically

ranges from 200-1000µm/h and the size of the crystals that have an extended diameter. The

disadvantages include the induced stress of material due to thermal gradient, polytype stability

and defects. Recent studies showed an improvement of the PVT capabilities due to simulation

and in-situ growth monitoring of the crystal interface and shape by 2D and 3D x-ray imaging

[36]. All this improvements allow a better process control and doping [37], polytype stability and

defects reductions (such as micropipes, stacking faults, different dislocations) [38]–[40].

Fig.1.9 Schematic representation of a PVT reactor with representative axial temperature

profile [29]

Other vapor growth techniques of bulk SiC crystal have been developed aiming to tackle

different technological challenges that PVT methods present. Growth techniques such as

Modified PVT (M-PVT) [41] and Continuous Feed PVT (CF-PVT) [42] were developed.

However, their technological processes are still in early stage of development finding their use

mostly at academic level.


20

1.4.1.2. High Temperature Chemical Vapor Deposition method

Another alternative method is the so called High Temperature Chemical Vapor

Deposition (HT-CVD) in which high purity SiC crystals are obtained due to high purity gases

introduced during growth [43]. Comparable with the sublimation technique, the HT-CVD method

takes advantage of the high purity of the precursor gases such as silane (SiH4) and ethylene

(C2H4), and of the more homogeneous temperature ( =1800-2300°C) profile for growing the

SiC crystal. Another key component of the method is the continuous supplying of source material

which gives direct control over the C/Si ratio and doping incorporation. Figure 1.10 shows the

representative geometry of a HT-CVD reactor with the corresponding temperature profile and

chemical stages.

Fig.1.10 Schematic representation of a HT-CVD reactor with representative temperature

profile and corresponding chemical stages [29]

The HT-CVD bulk growth process has been implemented at industrial level by

NORSTEL company which patented this growth technique for commercially purposes, producing

high resistivity semi-insulating 4H-SiC up to 100mm in diameter [44].

The Halide CVD (H-CVD) [45] is a recently introduced and promising bulk growth

technique (though it can be used for epitaxial growth also) for obtaining high purity SiC crystal

with low impurity levels and high electrical resistivity (producing semi-insulating wafers). This

method uses for growth precursor gases such as silicon tetrachloride (SiCl4) and propane (C3H8)

or methane (CH4). Although, compared with HT-CVD method, the H-CVD is in its early stage of

development, being used mostly at academic level.


21

1.4.1.3. Liquid phase growth

The liquid phase growth is a very promising growth technique for producing high quality

crystal SiC boules, once a controlled over the supersaturation at the grown surface is obtained.

Early efforts for liquid phase route to achieve SiC crystals were limited and not much progress

was made, but recently (in last 10 years) this method started to attract more attention. The most

feasible liquid phase method is the Top Seeded Solution Growth (TSSG) [46], [47] from the

perspective of large diameters and length of SiC boules. In the TSSG the graphite crucible acts

also as a carbon source that will be dissolved in the Si-based melt, and transported by diffusion

and convection towards the SiC seed. The driving force behind the growth is given by the

difference in temperature between the seed and the melt. Growth temperatures employed are in

the range of 1750-2100°C and an inert atmosphere (argon gas) is used. The difficulties

encountered are related to the lack of stoichiometric SiC liquid phase at atmospheric pressure, the

low solubility of carbon in the Si melt even at high temperature of 2800°C and the control over

the fluid convection in the melt. An alternative approach is the use of transition metal to Si melt

in order to increase the C solubility. However there is a drawback to this since the material gets

contaminated, resulting in lower purity of the crystal.

Apart from all the challenges the progress is quite rapid and advantages such as

production of crystals with reduce stacking fault densities, diameter enlargement and high

structural quality are already visible. The liquid phase growth is still in the early stage of

development which makes it viable at the level of academia, but the intensive progress places this

method towards the road for bulk SiC growth process [29], [48].

1.4.2. Current status of 4H-SiC substrates

The important factors for reducing the cost of SiC devices for different microelectronic

applications are the enhancement of the size and quality of SiC crystals.

1.4.2.1. Size evolution

Within the years, the continuous research development efforts from both industry and

academia have lead SiC wafer technology to new heights in terms of quality and size. During the

last two decades, viable process for substrates production (mature PVT growth process) was

available which enabled the progression of commercially SiC wafers with different diameters. In


22

fig. 1.11 it is presented the diameter progress of SiC crystals. Currently, 150mm diameter n-type

4H-SiC 4° off-axis substrates are commercially available [49], [50], [51], [52].

In 2015, at the International SiC Power Electronics Applications Workshop (ISiCPEAW-

2015) in Stockholm, Sweden and at the International Conference on Silicon Carbide and Related

Materials (ICSCRM-2015) the II-VI Advanced Materials and Cree Inc. have announced the

development of 200mm diameter SiC wafers [53]. It is clear that along the years the SiC wafer

technology has made tremendous progress that brought cost reduction of material and devices,

and moreover a boost of the SiC from R&D to volume production. The cost per SiC wafer

depends strongly on the supplier, material quality and the quantity ordered. Consequently, a 6

inches SiC substrate can reach prices as high as 2000 euro, but are continuing to decrease with

the increase in development.

Fig. 1.11 Development of 4H-SiC wafer diameter along the years

1.4.2.2. Defect density

Apart from increasing the SiC wafer diameter, the industry and academic attention was

directed towards defect reduction, as very low defects densities are required for power devices.

For achieving low defect density on a wafer, precise control of the heat and mass-transport

conditions during crystal growth are very important as described above. Micropipes which are

hollow core dislocation are an instant killer for SiC devices. Progress has been made for n-type

4H-SiC wafers with 100 and 150mm and now, wafers grown with low micropipe densities

0.1cm-2

or zero micropipes per cm2 offered by Cree Inc. are commercially available. Figure


23

1.12 displays the reduction of median micropipes density for 100mm diameters SiC wafers

grown by PVT method along the years.

Furthermore, studies of reducing the other defects such as Basal Plane Dislocations (BPDs),

Treading Edge (TEDs) or/and Screw Dislocations (TSDs) continue since new available wafers

with larger diameter are being commercialized [54]. Dow Corning has shown that the dislocation

density such as BPDs and TSDs are superior for 150mm compared to 100mm 4H-SiC wafers

(e.g. BPDs density in 100mm is 5.2x103cm

-2 compared to 11.7x10

3cm

-2 for 150mm wafer) [55].

Fig. 1.12 Reduction of median micropipe density on n-type 4H-SiC 100mm wafers along the

passing years [56]

If we compare the reduction of defects in SiC crystals with the ones in Si crystals well, it

has to be understood that SiC technology is still considered by the people from Si community as

“dirty” and “unstable”. Even though, SiC technology is not as mature as the Si, it is yet

competitive with the perfect Si semiconductor.

1.4.3. Epitaxial growth methods

The production of SiC devices for different electronic applications is determined by the

quality of the SiC material, since the active layers are obtained by epitaxial growth. This step is

2004 2005 2006 2007 2008 2009 2010 2011 20120

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su

bs

tra

tes

(c

m-2)

Year


24

required since the substrate has a high background doping and defects that can lead to electrical

quality limitation of the SiC for device performances. As a consequence, the development of SiC

devices relies upon the performance of epitaxial growth that can give rise to high volume

production capability of high quality and reproducible epitaxial SiC layers. The SiC epitaxial

growth can be performed homoepitaxially (i.e. similar substrate/layer) or heteroepitaxially (i.e.

substrate/layer different including 3C-SiC/ -SiC or Si substrates). Several techniques for

epitaxial growth of SiC that exist today are the CVD, sublimation, Liquid Phase Epitaxy (LPE),

Vapor-Liquid-Solid (VLS) epitaxy and Molecular Beam Epitaxy (MBE).

The SiC epitaxial process is usually performed by CVD [57] that is a Vapor Phase

Epitaxy (VPE) technique. In a CVD reactor, the growth is carried out under a mixture of gases

composed of precursors (such as silicon and carbon-containing gases) and a carrier gas (such as

hydrogen +/-argon) which are flowing inside a heated chamber. The mixture of gases inside the

chamber has a laminar flow, and over the heated susceptor a boundary layer is formed. As the

gases are heated they will start to decompose into reactive species that will diffuse through the

boundary layer and onto the substrate where, due to additional chemical reaction, the growth of

the epitaxial layer will take place. In CVD reactors, the growth temperature can be varied in the

range 1300-1800°C. At lower temperature, the growth of 3C-SiC on silicon or -substrates

(heteroepitaxy) can take place, while at growth temperature above 1500°C, the growth of 4H/6H-

SiC on similar substrates (homoepitaxy) is performed. The CVD epitaxy can take place under

Atmospheric Pressure (AP-CVD) or under Low Pressure (LP-CVD) [27]. The epitaxial growth of

SiC layers occurs at higher temperatures compared with the other WBG semiconductors.

Sublimation Epitaxy (SE) has a similar growth principle as standard PVT technique,

though the distance in epitaxy between the source and the substrate is reduced. The epitaxial

process with high growth rates can be achieved (e.g. 200µm/h), but a lack of precise control

makes difficult to achieve thickness and doping uniformities [58].

LPE is a technique that was widely used in the 1990s, but recently attracted the attention

since it can be used for eliminating or filling the micropipes. The growth process takes place due

to a temperature gradient applied between silicon melt (usually with and additional chemical

element to increase solubility) and the substrate [29]. Unfortunately, the inability to control the

doping level limits its use. VLS is a new technique compared with the previous ones described

above. The process takes place by decomposing the supplied hydrocarbon at the liquid vapor


25

interface that diffuses through the Si melt to the solid liquid interface leading to the growth. Even

though VLS method has its benefits, such as lower temperatures (<1400°C) and/or higher

aluminum incorporation, it’s still at the beginning and cannot compete yet with the CVD method

(i.e. a laboratory technique) [59]. MBE is another technique that can be used for growing very

thin layers under high vacuum conditions (precursors supplied by sputtering or solid source

heating) [60]. Even though a more precise control is possible due to in situ characterization, the

technique remains at laboratory scale due to reduced growth rate (<1µm/h).

1.4.3.1 Geometry of reactors

Presently, the technique of choice for the production of high-quality and reproducible

epitaxial layers for SiC power devices, in a commercial way, remains the CVD also called VPE.

Reactors used for CVD process can have different configurations that depend on the reactant gas

flow across the SiC substrate, namely the flow can be directed parallel or perpendicular to the

surface of the wafer. Reactors with vertical (chimney) or horizontal configurations have been

implemented and used for the CVD process. The horizontal configuration is the first

configuration developed for CVD reactors.

The chimney reactor is actually a vertical system who was developed by Ellison et al. [61]

and attracted attention due to advantages of the epitaxial process such as high growth rates. The

vertical geometry provides the mounting of the substrate on both inner walls of the reactor that

gives a minimization of thermal gradient close to growth area. The gas flow is added upward into

chamber, and transported by free convection due to a high temperature process.

1.4.3.2 Chamber heating

The CVD reactors configuration can be further categorized by the implemented

temperature that surrounds the SiC wafer dividing them into different groups: “cold-, hot- and

warm-walls” reactors. Up to mid-1990s the most common configuration for SiC process was the

cold-wall reactor since it was easier to adapt such reactors that were already used for III-V

semiconductor growth to the SiC technology. In such reactors, the process took place with the

classical precursor gases such as silane and propane, hydrogen as carrier gas, and with growth

temperatures above 1500°C. However, shortcomings which were related with the thermal

gradient (as the area above the substrate is not heated) resulted in non-uniform thickness and


26

doping concentrations, that are detrimental for industrial process [62]. Nowadays, the cold-wall

reactors are usually used in academic environments for further development of the SiC process

and quality of the layers.

In 1993, another reactor concept was developed i.e. the hot-wall reactor with a horizontal

geometry in which all the susceptor sides and reactor walls are heated. Thermal insulation

surrounds the susceptor to reduce the thermal losses by radiation resulting in a lower input power

compared with the cold-wall reactors. To prevent any additional impurity doping (e.g. with N2) of

the substrate during growth, the graphite reactor sidewalls and susceptor are covered by

polycrystalline SiC or eventually a Tantalum Carbide (TaC) coating. This type of reactor proved

to be capable of producing high quality thick layers (>50µm) with low background doping

(~1013

cm-3

) and good morphology. The high temperature used during the epitaxial process drives

a high cracking efficiency of the reactive gases which results in a uniform epitaxial layer. The

rotation of the substrate during the epitaxial process assures a higher uniformity of the layer.

Nowadays, due to advantages given by its geometry, the horizontal hot-wall reactor is one of the

most widely used reactors in industry and at academic level. In order to overcome different

challenges of the hot-wall CVD process such as: temperature control, large scale-area doping and

thickness uniformity of the layers, efforts have be made both experimentally and with numerical

modeling by different research groups to understand the deposition mechanism [63]–[67].

The planetary warm-wall reactors have been developed for SiC after Frijlink’s concept

[68] of III-V compound semiconductor growth based on gas foil rotation technique.

In the configuration of warm-wall reactors, a passively heating insulating ceiling that

reaches temperatures of 1500-1600°C is present. The principle of this reactor involves an

entrance of the reactive gases from reactor center that flow outward radially. For an improved

uniformity [69] the reactor design has a unique feature in which the gas (e.g. Argon gas) is used

to levitate the susceptor and rotate individually the wafers by viscous shear forces. Because of the

susceptor platform rotation and individually wafer rotation, highly uniform layers can be

obtained. Initially, in the warm-wall reactors configuration growth rates between 2-5µm/h could

be obtained but nowadays, growth rates up to 30µm/h can be achieved. For a more precise

control over temperature, hot-wall planetary configuration have been developed in which the

ceiling is heated with a separate induction coil enabling an improvement of layer quality,

thickness/doping uniformity, and defect density reduction [70], [71]. In fig. 1.13 schematic


27

illustrations of the different CVD reactors configurations employed for SiC epitaxial process

discussed above are represented.

Fig. 1.13 Schematic illustrations of the typical CVD reactors employed in the SiC

epitaxial process: (a) horizontal cold-wall, (b) vertical cold-wall, (c) horizontal hot-wall, (d)

warm-wall (hot-wall) planetary reactor and (e) chimney type-reactor [48]

The CVD reactors can run the epitaxial process under atmospheric or low pressure. Until

the early 1990s, the reactors ran mainly under atmospheric pressure, though soon the research

focused on low pressure reactors since a more precise control over the gas phase nucleation was

possible. Under reduced pressure conditions the thickness of the stagnant layer is decreased while

the gas velocity is increased, which leads to a more precise control over the deposition and higher

growth rates. A more detailed explanation of the epitaxial growth under low pressure conditions

is given in Chapter 2 of the thesis (section 2.2.2.1).

1.4.3.3 Capacity of the process chamber

Initially the CVD reactors were designed for single-wafer use since the research was

carried out at academic level. The capacity of the process chamber has been extended with the

increase of substrates diameters from 2 to 6 inches along the years. At industrial level, both cold-

wall and hot-wall CVD reactors have been extended to multi-wafer configuration as the batching

of wafers helps in reducing the overall cost for device processing. Moreover, multi-wafer reactors


28

are efficient to minimize the time per cycle which includes loading, heating, cooling, unloading,

enabling a throughput gain. The horizontal hot wall reactor has increased its growth capacity with

the evolution of bulk wafers diameters. For a high wafer uniformity of thickness and doping, the

rotation of the susceptor is applied.

From a commercial point of view the multi-wafer planetary warm-wall or hot wall

reactors have been proven to give excellent uniformity and reproducibility, thus making them

suitable for industrial applications. The planetary reactor configuration was scaled along the years

from 7x2 inches configuration [71] towards larger capacity configuration, resulting in 10x100mm

and further to 6x150mm capacity to match the evolution of substrate production.

It has been demonstrated by Burk and coworkers [72] that for a 4° off-axis 4H-SiC

150mm substrates grown in a warm-wall planetary reactor, a specular morphology and low

defects densities (e.g. <0.4cm-2

) of the epitaxial layers can be obtained.

Table 1.3 summarizes the commercially available CVD reactors for SiC epitaxial growth

with different configuration/chamber capacities and the main competitors on the market.


29

Company Reactor

name

Geometry Chamber

heating

Capacity

(inch)

Design

highlights

Aixtron AIX G5 WW Horizontal

Warm-wall

Planetary/Gas

foil rotation

12x4/8x6/5x8 Triple flow gas

injector

Sandwich

susceptor

Tokyo

Electron

Limited

(TEL)

Probus-SiC Horizontal

Hot-wall

Induction 10x3/7x4/3x6

Configured

with 1 or 2

process

chambers

Auto-loader

LPE PE106

ACiS M8

M10

Horizontal

Hot-wall

Induction 1x6

6x2/3x3/1x4

9x2/5x3/3x4

3 zone gas

injection

system

Growth rate

≤ 90µm/h

High vacuum

<1x10-7

mbar

p=50-400mbars

Epiluvac EPI-1000

EPI-1000X

Hot-wall unique Single wafer Different

configurations

High-growth

rate

NOVASiC1 R2

R3

Horizontal

Hot-wall

Resistive 1x2 to 1x4

1x6/2x4

p=50-800mbars

Equipped for

both n- and p-

type doping

Table 1.3 State of the art of the nowadays available CVD reactors and our reactors (non-

commercial) and their design highlights

1 non-commercial reactors


30

1.4.4. Epitaxial process (CVD)

1.4.4.1. Chemistry/ deposition rates

In order to enhance the commercialization and development of SiC power devices, much

effort has been made to improve the quality of SiC epitaxial layers and reduce the final cost of

device. For obtaining high power devices that can reach blocking voltages of several kV, thick

epilayers with low background doping concentration are required. Therefore, a high growth rate

is necessary during the epitaxial process. Higher growth rate process give rise to throughput

benefits thus, reducing the final cost of devices.

The main issue with increasing the growth rate is related to the amount of precursors used

in the gas mixture, resulting in the presence of homogeneous nucleation in the gas phase (i.e.

formation of silicon droplets due to high silicon concentration existent in the gas phase). Under

standard chemistry (i.e. precursors SiH4+C3H8 or ethylene (C2H4) and H2 as carrier gas), a

method that can be applied in order to avoid efficiently the Si cluster formation while increasing

the growth rate, is the reduction of the reactor pressure and/or the increase of carrier flow rate

[73], [74].

Another method that can lead to a further enhance of the growth rate, is the use of

chloride-based chemistry [75] which proved to be successful for growing fast and thick SiC

layers on different off-angle substrates [76]–[79]. The chloride-based chemistry includes the

addition of HCl to standard gas flow mixture or the use of chlorine containing precursors,

preventing the Si clusters formation since the silicon atoms will bind preferentially to chloride

atoms i.e. the dominating Si-specie for the growth will be the SiCl2 molecule.

Tanaka et al. [80] have studied the influence of growth pressure on growth rate of 4H-SiC

layers grown under standard chemistry and with addition of HCl gas. Figure 1.14 displays the

growth rate as a function of Si/H2 ratio at various pressures. As the Si/H2 ratio reflects the

precursor concentration, a direct proportionality with the growth rate can be obtained when all the

other process parameters are kept constant. Figure 1.14 reveals that a reduction in reactor

pressure from 250mbars to 50mbars under standard chemistry (no HCl addition) brings an

increase in growth rate. However, further increase of Si/H2 ratio at a given pressure results in a

saturation of growth caused by Si nucleation in the gas phase [81]. On the other hand, for a

constant pressure of 250mbars the addition of HCl gas compared to the one without, results in an

increase in growth rate (almost one order of magnitude difference) and no growth rate saturation.


31

Fig 1.14 Growth rate as a function of Si/H2 ratio for different pressures without and with HCl gas

(open symbol) [80]

The high growth rate achieved by chloride based growth process requires an increase in

growth temperature thus, enhancing the adatoms mobility on the surface. The higher temperature

can lead to a reduction of point defects and, to avoid the formation of 3C-SiC inclusions.

Nevertheless, a higher growth temperature delivers the disadvantage of step bunching formation

[79].

The most chosen approach for chloride based chemistry is the addition of HCl flow to the

standard precursors even though the other chloride precursors proved to be much more effective.

The growth process with hydrogen chloride addition (HCl) resulted in growth rates higher than

100µm/h [82], [83] while in standard process the growth rates can reach few tens of µm per hour

(8-10µm/h). Although, using different chlorine precursors it has been demonstrated an increase in

growth rate up to 350µm/h [84]. In terms of production cost and reproducibility for industry,

chloride based growth process seems to be very advantageous (e.g. lower energy consumption,

faster heat up time, reduction of hydrogen consumption).

1.4.4.2. Doping of SiC

SiC doping is usually carried out in-situ during the growth, by introducing n- or p-type

gas sources inside the reactor. For specific defined doping, the ion implantation method can be

applied, which leads to unwanted defects. In SiC, the n-type doping can be performed by using


32

Nitrogen (N) or Phosphorus (P) as dopant sources, while dopants such as Aluminum (Al) or

Boron (B) are used for p-type doping. The most common n- and p-type dopants gas sources are

respectively N2 and Trimethylaluminium (TMA). The ionization energies of the dopants applied

in SiC are dependent on the inequivalent lattice site, i.e. hexagonal or cubic site. The ionization

energies of the main dopants as a function of SiC polytypes are listed in table 1.4. In the SiC

lattice matrix, the nitrogen substitutes at carbon (C) lattice sites, which consist with the fact that

the atomic size of N (0.065nm) [85] is closer to that of C (0.07nm). In contrast to nitrogen,

aluminum substitutes at silicon (Si) lattice sites as the atomic size of Al (0.125nm) is closer to the

one of Si (0.11nm). However, compared with Aluminum and Boron, Nitrogen and Phosphorus

have relatively shallower ionization energies.

Impurity Polytype Ionization energy (meV)

Hexagonal/cubic sites

Nitrogen 4H-SiC 61/126

6H-SiC 85/140

3C-SiC 55

Phosphorus 4H-SiC 53/93

6H-SiC 80/110

Aluminum 4H-SiC 198/201

6H-SiC 240

3C-SiC 250

Boron 4H-SiC 300

6H-SiC 330

Table 1.4 Ionization energies of the main dopants employed in SiC (keep in mind that depending

on the work the reported values in literature can be different) from [48], [86], [87]

During dopant incorporation in SiC epilayers by different growth techniques, it has been

detected a polarity effect: higher nitrogen incorporation on the C-face (000 ) than on Si-face

(0001), while aluminum shows the inverse tendency. This polarity effect has been explained by

the surface kinetics during growth (different bonding of dopant atoms adsorbed onto Si- or C-

face) and elaborated further in chapter 3 of this manuscript.


33

In the early stages of SiC epitaxial growth studies, difficulties of doping control existed as

the background level of nitrogen was elevated. In the mid ‘90s Larkin and coworkers [88], [89]

have introduced the “site-competition effect”, which allows the control of the SiC doping in a

wide-range. The site competition effect is demonstrated between the dopant and the relaxed

matrix atom during the epitaxial growth. Therefore, at constant nitrogen flow rate, an increase of

carbon supply (C/Si>1) during the growth may reduce the probability of nitrogen incorporation

on the carbon site. Same principle can be applied for residual nitrogen, i.e. to reduce the

background concentration of non-intentionally doped epilayers. Such low n-type doped epilayers

are necessary for fabrication of vertical devices.

For the voluntary aluminum doping, an increase in silicon supply (C/Si<1) with a

constant Al flow rate should reduce the probability of aluminum incorporation on silicon site.

Applying the site competition effect during in situ N and Al doping by CVD growth, the doping

concentration on Si-face can be varied in the range 1x1014

to 2x1019

cm-3

, and respectively, in the

range 2x1014

to 5x1020

cm-3

[48], [86]. Such wide control over the n- and p-type doped SiC

epilayers helps to achieve abrupt junctions. It is noted that when trying to grow low doped n-type

or highly doped p-type epilayers, the task is more difficult on C-face than on Si-face.

1.4.4.3. Homoepitaxial growth of -SiC

In the early days of SiC epitaxy by CVD technique, the growth was performed on well

oriented available -SiC substrates where the polytype mixing turned out to be a serious problem.

To eliminate the polytype inclusions namely, the 3C-SiC presence onto the on-axis (0001) Si-

face 6H-SiC substrates, it was required to apply high temperatures above 1800°C [90].

Unfortunately, a high growth temperature conducts to different disadvantages such as higher H2

etching rate (i.e. lower growth rate), no precise doping control and elevated energy consumption

which results in higher cost per epilayer. However, in 1986 Matsunami et al. [17] found a key

method to grow single crystalline homoepitaxial 6H-SiC at lower temperatures (1400-1500°C) by

using off-angled (vicinal) 6H-SiC substrates. The method was named “step-controlled epitaxy”

since the polytype could be replicated by substrate step density. Therefore, the surface steps

could be used as templates for forcing the substrate polytype replication in the epilayer.

The SiC 0001 faces consist of terraces and steps, and for vicinal 4H-SiC or 6H-SiC

(0001) substrate the steps reflect the stacking sequence of the crystal. In the CVD growth the


34

adsorbed species are migrating onto the surface and will incorporate where a low potential is

present, i.e. at steps. A competitive two-dimensional growth process takes place on terraces when

high supersaturation is attained. On the on-axis 0001 faces, two-dimensional nucleation takes

place onto the terraces since the surface consists of low step density and wide terraces. Assigning

a low process temperature, the migration length of the adsorbed species is considerably reduced,

facilitating the 3C-SiC two-dimensional nucleation onto the terraces as shown in fig 1.15(a).

During the heteroepitaxy process the systematic formation of the defect called Double

Positioning Boundaries (DPBs), a particular case of twin boundaries, can be seen. The

appearance of the defect is based on two types of stacking sequence that 3C-SiC presents (i.e.

ABC and ACB), rotated by 60° to each other when the growth takes place on the basal plane of

hexagonal polytype.

Fig 1.15 Schematic representation of the step-controlled growth mode on: (a) on axis 6H-

SiC (0001) and (b) off-axis 6H-SiC (0001) substrate [48]

Contrarily, on off-angle 0001 faces even at low temperature, the migration of adsorbed

species towards steps is favored since the surface consists of high density of steps and narrow

terrace widths [91]. Thus, homoepitaxy can take place by lateral growth from steps along c-axis,

replicating the polytype stacking (step flow growth) as shown in fig 1.15(b).

The off-angle SiC substrates used for facilitating the homoepitaxy growth are usually

tilted several degrees towards low index direction such as <11 0>, though alternative

orientations such as <1 00> direction [92] can be applied. In the early stages, the SiC wafers


35

presented standard 8° and 3.5° off-angle toward <11 0> direction for 4H and 6H-SiC,

respectively. The 4H-SiC is the material of choice for electronic device applications and the 8°

off-cut provides the advantage of wider process window resulting in surfaces with low defect

density. However, larger off-angles signify reduced number of sliced wafers per boule i.e. more

waste, larger cost and moreover the presence of BPDs that cause a degradation in bipolar devices

[93], [94]. Nowadays, 150mm 4H-SiC substrates with 4° off-axis toward <11 0> direction are

commercially available and, efforts are made to reduce further the off-cut to 2°. Reducing the

substrate off-angle leads to rougher surfaces and extensive step-bunching due to narrower

window for the step-controlled growth (i.e. necessity to optimize the C/Si ratio and growth rate)

[95], [96].

The required conditions for SiC step-flow growth have been predicted by Kimoto and

Matsunami [98] by adapting the simple surface model described by Burton, Cabrera and Frank

(BCF) theory [99]. The conditions necessary for step-controlled epitaxy have been examined in

ref [98], in which the off-angles and growth temperature were varied in the range 0.2°-10°

(miscut towards <11 0> direction), and 1100-1500°C respectively. Figure 1.16 shows the

determined critical growth conditions for both (0001) Si and (000 ) C-face as a function of

growth temperature, growth rate and substrate off-angle.

The regions corresponding to the step-flow growth and two-dimensional nucleation

conditions are separated by the individual upper-left and lower-right curves, respectively. A step

flow growth with low off-angle substrates require high temperature and high growth rate,

respectively.


36

Fig. 1.16 Critical growth rate as a function of growth temperature determined by the BCF

theory for substrates with different off-cut angles [97]

For a substrate with a 0.2° off-angle, the model predicted a temperature of 1700°C and a

growth rate of ~6µm/h for obtaining step flow growth mode, as can be seen from fig 1.16.

However, for larger off-angle substrates i.e. more than 6°, lower growth temperature of ~1200°C

and growth rate as ~1µm/h were predicted for the step flow growth mode.

Although, the predicted conditions for step-controlled growth mode are a good theoretical

indication experimentally, the predicted curves are shifted towards higher temperatures (e.g. for

an 8° off-cut substrate the minimal growth temperature was found around 1450°C).

1.4.4.4. Si-face versus C-face

The SiC 0001 is a polar face as explained in section 1.3 that can be either (0001) Si

face or (000 ) C-face. The research and production of power devices currently, is mostly based

on Si-face 4H-SiC substrates because reproducible epilayers with low background doping levels

of less than 1015

cm-3

can be obtained by using the site-competition technique [88].

However, Kojima et al. [100] have shown that the site competition effect is present also

on C-face 4H-SiC epilayers under low pressure growth conditions thus being capable of reducing

the background doping level up to 2x1014

cm-3

. Furthermore, it was found that on C-face the

growth window for obtaining specular surface morphology is narrower than on Si-face. By


37

reducing the off-angle of the Si-face substrate, a degradation of the surface morphology was

observed as a result of extensive step-bunching and high surface roughness [95], [101]. The step-

bunching generation on Si-face of the grown epilayers is attributed to the surface free energy

[102], since along the (0001) direction the surface free energy on C-face -SiC is smaller

compared with that on Si-face [103]. On the other hand, on C-face specular surface morphology

can be obtained with straight-step structure along the off-cut direction and low surface roughness.

In contrast with Si-face, C-face substrates with low off-angle can be used to successfully grow

homoepitaxial layers without issues such as step-bunching, high roughness or polytype inclusions

[104]. For the heteroepitaxial growth, it has been demonstrated that the nominally on-axis C-face

SiC substrates provides favorable conditions for more uniform 3C-SiC nucleation when

sublimation epitaxy technique was used [105].

The oxidation rate on C-face 4H-SiC is almost ten times faster than on Si-face [106]

therefore, shortens the time for oxidation process in the fabrication of MOSFET devices where

the trap density ( ) reduction at the SiC-SiO2 interface is necessary. Fukuda et al. [107] showed

that by using H2 annealing, the threshold voltage can be improved and the can be decreased.

Moreover, C-face MOSFETs shows an increase channel mobility when a wet oxidation is applied

[108]. This exposes that even though Si-face is widely used, a potential for C-face especially for

SiC MOS-based devices for different applications exists.

1.4.4.5. Heteroepitaxial growth of -SiC

The opportunity to take advantages of the material properties at a lower cost compared

with the -SiC polytypes, motivated many works for performing the growth of cubic polytype

(called 3C-SiC or -SiC). However, no bulk 3C-SiC can be grown by Lely method due to the

high temperature applied during process. Furthermore, cubic SiC seeds were almost never

obtained, making the optimization of the growth process difficult. -SiC is a polytype that

appears spontaneously on -SiC and Si substrates. A grown interest was demonstrated for the

heteroepitaxial growth of 3C-SiC on large Si (111) or (100) substrates. For achieving higher

crystal quality, 3C-SiC epilayers were also grown on different substrates such as: silicon on

insulator (SOI), ‘‘undulant’’ <001> Si and on -SiC substrates.

The -SiC heteroepitaxial growth allows to obtain large size 3C-SiC layers which are

compatible for different device applications such as MOSFETs or MEMS.


38

Due to its stability at low temperature, the 3C-SiC can be grown on Si substrates below

the melting point of silicon ( =1414°C). The heteroepitaxial growth of 3C-SiC on Si substrates

cause several disadvantages since the two materials own different bulk lattice parameters (at

1350°C: 5.4598 vs 4.3850 see table 1.1) and thermal expansion coefficients (mismatch of

~8%) leading to induced strain and defects in the layers.

A solution to the mismatch problem was proposed by Nishino et al. [16], [109],

suggesting the use of a carburation step to create a buffer layer that acts as a seed/starting point

for the 3C-SiC growth by CVD. During the carburation step, 3C-SiC islands are created under a

carbon precursor (e.g. propane) diluted in hydrogen mixture which reacts with silicon at a given

temperature. The buffer layer is usually grown at lower temperature than the one use for CVD

epitaxy, and has a thickness of only few nanometers. The carburation step plays a decisive role in

the crystalline quality of the epitaxial layer. However, it has been observed that at Si/SiC

interface a high density of voids are present, which are considered as preferential areas for

defects nucleation such as Stacking Faults (SF), micro-twins or Anti-Phase Boundaries (APB).

For reducing the defect generation, most of the work was conducted towards optimization of the

carburation step [110], [111], but also compliant substrates were used [112], [113] 3C-SiC layers

can be grown on both Si(100) and (111) orientation substrate though, 3C-SiC(111) oriented

layers are limited in thickness to few µm due to the formation of cracks. Furthermore, after the

growth, the 3C-SiC(111) and (100) present a considerable roughness that is not suitable for

device fabrication. NOVASiC has developed the Chemical-Mechanical Polishing (CMP) [114]

process that helps reducing the unwanted surface roughness. Also, NOVASiC commercialize 3C-

SiC(100) [115] and (111) [116] epilayers with diameters from 2-4 inches.

Hoya Advanced Semiconductors Technology have demonstrated that using a compliant

substrate such as “Undulant-Si” substrate and apply the carburation step, 3C-SiC epitaxial layers

with reduced planar defects can be grown [112]. They have shown that an undulant substrate, i.e.

a Si substrate with a 4° off-angle towards [110] direction on which continuous undulations with

ridges aligned in direction are formed, promotes the elimination of planar defects (SFs and

APBs) [113]. Moreover, performing the epitaxial growth of 3C-SiC on undulant-Si substrates in

cold-wall, low pressure CVD system, under gas chemistry formed by dichlorosilane (SiH2Cl2),

acetylene (C2H2) and H2, almost 200µm thick 3C-SiC epitaxial layer was achieved. A solution


39

formed by hydrofluoric acid (HF) and nitric acid (HNO3) facilitated the removal of Si substrate,

therefore obtaining free-standing 3C-SiC substrates.

Another method for growing 3C-SiC layers is the use of commercially available on-axis

-SiC substrates (e.g. 4H or 6H-SiC substrates). This method is the best choice from the point of

view of thermal expansion coefficient compatibility and lattice constant matching (see table 1.1).

The increase in crystalline quality and size of -SiC substrates makes it more attractive for 3C-

SiC growth. The crystal structure of 4H-SiC and 6H-SiC 0001 face is similar to the 3C-SiC

(111) face therefore, the growth of 3C-SiC on (0001) face -SiC substrates by CVD technique is

assumed easier.

The growth of 3C-SiC on on-axis -SiC substrates requires two-dimensional nucleation

on terraces as explained in section 1.4.4.3, i.e. applying growth conditions that reduce the surface

mobility of adatoms (temperature or chemistry). The biggest challenge here remains the

elimination of the DPBs also called Twin Boundaries (TBs) during the growth since they are

killer defects for device performance [117].

Another issue is the reproducibility of free TBs 3C-SiC layers due to the fact that

perfectly on-axis substrates do not exist, but have generally an orientation uncertainty of +/- 0.5°

along the c-axis. An approach to tackle such issues involves the control of the 3C nucleation

stage by optimizing the growth conditions [118] (temperature, gas phase composition, pressure

and other growth parameters) or using different surface preparation prior to growth. It has been

demonstrated that in-situ etching by H2 or H2+HCl of the (0001) 6H-SiC surface prior to growth

reveal a step presence which contributes in obtaining DPB-free 3C-SiC layers [119].

Henry and co-workers [120]–[122] have reported the use of carbon-rich surface

preparation treatment during heating up or prior to the growth for obtaining high quality DPB-

free 3C-SiC epilayers grown on (0001) 4H-SiC substrate. Also, Alassaad and co-workers [123]

have reported that applying a GeH4 or SiH4 preparation treatment to slightly off-axis -SiC

surfaces, DPB-free 3C-SiC epilayers are promoted. It has been shown that despite the issues

encountered, high quality 3C-SiC epilayers can be grown on -SiC substrates by different growth

techniques such as VLS [124], sublimation epitaxy [125] and CVD [121].


40

1.4.5. Current status of SiC epilayers

Development of SiC power devices has been remarkably boosted since the enhancement

of substrate quality and size. The power electronics market is currently based on the mature Si

technology. However, due to the silicon material limitations, the SiC becomes a very attractive

candidate.

The research and development of SiC power devices is continuing to push their

performances towards higher current and voltage rating in order to access new applications areas.

SiC power devices can be characterized by high voltage blocking capability feature that is

strongly dependent on doping density. When considering the design of the SiC device structures,

the ideal Breakdown Voltage (BV) needs to be estimated. The desired BV can be determined

based on the relationship between the drift layer thickness and doping density as represented in

fig. 1.17.

Fig.1.17 Breakdown voltage as a function of both doping and thickness of 4H-SiC

epilayer. Two mechanisms are represented: punch-through and avalanche breakdown (red line)

[127]


41

As can be observed, the BV can be maintained when a certain thickness of the drift layer

is given while the doping concentration is increased. In this regime, the drift layer is fully

depleted and the breakdown occurs in the highly doped contact layer.

For example, to fabricate a 7000V 4H-SiC device, a thickness of 20µm is necessary for

the drift layer with a doping density in the range 1013

to ~5x1014

cm-3

. At higher doping

concentrations, an avalanche breakdown (red line fig. 1.17) occurs in the drift layer [126] leading

to a breakdown voltage reduction if the thickness is kept unchanged. As a consequence, for a

desired drift layer with a doping concentration above 5x1014

cm-3

, an increase of layer thickness is

necessary in order to maintain the 7000V breakdown voltage.

1.5. State of the art of SiC devices

SiC power devices are key components in power conversion systems and can be classified

in two main categories as represented in fig. 1.8 in section 1.3.2 for conductive SiC layers:

-Rectifiers (diodes) that regulate the electrical power, including unipolar (SBDs) and

bipolar junction (PN or PiN diodes) devices. Also the hybrid rectifier can be added, a merging p-

i-n and Schottky structure (Junction Barrier Schottky-JBS) that combines the benefits of high

blocking voltage capabilities, commutation speed and on state resistance of two involved diodes.

-Switches consist of unipolar transistors (e.g. MOSFETs and JFETs) and bipolar

transistors (e.g. Bipolar Junction Transistors-BJTs, Gate Turn-Off Thyristors-GTOs and Insulate

Gate Bipolar Devices-IGBTs).

To determine whether a unipolar or bipolar device is required for a given application, the

characteristic parameters of the device such as reverse blocking voltage, forward current density,

maximum allowable reverse current density, operating temperature and switching frequency are

considered.

Schottky barrier diodes

The SBDs are popular unipolar devices which are based as rectifying metal-

semiconductor junctions, and are very attractive due to their fast switching speed and relatively

high breakdown field. In SBDs, the forward current transport is completed by the majority

carriers injected from the semiconductor into the metal, allowing the device to exhibit fast reverse

recovery. Because of a fast reverse recovery in SiC SBDs, the switching losses can be reduced


42

which in turn will lead to higher switching frequency and furthermore, to a shrinkage of the

passive components. The schematic illustration of a SBD structure is represented in fig. 1.18. At

industrial level, an attention towards this emerging technology of SiC power devices started with

the first commercialized SiC SBDs by Infineon in 2001 (initial 300V/10A and 600V/6A).

Since then, SiC SBDs have shown a continuous increase in BV and conduction current,

showing at present BV ratings from 300-1700V and current ratings as high as 100A [128], [129].

Meanwhile, the applications for SiC SBDs have rapidly grown and extended to a variety of

power suppliers, building blocks of solar inverters, air conditioning, and industrial equipment.

SiC SBDs are well established components for Power Factor Correctors (PFC) circuits and

switching mode power suppliers. The highest BV of a SiC Schottky rectifier at industrial level is

1.7kV though, in research and development it exceeds 2.5kV.

Fig. 1.18 Schematic illustration of a cross-section 3.3kV Schottky barrier diode [130]

Bipolar PiN diodes

In PiN diodes, the conduction transport takes place by the majority and minority carriers.

Besides, PiN diodes demonstrated low leakage current, enabling higher blocking voltage

operation that requires reverse recovery charging during switching operation [29]. The schematic

illustration of a PiN diode structure (p+/i/n

+ structure) is represented in fig. 1.19. In this type of

bipolar diodes the electrical parameters depend on the difference in epilayers thickness, doping

concentration (an intrinsic layer n- doped between the p

+ and n

+ layers) and deposited ohmic

contacts. Applying large voltages to the structure, the conduction is dominated by injection of

minority carriers into the intrinsic layer (i-region).


43

As in SBD, the conductivity of the drift layer can be modulated through the choice of the doping

level, determining the forward voltage drop and the reverse breakdown voltage. However,

beyond the turn-on voltage of the junction, the conduction significantly increases due to the

injection of a large number of electrons and holes from the doped regions. Consequently, lower

static power losses can be achieved by PiN diodes, at the expense of a larger turn-on voltage.

The major issue in SiC PiN bipolar devices is the forward voltage degradation which can

be described by the larger voltage drop after current injection [131]. Degradation of the PiN

diodes is attributed to the nucleation and expansion of the Shockley-Stacking Faults (SSFs)

located on BPDs or other dislocations. The expanded SSFs activate a reduction of carrier lifetime

and formation of a barrier potential for carrier transport, resulting in the increase of forward

voltage drop. Driving force behind the expansion of these SSFs under forward bias has been

intensely studied and found to be caused by the material itself [132].

Fig. 1.19 Schematic illustration of a cross-section 3.3kV PiN diode [130]

The blocking voltage for SiC PiN junction diodes is desired to be higher than 5kV, though

the highest BV reported is of 21.7kV for a PiN diode with low reverse leakage current density

and high forward voltage drop [133]. Due to the challenges regarding the forward voltage

degradation, the development and commercialization of the diodes have been hindered.

To maintain low leakage current when the junction is in reverse mode and thus, reducing

the operating temperature resulted in the development of hybrid rectifiers such as JBS. In the last

decade, the research was oriented towards high-voltage and high-temperature applications,

leading to the development of SiC power switching devices like the MOSFET or JFET.


44

Development of power MOSFETs devices have been hampered by the degradation of channel

specific on-resistance due to low inversion channel mobility obtain on 4H-SiC. Another

challenging aspect in this technology is the gate oxide reliability. Despite the MOS technology

challenges, in 2011 a 1.2kV 4H-SiC MOSFETs devices have been commercialized by CREE and

a next generation of SiC MOSFETs with voltage ratings from 900V-15kV operating up to 175°C

has been demonstrated [134].

In recent years, bipolar devices such as IGBTs have attracted attention due to their

potential for high-voltage applications. This is the result of superior on-state performance. The

achievements of the IGBTs nowadays, reveal capabilities of blocking voltage higher than 12kV,

although it is expected to see an increase of their BV capabilities up to 20-30kV in near feature

[135]. Other promising bipolar devices are the BJTs which have been developed and

commercialized. The highest BV reported for a SiC BJT is 21kV, with a current gain of 63 and a

specific on-resistance of 321mΩ-cm2 [136]. The studies revealed that the bipolar devices suffer

from reliability problems, therefore reducing their commercial production.

Table 1.5 presents the state of the art of the current situation of SiC power devices on the

market, and the companies which are commercializing the devices.


45

Device type Unipolar Bipolar

Rectifiers (Diodes) SBD PIN

Company CREE

INFINEON

ROHM

GeneSiC

ST Microelectronics

Microsemi

TOSHIBA

GeneSiC

Voltage range (V) 600-1700V 8000-15000

Switches (Transistors) MOSFET JFET BJT GTO

Company CREE

INFINEON

ROHM

INFINEON

GeneSiC

GeneSiC

Voltage range (V) 400-1700 1200 1200-1700 6500

Table 1.5 The state of the art of the commercially available SiC devices and companies involved

in the market of power devices

1.6. Conclusions

The SiC semiconductor gains more and more attention for device fabrication due to the

rapid development at bulk and epitaxy level. Among the numerous SiC polytypes, 4H-SiC has

been the choice at the industrial levels due to its physical properties and the possibility to achieve

good quality epilayers. Nowadays, 4H-SiC wafers with 150mm diameter are commercially

available. For the industry, the development of CVD reactors with multiwafer configuration and

high growth rates offer benefits over the final device cost. Since the discovery of step-controlled

epitaxy and site-competition effect, high-quality -SiC (hexagonal polytypes) epilayers without

polytype mixing and with uniform doping density (in the range 1013

-1020

cm-3

) are obtained.

Compare with the hexagonal polytypes, the technological development of 3C-SiC is hindered due

to the presence of twin boundaries and stacking fault defects. Efforts are made by different


46

research groups to grow single domains 3C-SiC layers by heteroepitaxy, since no 3C-SiC seed is

available for bulk growth.

Furthermore, the recent developments in SiC power devices display the SBDs as the most

well established and mature power devices. However, SiC technology offers advancements in

power and current densities of the both unipolar and bipolar devices. Unfortunately, the bipolar

devices reveal reliability problems and their commercialization is somehow limited.

1.7. Bibliography

[1] http://www.energy.gov/eere/efficiency

[2] http://spectrum.ieee.org/semiconductors/materials/silicon-carbide-smaller-faster-tougher

[3] https://www.mesago.de/download/PCC/PCIM_Asia2015_Presentation_2015_Lin_Hong.pdf

[4] http://electronics360.globalspec.com/article/6286/iedm-divulges-advances-in-wide-bandgap-

devices

[5] T. Kimoto, “Material science and device physics in SiC technology for high-voltage power

devices,” Jpn. J. Appl. Phys., vol. 54, no. 4, p. 40103, Mar. 2015

[6] J. Millan, P. Godignon, X. Perpina, A. Pérez-Tomas, and J. Rebollo, “A survey of wide

bandgap power semiconductor devices,” IEEE Transactions on Power Electronics, vol. 29,

no. 5, 2014

[7] http://www.geaviation.com/press/systems/systems_20160315.html

[8] J. J. Berzelius, Annals of Physics, vol. 77, p. 169, 1824

[9] A.G. Acheson, English Patent, vol. 17911, 1892

[10] H. Moissan, “Etude du siliciure de carbone de la météorite Canon Diablo,” Compte-rendu de

l’Académie des Sciences de Paris, vol. 140, p. 405, 1905

[11] http://www.charlesandcolvard.com/our-moissanite/

[12] H.J. Round, “A note on Carborundum,” Electrical Word, vol. 19, p. 309, 1907

[13] J.A. Lely, Ber. Deut. Keram. Ges, vol. 32, p. 229, 1955

[14] W.J. Choyke, D.R. Hamilton, and L. Patrick, “Optical properties of cubic SiC:

Luminescence of Nitroge-exciton complexes and interband absorption,” Physical Review,

vol. 133, no. 4A, 1964

[15] Y.M. Tairov and V.F. Tsvetkov, “Investigation of growth process of ingots of silicon carbide

single crystals,” Journal of Crystal Growth, vol. 43, p. 29, 1978

[16] S. Nishino, J. A. Powell, and H. A. Will, “Production of large‐area single‐crystal wafers of

cubic SiC for semiconductor devices,” Applied Physics Letters, vol. 42, no. 5, pp. 460–462,

Mar. 1983

[17] N. Kuroda, K. Shibahara, W. Yoo, S. Nishino, and H. Matsunami, “Step-controlled VPE

growth of SiC single crystals at low temperatures,” Ext. Abstr. 19th Conf. on Solid State

Devices and Materials, Tokyo, p. 227, 1987

[18] R. Rupp, C. Miesner, and I. Zverev, “Combination of new generation CoolMOS and thinQ!

Silicon carbide Schottky diodes: new horizons in hard switching applications,” Proc. Power

Conversion Intelligent Motion Conf. 2001


47

[19] F. Bechstedt, P. Kackell, A. Zywietz, K. Karch, B. Adolph, K. Tenelsen, and J. Furtmuller,

“Polytypism and Properties of Silicon Carbide,” physica status solidi (b), vol. 202, p. 35,

1997

[20] S.E. Saddow and A. Agarwal, Advances in Silicon Carbide: Processing and Applications.

Artech House, Inc., 2004

[21] W. F. Knippenberg, Philips Research Reports, vol. 18, p. 161, 1963

[22] U. Starke, “Atomic Structure of SiC Surfaces,” Silicon Carbide Recent Major Advances ed.

W.J.Choyke, H.Matsunami, G.Pensl, p. 281, Springer-Verlag, Berlin 2004

[23] L. S. Ramsdell, American Mineralogist, vol. 32, p. 64, 1947

[24] http://www.iue.tuwien.ac.at/phd/ayalew/node20.html

[25] C. M. Zetterling, Process Technology for Silicon Carbide devices, Institution of Electrical

Engineers (INSPEC)-London, 2002

[26] J. Fan and P.K. Chu, Silicon Carbide Nanostructures, Engineering Materials and Processes.

Springer, Switzerland, 2014

[27] P. G. Neudeck, The VLSI handbook/ Chapter 5 SiC Technology, CRC Press. 2006

[28] A. Schoner, M. Krieger, G. Pensl, M. Abe, and H. Nagasawa, “Fabrication and

characterization of 3C-SiC based MOSFETs,” Chemical Vapor Deposition, vol. 12, p. 523,

2006

[29] M.B.J. Wijesundara and R.G. Azevedo, Silicon Carbide Microsystems for Harsh

Enviroments. Springer New York, 2011

[30] http://newsroom.toyota.co.jp/en/detail/2656842

[31] http://www.mitsubishielectric.com/news/2013/0326-a.html

[32] C. A. Zorman and R. J. Parro, “Micro- and nanomechanical structures for silicon carbide

MEMS and NEMS,” phys. stat. sol. (b), vol. 245, no. 7, pp. 1404–1424, 2008

[33] S. Jiao, J. F. Michaud, M. Portail, A. Madouri, T. Chassagne, M. Zielinski, and D. Alquier,

“A new approach for AFM cantilever elaboration with 3C-SiC,” Materials Letters, vol. 77, p.

54–56, 2012

[34] S. E. Saddow, “Silicon Carbide materials for biomedical applications”, Silicon Carbide

Biotechnology A Biocompatible Semiconductor for Advanced Biomedical Devices and

Applications, Elsevier Inc., p.1-15, 2012

[35] R.C. Glass, D. Henshall, V.F. Tsvetkov, and C.H. Carter, Jr., “SiC Seeded Crystal Growth,”

Physica Status Solidi (b), vol. 202, p. 149, 1997

[36] P. Wellman, G. Neubauer, L. Fahlbusch, M. Salamon, and N. Uhlmann, “Growth of SiC

bulk crystals for application in power electronic devices-process design, 2D and 3D X-ray in

situ visualisation and advanced doping,” Crystal Research and Technology, vol. 50, no. 1, p.

2, 2015

[37] T.L. Straubinger, M. Bickermann, M. Grau, D. Hofmann, L. Kadinski, S. Muller, M. Selder,

P.J. Wellmann, and A. Winnacker, “Growth rate control in SiC-Physical Vapor Transport

method through heat transfer modeling and non-stationary process conditions”; Materials

Science Forum, vol. 338–342, p. 39, 1999

[38] J. Heindl, H.P. Strunk, V.D. Heydemann, and G.Pensel, “Micropipes: Hollow Tubes in

Silicon Carbide,” Physica Status Solidi (a), vol. 162, p. 251, 1997

[39] B.M. Epelbaum, O. Filip, and A. Winnacker, “Bulk and epitaxial growth of micropipe-free

silicon carbide on basal and rhombohedral plane seeds,” Physica Status Solidi (b), vol. 245,

no. 7, p. 1257, 2008


48

[40] D. Nakamura, I. Gunjishima, S. Yamaguchi, T. Ito, A. Okamoto, H. Kondo, S. Onda, and K.

Takatori, “Ultrahigh-quality silicon carbide single crystals”, Nature Journal, vol. 430, p.

1009, 2004

[41] P. Wellman, P. Desperrier, R. Muller, T. Straubinger, A. Winnacker, F. Baillet, E. Blanquet,

J.M. Dedulle, and M. Pons, “SiC single crystal growth by a modified physical vapor transport

technique,” Journal of Crystal Growth, vol. 275, p. e555, 2005

[42] D. Chaussende, F. Baillet, L. Charpentier, E. Pernot, M. Pons, and R. Madar, “Continuous

Feed Physical Vapor Transport: towards high purity and long boule growth of SiC,” Journal

of Electrochem. Society, vol. 150 (10), p. G653, 2003

[43] O. Kordina, C. Hallin, A. Henry, J.P. Bergman, I. Ivanov, A. Ellison, N.T. Son, and E.

Janzén, “Growth of SiC by Hot-Wall CVD and HTCVD,” Physica Status Solidi (b), vol. 202,

p. 321, 1997

[44] http://www.norstel.com/products/semi-insulating-sic-substrates/

[45] M. Fanton, D. Snyder, B. Weiland, R. Cavalero, A. Polyakov, M. Skowronski, and H.

Chung, “Growth of nitrogen-doped SiC boules by halide chemical vapor deposition” Journal

of Crystal Growth, vol. 287, p. 359, 2006

[46] K. Kusunoki, K. Kamei, N. Okada, K. Moriguchi, H. Kaido, H. Daikoku, M. Kado, K.

Danno, H. Sakamoto, T. Bessho, and T. Ujihara, “Top-seeded solution of 3 inch diameter

4H-SiC bulk crystal using metal solvents,” Materials Science Forum, vol. 778–780, p. 79,

2014

[47] T. Ujihara, K. Seki, R. Tanaka, S. Kozawa, Alexander, K. Morimoto, K. Sasaki, and Y.

Takeda, “High-quality and large-area 3C–SiC growth on 6H–SiC(0 0 0 1) seed crystal with

top-seeded solution method,” Journal of Crystal Growth, vol. 318, no. 1, pp. 389–393, Mar.

2011

[48] T. Kimoto and J.A. Cooper, Fundamentals of Silicon Carbide Technology. John Wiley &

Sons Singapore Pte. Ltd, 2014

[49]http://www.cree.com/News-and-Events/Cree-News/Press-Releases/2012/August/150mm-

wafers

[50] http://www.iiviadvmat.com/SiC-products/SiC-Products.html

[51] http://www.dowcorning.com/applications/search/default.aspx?R=12686EN

[52] https://sicrystal.de/index.php/en/produkte-en

[53]http://globenewswire.com/news-release/2015/07/16/752880/10142036/en/II-VI-Advanced-

Materials-Demonstrates-World-s-First-200mm-Diameter-SiC-Wafer.html

[54] St. G. Muller, E.K. Sanchez, D. Hansen, R.D. Drachev, G. Chung, B. Thomas, J. Zhang,

M.J. Loboda, M. Dudley, H. Wang, F. Wu, S. Byrappa, B. Raghothamachar, and G. Choi,

“Volume production of high quality SiC substrates and epitaxial layers: Defect trends and

device applications,” Journal of Crystal Growth, vol. 352, p. 39, 2012

[55] J. Quast, D. Hansen, M. Loboda, I. Manning, K. Moeggenborg, S. Muller, C. Parfeniuk, E.

Sanchez, and C. Whiteley, “High quality 150 mm 4H-SiC wafers for power device

production,” Materials Science Forum, vol. 821–823, p. 56, 2015

[56] R. Rupp, “SiC power electronic devices: historical development, current and future

concepts/ presentation,” Tutorial ICSCRM conference, 2013

[57] O. Kordina, C. Hallin, A. Ellison, A. S. Bakin, I. G. Ivanov, A. Henry, R. Yakimova, M.

Touminen, A. Vehanen, and E. Janzén, “High temperature chemical vapor deposition of

SiC,” Applied Physics Letters, vol. 69, no. 10, p. 1456–1458, 1996

[58] R. Yakimova and E. Janzén, “Current status and advances in the growth of SiC,” Diamond

and Related Materials, vol. 9, no. 3–6, p. 432–438, 2000


49

[59] G. Ferro, D. Chaussende, and C. Jacquier, “VLS growth of SiC epilayers,” Research sign-

post, p. 91, 2006

[60] R.S. Kern, K. Jarrendahl, S. Tanaka, and R.F. Davis, “Homoepitaxial SiC growth by

Molecular beam epitaxy,” Physica Status Solidi (b), vol. 202, p. 379, 1997

[61] A. Ellison, J. Zhang, A. Henry, and E. Janzén, “Epitaxial growth of SiC in a chimney CVD

reactor,” Journal of Crystal Growth, vol. 236, no. 1–3, p. 225, 2002

[62] J.A. Powell, D.J. Larkin, L.G. Matus, W.J. Choyke, J.L. Bradshaw, L. Henderson, M.

Yoganathan, J. Yang, and P. Pirouz, “Growth of high quality 6H-SiC epitaxial films on

vicinal (0001) 6H-SiC wafers,” Applied Physics Letters, vol. 56, p. 1442, 1990

[63] G. Wagner, D. Schulz, and D. Siche, “Vapour phase growth of epitaxial silicon carbide

layers,” Progress in Crystal Growth and Characterization of Materials, vol. 47, no. 2–3, p.

139–165, 2003

[64] Ö. Danielsson, A. Henry, and E. Janzén, “Growth rate predictions of chemical vapor

deposited silicon carbide epitaxial layers,” Journal of Crystal Growth, vol. 243, no. 1, p.

170–184, 2002

[65] Ö. Danielsson, U. Forsberg, A. Henry, and E. Janzén, “Investigation of the temperature

profile in a hot-wall SiC chemical vapor deposition reactor,” Journal of Crystal Growth, vol.

235, no. 1–4, p. 352–364, 2002

[66] J. Meziere, M. Ucar, E. Blanquet, M. Pons, P. Ferret, and L. Di Cioccio, “Modeling and

simulation of SiC CVD in the horizontal hot-wall reactor concept,” Journal of Crystal

Growth, vol. 267, no. 3–4, p. 436–451, 2004

[67] S. Nishizawa and M. Pons, “Growth and Doping Modelling of SiC-CVD in a horizontal hot

wall reactor,” Chemical Vapor Deposition, vol. 12, p. 516–522, 2006

[68] P. M. Frijlink, “A new versatile, large size MOVPE reactor,” Journal of Crystal Growth,

vol. 93, no. 1–4, p. 207–215, 1988

[69] J.J. Sumakeris, M.K. Das, S. Ha, E. Hurt, K. Irvine, M.J. Paisley, M.J. O’Loughlin, J.W.

Palmour, M. Skowronski, H.McD. Hobgood, and C.H. Carter, Jr., “Development of epitaxial

SiC processes suitable for bipolar power devices,” Materials Science Forum, vol. 483–485, p.

155, 2005

[70] A. A. Burk Jr., M. J. O’Loughlin, and H. D. Nordby Jr., “SiC epitaxial layer growth in a

novel multi-wafer vapor-phase epitaxial (VPE) reactor,” Journal of Crystal Growth, vol. 200,

no. 3–4, p. 458–466, 1999

[71] A. A. Burk, “Development of multiwafer Warm-Wall planetary VPE reactors for SiC device

production,” Chemical Vapor Deposition, vol. 12, p. 465, 2006

[72] A. A. Burk, D. Tsvetkov, M.J. O’Loughlin, S. Ustin, L. Garrett, A.R. Powell, J. Seaman, and

N. Partin, “Latest SiC epitaxial layer growth results in a high-throughput 6x150mm Warm-

Wall Planetary reactor,” Materials Science Forum, vol. 778–780, p. 113, 2014

[73] H. Tsuchida, I. Kamata, T. Jikimoto, and K. Izumi, “Epitaxial growth of thick 4H–SiC layers

in a vertical radiant-heating reactor,” Journal of Crystal Growth, vol. 237–239, Part 2, p.

1206–1212, 2002

[74] M. Ito, L. Storasta, and H. Tsuchida, “Development of 4H–SiC Epitaxial Growth Technique

Achieving High Growth Rate and Large-Area Uniformity,” Appl. Phys. Express, vol. 1, no. 1,

p. 15001, 2008

[75] O. Kordina, “Chloride based chemistry for SiC growth,” US Patent 7,247,513, 2007

[76] F. La Via, G. Galvagno, G. Foti, M. Mauceri, S. Leone, G. Pistone, G. Abbondanza, A.

Veneroni, M. Masi, G L. Valente, and D. Crippa, “4H-SiC epitaxial growth with chlorine

addition,” Chemical Vapor Deposition, vol. 12, p. 509, 2006


50

[77] H. Pedersen, S. Leone, A. Henry, F. C. Beyer, V. Darakchieva, and E. Janzén, “Very high

growth rate of 4H-SiC epilayers using the chlorinated precursor methyltrichlorosilane

(MTS),” Journal of Crystal Growth, vol. 307, no. 2, p. 334–340, 2007

[78] S. Leone, H. Pedersen, A. Henry, O. Kordina, and E. Janzén, “Thick homoepitaxial layers

grown on on-axis Si-face 6H- and 4H-SiC substrates with HCl addition,” Journal of Crystal

Growth, vol. 312, no. 1, p. 24–32, 2009

[79] S. Leone, H. Pedersen, A. Henry, O. Kordina, and E. Janzén, “Improved morphology for

epitaxial growth on 4° off-axis 4H-SiC substrates,” Journal of Crystal Growth, vol. 311, no.

12, p. 3265–3272, 2009

[80] T. Tanaka, N. Kawabata, Y. Mitani, M. Sakai, N. Tomita, M. Tarutani, T. Kuroiwa, Y.

Toyoda, M.Imaizumi, H. Sumitani, and S. Yamakawa, “Influence of growth pressure and

addition of HCl gas on growth rate of 4H-SiC epitaxy,” Materials Science Forum, vol. 821–

823, p. 133, 2015

[81] S. Nakamura, T. Kimoto, and H. Matsunami, “Rate-determining process in chemical vapor

deposition of SiC on off-axis -SiC,” Journal of Crystal Growth, vol. 270, no. 3–4, p. 455–

461, 2004

[82] F. La Via, G. Izzo, M. Mauceri, G. Pistone, G. Condorelli, L. Perdicaro, G. Abbondanza, L.

Calcagno, G. Froti, and D. Crippa, “4H-SiC epitaxial layer growth by trichlorosilane (TCS),”

Journal of Crystal Growth, vol. 311, p. 107, 2008

[83] H. Fujibayashi, M. Ito, H. Ito, I. Kamata, M. Naito, K. Hara, S. Yamauchi, K. Suzuki, M.

Yajima, S. Mitani, K. Suzuki, H. Aoki, K. Nishikawa, T. Kozawa, and H. Tsuchida,

“Development of a 150 mm 4H-SiC epitaxial reactor with high-speed wafer rotation,” Appl.

Phys. Express, vol. 7, no. 1, p. 15502, 2013

[84] S. Leone, F. C. Beyer, H. Pedersen, O. Kordina, A. Henry, and E. Janzén, “High Growth

Rate of 4H-SiC Epilayers on On-Axis Substrates with Different Chlorinated Precursors,”

Crystal Growth & Design, vol. 10, no. 12, p.5334, 2010

[85] http://www.webelements.com/periodicity/atomic_radius_empirical/

[86] G. Dhanaraj, K. Byrappa, V. Prasad, and M. Dudley, Handbook of Crystal growth, Springer.

2010


[88] D. J. Larkin, “SiC Dopant Incorporation Control Using Site-Competition CVD,” phys. stat.

sol. (b), vol. 202, no. 1, p. 305–320, 1997

[89] D. J. Larkin, P. G. Neudeck, J. A. Powell, and L. G. Matus, “Site-competition epitaxy for

superior silicon carbide electronics,” Applied Physics Letters, vol. 65, no. 13, p. 1659, 1994

[90] S. Nishino, H. Matsunami, and T. Tanaka, “Growth and morphology of 6H-SiC epitaxial

layers by CVD,” Journal of Crystal Growth, vol. 45, p. 144–149, 1978

[91] H. Matsunami and T. Kimoto, “Step-controlled epitaxial growth of SiC: high quality

homoepitaxy,” Materials Science and Engineering R, vol. 20, p. 125, 1997

[92] B. E. Landini and G. R. Brandes, “Characteristics of homoepitaxial 4H-SiC films grown on

c-axis substrates offcut towards 〈11 00〉 or 〈112 0〉,” Applied Physics Letters, vol. 74,

no. 18, p. 2632–2634, 1999

[93] W. Chen and M. A. Capano, “Growth and characterization of 4H-SiC epilayers on substrates

with different off-cut angles,” Journal of Applied Physics, vol. 98, no. 11, p. 114907, 2005

[94] H. Jacobson, J. Birch, R. Yakimova, M. Syväjärvi, J. P. Bergman, A. Ellison, T. Tuomi, and

E. Janzén, “Dislocation evolution in 4H-SiC epitaxial layers,” Journal of Applied Physics,

vol. 91, no. 10, p. 6354–6360, 2002.


51

[95] B. Kallinger, P. Berwian, J. Friedrich, and B. Thomas, “Step-controlled homoepitaxial

growth of 4H–SiC on vicinal substrates,” Journal of Crystal Growth, vol. 381, p. 127–133,

2013.

[96] X. Li, E. Janzén, and A. Henry, “Homo-epitaxial growth on low-angle off-cut 4H-SiC”,

Materials Science Forum, vol. 778–780, p. 131, 2014.

[97] F. L. Via, M. Camarda, and A. L. Magna, “Mechanisms of growth and defect properties of

epitaxial SiC,” Applied Physics Reviews, vol. 1, no. 3, p. 31301, 2014

[98] T. Kimoto and H. Matsunami, “Surface kinetics of adatoms in vapor phase epitaxial growth

of SiC on 6H‐SiC0001 vicinal surfaces,” Journal of Applied Physics, vol. 75, no. 2, p.

850–859, 1994

[99] W. K. Burton, N. Cabrera, and F. C. Frank, “The Growth of Crystals and the Equilibrium

Structure of their Surfaces,” Philosophical Transactions of the Royal Society of London A:

Mathematical, Physical and Engineering Sciences, vol. 243, no. 866, p. 299–358, 1951

[100] K. Kojima, S. Kuroda, H. Okumura, and K. Arai, “Homoepitaxial growth on a 4H-SiC C-

face substrate,” Chemical Vapor Deposition, vol. 12, p. 489, 2006

[101] K. Kojima, K. Masumoto, S. Ito, A. Nagata, and H. Okumura, “4H-SiC Homoepitaxial

Growth on Substrate with Vicinal Off-Angle Lower than 1°,” ECS J. Solid State Sci.

Technol., vol. 2, no. 8, pp. N3012–N3017, 2013

[102] T. Kimoto, Z.Y. Chen, S. Tamura, S. Nakamura, N. Onojima, and H. Matsunami, “Surface

morphological structures of 4H-, 6H- and 15R-SiC (0001) epitaxial layers grown by CVD,”

Japanese Journal of Applied Physics, vol. 40, p. 3315, 2001

[103] M. Syväjärvi, R. Yakimova, and E. Janzén, “Growth of SiC from the liquid phase: wetting

and dissolution of SiC,” Diamond and Related Materials, vol. 6, no. 10, p. 1266–1268, 1997

[104] K. Kojima, H. Okumura, S. Kuroda, and K. Arai, “Homoepitaxial growth of 4H-SiC on on-

axis (0 0 0 1 ) C-face substrates by chemical vapor depositon,” Journal of Crystal Growth,

vol. 269, no. 2–4, p. 367–376, 2004

[105] R. Vasiliauskas, S. Juillaguet, M. Syväjärvi, and R. Yakimova, “Cubic SiC formation on

the C-face of 6H–SiC (0001) substrates,” Journal of Crystal Growth, vol. 348, no. 1, p. 91–

96, 2012

[106] A. Golz, G. Horstmann, E. Stein von Kamienski, and H. Kurz, Inst. Phys. Conf. Ser., vol.

142, p. 633, 1996

[107] K. Fukuda, M. Kato, K. Kojima, and J. Senzaki, “Effect of gate oxidation method on

electrical properties of metal-oxide-semiconductor field-effect transistors fabricated on 4H-

SiC C(0001 ) face,” Applied Physics Letters, vol. 84, no. 12, pp. 2088–2090, 2004

[108] T. Hatakeyama, T. Shimizu, T. Suzuki, Y. Nakabayashi, H. Okumura, and T. Kimoto,

“Deep level transient spectroscopy characterization of mobility-limiting traps in SiO2/SiC

interfaces on C-face 4H-SiC,” Materials Science Forum, vol. 740–742, p. 477, 2013

[109] S. Nishino, H. Suhara, H. Ono, and H. Matsunami, “Epitaxial growth and electric

characteristics of cubic SiC on silicon,” Journal of Applied Physics, vol. 61, no. 10, pp.

4889–4893, 1987

[110] M. Zielinski, A. Leycuras, S. Ndiaye, and T. Chassagne, “Stress relaxation during the

growth of 3C-SiC∕Si thin films,” Applied Physics Letters, vol. 89, no. 13, p. 131906, 2006

[111] G. Ferro, T. Chassagne, A. Leycuras, F. Cauwet, and Y. Monteil, “Strain tailoring in 3C-

SiC heteroepitaxial layers grown on Si(100),” Chemical Vapor Deposition, vol. 12, p. 483,

2006


52

[112] M. Zielinski, C. Moisson, S. Monnoye, H. Mank, T. Chassagne, S. Roy, A.E. Bazin, J.F

Michaud, and M. Portail, “Recent advances in surface preparation of silicon carbide and

other band gap materials,” Materials Science Forum, vol. 645–648, p. 753, 2010

[113] http://www.novasic.com/services/epitaxy/specs_3C-SiC-100.php

[114] http://www.novasic.com/services/epitaxy/specs_3C-SiC-111.php

[115] H. Nagasawa, K. Yagi, and T. Kawahara, “3C-SiC hetero-epitaxial growth on undulant

Si(0 0 1) substrate,” Journal of Crystal Growth, vol. 237–239, Part 2, pp. 1244–1249, 2002.

[116] H. Nagasawa, K. Yagi, T. Kawahara, and N. Hatta, “Reducing planar defects in 3C-SiC,”

Chemical Vapor Deposition, vol. 12, p. 502, 2006

[117] A. A. Lebedev, P. L. Abramov, S. P. Lebedev, G. A. Oganesyan, A. S. Tregubova, and D.

V. Shamshur, “Influence of the defect density (twins boundaries) on electrical parameters of

3C-SiC epitaxial films,” Physica B: Condensed Matter, vol. 404, no. 23–24, pp. 4758–4760,

2009

[118] M. Soueidan, G. Ferro, B. Nsouli, F. Cauwet, J. Dazord, G. Younes, and Y. Monteil,

“Effect of growth parameters on the heteroepitaxy of 3C-SiC on 6H-SiC substrate by

chemical vapor deposition,” Materials Science and Engineering: B, vol. 130, no. 1–3, pp.

66–72, 2006

[119] Z. Y. Xie, J. H. Edgar, B. K. Burkland, J. T. George, and J. Chaudhuri, “DPBs-free and

polytype controlled growth of SiC via surface etching on on-axis 6H-SiC(0 0 0 1),” Journal

of Crystal Growth, vol. 224, no. 3–4, pp. 235–243, 2001

[120] A. Henry, X. Li, H. Jacobson, S. Andersson, A. Boulle, D. Chaussende, and E. Janzén,

“3C-SiC Heteroepitaxy on Hexagonal SiC Substrates,” Materials Science Forum, vol. 740–

742, pp. 257–262, 2013

[121] X. Li, H. Jacobson, A. Boulle, D. Chaussende, and A. Henry, “Double-positioning

boundaries free 3C-SiC layers grown on-axis 4H-SiC,” Journal of Solid Science and

Technology, vol. 3, p. P75, 2014

[122] A. Henry, S. Leone, F. C. Beyer, S. Andersson, O. Kordina, and E. Janzén, “Chloride based

CVD of 3C-SiC on alpha-SiC substrates,” Materials Science Forum, vol. 679–680, p. 75,

2011.

[123] K. Alassaad, V. Soulière, F. Cauwet, D. Carole, B. Doisneau, and G. Ferro, “Understanding

of the growth mechansim leading to twin boundary elimination during 3C-SiC hetereoepitaxy

on alpha-SiC sybstrates by CVD,” Materials Science Forum, vol. 821–823, p. 209, 2015.

[124] M. Soueidan and G. Ferro, “A Vapor–Liquid–Solid Mechanism for Growing 3C-SiC

Single-Domain Layers on 6H-SiC(0001),” Advanced Functional Materials, vol. 16, no. 7, p.

975–979, 2006.

[125] R. Vasiliauskas, M. Marinova, M. Syväjärvi, R. Liljedahl, G. Zoulis, J. Lorenzzi, G. Ferro,

S. Juillaguet, J. Camassel, E. K. Polychroniadis, and R. Yakimova, “Effect of initial substrate

conditions on growth of cubic silicon carbide,” Journal of Crystal Growth, vol. 324, no. 1, p.

7–14, 2011

[126] B.J. Baliga, Fundamentals of Power Semiconductor Devices. Springer, Berlin, 2008


[128]https://www.infineon.com/cms/en/product/power/sicarbide-

sic/channel.html?channel=ff80808112ab681d0112ab6a50b304a0

[129] http://www.genesicsemi.com/commercial-sic/sic-schottky-rectifiers/

[130] A. Pérez-Tomas, P. Brosselard, J. Hassan, X. Jorda, P. Godignon, M. Placidi, A. Constant,

J. Millan, and J.P. Bergman, “Schottky versus bipolar 3.3 kV SiC diodes,” Semiconductor

Science and Technology, vol. 23, p. 125004, 2008


53

[131] H. Lendenmann, F. Dahlquist, N. Johansson, R. Soderholm, P.A. Nilsson, J.P. Bergman,

and P. Skytt, “Long term operation of 4.5kV PiN and 2.5kV JBS diode,” Materials Science

Forum, vol. 353–356, p. 727, 2001

[132] M. Skowronski and S. Ha, “Degradation of hexagonal silicon-carbide-based bipolar

devices,” Journal of Applied Physics, vol. 99(1), p. 011101, 2006

[133] H. Niwa, J. Suda, and T. Kimoto, “21.7 kV 4H-SiC PiN Diode with a Space-Modulated

Junction Termination Extension,” Appl. Phys. Express, vol. 5, no. 6, p. 64001, 2012

[134] S. Allen, V. Pala, E. VanBrunt, B. Hull, L. Cheng, S. Ryu, J. Richmond, M. O’Loughlin,

A. Burk, and J. Palmour, “Next-generation planar SiC MOSFETs from 900V to 15KV,”

Materials Science Forum, vol. 821–823, p. 701, 2015

[135] R.J. Callanan, A. Agarwal, A. Burk, M. Das, B. Hull, F. Husna, A. Powell, J. Richmond, S-

H. Ryu, and Q. Zhang, “Recent progress in SiC DMOSFETs and JBS diodes at Cree,” Proc.

34th Ind. Electron. Conf, p. 2885, 2008

[136] H. Miyake, T. Okuda, H. Niwa, T. Kimoto, and J. Suda, “21-kV SiC BJT with Space-

Modulated Junction Termination Extension,” IEEE Electron Device Letters, vol. 33 (11), p.

1598, 2012

Chapter 2. Growth process of SiC epilayers and characterization techniques

54

Contents

Chapter 2 Growth process of SiC epilayers and characterization techniques .................................... 55

2.1. Introduction ........................................................................................................................................ 55

2.2. Hot wall Chemical vapor deposition (CVD) reactor ....................................................................... 55

2.2.1. Operational principle ..................................................................................................................... 55

2.2.2. Growth of SiC in a hot-wall CVD reactor ..................................................................................... 56

2.2.2.1. Boundary layer model ............................................................................................................ 57

2.2.3 Kinetics influence ........................................................................................................................... 62

2.2.3.1. Temperature influence on growth rate.................................................................................... 62

2.2.3.2 Precursor supply influence on the growth rate ........................................................................ 65

2.2.3.3. C/Si ratio influence on growth rate ........................................................................................ 67

2.2.3.4. Growth pressure influence on growth rate ............................................................................. 69

2.2.4. Doping of the SiC layers in CVD reactor ...................................................................................... 71

2.2.4.1. Trimethylaluminum (TMA) as p-type precursor gas.............................................................. 71

2.2.4.2. Doping incorporation.............................................................................................................. 74

2.3. Characterization techniques .............................................................................................................. 76

2.3.1. Surface inspection techniques: ...................................................................................................... 76

2.3.1.1. Optical Microscopy (OM) ...................................................................................................... 76

2.3.1.2. Scanning Electron Microscopy (SEM) ................................................................................... 77

2.3.1.3. Atomic Force Microscopy (AFM) .......................................................................................... 78

2.3.2. Optical characterization techniques: .............................................................................................. 81

2.3.2.1. Fourier Transform Infrared Spectroscopy (FTIR) .................................................................. 81

2.3.2.2. Low Temperature PhotoLuminescence (LTPL) ..................................................................... 84

2.3.2.3. Raman scattering spectroscopy .............................................................................................. 89

2.3.3 Structural characterization: X-ray Diffraction ................................................................................ 92

2.3.4. Chemical characterization: Secondary Ion Mass Spectrometry (SIMS) ....................................... 95

2.3.5 Electrical characterization .............................................................................................................. 98

2.3.5.1 Capacitance-Voltage technique (C-V)..................................................................................... 99

2.3.5.2 Hall Effect measurements ...................................................................................................... 104

2.4. Conclusions ....................................................................................................................................... 108

2.5. Bibliography ..................................................................................................................................... 108


55

Chapter 2 Growth process of SiC epilayers and characterization techniques

2.1. Introduction

In this chapter, the first part is devoted to the discussion of the chemical vapor deposition

reactor and deposition mechanism of the SiC epitaxial layers. The second part describes different

characterization techniques used in order to get detailed information regarding the grown SiC

epilayers. The investigation techniques of the material properties presented is divided by

categories such as: surface inspection, optical, chemical, structural and electrical characterization.

2.2. Hot wall Chemical vapor deposition (CVD) reactor

The epitaxial growth of all the epilayers presented throughout this thesis, has been

performed in two home-made reactors based on an innovative adaptation of the hot wall reactor

concept patented by André Leycuras [1]. The reactors are developed in the frame of a

cooperation between NOVASiC and Centre de Recherche sur l'Hétéro-Epitaxie et ses

Applications (CRHEA-CNRS) located at Valbonne.

First reactor denoted hereafter R1 reactor (property of CRHEA-CNRS) has a capacity of 1x2”. In

this work, it was mostly used for the growth of aluminum doped epilayers. The second one

denoted R2 reactor (property of NOVASiC) has a capacity of 1x100mm. Initially not equipped

with an Al-doping line, R2 was used for the elaboration of nitrogen doped epilayers. It is worth

pointing out, that after the installation of TMA line in R2 reactor, the growth process of Al-doped

SiC layers was successfully up-scaled to 100mm wafers.

NOVASiC fabricated two more reactors with a capacity of, respectively, 2x4” and 1x6”. They

were not used in the present study.

2.2.1. Operational principle

The originality of the CVD system lies in the heating method. Contrarily to other hot wall

reactors that are using in general the inductive heating produced by radio-frequency (RF)

generators, our reactor chamber is heated by two independent graphite resistive elements, as

represented in fig. 2.1. The resistive heating elements are placed above and under the susceptor

area. This concept of resistive heating gives the opportunity to control also the vertical

temperature gradient. Moreover, the radiative heat losses are reduced due to a thermal screen


56

system that encloses the susceptor and the heating elements. The internal pieces and the thermal

screens are placed inside a double wall water cooled stainless steel chamber. Besides, the SiC

substrate is placed on a rotating graphite plate (susceptor) that is coated by polycrystalline SiC.

The gases are introduced inside the rectangular section (liner) of the chamber through an inlet

pipe and towards the heated susceptor as can be seen in fig.2.1.

Fig. 2.1 Historic sketch of the CVD reactor concept for SiC deposition

The temperature inside the reactor is measured at the center of the susceptor by the help

of an optical bichromatic pyrometer for T>500°C (wavelength =1.28-1.65µm). Additional

information about the growth temperature is obtained by two thermocouples system. During the

reactor heating, growth temperatures as high as 1850°C in R2 and 1650°C in R1 can be attained.

The growth temperature range is suited for both heteroepitaxy i.e. cubic polytype 3C-SiC on

silicon substrates ( ~1200-1400°C) and homoepitaxy i.e. hexagonal polytypes like 4H-SiC

epilayers ( ~1550-1850°C). The process pressure during growth can be varied in the range of 50

to 800mbars. Usually, the growth was performed at pressures of 100-200mbars.

2.2.2. Growth of SiC in a hot-wall CVD reactor

The epitaxial growth of SiC layers in our horizontal, low pressure resistivity heated hot

wall CVD reactor is performed using classical precursors: silane (SiH4) and propane (C3H8).

Hydrogen (H2) purified by a palladium membrane is used as a carrier gas. The gas sources used

for n- and p-type doping are the nitrogen (N2) and the trimethylaluminum (TMA) respectively.


57

When looking into the epitaxial CVD growth process system at a macroscopic scale,

several key steps are found to be involved: starting with transportation of the precursors by the

carrier gas into the deposition zone of the reactor. Once the gases arrive in the deposition area,

chemical reactions occur which are strongly temperature dependent. Moreover, due to thermal

losses, such chemical reactions appear also outside the deposition area. Nevertheless, the

chemical composition of the gas mixture allows the generation of reactants. The generated

intermediate species diffuse across a boundary layer to reach the growth surface. On the grown

surface the gaseous reactants will be involved in surface reactions such as adsorption,

dissociation of the molecules and surface diffusion towards the grown sites, which result in film

formation or in some cases on surface etching. The last steps involved in the process include the

byproducts desorption and their transport away from the deposition zone. These CVD growth

process deposition steps are illustrated in the fig. 2.2.

Fig. 2.2 Schematic illustration of some of the gas-phase and surface processes during CVD

deposition [2]

2.2.2.1. Boundary layer model

Inside the CVD chamber, the mass transport phenomena can be described by fluid

dynamics and diffusion of reactive species from gas phase towards crystal surface.


58

The fluid dynamics is characterized by the dimensionless parameter named Reynolds

number ( ), defined by:

(2.1)

where is the gas velocity, the gas density, the gas viscosity and is the tube diameter. The

number describes the flow regimes, i.e. laminar if <2100 or turbulent if >2100.

In general, CVD systems are characterized by a low number which indicates a laminar flow

regime (e.g. for the 4H-SiC process we calculated a number of ~12 in R1 and ~20 in R2

reactor) although, the heating of the walls can give rise in some cases to convective currents.

Since the flow regime is laminar, a stagnant boundary layer at the surface of the substrate is

developed. In the boundary layer, the gas velocity of the species is considerably reduced

compared to main gas stream and the reactants and products diffuse across this layer to reach the

deposition substrate as represented in fig. 2.3.

The boundary layer thickness can be considered as an important parameter to control the

growth rate (deposition mechanism) at different substrate positions along reactor axis. In 1970,

Eversteyn et al. [3] proposed a “stagnant layer” model for Si deposition in a CVD horizontal

reactor to determine the growth rate ( ) and improve the epilayer thickness uniformity.

Fig. 2.3 Schematic illustration of a boundary layer formation under a laminar flow

The authors presented a quantitative model explaining the growth rate evolution along the

reactor axis, assuming a simplified vision of the gas flow inside the reactor for a non-tilted


59

susceptor: (i) the gas velocity is constant over the height of reactor tube h except in the

boundary layer (stagnant layer) close to the surface with a thickness ; (ii) in the boundary layer

the gas velocity is zero and the reactants are diffusing across the layer towards the grown surface.

The SiC growth is usually performed under carbon-rich environment (C/Si ratio≥1), and thus the

growth rate is controlled by the supply of Si species in the gas phase [4]. Therefore, the proposed

model for Si growth described by Eversteyn et al. can be applied for the SiC system to determine

the growth rate. Keeping the assumptions described above for a non-tilted susceptor, the growth

rate along the reactor axis can be calculated as:

(2.2)

where represents the position along the reactor axis, e.g. =0 at the beginning of the reactor hot

zone (in our case the upstream limit of the heating element); is the gas velocity at ambient

temperature , is the susceptor temperature, represents the partial pressure of silane

calculated by the ratio of silane flow and the total flow rate; D0=0.2cm2s

-1 is the diffusion

coefficient of silane in hydrogen (across boundary layer), =8.314JK-1

mol-1

is the universal gas

constant, h is the free height above the susceptor and =12.491cm3 is the volume

of a mole of crystalline SiC, given by the ratio between the SiC molar mass (40.096gmol-1

) and

density (3.21gcm-3

). The only unknown variable consider as adjustable and variable along the

reactor axis is the thickness of the boundary layer.

The boundary layer thickness at point can be determined as:

(2.3)

A second model has been proposed by the same authors in the paper, which takes into

consideration this time the susceptor tilt, leading to a decrease of boundary layer thickness and

gas variation. As noticed from eq. 2.3 the boundary layer thickness is increasing as the square

root of the distance along the reactor axis. Therefore, for obtaining a more uniform deposition

rate along the substrate surface, the tilting of the substrate can be adopted in order to maintain a


60

constant . This method leads to a cross-section diminution, as a result of the gas velocity

increase along the substrate in order to achieve a more constant growth rate. However, the growth

rate determination by this model is more complex.

It is important to point out, that a decrease of results in an increase of growth

rate . The decrease of is given by an increase in gas velocity , which can be adjusted

by modifying the growth parameters such as: pressure, temperature and gas flow rate as described

by eq.2.4.

(2.4)

where is the carrier gas (hydrogen in our case) flow rate, is the cross-section area of the

reactor, it the substrate temperature, and is the process pressure. While and are the

standard room temperature and atmospheric pressure at which the mass flow rates are measured.

A reduction of process pressure brings an increase of due to a decrease of , while

an increase of growth temperature or gas flow rate leads also to the reduction.

The theoretical growth rate has been calculated by the first model described by eq. 2.2

for R1 reactor, even though it doesn’t take into account the complexity of the gas flow. The

theoretical growth rate has been determined along the reactor axis for various hydrogen flow

rates while maintaining constant the reactor pressure at 350mbar and the silane flow at

4sccm, as represented in fig. 2.4. The boundary layer thickness for the simulation of has been

adjusted to ~0.1cm, being similar with the one calculated with the formula exposed upper and in

ref. [3], i.e. =0.17cm. Figure 2.4 displays the simulated profile of the growth rate along reactor

axis for different flow rates of carrier gas . Regardless of the input , the growth

rate decreases exponentially along reactor axis. Yet, our consideration is to have a larger

deposition area to cover the substrate and furthermore, to obtain high quality epilayers.


61

Fig. 2.4 Predicted growth rates along the R1 reactor axis for different carrier gas flow

rates with a constant p=350mbar and a =4sccm. The masked rectangle indicates the position

of the substrate during the experiments.

The simulation profile along the reactor axis illustrates that using low carrier gas flow

rates (e.g. =4slm), the deposition takes place mostly in the upstream of the reactor. Increasing

the flow rate to about 14slm leads the deposition towards the downstream, while at substrate

level a higher growth rate is observed despite stronger dilution of precursors. A further increase

of flow rate towards 30slm results to a decrease on the substrate, which coincides with

expected tendency assuming the increase of precursor dilution.

The simple model proposed by Eversteyn et al. that doesn’t account the complexity of gas

flow inside the reactor, seems to indicate in a qualitative manner the growth kinetics of SiC inside

R1 reactor. Similar calculation for determining the (x) profile have been done for R2

(1x100mm) reactor, and similar tendencies with R1 reactor were observed.

0 5 10 150

5

10

15

20

25

30

Gro

wth

ra

te (

µm

/h)

Position along reactor axis x (cm)

H2=4slm

H2=10slm

H2=14slm

H2=30slm

substrate


62

2.2.3 Kinetics influence

2.2.3.1. Temperature influence on growth rate

The growth rate dependence on process temperature is described by eq. 2.2. The predicted

growth rate shows a gradual increase with the raise of process temperature, which can be

attributed to an enhancement of reactants diffusion across the boundary layer. In the gas phase

the decomposition of precursors, movement and reaction of reactants is influenced by system

temperature [5]. Therefore, in epitaxial growth the mass transport of reactants from gas phase to

surface is governed by diffusion, which is defined as the migration of reactive species caused by

a concentration gradient [6], [7]. The process temperature has a strong influence on the surface

processes giving rise to different growth regimes.

Fig. 2.5 Schematic representation of the deposition rate as a function of inverse temperature with

its different growth regimes

Figure 2.5 displays the rate-limiting mechanisms existent in a deposition system when the

temperature is varied. At low temperature the growth rate is limited by surface-reactions such as

chemical reaction, surface migration, lattice incorporation due to slow surface diffusion. At

higher temperatures (i.e. in an intermediate range of temperatures) the growth rate seems to be

independent of process temperature, and governed by a mass-transport limited regime. Increasing

further the process temperature, a decrease in growth rate is observed even though the process is


63

still governed by mass-transport. The obtained growth rate reduction in CVD system is driven by

some thermodynamic factors such as H2 etching [7] or enhance desorption rate.

For determining the growth regimes in our reactor, an analysis of the growth rate

dependence on growth temperature was performed on both polarities of 8° off-axis 4H-SiC

substrates. The growth temperature was varied in the range 1400-1800°C while all the other

process parameters were kept constant. The set process conditions were: reactor pressure

400mbar, flow rate 35slm, flow of 12sccm and a flow of 6sccm. A C/Si ratio

(i.e. the ratio of carbon and silicon atoms of supplied precursors) of 1.5 was applied. The

obtained growth rate values for both Si- and C-face of 4H-SiC epilayers when the growth

temperature ( ) was varied are represented in fig.2.6.

Fig. 2.6 Growth temperature dependence of growth rates for Si-face (blue symbols) and C-face

(red symbols) of off-axis 4H-SiC epilayers

When varying the temperature in the range 1400-1600°C ( registered by pyrometer) on

both polarities, the growth rate shows a weak increase. The slight dependence of the growth rate

on temperature on Si-face reveals a mass-transport limited regime, which is in agreement with

reported results in literature [8], [9]. Increasing further the temperature in the range 1650-1800°C,

4.8 5.0 5.2 5.4 5.6 5.8 6.0 6.20

2

4

6

8

10

12

14

C-face

Gro

wth

ra

te (

µm

/h)

1/Tg (10-4 K-1)

SiH4=12sccm

Si-face

1800 1700 1600 1500 1400

Tg (°C)


64

the growth rate starts to decrease for both Si and C-face, respectively. A negative activation

energy has been calculated from the Arrhenius law on both polarities, i.e. on Si-face equal to -

10.8kcal/mol (-0.47eV), and on C-face equal to -11.9kcal/mol (-0.52eV).

If we take into consideration that the growth rate is determined by deposition rate and

etching rate as shown by eq.2.5, the negative activation energy can be explained.

(2.5)

This indicates the presence of two competing mechanisms during the growth process, in

which one becomes more relevant with increasing temperature. Thus, in the range 1650-1850°C

in our CVD reactor the enhanced mechanism is the etching, and as a result a drop in growth rate

is observed. Consequently, the calculated on both substrate polarities becomes positive for

etching mechanism.

It has been confirmed in the literature that the etching rate of hydrogen in SiC growth

increases exponentially with temperature and proportionally with the hydrogen flow rate [10]. In

order to determine the contribution of hydrogen etching on SiC surface, an etching run was

performed using only a hydrogen atmosphere in R2 reactor. A 4H-SiC wafer was placed on the

non-rotated susceptor in order to study the effect of hydrogen flow along the reactor axis. For this

experiment, similar conditions were applied as for the 4H-SiC deposition process, except for the

precursors flow rate. During one hour the sample was heated at a temperature of 1800°C

(temperature indicated by pyrometer) under a hydrogen flow rate of 35slm. A reactor pressure of

100mbars was applied.


65

Fig. 2.7 Hydrogen etching rate determined at a temperature of 1800°C along the reactor

axis under: (a) mapping form and (b) x-y form

Figure 2.7 shows the etching/deposition rate deduced from thickness measured along the

substrate diameter, parallel and perpendicular to the hydrogen flow by Fourier Transform

Infrared Spectroscopy (FTIR). As can be noted, an etching (negative growth rate) is present on

the wafer, with a maximum rate of ~7µm/h in the center of the wafer. Along the gas flow

(parallel axis) i.e. in the left side of the wafer, an increase in growth rate as high as ~7µm/h is

seen. This is proportional with the etching rate, indicating a redeposition of material in the

downstream of the reactor. A similar profile of hydrogen etching along the reactor axis was

observed by Meziere et al. in ref. [11]. Perpendicular to the gas flow, a symmetric profile is seen.

Moreover, a similar test was carried out on a 4H-SiC substrate under an argon (Ar) flow of 5slm,

while keeping the same growth conditions described above. The FTIR analysis indicated an

etching rate of ~0.04µm/h along the reactor axis (i.e. no influence of Ar gas on surface). This

confirms that under the range of temperatures used for the 4H-SiC process, the observed decrease

in growth rate can be attributed to an increase hydrogen etching rate.

2.2.3.2 Precursor supply influence on the growth rate

As described by eq. 2.2 the growth rate shows a dependence on the precursor flow

determined by the silane (SiH4) partial pressure. The silicon supply from the gas phase

determines the deposition rate at the surface of the substrate under C-rich conditions.

0

10

20

30

40

0

90

180

270

0

10

20

30

40

Po

sit

ion

(m

m)

-8-6-4-202468

gas flow

µm/h

a)-40 -20 0 20 40

-8

-4

0

4

8

Etc

hin

g r

ate

(µ

m/h

)

Position (mm)

// axis axis

Gas flow


66

The SiH4 partial pressure (i.e. ) is defined by the following equation:

(2.6)

where is the reactor pressure, and are the flow rate of silane and propane

(precursors gases), and is the hydrogen (carrier) gas, respectively.

An analysis to determine the growth rate evolution as a function of the input silane flow

rate was made in R2 reactor on 2° off-oriented 4H-SiC substrates, and represented in

fig.2.8.

Fig. 2.8 Input silane flow rate dependence of the growth rate for 4H-SiC epilayers growth at

1770°C. The solid line is guidance for the eyes.

The flow was varied in the range 5-40sccm, while the process conditions were kept

stable: growth temperature of 1770°C, a flow equal to 35slm and a reactor

pressure=150mbar. During the growth, the C/Si flow rate ratio of 1.2 was fixed by varying the

input flow of propane. It is important to note that the growth rate may be dominated by the

0 5 10 15 20 25 30 35 40 450

5

10

15

20

25

30

35

40

Gro

wth

ra

te (

µm

/h)

Silane flow rate SiH4

(sccm)

H2=35 slm

p=150 mbar Si-face


67

carbon supply when the C/Si ratio <1 (Si-rich conditions) whereas under higher C/Si ratio>1 (C-

rich conditions) the growth rate may be dominated by the silicon supply.

From the graph it is obvious that the growth rate increases proportional with the input

flow rate introduced in the gas phase i.e. the growth rate is fed by the Si-supply. Moreover,

in literature [12] it was reported a saturation of growth rate at high Si-supply, due to a nucleation

(polymerization) of Si clusters in the gas phase as described previously in Chapter 1.

The growth rate was determined from thickness measurements by FTIR. A growth rate of

~33µm/h can be obtained for a =40sccm, which indicates the possibility of obtaining a

wide range of deposition rates by adjusting the silane flow and process conditions. However,

when a low flow is used during the growth, the competition between deposition and

etching starts to be visible and, the mean etching rate can be predicted (e.g. at =5sccm an

etching rate of ~3µm/h was displayed).

2.2.3.3. C/Si ratio influence on growth rate

Besides the influence, the input C/Si ratio has also an impact on the deposition rate

of SiC layers. Even though, the eq. 2.2 doesn’t describe the C/Si ratio as an input parameter, it is

impacting the gas phase directly. As described above, the C/Si ratio describes the ratio of C and

Si atoms available in the supply of precursors thus, influencing the silane partial pressure

(eq.2.6). Therefore, the input C/Si ratio is defined as the ratio between a triple of propane

flow rate and silane flow rate as described by eq. 2.7 [13].

(2.7)

Adjusting the C/Si ratio, the deposition can appear theoretically under two limited

deposition regimes: (i) under Si-supply limited (for C/Si>1) or (ii) under C-supply limited (for

C/Si<1). Usually, during the process the growth rate is maintained constant by keeping fixed the

SiH4 flow while varying the C3H8 flow rate. For obtaining a good surface morphology, the

epitaxial process should be performed at C/Si ratios that offer near-stoichiometric conditions.

Low C/Si ratio (Si-rich conditions) gives rise to Si droplets that affect the surface morphology. In

the case of high C/Si ratio (C-rich conditions), an excess of carbon is present which can cause


68

different defects and surface degradation. The study of the growth rate as a function of the C/Si

ratio was investigated in both reactors by varying the ratio of C/Si in the range 0-3 while keeping

constant all the other growth parameters. As mention a change in C/Si ratio can influence the

growth rate Gr during the growth process [13], [14]. The C/Si ratio was adjusted in our

experiment by two methods: (i) changing the C3H8 flow rate as the SiH4 flow was maintained

constant and, (ii) varying the SiH4 flow rate while the C3H8 flow was fixed. The results of the

growth rate as a function of C/Si ratio for both reactors are represented in fig. 2.9(a)-(b). The

growth rate is represented as the normalized with the obtained at C/Si=1 (i.e. the theoretical

point of stoichiometry).

Fig. 2.9 Dependence of normalized growth rate on C/Si ratio when the C3H8 flow rate

(SiH4 flow rate) was varied and the SiH4 flow (C3H8 flow rate) was fixed for: (a) R2 and (b) R1

reactor

Let’s start our discussion with the first method applied, i.e. changing the propane flow

rate while the silane flow rate was fixed. It can be seen from fig. 2.9(a), that up to a C/Si=1.5, the

increases with the increase of C/Si ratio. It suggests that in this range, the growth rate is

limited by the carbon supply. The input C/Si ratio of 1.5 corresponds to “stoichiometric”

0.1 1 100.1

1

10

C-rich conditionsSi-rich conditions

Gr /

Gr (

C/S

i=1

)

C/Si ratio

C3H

8 variation

SiH4 variation

R2 reactor

a)

0.1 1 100.1

1

10


b)

R1 reactor

Gr /

Gr (

C/S

i=1

)

C/Si ratio

C3H

8 variation

SiH4 variation


69

conditions in the deposition zone. For C/Si>1.5, a saturation of the growth rate is observed

indicating that the deposition regime is changing to Si-supply limited.

To avoid the decrease of deposition rate in the C-supply limited regime, a second method

was applied in which the silane flow rate was varied, and the propane flow rate was fixed. As can

be observed, using such method leads to almost constant growth rates in the given range of C/Si

ratios. However, the inconvenient of working at low C/Si ratio remains the apparition of Si

droplets that results in a deterioration of the surface morphology.

In the case of R1 reactor (fig. 2.9b) similar deposition regimes are found. The

“stoichiometric” conditions were noticed at the input C/Si of 0.5 for this deposition zone. Under

C/Si ratio<0.5, an almost constant growth rate was noticed compared to R2 reactor when the

propane flow rate was varied. Moreover, the SiC deposition at an input C/Si ratio equal to 0 (no

C3H8 flow rate added to the gas phase) was obtained. This suggests the existence of a carbon

supply source inside the reactor. When the silane flow rate was varied while maintaining the

propane flow rate constant, slight variation of the growth rate was seen. This effect indicates the

presence of some fluctuations of the carbon supply in the gas phase given by the flow meter as

minimum set point was reached. Therefore, in our study we will need to consider besides the C/Si

ratio, also the influence of growth rate variation on dopant incorporation.

2.2.3.4. Growth pressure influence on growth rate

Reactor pressure is not explicitly present in eq. 2.2, however it influences different

process parameters during growth such as: gas velocity (eq.2.4), boundary layers thickness (eq.

2.3), and precursor partial pressure (eq.2.6). Furthermore, the proportions of various reactants and

the surface mobility of adatoms [15] are influenced by process pressure. In the CVD process of

SiC layers, low pressures are often applied in order to obtain higher growth rate and better

uniformity of the layers [12].


70

Fig. 2.10 Experimental dependence of mean growth rate on reactor pressure at fixed flow

rates of precursor and carrier gas and line is present as guidance for the eyes

Our CVD reactors are designed to work at low pressures, thus the study of the growth rate

dependence on reactor pressure was performed by varying the pressure in the range 50-800mbars.

As substrates we used 4° off axis Si-face 4H-SiC wafers. The set conditions for the growth were:

a =1700°C, =35slm, a C/Si=2 with a SiH4 and C3H8 flow rates of 12sccm and 8sccm,

respectively. Figure 2.10 represents the obtained as a function of reactor pressure. From the

graph, we observe a maximal growth rate at a reactor pressure in the range 100-200mbar.

When the total pressure is increased up to 800mbars, an important decrease of deposition rate is

seen. This can be attributed to two effects: firstly, to an increase of boundary layer thickness at

high pressures which limits the diffusion of reactants towards the substrate. Secondly, to a

reduction of gas velocity which lead the deposition zone mostly towards the upstream part of the

reactor (i.e. a reduction of reactive species above the substrate).

Moreover, decreasing the reactor pressure below 100mbar results also to a reduction of

growth rate even though the boundary layer thickness is already diminished. Low pressures give

rise to higher gas flow velocity and, as a consequence, the deposition zone moves towards the

100 10000

4

8

12

16

Gro

wth

ra

te

(µm

/h)

preactor (mbar)

SiH4=12sccm

H2 =35sccm

Tg =1700°C

Si-face


71

downstream part of the reactor. Additionally, it reduces the residence time of reactant species

which becomes also a limiting factor for the deposition rate.

2.2.4. Doping of the SiC layers in CVD reactor

The n- and p-type doping of SiC layers in our reactors are performed using molecular

Nitrogen (N2) and respectively, organometallic compound Trimethylaluminum (i.e. TMA or

Al(CH3)3) as precursors sources.

2.2.4.1. Trimethylaluminum (TMA) as p-type precursor gas

The TMA is a liquid organometallic compound that needs a more complex operational

mode for the doping flow compared with the nitrogen precursor. The TMA compound is

available in a stainless steel bubbler. The TMA bubbler is installed under a hydrogen line that

needs to be maintained under constant hydrogen pressure and temperature . The

temperature is controlled by a heated bath in which the TMA bubbler is immerged as shown in

fig. 2.11. A hydrogen gas flow passes through the bubbler where it saturates in

organometallic compound. The transported TMA flow can be determined using the following

expression:

(2.8)

where is the hydrogen flow in the bubbler, is the saturated vapor pressure of

TMA, which depends on the bubbler temperature and bubbler total pressure (i.e. usually set at

~2bars).

The saturated vapor pressure of the TMA depends strongly on the bubbler temperature

and can be expressed following the Clapeyron approximation [9]:

(2.9)


72

where the is given in mmHg and the is the bubbler temperature in Kelvin. When the

bubbler temperature is in the range 18-60°C, the and coefficients given in the eq. 2.8 are:

=2134 and =8.22 [15].

Our doping experiments were carried out using a TMA bubbler temperature =20°C

resulting in a saturation vapor pressure of TMA: ~8.72mmHg. It has to be pointed out

that any variation in bubbler temperature or pressure can result in unwanted variation of the TMA

doping flow during the growth.

Figure 2.11 shows the TMA line installed for both (a) R1 and (b) R2 reactors used during

this thesis. The TMA line is heated in order to avoid any undesired TMA condensation that may

block the flow towards the reactor. In a first step, the H2 enters the bubbler where it gets saturated

in aluminum compound (H2b). The gas mixture exiting the bubbler is further diluted with pure

hydrogen (H2d1 and H2d2). However, from the quantity of obtained mixture, only a small amount

is injected in the reactor during growth (i.e. minimum amount of TMA obtained at H2P2=1.5x10-

7sccm). For stronger Al concentration of the epilayers, the doping can be made directly with the

TMA amount extracted after the bubbling step (H2b).


73

Fig. 2.11 Schematic representation of the TMA line in (a) R1 and (b) R2 reactor


74

Any excess of TMA gas flow is sent towards the evacuation line (vent line in fig 2.11)

through the regulation valves. As one may notice from the schematic representation, a difference

between the two TMA lines installed on R1 and R2 reactors is visible. The distinction between

the two TMA lines is represented by an additional line of dilution (H2d2) in R2 reactor. The

second TMA dilution line was added to achieve a wider control range over the Al concentration

(i.e. from 3x1016

cm-3

to 2x1020

cm-3

) during the epitaxial growth. Unfortunately, the TMA

precursor presents also the disadvantage of a strong “memory effect”. This can be reduced only

by an annealing of the reactor under hydrogen flow prior to growth or using an HCl flushing [17].

Also, a slow stabilization step of the gas flow in the reactor was found to be necessary.

On the other hand, for the case of N2 doping, the control of dopant supply is easier than

for TMA. The nitrogen is diluted in pure hydrogen and, further the mixture is injected in the

reactor for the epitaxial growth. The nitrogen/hydrogen mixture flow injected in the reactor can

be determined as follows:

(2.10)

where the is the diluted nitrogen flow rate, the is the pure nitrogen flow rate and the

is the pure hydrogen flow rate. The nitrogen doping in our reactors can be controlled in the

range of ~1x1014

cm-3

to 4x1019

cm-3

.

2.2.4.2. Doping incorporation

The doping experiments that are presented within this thesis have been accomplished in the

two reactors described above as follows:

a) The p-type doping experiments on different SiC substrates have been carried out in the R1

reactor (1x2”). The doping process was then up-scaled to 100mm in R2 reactor,

b) The n-type doping experiments have been carried out only in the R2 reactor.

As presented in the previous section, when the susceptor is immobile during the growth, a

variation of the growth rate along the reactor axis is created and the resulting layer is non-

uniform in term of thickness. Also the dopant incorporation follows the same non-uniformity.

The variation of the voluntary doping incorporation (the case of nitrogen and aluminum) during


75

the growth of 4H-SiC and 3C-SiC epilayers with no substrate rotation has been studied. The

obtained results are represented in fig. 2.12.

Fig. 2.12 Aluminum and nitrogen doping profile on SiC epilayers deposited without substrate

rotation. Measurements performed (a) along the reactor axis and (b) perpendicular to the reactor

axis. The 0 position indicates the sample center and the arrow indicates the gas flow direction.

The doping concentration has been obtained for the case of 4H-SiC layers by

Capacitance-Voltage (C-V) measurements, while for the 3C-SiC layers the concentration was

deduced from FTIR spectra [18]. In 3C-SiC epilayers the doping concentration was not

determined by C-V measurements since the surface roughness and the defect density leads to

poor rectifying contact that give rise to measurement errors.

As can be seen, the aluminum incorporation decreases along the reactor axis while the

nitrogen incorporation shows an opposite tendency, i.e. increases along the axis. This

experimental observation for the nitrogen incorporation may be associated to the stronger

bonding of N2 molecule which needs more time/temperature to decompose [11], [19]. In the case

of aluminum, the TMA molecule decomposes much faster (at temperature above 100°C) which

can explain the decrease trend obtained experimentally.

-50 0 50

1015

1016

1017

1018

1019

Al

N

N

axis// axis

Dopin

g (

cm

-3)

position (mm)

gas flow

-50 0 50

1015

1016

1017

1018

1019

N in 3C (FTIR, R2) N in 4H (C-V, R2) Al in 4H (C-V, R1)

b)a)


76

2.3. Characterization techniques

2.3.1. Surface inspection techniques:

2.3.1.1. Optical Microscopy (OM)

Considering that for device fabrication the requirement is a high quality of epilayer

surface, an attention needs to be directed towards the morphology. Optical Microscopy (OM) is a

fast and non-destructive method, which was used after each growth to determine the morphology,

uniformity and epilayer defects. Moreover, OM was carried out also to reveal any presence of

3C-SiC domains when the growth was performed on -SiC substrates.

In OM, the essential elements are the magnification, the resolution and the contrast which define

the quality of the specimen image. The magnification and resolution of the OM are somewhat

limited by the light wavelength thus reaching usually a maximum magnification of x1500 and

resolution of 200nm [20].

In the study of our epilayers, Differential Interference Contrast (DIC) also known as

Nomarski DIC (NDIC) configuration was used due to its advantages of improving the contrast of

morphological features like epitaxial defects or surface irregularities. The contrast is generated by

any refractive index gradient on the sample. A difference in brightness is visualized due to

amplitude variations recorded by the reflected light beams. Other contrast configurations such as

Dark Field (DF) and Bright Field (BF) can be applied. Therefore, the OM is a fast operating tool

to assess the epilayer surface quality. This helps to improve the growth process in view of

eliminating the defects of future epilayers.

To quantify the density of defects existing on the epilayer surface, the OM is a good tool but a

long time for image acquisition is necessary. Therefore, when a full surface mapping to estimate

the defect density was required, we have taken advantage of the full automated DIC imaging. In

this manner, a faster and more productive characterization was achieved.

Additionally, the microscope can operate under BF or DF configurations as desired.

Selecting the objective x2.5 in DIC configuration, the mapping of a 100mm wafer takes only 15

minutes (fig. 2.13a). Regarding the defects detection, it is based on the analysis of image contrast

and intensity. The use of such method leads to reliability losses in the favor of smaller acquisition

time. However, the analysis on SiC epilayers by DIC imaging became quickly a standard

characterization technique, since it offers a fast report on the grown surface quality.


77

Figure 2.13 shows a schematic representation of the available optical characterization

techniques for the epilayer surface morphology. As it is noticed, the surface morphology

assessment can be downscaled from macroscopic (fig. 2.13a) to nanometric level (fig. 2.13b).

The OM offers large scale images of the surface morphology with indications on defects and

step-bunching formation. Yet, for more details on defects or macro-steps, higher resolution

microscopy techniques are required as described below.

Fig. 2.13 Schematic representation of the epilayers surface morphology characterized by

different microscopy techniques: (a)-(b) OM, (c) Scanning Electron Microscopy (SEM) and (d)

Atomic Force Microscopy (AFM)

2.3.1.2. Scanning Electron Microscopy (SEM)

Another characterization technique that was used for the study of surface morphology was

the Scanning Electron Microscopy (SEM), since no complex sample preparation is needed except

for a clean surface. This technique was used mostly for detailed information about the epilayer

surface, since the SEM provides a higher magnified image of the surface (fig. 2.13c). The SEM

uses electrons instead of light to scan the surface enabling a higher resolution.

The basic SEM principle includes the focus of an electron beam by different lenses and

apertures towards the surface where it generates various signals due to electron-sample


78

interactions. Signals generated by secondary electrons are collected by detectors, where the

imaging of the sample morphology is formed [20]. Usually the data are collected from a defined

area on the sample, where a magnification ranging from x20 to approximately x30,000 can be

obtained. During our study, a Carl Zeiss Supra40 SEM instrument was used to observe the

surface morphology and the thickness of the grown epilayer in cross-section. Throughout the

work, the SEM was used both for plan-view and cross-section analysis of the aluminum doped

4H-SiC epilayers (section 4.3 of Chapter 4). Also, the SEM was used to monitor the power

device fabrication that will be discussed later in Chapter 5 of this thesis.

2.3.1.3. Atomic Force Microscopy (AFM)

Atomic Force Microscopy (AFM) technique supplies detailed information regarding the

atomic scale roughness of the thin films (fig. 2.13d). It was used to evaluate the height of the

atomic growth steps of the crystal, which potentially can be used for polytype identification. The

measurements were performed in tapping mode, in which the tip comes in intermittent contact

with the sample. In this mode, the scanning is obtained by oscillating the cantilever at a

frequency close to the cantilever resonance using a piezoelectric crystal. The amplitude of

oscillations influenced by the surface irregularities is used as feedback to form the image. The

tapping mode presents the main advantage of allowing higher resolution of the topographic

imaging of sample surface in ambient air. The limiting factor for good resolution is given by the

tip sharpness. During AFM measurements, any observed defect can disturb the tip sharpness and

the recorded image. Nevertheless, compared to the contact mode, the tapping mode allows a

limitation of the tip degradation.

Throughout the work, it was noticed that the homoepitaxially grown layer show very

smooth roughness with Root Mean Square (RMS) values of only few nm. A high incorporation

of nitrogen or aluminum in the epilayers can result to an increase of surface roughness. Of

course, the selected process parameters and/or the substrate off-cut can also play a role in the

surface morphology, i.e. poorly chosen growth parameters can lead to a surface deterioration.

Examples of tapping mode: AFM images recorder from a 4° and 8° off-cut 4H-SiC samples with

a nitrogen concentration variation are represented in fig. 2.14(a)-(b).

The 5x5µm² AFM images recorded on 4H-SiC layers with low nitrogen concentration

[N]=7x1014

cm-3

and high nitrogen concentration [N]=3.8x1019

cm-3

show a step-bunching


79

morphology. The surface roughness increases with the nitrogen concentration enhancement, from

RMS values of ~1.8nm (fig 2.14a) to ~6nm (fig 2.14b). Of course, we may consider that the off-

cut difference may account somehow to the observed roughness difference.

Fig. 2.14 AFM images obtained on:(a) 5x5µm² scan of a 4H-SiC 4° off-axis with a

[N]=7x1014

cm-3

, (b) 5x5µm² scan of a 4H-SiC 8° off-axis with a [N]=3.8x1019

cm-3

and (c)

50x50µm² scan of a 3C/4H heterostructure where the inset shows an 10x10µm² image obtained

on the same sample representing a spiral growth

Figure 2.14(c) displays a 50x50µm² AFM scan of a 3C-SiC epilayers grown on an on–

axis 4H-SiC substrate. The scan revealed triangular patterns that are attributed to 3C-SiC

polytype, and roughness RMS value of ~3nm was found. The inset image shows a 10x10µm²

AFM scan of a spiral step, which is characteristic to an -SiC growth [21].

From the AFM characterization of 3C-SiC layers grown on on-axis 6H-SiC/4H-SiC substrates,

we have found the presence of different domains associated to hexagonal, cubic and even to

polytype mixing.

a)

c)


80

Fig. 2.15 AFM image and the height profile steps obtained after heteroepitaxial growth on

different samples and regions: (a) 6H-SiC, (b) 4H-SiC and (c) 3C-SiC regions

Thus, in some areas based on the height profile from AFM data represented in fig.

2.15(a)-(b), steps with heights of 0.75nm and 1.5nm associated to 6H-SiC polytype or 0.5nm and

1.0nm related to 4H-SiC polytype were observed. In the regions assigned as 3C-SiC polytype

(yellow areas), steps with ~0.25nm height were usually dominant (not illustrated here). This

illustrates a slight dilemma in connecting the step height composed of one Si-C bilayer to 3C-SiC

film formation. Although, a good indicator of 3C-SiC domains formation as represented in fig.

2.15(c) is the presence of planar defects with 60° angles such as Stacking Faults (SFs). AFM

technique proved to be useful to characterize macrosteps formation of hexagonal polytypes.


81

2.3.2. Optical characterization techniques:

2.3.2.1. Fourier Transform Infrared Spectroscopy (FTIR)

Along the years, different works have shown the reliability of Fourier Transform Infrared

Spectroscopy (FTIR) method. Moreover, it is a non-destructive, low-cost, contactless and easy to

use technique that became fast a routinely method to determine epilayer thickness [22].

During this study, we used the technique to determine the thickness and uniformity of full SiC

wafers. The FTIR measurements were performed at room temperature in the wavelength range of

400-5000cm-1

.

A reflectance spectrum can be registered if a smooth air-epilayer and epilayer-substrate

interfaces are present. Nevertheless, for accurately determination of epilayer thickness, the

presence of interference fringes is required. The interference fringes are detected in the case of

homoepitaxy due to a difference in refractive index between the epilayer and the substrate. The

difference can be achieved by strong doping concentration discrepancy between the epilayer and

substrate.

The complex refractive index n of the SiC material in infrared (IR) region can be derived

from the standard form of dielectric constant as follows [23]:

(2.11)

where: is the refractive index, is the extinction index, is the high-frequency dielectric

constant (i.e. for 4H-SiC=6.52), and are the transversal and longitudinal phonon

frequencies of the crystal ( =793cm-1

and =974cm-1

), is the crystal damping parameter

( =5cm-1

), is the measurement frequency, is the damping parameter for the free

electrons/holes which depends on the free-carrier mobility, is the plasmon frequency denoted

as a function of doping concentration and expressed as:

(2.12)


82

where is the Dirac constant, the free-carrier concentration, the electron charge, is the

effective mass, is the vacuum permittivity constant.

Fig. 2.16 FTIR reflectance spectra recorded on a 12 µm thick 4H-SiC epilayer

The dielectric function can be used as model for thickness calculation, and furthermore,

the calculated spectra can be compared with the measured spectrum as reported in ref [22].

Figure 2.16 displays the FTIR spectra collected in the 400-5000cm-1

energy range on a

homoepitaxial grown 4H-SiC sample. As can be observed, the reflectance spectrum exhibits the

interference fringes necessary for thickness evaluation [22], [24]. The fringes were generated by

the difference in doping concentration among the 4H-SiC epilayer ( - ~1016

cm-3

) and the 4H-

SiC substrate, which has usually a nitrogen doping concentration in the range 1018

-1019

cm-3

.

The FTIR spectrum is dominated by the “reststrahlen” (RS) band in the energy range

790-970cm-1

, in which the transversal and longitudinal optical phonon frequencies are located.

From the RS band, information regarding the doping concentration of the grown layers can be

extracted as reported by Portail et al. [18].

The graph inserted in fig. 2.16 shows a selected range of the spectrum, where oscillation

of the fringes with a certain wavelength can be used to calculate the epilayer thickness :

1000 2000 3000 4000 50000

20

40

60

80

100

400 450 500 550 600 6500

10

20

30

40

50

60

Re

fle

cta

nc

e (

%)

Wavenumber (cm-1)


83

(2.13)

where is the number of interference periods that separate the two maxima or minima, and

are the refractive index and wavelength of the two maxima/minima, respectively.

The thickness uniformity of the layer was also investigated with the help of FTIR

technique. The thickness uniformity of a 4H-SiC epilayers grown in our reactors using the

susceptor rotation, shows a typical standard deviation over mean value of around 2%, although

lower value can be attained (~0.7%) while excluding ~3mm from the wafer edges.

Figure 2.17 shows the collected thickness map on a 4° off-axis 4H-SiC epilayer grown

with substrate rotation in the R2 reactor. The growth conditions involved a temperature of

1820°C, reactor pressure of 150 mbar, C/Si ratio=1.8 and a carrier flow rate of 38slm. It was

found that the thickness uniformity of the epilayer with ~3mm edge exclusion showed a standard

deviation over mean value of 2.4%.

Fig. 2.17 Thickness uniformity map of an Al doped 4H-SiC epilayer grown using the substrate

rotation ( - of ~3x1019

cm-3

). The standard deviation over mean with 3mm edge exclusion

was found to be 2.4%.

0

10

20

30

40

0

90

180

270

0

10

20

30

40

Po

sit

ion

(m

m)

19.01

19.37

19.72

20.08

20.44

20.80

21.15

21.51

µm/h


84

2.3.2.2. Low Temperature PhotoLuminescence (LTPL)

Low Temperature Photoluminescence or LTPL is an optical spectroscopy technique

considered as the most informative tool for qualitative characterization of semiconductor thin

films. In the case of SiC layers, LTPL spectrum allows the polytype identification and provides

information about the impurities and defects presented into the material. It is considered as a non-

destructive technique that does not require any sample preparation. Although, if the cryostat size

is smaller than the wafer diameter, the cut of the wafer is necessary, disregarding it as a routine

technique. Also, a quantitative evaluation of the impurity concentration can be extracted after the

correlation between intrinsic and extrinsic emissions.

In this work, the LTPL was mostly used for polytype identification of the 3C-SiC layers

grown on -SiC substrates, which are presented in Annex I. Moreover, in collaboration with

Laboratoire Charles Coulomb in Montpellier, the LTPL technique was used to determine the

dopant concentrations in 4H-SiC epilayers, which is described in detail in the PhD thesis of P.

Kwasnicki [25].

Low Temperature PhotoLuminescence (LTPL) principle

Photoluminescence (PL) is one of many forms of light emission from the material. In PL,

the light emission is initiated by photoexcitation: the electrons are excited from their ground state

by an excitation source (usually a laser) with certain photon energy, creating electron-hole pairs

or excitons. To reach back the equilibrium state, radiative and non-radiative recombinations of

the electron-hole pairs take place.

Since the SiC is an indirect band gap semiconductor, the recombination of electron-holes

pairs occurs only assisted by one or multiple momentum conserving phonons. Furthermore,

contrarily to direct band gap semiconductors where the PL measurements can be performed at

room temperature, the analyses on SiC layers have to be performed at low temperatures (e.g. in

the range 2-5K) in order to obtain the resolution and intensity high enough to correctly interpret

the spectra [26].

The radiative recombinations of electron-holes pairs include:

i) Intrinsic radiative recombinations observed in high purity materials,

ii) Extrinsic radiative recombinations related to the presence of impurities and/or defects

in the epilayers.


85

Intrinsic radiative recombination

During excitation process, an electron from the valence band is displaced to conduction

band and, a positively charge (hole) is created. The process generates the creation of electron-

hole pairs bounded by Coulomb interaction called Free Excitons (FEs).

In this case, when an electron and a hole recombine, a photon is emitted with lower

energy than the band gap, as:

(2.14)

where is the Coulomb energy of the exciton and is the excitonic bandgap energy of the

material.

In LTPL, the intrinsic transition is important since it is a typical indicator of

epilayers/substrate quality. The lines attributed to radiative FEs replicas are labeled on the

LTPL spectrum, where x= and designates the phonon replica energy. However, in SiC material

the most intense FE replica on the spectrum is for 4H-SiC and for 6H-SiC, respectively.

The FE replicas intensities are decreasing with the increase in dopant concentration.

Extrinsic radiative recombination

During extrinsic recombination, the created excitons can bind to impurities (e.g. donor or

acceptor) or to defects, becoming Bound Excitons (BE). In the BE recombination, either the

impurity or the phonon can accomplish the momentum conservation, and the energy of the

emitted photon becomes:

(2.15)

where is the binding energy of the exciton to the impurity/defect.

The transition without the phonon participation gives rise to Zero Phonon Line (ZPL) on

the LTPL spectrum. The quantity of ZPL depends on the number of non-equivalent sites for

impurity incorporation, related to the investigated polytype. Sites can be of cubic (“k”) or

hexagonal (“h”) type as represented in fig. 2.18.


86

Fig. 2.18 Schematic representation of the available sites in the SiC unit cell: (a)

hexagonal “h” site and (b) cubic “k” site

To illustrate this, we can consider the case of exciton bound to nitrogen–donor impurity

which substitutes at carbon sites in the SiC matrix. For 3C-SiC, where all the sites are of the same

type, a single ZPL related to the nitrogen bound exciton can be observed. It is assisted by four

phonon replicas labeled: Transverse and Longitudinal Acoustic (TA and LA) and, Transverse and

Longitudinal Optical (TO and LO). For 4H-SiC, the lines attributed to the nitrogen BE are the P0

(N at hexagonal site) and Q0 (N at cubic site), followed by 24 phonon replicas identified by their

energy position. Finally, for 6H-SiC, three types of available sites are present: one hexagonal plus

two cubic, and the ZPLs are denoted as P0 (N at hexagonal site) and R0, S0 (N at cubic site), each

of them assisted by 36 phonon replicas.

In the case of exciton bound to aluminum (acceptor impurity), the ZPLs related to

polytypes can be observed only at high spectral resolution. Therefore, we observe in the collected

spectra the transition of the ZPLs to 4Al0 complex [27]. In the case of 4H-SiC, the 4Al0 line can

be found between the P0 and Q0 bound to nitrogen as shown in fig 2.19.

Figure 2.19 displays the typical LTPL spectra collected in the Near Band Edge (NBE)

range for 4H-SiC, 6H-SiC and 3C-SiC polytype. The SiC polytype identification is based on the

characteristics lines related to FE, nitrogen or aluminum bound excitons in the NBE range. The

distinction between the polytypes can be observed in the collected spectrum in terms of energetic

position of the characteristics lines and the number of peaks.


87

Fig. 2.19 Typical photoluminescence spectra of near-band recombination for 4H-SiC (blue), 6H-

SiC (red) and 3C-SiC (black) polytype obtained at temperature of 4-5K

Donor Acceptor Pairs (DAP) recombination

When a semiconductor is simultaneously doped (intentionally or unintentionally) by both

donor and acceptor impurities, the electrons can be trapped forming neutral donors (D0) and, the

holes can be trapped forming neutral acceptors (A0). Most of the sites will form Donor-Acceptor

Pairs (DAP) excitons when the compensation (~similar donor and acceptor concentration) of the

impurity concentration is presented in the sample. The DAP recombination emits a photon with

the energy:

(2.16)

where is the band gap, and are the donor and acceptor binding energies, and the

last term is the Coulomb interaction energy between donor and acceptor ions separated by

distance . In SiC films, the DAP recombination leads to the presence of a broad band on the

LTPL spectra [16], [28].

380 400 420 520 540 560

380 381 382 383

412 413 414 415 416 417

3C-SiC6H-SiC

LO

TO

LA

TA

Inte

ns

ity

(a

.u)

Wavelength (nm)

ZPL

Ph

on

on

re

pli

ca

s (

I x a

nd

Px

)

4H-SiC

Q0

4Al0

P0Al0

S0R0

P0

3.3 3.2 3.1 3.0 2.9 2.35 2.30 2.25

Energy (eV)


88

Doping concentration evaluation

The intensity of BE recombination lines depends on the impurity concentration in the

matrix. Consequently, as already demonstrated in literature [29], calibration curves can be

proposed to allow the estimation of doping concentration from the intensity of related NBE

features. Ivanov et.al [30] have demonstrated a calibration curve for the low (non-intentional)

nitrogen doping 4H-SiC layers, taking into consideration a ratio related to the recombination of

exciton bound to nitrogen donors and the recombination of the FE as:

(2.17)

where the and represent the area of the BE and FE lines.

They demonstrated with a series of 4H-SiC and additionally with 6H-SiC (the FE= )

epilayers that the intensity ratio scales linearly in log-log scale, and can be used as reference to

which the emissions from impurities can be compared to determine the nitrogen concentration.

The FE lines are only visible when the nitrogen doping is below 2x1016

cm-3

. Hence, increasing

the nitrogen doping in the range 2x1016

cm-3

to ~1018

cm-3

(i.e. medium doped layers), the intensity

of line and its phonon replicas are decreasing with respect to the line intensity. Due to this

phenomenon, the nitrogen concentration determination was carried out by Henry et al. [31] from:

(2.18)

where and are the intensity corresponding to the and line from the measured

LTPL spectra. At doping concentrations higher than 1018

cm-3

, the intensity of the phonon replica

decreases until no trace of the line is visible on the spectra. In this case the nitrogen

concentration is estimated from the energy position of line, since the strong nitrogen presence

induces a shift towards lower energies.

In the case of aluminum doping (1015<[Al]<10

17cm

-3) in 4H-SiC layers, the concentration

estimation by LTPL involves the intensities ratios of the corresponding and lines [28]:


89

(2.19)

When the Al incorporation is above 1018

cm-3

, on the LTPL spectrum a Broad Band (BB)

is formed and the line is shifted towards lower energies. This effect can be correlated to the

band gap shrinkage that appears at high dopant concentration. For a correct estimation of the

dopant concentration, the compensation effect needs to be taken into consideration since DAP

transitions appear when Al and N concentrations are high enough [28]. It has been shown that Al

concentration can be calculated either from the Full Width at Half Maximum (FWHM) or from

the line shift [25].

Experimental set-up

The LTPL measurements were performed on nitrogen/ aluminum doped 4H-SiC layers

and 3C/6H-SiC heterostructure grown in our reactors. The excitation source was a frequency

double argon ion (FRED) laser with a 244nm wavelength and 30mW power. For the LTPL

spectrum acquisition, the samples were immersed under a gas Helium (He) flow in an Oxford

cryostat and cooled at 4K. The laser beam spot covers few hundred micrometers on the sample

surface, and penetrates to about few µm in 4H-SiC samples. A Jobin-Yvon Horiba TRIAX550

spectrometer with different gratings (e.g. 600, 1200 and 2400gr/mm) providing various spectra

resolutions was also used.

2.3.2.3. Raman scattering spectroscopy

Raman spectroscopy is a very effective, non-destructive tool for sample characterization,

especially since it doesn’t require any sample preparation. Raman technique is based on inelastic

scattering of monochromatic laser light, where the measured light after interaction is shifted in

frequencies that correspond to energy levels of different vibrations. The Raman scattering

involves the exciting vibrational modes: a diminished energy of photons scattering corresponding

to Stokes lines, or when the scatter photons gain energy they correspond to anti-Stokes lines.

The Raman scattering spectroscopy is usually used to characterize the quality of the

grown layers or substrates, the polytype identification and the estimation of the carrier

concentration. Raman spectrum is sensitive to the stress within the epitaxial layer.


90

Polytype identification

In this study, Raman technique was mainly used to identify the SiC polytypes within the

layers (see Annex I). Raman is a very useful and easy tool to determine the SiC polytypes. The

polytypes identification is actually based on the position lines of the phonon frequencies, such as,

Transverse Acoustic (FTA) and Optical (FTO) Folded modes. The periodic structure of each

polytypes corresponds to a given number of folded modes arriving from the Transverse Acoustic

(TA) and Transverse Optical (TO) branches. Therefore, the scattering allows the observation of

the phonons E1-giving the TO and A1-giving the Longitudinal Optical (LO) modes collected by

Raman bands. The Raman frequencies of the folded modes of typical polytypes are summarized

in the table 2.1

Polytype Frequency (cm-1

)

Planar acoustic Planar optic Axial acoustic Axial optic

x=q/qB FTA FTO FLA FLO

3C-SiC 0 - 796 - 972

4H-SiC

0 - 796 - 972

2/4 196, 204 776 - -

4/4 266 - 610 838

6H-SiC

0 - 797 - 965

2/6 145, 150 789 - -

4/6 236, 241 - 504, 514 889

6/6 266 767 - -

Table 2.1 Raman frequencies of the folded modes for the typical polytypes used along this study.

Only the FTA and FTO modes with E1/2 and A1 symmetry are shown [32].

A simple and more direct way is to compare the obtained spectrum (i.e. the frequency

position lines) with the already calculated spectrum for a given polytype. The Raman spectra

collected from 4H-SiC and 6H-SiC polytypes are shown in the fig. 2.20.


91

Fig 2.20 Raman spectra collected at 300 K for 4H-SiC and 6H-SiC polytypes

Furthermore, information regarding any SiC defects or stacking disorder can be acquired

from the Raman bands since a broadening of the peaks is induced. The exhibit broadening

attributed to defects is usually seen by the FTA modes. The stress and strain can also be

identified by the shift brought to the phonon modes lines.

Doping level evaluation

The room temperature Raman spectroscopy is a good candidate to evaluate the n- or p-

type carrier concentration in the SiC epilayers. The doping concentration can be estimated from

the shape of either, Longitudinal Optical Phonon Coupled (LOPC) mode which is given by

FLOx(0) line on the Raman spectra, or by the acoustic modes (FTA line). An evaluation of the

doping level by Raman spectroscopy was performed by P. Kwasnicki, and explained in his PhD

manuscript for both nitrogen and aluminum doped 4H-SiC epilayers grown in our CVD reactors

[25].

200 400 600 800 1000

6H-SiC

FTAx(1/3)

FLAx(2/3)

FTOx(1)

FLOx(2/3)

FTOx(1/3)

FLOx(0)

FTOx(1/2)

FLAx(1)

FTAx(1)

FTAx(1/2)

Inte

ns

ity

(a

.u)

Raman shift (cm-1)

T=300K

4H-SiC

FTOx(0)


92

Experimental set-up

The Raman measurements have been performed in order to determine the polytype of the

grown layers using a Jobin Yvon T64000 system. The excitation source was a He-Ne laser with a

532nm wavelength and 30mW power. The signal was collected by a CCD camera cooled with

liquid nitrogen. The measurements were done at room temperature in the range 100-1200cm-1

.

All Raman spectroscopy measurements have been performed at Laboratoire Charles Coulomb by

P. Kwasnicki.

2.3.3 Structural characterization: X-ray Diffraction

The XRD is an important characterization technique to determine the structural

information about the grown crystal (e.g. crystalline structure and quality, lattice parameters,

phase composition, strain and grain size). Diffraction occurs when an incident beam of

monochromatic x-ray of wavelength scatters a set of planes, and the scattering phenomena

produces constructive interference at specific angles. For a set of parallel planes, the diffraction

occurs if the incidence angle fulfills the specific conditions defined by Bragg’s law [33]:

(2.20)

where is an integer number, is the wavelength of the incident X-ray beam, is the

interplanar spacing of the crystal planes having Miller indices (hkl) [34].

Experimental details

The equipment involved in the SiC structural characterization is a PANalytical X’Pert Pro

MRD diffractometer available in the lab that uses a Cu X-ray source emitting at a wavelength

K 1 of 0.15406nm with a Ge(220) crystal monochromator. The XRD analyses of the epilayer

crystalline quality can become a routine investigation due to fast measurements and no sample

preparation.

An example of 2 - symmetrical scan recorded at different angles on 3C-SiC epilayer

grown on an on-axis 4H-SiC substrate, is represented in fig. 2.21 The 2 - scan displays the

presence of two peaks associated to: 3C-SiC111 reflections from the epilayer and to 4H-

SiC0001 reflections from the substrate, respectively. Moreover, the out of plane lattice


93

parameter can be calculated from the 2 - scan. The crystalline quality of the SiC material can

be obtained by performing an -scan (or “rocking-curve”).

Fig. 2.21 2 - symmetrical scan of a 3C-SiC epilayer grown on a well-oriented

4H-SiC substrate

A low value of FWHM indicates the presence of a good material quality while a crystal

structure distortion by defects is associated with high FWHM values (e.g. highly Al doped 4H-

SiC). To enhance the XRD measurement, a triple-axis analyzer can be used in order to obtain

higher resolution.

XRD study of 3C-SiC grown on -SiC substrates

In this work, X-ray diffraction was used not only to control the structural quality of the

layers but also as an effective tool for polytype identification, especially when 3C-SiC films were

grown on -SiC substrates. We can distinguish two cases. The first concerns the films composed

mainly of 3C polytype – in this case the symmetrical 2 - scans show clearly the 3C-SiC(111),

(222) and (333) peaks in the vicinity of 4H-SiC(004), (008) and (0012) peaks of the substrate (i.e.

fig. 2.21). The second case concerns the layers having sporadic presence of 3C-SiC as isolated

domains or local inclusions, for which the intensity of 3C-SiC(111), (222) and (333) reflections is

35 36 37 75 76 132 133 134

3C

-SiC

(1

11

)

Inte

ns

ity

(a

.u)

4H

-SiC

(00

4)

3C

-SiC

(2

22

)

2 (°)

4H

-SiC

(0

08

)

3C

-SiC

(3

33

)

4H

-SiC

(0

01

2)


94

too low to make them visible as independent peaks on the symmetrical 2 - scans. The

identification of 3C-SiC has to be based on the detection of an X-Ray peak specific for cubic

polytype. 3C-SiC113 reflections at 2 =71.85° are a potential candidate. Even if the

4H-SiC2 4 peaks have a similar Bragg angle, 2 =71.75°, the corresponding reflection is

observed at different angles than 3C-SiC113 so that no confusion is possible2. Thus, the

presence/absence of 3C-SiC(113) peak is a unequivocal indication of presence/absence of cubic

polytype in the layer.

A particular care has to be taken while analyzing the films grown on vicinal substrates.

Usually, the presence of sample off-cut is compensated by an adequate disorientation of the

sample holder before measurement. However, in the equipment used for this study, such

compensation was allowed only for the angles up to 4°. Since an important quantity of layers was

deposited on 8°off 4H-SiC substrates, the off-cut had to be compensated by adjusting and

angles. The adjustment depends on the azimuth angle between the off-cut axis and the

diffraction plane (=off-anglecos ; =-off-anglesin ). Considering that the SiC peaks are

narrow (FWHM<<0.1°), an imprecise choice of may prevent the 3C-SiC113 peak detection

and lead to wrong conclusions. To avoid such a misinterpretation, the compensation adjustment

has to be validated by successful detection of a substrate-related peak at same azimuth angle, e.g.

4H-SiC108.

Figure 2.22 shows the schematic representation of the crystal structure and its

corresponding lattice vectors for: (a) an 8° off-cut towards (11 0) 4H-SiC sample and (b-c) 3C-

SiC with double domains rotated by 60° around its symmetry axis. A simulation of the

corresponding hexagonal and cubic structures projected on the (0001) and respectively (111)

plane obtained with Carine software are represented also. The crystallographic representation of

the structures is constructed from Si-C tetrahedral, where the Si and C atoms are defined by blue

and red symbols. The red arrows indicate the off-cut direction of the vicinal substrate.

2 4H-SiC 2 - 4 can be observed at =0, 60, 120, 180, 240, 300°. 3C-SiC113 can be observed at =90, 210,

330° and 30, 150, 270°


95

Fig. 2.22 Schematic representation of the crystal structure and its corresponding lattice vectors

for: (a) 4H-SiC with 8° off-cut towards (11 0), (b) first 3C-SiC domain and (c) second 3C-SiC

domain rotated by 60° with respect to the first. Si-C tetrahedral (Si –blue and C- red symbols) are

shown in the projected crystal structure

This representation acts as a model to validate the planes and positions at which the

diffraction can occur for the two structures when a vicinal substrate is used during the growth

process. Another X-ray diffraction technique used during our study to analyze the properties of

the 3C-SiC films grown on -SiC substrates was the Reciprocal Space Mapping (RSM). This

involves a series of -2 scans performed at successive values, and they will be discussed

later in chapter 4.

2.3.4. Chemical characterization: Secondary Ion Mass Spectrometry (SIMS)

Secondary Ion Mass Spectrometry (SIMS) is a highly sensitive technique used to

determine the semiconductor material chemistry, i.e. the composition and doping incorporation of

the thin films [35].


96

The analyzed sample is placed under vacuum, and the sample surface is bombarded by a primary

ions beam (e.g. Cs+ O2

+, O

- or Ar

+ ions) with energies in the range of 5-17keV. The incident

beam has a 20-50µm diameter and can be scanned over a 100µm-1mm size rectangular raster.

The collisions of energetic primary ions with sample induce the surface sputtering, leading to

formation of rectangular craters. Sputtered matter is composed mainly from neutral particles, but

contains also ionized atoms or molecules (secondary ions) released from the material [36] as

displayed in fig. 2.23.

Fig. 2.23 Schematics of the material sputtering when the energetic primary beam is applied in

SIMS investigation [37]

The extraction of the secondary ions from the sample surface is performed by an applied

electric potential in the range 1-10kV. Accelerated secondary ions are afterwards subjected to

mass filtration in the mass spectrometer [38]. The mass filter includes electrostatic and magnetic

analyzer that allows appropriate selection of secondary ions with specific energy and momentum

(i.e. ions having given mass to charge ratio). The intensity of the detected beam is proportional to

the concentration of element in the sample and its sensitivity limit.

SIMS is a versatile technique that allows identifying the elements composing the material,

to determine the concentration depth profiles of different elements or the ion imaging of the

sample. It is important to mention that all the elements from the periodic table can be detected by

SIMS, but different sensitivity levels need to be applied. However, the ionization efficacy varies

depending on the element, resulting in different low detection limits.


97

In order to determine the element concentration from the intensity of related secondary

ions, one has to know the Relative Sensitivity Factor (RSF) of the ion in the analyzed matrix [39].

The concentration of the element can be expressed as:

(2.21)

where is the investigated ion intensity, and is the intensity of ion matrix chosen as reference.

The RSF value comes from the calibration process, i.e. from a measurement of the reference

sample of the same compound, having known the studied element concentration. The RSF value

depends, obviously, on the choice of reference and investigated ions, but also on the nature of

primary beam and secondary ions polarity. Moreover, it shows an environment dependence, i.e.

on the type of instrumentation used during the analysis. Although, the SIMS technique gives

competent results, the destructive preparation of the sample and the high cost of the analysis,

makes it usage difficult as a routine characterization tool.

Experimental investigation of 4H-SiC epilayers

The detection of N and Al dopants from the SiC matrix by SIMS has been performed on a

CAMECA IMS 5f spectrometer available at Laboratoire Charles Coulomb. During the analysis,

O2+

primary beam was accelerated at 15keV and positive secondary ions were collected: 27

Al+,

14N

+ for dopant,

28Si

++,

13C

+ and

12C2

+ were used as reference signal. The secondary beam voltage

was set at 4.5kV.

For this apparatus, the Detection Limit (DL) of nitrogen and aluminum were found to be

approximately: DL(N)=1x1017

cm-3

and DL(Al)=3x1013

cm-3

, respectively. The craters size was

150x150µm2 and, a Dektak profilometer was used to measure the depth. Typical sputtering rate

was 2-5nm/s. The RSFs values are given in table 2.2.

14N

+ /

12C2

27Al

+ /

12C2

RSF 1.01×1021

1.08×1018

Table 2.2 RSF values used in O2+ configuration for p-type 4H-SiC samples


98

Figure 2.24 shows an example of a Al (pink) and N (green) concentration profile as a

function of sputtering time, obtained on an intentionally aluminum doped 4H-SiC layer. A quite

constant Al concentration level of ~3x1018

cm-3

can be observed until 2000s which corresponds to

~10µm. After this depth, the concentration level changes corresponding to the interface between

the epilayer and substrate leading to the measurement of a residual Al concentration. The

substrate N concentration is ~5x1018

cm-3

while a residual N concentration coming from the SIMS

chamber was measured in the epilayer.

Fig. 2.24 Doping profile obtained by SIMS analysis on aluminium doped 4H-SiC

epilayers grown on homoepitaxially on an n+ type substrate

2.3.5 Electrical characterization

Secondary Ion Mass Spectrometry discussed in the previous paragraph supplies

information about principal dopants (N, Al) incorporation in the SiC matrix. However, SIMS

does not indicate whether the impurity is electrically active or not. Moreover, the high detection

limit of nitrogen (DL(N)~1e17cm-3

) does not allow an effective control of n-type layers for

electronic applications (n1e16cm-3

).

A routine tool to determine the net concentration of electrically active impurities (donors or

acceptors) in the epilayers is the Capacitance-Voltage (C-V) technique. Still, it does not indicate

the impurity ionization rate, i.e. the amount of free charge carriers. In order to evaluate the carrier


99

concentration, Hall effect measurements have to be performed. Both techniques are presented in

this section.

2.3.5.1 Capacitance-Voltage technique (C-V)

Principle of the method

The first investigation of doping impurities profile of a semiconductor by C-V

measurement was first described by Hillibrand and Gold in 1960 [40]. Typically, the C-V method

requires the presence of PN junctions or Schottky contacts for correct measurements. A faster and

non-destructive alternative is the use of a mercury (Hg) probe, where the contact is formed

locally by liquid Hg at the epilayer surface. Moreover, any residuals can be clean away after the

measurements, leaving the surface ready for additional characterization. The C-V measurements

determine the net doping concentration of the layers, i.e. the number of acceptors minus the

number of donors ( ) for p-type layers, or the number of donors minus the number of

acceptors ( ) for n-type layers that can be ionized. The C-V technique proves to be

capable to evaluate the doping profile up to a concentration of about 5x1019

cm-3

due to carrier

tunneling effect [41], [42].

The mercury probe C-V measurement requires a sample with sufficient size and low roughness to

cover the ring of the Hg contacts as shown in fig. 2.25(a). The contacts are made between the Hg

and the epilayer surface therefore, a flip of the wafer is necessary for a correct position on the

substrate holder of the equipment. Under an applied vacuum, the Hg from the two vials is coming

in contact with the epilayer surface forming two Schottky contacts with different diameters. The

first contact ring with larger diameter has a capacitance C1 while the second one presents a

smaller diameter (e.g. in our case 792µm) with a capacitance C2 as display in the fig. 2.25(b). As

C1 is noticeably larger than C2, the total capacitance is approximated to C2. The extraction of the

capacitance from the impedance measurement can be performed in series or parallel mode in case

of electrical leakage.


100

Fig. 2.25 (a) Diagram of the mercury probe C-V measurement and (b) equivalent circuit under

measurement between the sample and the mercury contacts [36], [43]

Considering a metal-semiconductor structure of n-type where the metal work function

is higher than the semiconductor work function ( ), or for the case of p-type-

, the electrons/holes from the semiconductor side will move towards the metal side leaving

behind positively charged fix donors ( ) or negatively charged fixed acceptors ( ). Under

thermodynamic equilibrium, a space charge region (scr) of width is created at the structure

interface giving rise to a potential barrier and, an electrical field [41], [44].

When a voltage is applied to the metal-semiconductor junction, a variation of the scr width

of the two contacts is obtained as represented in figure 2.25(b).

Therefore, the variation of the depletion region width with the applied voltage is expressed

as:

(2.22)

where is the vacuum permittivity ( =8.854x10-12

F/m), is the dielectric constant ( =9.6),

is the barrier potential, is the elementary charge ( =1.6022x10-19

C), - is the net

ionized dopant concentration, is the contact surface and is the capacitance (in F).

The differential capacitance of the structure considerate as a parallel-plate capacitor of width

is determined as follows [41]:


101

(2.23)

The doping density profile ( or ) for a constant doping can be obtained

as:

(2.24)

To obtain the net doping concentration, the slope of the curve can be used.

Characterization of the epilayers

In our study, the doping profile of 4H-SiC grown epilayers was evaluated by C-V

measurements using an Hg-probe with an Agilent LCR meter (4284A). The C-V analysis on 4H-

SiC epilayers was performed usually under an applied positive bias up to +40V for p-type layers,

or negative bias up to -40V, for n-type layers.

During C-V measurements, an additional alternative current (ac) bias is applied for which

the magnitude and frequency are fixed compared to the direct current (dc) bias which is varied.

Thus, the applied frequency (ac bias) provides the time necessary during the measurement for the

majority carriers to respond.

Figure 2.26(a) displays the capacitance as a function of forward bias at various

frequencies (from 20Hz to 1MHz), recorded on an non-intentionally doped 4H-SiC epilayer with

[Al]SIMS~2.3x1016

cm-3

. As can be seen, the capacitance is sensitive to frequency change, leading

to a decrease of capacitance value with increasing frequency. In fig. 2.26(b) the effective dopant

concentration profile derived from C-V measurements performed at different

frequencies and forward biases are illustrated. The depletion width varies over ~1µm under the

applied biases and frequencies.


102

Fig. 2.26(a) Capacitance-voltage measurements performed on a non-intentional doped

4H-SiC sample containing 2.3x1016

cm-3

Al atoms at different frequencies [20Hz-1MHz] and

(b) effective acceptor concentration profile at various frequencies.

0 5 10 15 20 25 30 35 40

5x10-11

1x10-10

2x10-10

2x10-10

Ca

pa

cit

an

ce

(F

)

Voltage (V)

20 Hz 100 Hz 100 Hz 1 kHz 10 kHz 20 kHz 50 kHz 100 kHz 1 MHz

Na-N

d=2.3x10

16 cm

-3

a)

0.0 0.4 0.8 1.2 1.6 2.010

15

1016

1017

1018

b)

Na-N

d (

cm

-3)

Depletion width W (µm)

20 Hz 100 Hz 100 Hz 1 kHz 10 kHz 20 kHz 50 kHz 100 kHz 1 MHz


103

However, at low frequencies (20-100Hz) the measurement is somehow compromised by

the carrier response (noise), while at high frequency, e.g.1MHz, an underestimated net dopant

concentration value was observed.

Reliable dopant concentration values can be achieved when the frequency is varied in the

range 100Hz–100kHz. An increase in dopant concentration results to a reduction of the depletion

width, i.e. the holes/electrons are favored to pass the interface by tunneling effect. This creates an

electrical leakage through the space charge region. For a given doping concentration, the

analyzed depth can be varied from~5μm (for films doped at ~1×1015

cm-3

) to ~0.1μm (for films

doped at ~5×1018

cm-3

).

Fig. 2.27 Wafer mapping of the net dopant concentration obtained on an Al p-type doped 4H-SiC

epilayer grown with a substrate rotation

The uniformity of the doping concentration is a necessary condition for fabrication of

reliable power devices. Therefore, the C-V technique can be used to determine the effective

dopant concentration along the sample diameter. Figure 2.27 shows the effective doping

uniformity map of an Al doped 4H-SiC epilayer. During growth, a substrate rotation was used.

The net doping concentration of the layer was determined at ~3x1019

cm-3

, and excluding 3mm

from the wafer edge, a standard variation over the mean value of 33.2% was found.

0

10

20

30

40

0

90

180

270

0

10

20

30

40

Po

sit

ion

(m

m)

5.0E18

1.1E19

1.8E19

2.4E19

3.1E19

3.7E19

4.4E19

5.0E19

cm-3


104

A comparison between SIMS and C-V technique was performed to determine the

accuracy in the measurements of Al dopant concentrations in 4H-SiC layers. A satisfactory

agreement between the two techniques was found as shown in fig. 2.28 for Al doped epilayers in

range ~1017

up to ~3x1018

cm-3

. The agreement between the two techniques was observed

previously in literature [45]. Moreover, it emphasizes the necessity of epitaxial layers with low

defect density and impurity compensation for a good characterization.

Fig. 2.28 Net dopant concentration determined by C-V measurements as a function of

the Al concentration measured by SIMS technique

2.3.5.2 Hall Effect measurements

Electrical properties of the grown epilayers can be acquired by Hall effect measurements.

The principle of the method involves placing the sample into a uniform magnetic field (B)

perpendicular to the electrical current (I), where the Lorentz force (F) will act on the carrier

charges as:

(2.25)

1017

1018

1019

1017

1018

1019

Na-N

d (

cm

-3)

Al concentration (cm-3)


105

where is the elementary charge, is the particle velocity and is the applied magnetic field.

The potential difference created by the electrical field is called the Hall voltage and can be

written as [34]:

(2.26)

where is the Hall coefficient, is the electrical current and is the epilayer thickness. The

Hall coefficient can be calculated by:

(2.27)

where is the Hall concentration for p-type materials; or is the electron concentration for n-

type materials and is the Hall scattering factor which is temperature dependent as shown by

Pensl et al. [46]

The Hall concentration can be determined as:

(2.28)

where is the drift carrier concentration (i.e. true hole density).

Moreover, the resistivity and mobility can also be derived from the Hall Effect

measurement using the given expressions:

(2.29)

(2.30)

Experimental details

The electrical properties which involve the free carriers concentration, mobility and

resistivity of the Al doped 8° off-axis 4H-SiC epilayers were characterized by Hall effect

technique at Laboratoire Charles Coulomb in Montpellier, but also at CNR-IMM in Catania.


106

The electrical isolation between the epilayer and the conductive 8°off-axis n+-type 4H-

SiC substrate was obtained by an epitaxial p-n junction. Ti/Al metal was evaporated in Van der

Pauw configuration (i.e. at the sample edges) on the epilayers surfaces to create the contacts. To

obtain the ohmic behavior, the contacts were annealed at temperature of 700°C under N2

atmosphere. The electrical connection for the measurements was done by gold wires realized

with silver epoxy. Hall measurements were performed in the temperature range of 300-850K

under a He atmosphere, while a platinum Pt resistor was used for temperature control. The

maximum applied magnetic field B was of 1Tesla.

The contacts ohmicity has been checked by Current-Voltage (I-V) characteristics, before

any Hall measurements. The resistivity measurement principle by Van der Pauw method involves

applying a current between two contacts (e.g. 1 and 2), followed by a voltage measurement at the

opposite contacts, as represented in fig. 2.29(a). The Van der Pauw voltages have been measured

for the four deposited contacts at positive current I+, negative I- and I=0.

Fig. 2.29 Schematic representation of the Van der Pauw configuration for measurement of (a)

resistivity and (b) Hall-effect

The contact resistance and epilayer resistivity can be calculated by the following expressions:

(2.31)

(2.32)


107

The Hall-effect measurement in Van der Pauw configuration consists of measuring the

Hall voltage VH at an opposite pair of contacts than the ones where the current and the

perpendicular magnetic field B are applied, as displayed in fig. 2.29(b). For our samples, the

average was obtained from the analysis of the contacts at both I+ and I-.

Figure 2.30 shows the Hall concentration derived from Hall effect measurements on

Al doped 4H-SiC epilayer with an [Al]=6.8x1018

cm-3

determined by SIMS, as a function of

temperature. As can be seen, the increases with the temperature. However, above 600K the

starts to increase up to ~2x1019

cm-3

which is over the concentration obtained by SIMS. Such

behavior may suggest a leakage current at the interface between the epilayer and the substrate or

due to the value used. More details about the Hall measurements on Al doped 4H-SiC layers

will be discussed further in Chapter 4. It is worth mentioning that the electrical properties of Al

doped 4H-SiC layers grown in our CVD reactors were investigated by Hall effect for the first

time.

Fig. 2.30 Hall concentration as a function of reciprocal temperature measured on a Al-doped 4H-

SiC layer at 6.8x1018

cm-3

1.0 1.5 2.0 2.5 3.0 3.510

17

1018

1019

1020

Ha

ll c

on

ce

ntr

ati

on

(c

m-3)

1000/T (K-1)

[Al]=6.8e18 cm-3

800 700 600 500 400 300

T (K)


108

2.4. Conclusions

The first part of this chapter was dedicated to the deposition system and growth details of

heteroepitaxial and homoepitaxial SiC epilayers. The boundary layer model proposed by

Eversteyn et al. [3] has showed that can describe in a qualitatively manner the growth kinetics

inside the CVD reactor. Moreover, the deposition rate was found to be influenced by various

process parameters such as growth temperature, silane flow rate, reactor pressure and C/Si ratio.

In the high temperature range, an important impact of H2 etching on growth rate was registered.

Moreover, due to the dilution lines installed, the dopant source can be widely varied, resulting in

a tuning of p-type doping in the range 1016

cm-3

to 1020

cm-3

and of n-type in the range 1014

cm-3

to

1019

cm-3

in 4H-SiC epilayers.

The second part of the chapter was assign to the various characterization techniques

employed to analyze different features or properties of the grown layers. An elaborate study

regarding the use of XRD technique as an effective tool for polytype identification even for the

case of off-angles substrates was presented. For fast polytype identification, the LTPL and

Raman techniques have be used to determine the presence of 3C-SiC polytype on -SiC

substrates. The dopant concentration of the layers was determined by Hg probe C-V

measurements and SIMS analyses. Due to the limitation of dopant detection (especially for [N]

where DL(N)=1x1017

cm-3

), and the high-cost of the measurement, SIMS technique remains

usually the second option. On the other hand, the mercury probe C-V technique is a fast and easy

method to determine the net doping concentration ( - or - ) of the layers up to

~5x1019

cm-3

. The electrical properties of Al doped 4H-SiC layers have been investigated for the

first time by Hall Effect measurements. It allowed gaining valuable information regarding the

Hall holes concentration and their mobility.

2.5. Bibliography

[1] A. Leycuras, “Réacteur pour dépot chimique en phase vapeur et procédé pour sa mise en

oeuvre, brevet française (1998),” no. FR2 786 203, 1998

[2] U. Forsberg, “CVD Growth of Silicon Carbide for High Frequency Applications,” Linköping

University, Department of Physics and Measurement Technology, 2001

[3] F.C. Eversteyn, P.J.W. Severin, C.H.J. v.d. Brekel, and H.L. Peek, “A Stagnant Layer Model

for the Epitaxial Growth of Silicon from Silane in a Horizontal Reactor,” J. Electrochem.

Soc., vol. 117, p. 925, 1970


109

[4] T. Kimoto, H. Nishino, W. S. Yoo, and H. Matsunami, “Growth mechanism of 6H-SiC in

step-controlled epitaxy,” Journal of Applied Physics, vol. 73, no. 2, p. 726, 1993

[5] K.L. Choy, “Chemical vapour deposition of coatings,” Progress in Materials Science, vol.

48, p. 57, 2003

[6] O. Danielsson, “Simulations of silicon carbide chemical vapor deposition,” Dissertation no.

773, Linköping University, Sweden, 2002

[7] M. Ohring, “Chapter 6 -Chemical vapor deposition,” in Materials Science of Thin Films

(Second Edition), M. Ohring, Ed. San Diego: Academic Press, p. 277–355, 2002

[8] T. Kimoto, “Step-controlled epitaxial growth of alpha-SiC and device applications,” Kyoto

University, Japan, 1995

[9] P. Aboughe-nze, “Epitaxie du carbure de silicium cubique. Utilisation de

l’hexamethydisilane.Realisation de pseudo substrats pour la croissance de GaN cubique”,

Université Claude Bernard-Lyon1, 2001

[10] M. Kumagawa, H. Kuwabara, and S. Yamada, “Hydrogen etching of Silicon carbide,”

Japanese Journal of Applied Physics, vol. 8, no. 4, p. 421, 1969



Growth, vol. 267, no. 3–4, p. 436–451, 2004

[12] T. Tanaka, N. Kawabata, Y. Mitani, M. Sakai, N. Tomita, M. Tarutani, T. Kuroiwa, Y.

Toyoda, M.Imaizumi, H. Sumitani, and S. Yamakawa, “Influence of growth pressure and

addition of HCl gas on growth rate of 4H-SiC epitaxy,” Materials Science Forum, vol. 821–

823, p. 133, 2015

[13] K. Kojima, S. Kuroda, H. Okumura, and K. Arai, “Nitrogen incorporation characteristics

on a 4H-SiC epitaxial layer,” Applied Physics Letters, vol. 88, no. 2, p. 21907, 2006

[14] W. Chen, and M.A. Capano, “Growth and characterization of 4H-SiC epilayers with

different off-cut angles,” Journal of Applied Physics, vol. 98, no. 11, p. 114907, 2005

[15] Ö. Danielsson, A. Henry, and E. Janzén, “Growth rate predictions of chemical vapor

deposited silicon carbide epitaxial layers,” Journal of Crystal Growth, vol. 243, no. 1, p.

170–184, 2002

[16] C. Sartel, “Homoépitaxie du SiC-4H a partir de différents précurseurs. Realisation du

dopage p,” Université Claude Bernard-Lyon1, 2003

[17] S. Ji, K. Kojima, Y. Ishida, S. Saito, S. Yoshida, H. Tsuchida, and H. Okumura,

“Suppressing Al memory effect on CVD growth of 4H-SiC epilayers by adding hydrogen

chloride gas,” Japanese Journal of Applied Physics, vol. 53, no. 4S, p. 04EP07, 2014

[18] M. Portail, M. Zielinski, T. Chassagne, H. Chauveau, S. Roy, and P. De Mierry, “Highly

sensitive determination of N+ doping level in 3C–SiC and GaN epilayers by Fourier

transform infrared spectroscopy,” Materials Science and Engineering: B, vol. 165, no. 1–2, p.

42–46, 2009

[19] Ö. Danielsson, U. Forsberg, and E. Janzén, “Predicted nitrogen doping concentrations in

silicon carbide epitaxial layers grown by hot-wall chemical vapor deposition,” Journal of

Crystal Growth, vol. 250, no. 3–4, pp. 471–478, 2003

[20] C.R. Brundle, C. A. Evan Jr., and S. Wilson, Encyclopedia of Materials characterization,

Manning Publication Co. 1992


the C-face of 6H–SiC(0001) substrates,” Journal of Crystal Growth, vol. 348, no. 1, p. 91–

96, 2012


110

[22] M. F. Macmillan, A. Henry, and E. Janzén, “Thickness determination of low doped SiC

epi-films on highly doped SiC substrates,” Journal of Elec Materi, vol. 27, no. 4, p. 300–303,

1998

[23] R. T. Holm, P. H. Klein, and P. E. R. N. Jr, “Infrared reflectance evaluation of chemically

vapor deposited β‐SiC films grown on Si substrates,” Journal of Applied Physics, vol. 60, no.

4, p. 1479–1485, 1986

[24] J. Camassel, J. Pernot, H. Y. Wang, and H. Peyre, “Modeling the infrared reflectance of

n–/n+ SiC layers on top of n+ SiC substrates for epitaxy control application,” physica status

solidi (a), vol. 195, no. 1, p. 38–43, 2003

[25] P. Kwasnicki, “Evaluation of doping in 4H-SiC by optical spectroscopies,” Université

Montpellier II Sciences et Techniques du Languedoc, 2014

[26] W.J. Choyke and R.P. Devaty, “Optical Properties of SiC: 1997-2002,” Silicon Carbide

Recent Major advances, Springer-Verlag Berlin, p. 413–433, 2004

[27] L. L. Clemen, R. P. Devaty, M. F. MacMillan, M. Yoganathan, W. J. Choyke, D. J.

Larkin, J. A. Powell, J. A. Edmond, and H. S. Kong, “Aluminum acceptor four particle bound

exciton complex in 4H, 6H, and 3C SiC,” Applied Physics Letters, vol. 62, no. 23, p. 2953–

2955, 1993

[28] Balloud Carole, “Méthodes de caractérisation optique de SiC,” Université Montpellier II

Sciences et Techniques du Languedoc, 2005

[29] J. Camassel, S. Juillaguet, M. Zielinski, and C. Balloud, “Application of LTPL

Investigation Methods to CVD-Grown SiC,” Chemical Vapor Deposition, vol. 12, no. 8–9, p.

549–556, 2006

[30] I. G. Ivanov, C. Hallin, A. Henry, O. Kordina, and E. Janzén, “Nitrogen doping

concentration as determined by photoluminescence in 4H–and 6H–SiC,” Journal of applied

physics, vol. 80, no. 6, p. 3504–3508, 1996

[31] A. Henry, O. Kordina, C. Hallin, C. Hemmingsson, and E. Janzén, “Photoluminescence

determination of the nitrogen doping concentration in 6H‐SiC,” Applied Physics Letters, vol.

65, no. 19, p. 2457–2459, 1994

[32] S. Nakashima and H. Harima, “Raman Investigation of SiC polytypes,” Physica status

solidi (a), vol. 162, p. 39, 1997

[33] R.J.D, Tilley, Crystals and Crystal Structures. John Wiley & Sons, Ltd, 2006

[34] C. Kittel, Introduction to Solid State Physics, Seventh edition. John Wiley & Sons, Inc,

1996

[35] H. W. Werner and H. A. M. De Grefte, “Investigation of surface layers by SIMS and

SIIMS,” Surface Science, vol. 35, p. 458–472, 1973

[36] A. E. Bazin, “Conception de diodes Schottky sur 3C-SiC épitaxie sur silicium,”

Université François-Rabelais de Tours, 2009

[37] http://www.nt.ntnu.no/users/floban/KJ%20%203055/Surface/SIMS.htm

[38] http://www.cameca.fr/instruments-for-research/sims.aspx

[39] R. G. Wilson and J. M. Zavada, “Secondary ion mass spectrometry of dopant and

impurity elements in wide bandgap semiconductors,” Materials Science and Engineering: R:

Reports, vol. 73, no. 11–12, p. 101–128, 2012

[40] Hilibrand J. and Gold R.D., “Determination of the impurity distribution in junction diodes

from capacitance-voltage measurements,” RCA Rev, vol. 21, p. 242–252, 1960

[41] Schroder D.K., Semiconductor material and device characterization, 3d ed. John Wiley &

Sons, Inc, 2006


111

[42] C.C. Hu, Modern Semiconductor Devices for Integrated Circuits. New Jersey:

Pearson/Prentice Hall, 2010

[43]https://www.google.fr/url?sa=t&rct=j&q=&esrc=s&source=web&cd=1&cad=rja&uact=8&v

ed=0CCgQFjAAahUKEwig0q64kfTGAhVkcNsKHe73A40&url=http%3A%2F%2Fgato-

docs.its.txstate.edu%2Fcos-department-of-

physics%2Fmanuals%2FCVanalyzer%2FCVanalyzer.pdf&ei=5WCyVeDLNeTg7Qbu74_oC

A&usg=AFQjCNF1Is32zx0_84Y37RQCMhZFZne11w&bvm=bv.98717601,d.ZGU

[44] Mishra U.K and Singh J., Semiconductor device physics and design. Springer, 2008

[45] W.J. Choyke, H. Matsunami, and G. Pensl, Silicon Carbide Recent Major Advances.

Springer-Verlag Berlin, 2004

[32] G. Pensl, F. Schmid, F. Ciobanu, M. Laube, S.A. Reshanov, N. Schulze, K. Semmelroth,

H. Nagasawa, A. Schoner and G. Wagner, “Electrical and optical characterization of SiC,”

Material Science Forum, vol. 433-436, p. 365, 2003

Chapter 3 Dopant incorporation in SiC epitaxial layers

112

Contents

Chapter 3 Dopant incorporation in SiC epitaxial layers: the case of nitrogen and aluminum .......................... 113

3.1. Introduction ....................................................................................................................................................... 113

3.2. Evolution of dopant incorporation ................................................................................................................... 114

3.2.1 Influence of polytype and off-angle ............................................................................................................. 114

3.2.1.1 Nitrogen case ........................................................................................................................................ 115

3.2.1.2 Aluminum case ..................................................................................................................................... 118

3.2.2 Influence of crystal orientation ..................................................................................................................... 121

3.2.2.1 Nitrogen case ........................................................................................................................................ 122

3.2.2.2 Aluminum case ..................................................................................................................................... 123

3.3 Dopant incorporation trends ............................................................................................................................. 125

3.3.1 Growth temperature dependence .................................................................................................................. 125

3.3.1.1 Nitrogen case ........................................................................................................................................ 125

3.3.1.2 Aluminum case ..................................................................................................................................... 127

3.3.1.3 Discussions ........................................................................................................................................... 129

3.3.2 Dopant supply dependence ........................................................................................................................... 131

3.3.2.1 Nitrogen flow rate ................................................................................................................................. 131

3.3.2.2 TMA flow rate ...................................................................................................................................... 132

3.3.2.3 Discussions ........................................................................................................................................... 133

3.3.3 Growth rate dependence ............................................................................................................................... 135

3.3.3.1 Nitrogen case ........................................................................................................................................ 135

3.3.3.2 Aluminum case ..................................................................................................................................... 136

3.3.3.3 Discussions ........................................................................................................................................... 137

3.3.4 Reactor pressure dependence ....................................................................................................................... 138

3.3.4.1 Nitrogen case ........................................................................................................................................ 138

3.3.4.2 Aluminum case ..................................................................................................................................... 139

3.3.4.3 Discussions ........................................................................................................................................... 140

3.3.5 C/Si ratio dependence ................................................................................................................................... 141

3.3.5.1 Nitrogen case ........................................................................................................................................ 142

3.3.5.2 Aluminum case ..................................................................................................................................... 144

3.3.5.3 Discussions ........................................................................................................................................... 147

3.4 Conclusions ......................................................................................................................................................... 150

3.5 Bibliography ........................................................................................................................................................ 152


113

Chapter 3 Dopant incorporation in SiC epitaxial layers: the case of nitrogen

and aluminum

3.1. Introduction

A key issue in production of high quality epitaxial layers for different electronic devices

remains the effective control of the voluntary dopant incorporation and the reduction of residual

impurity incorporation.

The latter was demonstrated by Larkin et al. [1], [2] as previously mentioned in Chapter

1, who identified the “site competition” mechanism. Larkin demonstrated that changing the

proportions between carbon and silicon precursors (C/Si ratio) while keeping constant all the

other parameters during the growth, allows a control of impurity incorporation over a wide

range. Moreover, it is well established that the incorporation of dopants in the SiC matrix

depends on the atomic size thus, nitrogen and aluminum atoms will preferentially incorporate on,

respectively, C and Si sites [3]. The site preference of the dopant in the SiC matrix has a direct

consequence on: (i) the site competition effect between dopant and the related matrix atom and

(ii) the substrate crystalline orientation since the dopant incorporation depends strongly on

substrate polarity.

Larkin’s finding proved that C/Si ratio is the process parameter for an effective doping

control, and it became a catalyst for subsequent investigations of doping mechanism. Various

groups intensively studied the doping mechanism especially for n-type material [4]–[10] widely

used in fabrication of unipolar devices. For p-type doping with aluminum, the studies are less

abundant [11]–[15]. Furthermore, experimental and theoretical studies have been extended also

for phosphorus (n-type) and boron (p-type) doping [11], [16]–[18]. Large range of doping

concentration (1014

-1020

cm-3

) can be achieved, as the dopant incorporation depends on process

parameters, crystal orientation and reactor geometry. The visible disparities between the results

of certain studies of dopant incorporation mechanisms have usually been attributed to differences

in reactor geometry.

In the present chapter, the discussion will be focused on the nitrogen and aluminum

incorporation in SiC layers. The trends of voluntary N and Al incorporation on SiC layers were

experimentally investigated. The influence of polytype and crystal orientation of the SiC

substrates on N and Al incorporation will be firstly discussed. Furthermore, the role of the


114

deposition parameters such as growth temperature, dopant flow rate, growth rate, growth

pressure and C/Si ratio in dopant incorporation process will be shown and discussed. A particular

attention was allocated to the role of C/Si ratio and related site competition effects.

3.2. Evolution of dopant incorporation

In order to obtain exhaustive information regarding the trends of N and Al incorporation in

the SiC layers, the growth experiments were performed in two reactors of different sizes, on

substrates having diverse nature /orientation, and explored a wide range of growth conditions. To

evaluate the role of each process parameter, we used two types of samples: multilayer stacks

grown for SIMS analyses (all process parameters except one are kept constant during growth),

and series of individual layers prepared for C-V measurements.

The substrates used for homoepitaxial growth were mainly Si-face and C-face off-axis (2°, 4°

and 8°) commercial 4H-SiC wafers (or pieces of wafers) planarized by NOVASiC. Some of the

growth runs were performed on on-axis 6H-SiC wafers. Heteroepitaxial 3C-SiC films grown on

Si(100) and Si(111) on-axis and 4° off-axis substrates were also under investigation. Some films

were deposited simultaneously on substrates of various nature (e.g. Si- and C-face 4H-SiC)

The Al incorporation in the grown films was determined by SIMS and C-V analyses,

techniques described previously in Chapter 2. Since the results of both measurement techniques

give a satisfactory agreement, it is here-after referred as Al incorporation [Al]. On the other

hand, the N concentration of the epilayers was determined only by C-V measurements and will

be referred herein as N incorporation [N]. It has to be mentioned that the [N] study was

performed only in R2 reactor (100mm), while the [Al] was initially study in R1 reactor (50mm)

and after the process transfer, in R2 reactor.

3.2.1 Influence of polytype and off-angle

For this work, substrates of different nature were used as the common polytypes of SiC

(4H-SiC, 6H-SiC and 15R-SiC), and substrates with various off-angles (on-axis, 2°, 4° and 8°

off). This study seemed appropriate to verify whether these parameters may lead to any influence

of the dopant incorporation. The literature results report that at given growth conditions, the

dopant incorporation (N or Al) does not differ significantly between polytypes or off-angles [7],


115

[8], [14], [19]–[21]. Forsberg et al. [14] proposed that Al will incorporate on the sites available

on Si-face terraces since Al atoms are relative immobile. However, other contributions suggested

that nitrogen and aluminum atoms should preferentially incorporate at steps sites [20]. In such

case, an increase of the substrate off-cut angle results in higher step density per unit area and

shorter terrace width [21], [22], which should lead to stronger dopant incorporation. Moreover,

we may also suppose that the difference in step height between various SiC polytypes should

influence the dopant incorporation [23], [24].

3.2.1.1 Nitrogen case

The influence of the SiC polytypes on dopant concentration in our setup was studied by

performing several growth experiments on a multi-polytype (4H-/6H-/15R-) SiC 4°off-angle

substrate. Before the epitaxial process, the sample was characterized by Raman spectroscopy to

identify the polytype from each region of the epiwafer. Therefore, from the recorded Raman

spectra collected at room temperature in the range 750-800cm-1

, the polytypes were identified

from the shape and position of the TO peaks as represented in fig. 3.1.

Fig. 3.1 Raman spectra collected in different regions of the multi-polytype (4H-/6H-

/15R-) SiC epiwafer


116

The study was carried out during one growth run under constant process conditions, i.e.

growth temperature of 1813°C, reactor pressure of 100mbar, C/Si ratio=1.2 given by 12sccm

C3H8 flow and 30sccm SiH4 flow rate. The N2 flow rate was set at 0.1sccm, and a growth rate of

~20µm/h was obtained. A substrate rotation was used during growth. Figure 3.2 shows the

normalized net donor concentration [N] obtained on 6H-SiC and 15R-SiC regions as a function

of the [N] obtained in 4H-SiC region. It can be observed that at the given process conditions, the

[N] does not differ significantly between the studied polytypes. This appears to be in accordance

with literature results [7], [8], and, thus we may consider similar incorporation tendencies for all

the SiC polytypes.

Fig. 3.2 Net donor concentration [N] in 6H- (blue) and 15R- (red) SiC regions reported to

doping of 4H-SiC region

The results were obtained from the doping profile of each polytype obtained by C-V

measurements along the wafer diameter (perpendicular to wafer primary flat) as shown in fig.

3.3. As a result, the points obtained at the same distance from the wafer center were compared,

from where more data points for 6H-SiC than 15R-SiC polytype.

1014

1015

1016

0.1

1

10

[N] 6

H,1

5R/[

N] 4

H

[N]4H (cm-3

)

15R-SiC

6H-SiC


117

Fig. 3.3 Net donor concentration [N] obtained along the wafer diameter for 4H-SiC, 6H-

SiC and 15R-SiC polytypes

The study to determine whether the off-cut angles have any influence on the N

incorporation was carried out on Si-face 4H-SiC substrates with various off-cuts (2°, 4° and 8°)

during several growth runs. Thereafter, the growth parameters were: =1770°C, C/Si ratio of

1.5, and a reactor pressure of 150mbar. The obtained growth rate was ~11µm/h. A slightly

different N2 flow rate was applied during the growth runs therefore the presented results are

normalized with respect to N2 flow. Figure 3.4(a) represents the normalized [N] with the N2 flow

rate profile obtained along the wafer axis (perpendicular to the primary flat of the wafer) on

4H-SiC films with various off-cut angles. In the studied range, the off-cut angle does not

influence the N incorporation in 4H-SiC layers.

-40 -20 0 20 4010

14

1015

1016

[N] (c

m-3)

Position along wafer (mm)

4H-SiC 6H-SiC 15R-SiC


118

Fig. 3.4 (a) Normalized donor concentration with N2 flow rate in 4H-SiC films with

various off-cut angles and (b) [N] in 0° 6H- (dark cyan) and 4° off-angle 4H-SiC (purple) versus

growth pressure

Additionally, the [N] dependence was investigated for two different polytypes (4H- and

6H-SiC) with various off-cut angles (4° and 0°, respectively) under constant process parameters.

The growth was performed at growth temperature of 1780°C, and a C/Si ratio of 1.5. The growth

pressure was varied in the range 100-250mbar, and a growth rate of ~7µm/h was achieved. The

results are displayed in fig. 3.4(b). Even if we may suppose that the growth mechanism is

different between off-axis and on-axis film (step flow versus nucleation on terraces) no

considerable [N] difference between the 6H-SiC 0° and 4H-SiC 4° off was seen, which

corroborates well with the results presented in fig. 3.2 and 3.4(a).

3.2.1.2 Aluminum case

The influence of the SiC polytypes on Al incorporation in our setup was also studied on a

multi-polytype (i.e. with 4H-/6H-/15R-) 4°off-angle SiC substrate. As represented in fig. 3.1, the

different polytype regions presented on the epiwafer were identified by Raman spectroscopy.

The study was performed within several growth runs under various conditions. The constant

102

103

1015

1016

1017

4° 4H-SiC

0° 6H-SiC

[N]/

N2

(c

m-3

/sc

cm

)

pgrowth (mbar)b)-40 -20 0 20 40

1015

1016

1017

[N]/

N2 (

cm

-3/s

cc

m)

Position (mm)

2° 4H-SiC

4° 4H-SiC

8° 4H-SiC

a)


119

process conditions between the growth runs where =1820°C and a C/Si ratio of 1.8. The

obtained growth rate was ~20µm/h. The growth pressure and TMA flow rate were varied

in the range 95-150mbar and 8.7x10-4

to 5.6x10-3

sccm, respectively.

The results are shown in fig. 3.5 where the Al incorporation in the 6H-SiC and 15R-SiC

regions are reported to the [Al] in the 4H-SiC regions in the same manner as explained for

nitrogen. The investigation reveals that, likewise for nitrogen, the [Al] is similar in 6H-SiC and

15R-SiC regions, and follows closely the values measured in 4H-SiC regions. The results

obtained in this study are consistent with the results reported previously in literature [14], [19],

which indicate no polytype influence on Al incorporation.

Fig. 3.5 Al incorporation in 6H- (blue) and 15R- (red) SiC regions reported to the doping

of 4H-SiC region

It would be interesting to include the cubic polytype, 3C-SiC, in this comparison.

Unfortunately, the heteroepitaxial growth of 3C-SiC on Si substrates and the homoepitaxial

growth of 4H-SiC layers are performed at different temperatures and it is impossible to directly

compare dopant incorporation under same growth conditions. The Al incorporation in 3C-SiC

films grown on Si(111) oriented substrates and on Si-face 4H-SiC films grown under different

growth temperatures are represented in fig. 3.6. The layers were grown at growth temperatures of

1015

1016

1017

0.1

1

10

6H-SiC

15R-SiC

[Al]6

H,1

5R/[

Al]4

H

[Al]4H (cm-3

)


120

1370°C and 1600°C, and a C/Si ratio of 0.8 and 1, respectively. The reactor pressure 200mbar

and a growth rate of ~8µm/h was obtained on both polytypes. It can be observed a comparable

[Al] between the two polytypes 4H-SiC(0001) and 3C-SiC(111). This observation can be

associated to the bonding of Al atoms on silicon sites in the SiC matrix and thus revealing a

stronger bonding on the (111) surface. Such a result leads us to the conclusion that 3C-SiC(111)

surface is equivalent to Si-face (rather than C-face) of 4H-SiC.

Fig. 3.6 Al incorporation in 4H-SiC(0001) and 3C-SiC(111) layers grown under various growth

temperature

The influence of 4H-SiC substrate off-cut angle on Al incorporation was investigated by

performing a simultaneous growth run on pieces of Si-face 4H-SiC substrates with 2°, 4° and 8°

off-cut angles. The growth conditions were: =1700°C, growth pressure 150mbar, and a C/Si

ratio=1.8. Figure 3.7 display the [Al] as a function of 4H-SiC substrate off-cut angle. The results

show that at the given growth conditions, the off-cut angle do not influence the incorporation of

Al in 4H-SiC films. Yamamoto et al. [20] reported similar results when they investigated the

[Al] dependence on C/Si ratio for 4° and 8° off-angle 4H-SiC substrates.

10-4

10-3

10-2

10-1

100

1017

1018

1019

1020

1021

[Al]

(c

m-3

)

TMA flow rate (sccm)

4H-SiC(0001)

3C-SiC(111)


121

Fig. 3.7 Al incorporation in 4H-SiC films with various off-cut angles

3.2.2 Influence of crystal orientation

In literature studies [6]–[8], [11], [12], [14], [19], [25]–[27] it is stated that the N

incorporation is systematically higher on C-face than on Si-face of -SiC substrates, while the Al

incorporation shows an opposite tendency (i.e. lower incorporation on C-face than on Si-face).

This lattice polarity dependence is attributed to the different bonding of impurity on the surface.

Considering that the N and Al atoms incorporate respectively on carbon and silicon sites in the

SiC matrix, N on C-face (and Al on Si-face) will present a stronger, three-fold, bonding than N

on Si-face (and Al on C-face) having only one bond with the surface. Stronger bonding reduces

desorption probability and leads to stronger incorporation of impurity.

This is schematically represented in fig. 3.8, which shows the projection of the SiC lattice

along (11 0) direction with the substitution of the impurity atoms in the SiC matrix on Si-face

and C-face.

2 4 6 810

17

1018

1019

[Al]

(c

m-3)

off-cut (°)

4H-SiC


122

Fig. 3.8 Schematic representation of the SiC matrix with the two surface polarities: (a)

Si-face and (b) C-face when the atoms are substitute by impurity atoms. Si and C atoms are

represented by red and black symbols, while Al and N atoms by violet and blue symbols.


The surface polarity influence of -SiC substrates on N concentration was also observed

in our experiments. The results displayed in fig. 3.9, obtained on simultaneously grown Si- and

C-face on-axis 6H-SiC films during various growth runs, are a typical example. The growth

conditions for the series of layers include a growth temperature 1780°C and a C/Si ratio equal to

1.5. A growth rate of ~6µm/h was attained. It can be seen from the graph, that the [N] is around

one order of magnitude higher on C-face [NC] compared to Si-face [NSi]. Some previous studies

reported for both 6H-SiC and 4H-SiC layers, a similar incorporation ratio [NC]/ [NSi] [28], [29].

Other groups have found a smaller difference in N incorporation between the two faces [7], [8],

[13], [30] compared to ours, which may be related to the growth conditions (i.e. C/Si ratio).


123

Fig. 3.9 Normalized N incorporation with the N2 flow rate measured on simultaneously

grown Si- and C-face on-axis 6H-SiC layers

In the case of 3C-SiC layers grown on Si substrates with (111) or (100) orientation it was

reported by Zielinski et al. [10] a higher N incorporation on (100) than on (111) surface. Using

the argument regarding the N atoms bond formation, weaker bonding is created on (111) surface

which results in stronger N desorption. This observation reintroduce the conclusion that 3C-

SiC(111) oriented surface is equivalent to Si-face of -SiC polytype.


The polarity dependence of Al incorporation was investigated through simultaneous

growth on Si- and C-face 8° off 4H-SiC substrates. The measured SIMS profiles are represented

in fig. 3.10(a). In the SIMS profiles the visible steps are the result of multiple layers grown on

the same sample by varying one growth parameter (in this case the TMA flow rate).

The deposition conditions of the multilayer samples were: growth temperature 1600°C,

reactor pressure 200mbar and a C/Si ratio~1, while a growth rate of ~8µm/h was obtained. As

expected, we observe stronger Al incorporation on Si-face. Al incorporation on Si-face [AlSi] is

~2 orders of magnitude higher than on C-face [AlC]. Similar results were described previously in

literature [13], [14], [19], [30]. However, the literature results reported a smaller incorporation

102

103

1015

1016

1017

1018

Si-face

[N]/

N2 (

cm

-3/s

cc

m)

pgrowth

(mbar)

C-face


124

ratio [AlSi]/[AlC]. Therefore, we may speculate that the [AlSi]/[AlC] ratio difference between our

results and literature are mainly related to particular growth conditions (e.g. C/Si ratio) and/or

reactor configuration.

Fig. 3.10 Examples of Al concentration depth profiles measured by SIMS on multilayer

stacks grown on: (a) Si- and C-face 4H-SiC substrates and (b) 3C-SiC(100) and(111) oriented

layers

0 2 4 6 8 10 1210

14

1015

1016

1017

1018

1019

1020

4H-SiC C-face 4H-SiC Si-face

[Al] (

cm

-3)

depth (µm)a)

0 2 4 6 8 10 12 14 16 1810

15

1016

1017

1018

1019

1020

1021

[Al]

(c

m-3

)

depth (µm)

3C-SiC (100) oriented 3C-SiC (111) oriented

b)


125

An additional study was performed on 3C-SiC layers grown simultaneously on 3C-

SiC(100) and (111) oriented template substrates. The multilayer stacks were deposited under

following conditions: =1370°C, reactor pressure of 200mbar, C/Si ratio ~0.8 and a growth rate

of ~8µm/h. The obtained SIMS profiles measured on both 3C-SiC(100) and (111) oriented layers

are displayed in fig. 3.10(b). The examination reveals ~20 times higher Al incorporation on 3C-

SiC(111) oriented layer [Al(111)] compared to 3C-SiC(100) oriented layer [Al(100)]. The visible

difference of [Al] is associated to the bonding configuration on both surface orientations.

Consequently, the Al atoms are bond to the (100) oriented surface only by two bonds [4], while

on (111) oriented surface due to stronger [Al] it may suggest the presence of three available

bonds. This indicates that the (111) oriented surface is equivalent to Si-face in 4H-SiC. In

literature it was reported only the study of Al incorporation on 3C-SiC(100) oriented layers [31].

3.3 Dopant incorporation trends

In this section, a detailed discussion about the influence of process parameters on dopant

incorporation will be presented. Major part of the observed tendencies (except temperature

dependence) can be fitted using a power function. Knowing the value of the function exponent

(slope in log-log representation), we can compare different datasets. Sometimes, the exponent

value gives us an indication on the dopant incorporation mechanism.

3.3.1 Growth temperature dependence


The temperature dependence of N incorporation was studied on both Si- and C-face 8°-

off-axis 4H-SiC substrates in R2 reactor. Various growth runs were performed for this study. In

order to investigate the influence of growth temperature on N incorporation, the temperature was

varied in the range 1400-1800°C, while all the other process parameters were maintained

constant. The growth conditions were set to: reactor pressure 400mbar, C/Si ratio equal to 2 and

a growth rate of ~8µm/h. The N2 flow rate was set slightly different for both polarities, and

therefore, the results will be presented under normalized value (i.e. [N]/ ). The experimental

results are represented under an Arrhenius plot in fig. 3.11. We found higher N incorporation on

C-face compared to Si-face, and their incorporation ratio [NC]/[NSi] varies with growth


126

temperature. Furthermore, the two polarities are showing an opposite dependence of [N] with

growth temperature.

Fig. 3.11 The normalized N incorporation as a function of growth temperature for Si- (blue

symbols) and C-face (red symbols) 4H-SiC layers. The line is used as guidance for the eyes.

Two N incorporation regimes are found for both 4H-SiC faces. A first regime is seen for

growth temperatures in the range 1400-1600°C, in which [N] increases on C-face and decreases

on Si-face. The second regime is seen at growth temperatures 1600-1800°C in which [N]

exhibits different dependence on SiC polarity. Consequently, with the increase in growth

temperature, the [N] decreases sharply on C-face and increases slightly on Si-face. An additional

series of epilayers were grown under lower C/Si ratio=1.5 (empty symbols in fig. 3.11) on Si-

face that display a similar temperature tendency. Slightly higher N incorporation was obtained

for this series. The [N] difference between both series of films grown under different C/Si ratio

can be explained by the site competition effect. A more detailed discussion will be attributed to

C/Si ratio dependence on dopant incorporation further in this chapter.

4.6 4.8 5.0 5.2 5.4 5.6 5.8 6.010

15

1016

1017

1018

1019

1020

Si-face

[N]/N

2 (

cm

-3/s

cc

m)

1/Tgrowth (10-4K)

C/Si=2 C/Si=1.5

C-face

1800 1700 1600 1500 1400

Tgrowth (°C)


127


Like for nitrogen, the influence of growth temperature on Al incorporation was

investigated on both Si- and C-face 4H-SiC layers in both R1 and R2 reactor. The layers were

grown simultaneously in R1 reactor with a variation of growth temperature in the range 1410-

1660°C. The growth was performed under fix conditions: C/Si ratio was set to 1, reactor pressure

200mbar, and TMA flow rate at 1.3x10-3

sccm. The growth rate was of ~8µm/h. The results are

presented in fig. 3.12. We observed that the Al incorporation decreases for both polarities with

increasing growth temperature. As expected, a higher [Al] was found on Si-face compared with

C-face.

Fig. 3.12 Al incorporation as a function of growth temperature on Si- and C-face 4H-SiC layers

grown in R1 and R2 reactor

The study was extended to the films grown in R2 reactor, in order to verify if a similar

dependence can be found. A multilayer stack was deposited on Si-face 4H-SiC and a variation of

growth temperature in the range 1500-1850°C was applied. Compared to the study performed in

R1, slightly different growth conditions were set: reactor pressure 150mbar, C/Si ratio equal to

1.5, and a growth rate of ~5.5µm/h. The TMA flow rate was fixed at 8.66x10-4

sccm. The results

4.6 4.8 5.0 5.2 5.4 5.6 5.8 6.010

15

1016

1017

1018

1019

1020

R1: Si-face

[Al]

(c

m-3

)

1/Tgrowth (10-4K)

R2 : Si-face

R1: C-face

1800 1700 1600 1500 1400

Tgrowth

(°C)


128

are also found in fig. 3.12. It can be seen that the Al incorporation tendency is similar with the

one observed previously. Thus, the [Al] decreases as the temperature increases. The [Al]

difference between the two series of grown films on Si-face can be attributed to different growth

conditions, where the most important factors seem to be the TMA flow rate, reactor pressure and

carrier gas flow rate.

Furthermore, the influence of growth temperature on residual Al incorporation was

investigated. For this study, 5µm thick Si-face 4H-SiC films grown using the gas flow rates and

reactor pressure described previously for voluntary Al incorporation in R1 reactor were used.

The growth temperature was varied between 1660-1410°C with 50°C step, while the aluminum

was not supplied in the gas phase. Consequently, the only possible Al source was the SiC

deposited on the reactor walls during previous growth runs and/or the traces of TMA remaining

in gas lines. Figure 3.13 shows the residual [Al] as a function of reciprocal growth temperature.

Fig. 3.13 Residual Al incorporation as a function of growth temperature on Si-face

4H-SiC films

It can be seen that decreasing the growth temperature results in slight decrease of residual

[Al] from 4.4x1016

cm-3

(at 1660°C) to 1.3x1016

cm-3

(at 1410°C). The observed trend can be

5.3 5.4 5.5 5.6 5.7 5.8 5.9 6.010

15

1016

1017

1018

[Al]

(c

m-3)

1/Tgrowth

(10-4K

-1)

1600 1550 1500 1450 1400

Tgrowth

(°C)

Si-face


129

attributed either to a reduction of Al desorption from the upstream part of reactor walls

(enhanced at higher growth temperatures), or to a progressive exhaustion of Al source (i.e. the

memory effect).

3.3.1.3 Discussions

In order to understand the trends obtained experimentally, let’s start our discussions with

the Al case. The experimental study revealed for both investigated polarities a decrease of Al

incorporation as the growth temperature increases (fig.3.12). According to the thermodynamic

calculations performed by Forsberg et al. [14], the Al incorporation process is governed mainly

by single Al-containing species. The mole fraction of the Al atoms increases continuously with

increasing temperature and thus, should result in higher Al incorporation. Yet, our study showed

an inverse tendency that is usually assigned to a strong Al desorption at high temperatures.

Comparable results have been reported by various research groups [12], [14], [30], [32].

The activation energies have been calculated from the Arrhenius plot (slope of log

[Al] versus 1/ curve) for both Si- and C-face layers grown in the two reactors. Hence, in R1

reactor the study revealed the presence of two activation energies on Si-face, while on C-face

only a single was found. The activation energies on Si-face were: in the temperature range

1410-1560°C estimated at -40kcal/mol (-1.8eV), while in the range =1560-1660°C was

estimated at -146kcal/mol (-6.3eV). On the C-face, the was estimated at around -104kcal/mol

(-4.5eV) in the temperature range 1410-1660°C. This is comparable to the energy obtained on Si-

face for temperatures above 1560°C, which indicates a similar thermal behavior of [Al]

incorporation between both SiC faces.

For the layers grown in R2 reactor on the Si-face, two activation energies were observed.

Thus, the was estimated at -81kcal/mol (-3.5eV) in the temperature range 1500-1650°C, and

around -54kcal/mol (-2.3eV) between 1700-1850°C. In literature, Kojima et al. [12] and

Forsberg et al. [14] have reported stronger Al desorption rates on Si-face, with an of ~ -200 to

-300kcal/mol. On C-face a smaller was reported at~ -50 to -80kcal/mol.

Moreover, the study of growth temperature dependence of N incorporation on both SiC

faces showed to be much complicated than for Al doping. The experimental results display a

strongly dependence of [N] on process temperature and SiC polarity (fig. 3.11). Hence, in the

temperature range 1450-1600°C the [N] increases on C-face with an estimated at ~45kcal/mol


130

(2eV). On Si-face an inverse tendency was seen in the given temperature range, i.e. the [N]

decreases with the increase in growth temperature. An activation energy of -50kcal/mol (-

2.2eV) was calculated. The observed trend may suggest that the N incorporation mechanism is

governed by an adsorption process on C-face and respectively, desorption process on Si-face in

this temperature range. A similar result was reported by Forsberg et al. [7], while Kojima et al.

[12] and Kimoto et al. [30] indicated a desorption limited process for both polarities. Increasing

further the growth temperature up to 1800°C, resulted to a tendency shift while maintaining the

opposite behavior between both polarities. Therefore, in the given temperature range, a strong

decrease of [N] on the C-face, and a slight increase of [N] on Si-face was observed. The

activation energy was estimated on C-face at -186kcal/mol (-8eV), while on Si-face at

58kcal/mol (2.5eV). The activation energies obtained on both SiC faces, in the range of

temperatures studied on both reactors, for the N and Al case are summarized in table 3.1.

for Al case (kcal/mol) for N case(kcal/mol)

(°C) Si-face C-face Si-face Si-face C-face

R1 reactor R2 reactor R2 reactor

1410-1560 -40

-104

1560-1660

-146

1500-1650 -81

1700-1850 -54

1450-1600 -50 45

1600-1800 58 -186

Table 3.1 Activation energies calculated from the Arrhenius plot for both Si- and C-

face layers grown in both reactors for the Al and N cases

According to this study, the dopant incorporation depends strongly on the applied process

temperature and SiC polarities. In order to understand the experimentally observed tendencies of

dopant incorporation, we have considered the crystal surface during the growth under the silane,

propane and dopant flow rates. In Annex II we proposed an explanation regarding the

mechanisms that could be responsible for the observed tendencies of N and Al incorporation.


131

3.3.2 Dopant supply dependence

3.3.2.1 Nitrogen flow rate

The nitrogen flow rate dependence on N incorporation was investigated on both C- and

Si-face of 4H-SiC layers. The series of layers were grown during various runs with different

process conditions. For the study, the input N2 flow rate was varied from 0.03-1000sccm while

the rest of process parameters were kept constant. On C-face, the set conditions for the growth

involved a growth temperature of 1700°C, reactor pressure of 400mbar and a C/Si ratio set at 2.

A growth rate of ~8µm/h was obtained. On the other hand, on Si-face the growth conditions were

slightly different as: growth temperature 1780°C, reactor pressure 100mbar and a C/Si ratio

equal to 1. A growth rate of 24µm/h was found.

The nitrogen concentration [N] as a function of the input nitrogen flow rate is displayed

in fig. 3.14. As can be seen, the [N] increases with the increase in nitrogen flow rate on both SiC

polarities.

Fig. 3.14 Nitrogen incorporation as a function of input N2 flow rate

10-2

100

102

1015

1016

1017

1018

1019

1020

C-face

[N] (c

m-3)

N2 flow rate (sccm)

Si-face


132

3.3.2.2 TMA flow rate

The influence of the TMA flow rate on the Al incorporation was studied for both Si- and

C-face of 4H-SiC layers. The study was performed in R1 reactor by growing multilayer stacks

simultaneously on Si- and C-face, while varying the TMA flow rate in the range 0.001-0.1sccm.

For the growth of the multilayer samples, the set conditions were: growth temperature 1610°C,

reactor pressure 200mbar, C/Si ratio ~1 and a growth rate of 8µm/h. In fig. 3.15(a) the Al

incorporation as a function of input TMA flow rate on both polarities is shown.

It can be noticed from the plot that on both faces the Al incorporation is proportional to

the TMA flow rate. A higher [Al] is found on Si-face as related to C-face. The observed

dependence was confirmed by an additional series of Si-face 4H-SiC layers grown during

various runs and presented also by the empty symbols in fig. 3.15(a). The difference between the

two series comes from process conditions applied during growth runs (e.g. temperature, pressure,

carrier gas flow rate).

The dependence on TMA flow rate of Al incorporation was studied also for a series of

films grown in R2 reactor. The investigation was performed also by growing multilayer stacks

simultaneously on Si- and C-face 4H-SiC as the TMA flow rate was varied in the range 6x10-4

to

1sccm. The deposition conditions were: growth temperature 1820°C, reactor pressure 150mbar,

C/Si ratio ~1.8 and a growth rate of 15µm/h. The obtained results are displayed in fig. 3.15(b).


133

Fig. 3.15 Al incorporation as a function of TMA flow rate for the series grown in: (a) R1

and (b) R2 reactor

It was found a similar trend as the one obtained in R1 reactor, i.e. the [Al] increases with

the increase of the TMA supply for both SiC polarities. However, for this study, two distinct

dependences were found on both faces.

3.3.2.3 Discussions

The dopant incorporation is often considered as a power function of the dopant partial

pressure, which in turn is proportional to the precursor flow and reactor pressure, since the

dopant incorporation is governed by Henry’s law [33]:

(3.1)

where is the dopant concentration, is the partial pressure of dopant containing species and

the is an exponent equal to the reciprocal of the number of dopant atoms presented in the

dopant containing species. For example, if we consider that for Al or N doping the contribution

10-3

10-1

1014

1015

1016

1017

1018

1019

1020

1021

R1:C-face

[Al]

(c

m-3)


R1: Si-face

Tg: 1600°C

a)

R1: Si-face

10-5

10-3

10-1

1014

1015

1016

1017

1018

1019

1020

1021

R2 :C-face

Tg: 1820°C

[Al]

(c

m-3)


R2 :Si-face

b)


134

is made by single atoms in the Al or N-containing species, the =1 while for the molecules with

two atoms (e.g. N2 molecule), the =

.

The dopant partial pressure can be obtained as:

(3.2)

where is the reactor pressure and the , , and are the dopant, silane,

propane and hydrogen flow rates, respectively.

The results on both Si- and C-face showed an increase of N incorporation with increasing

N2 supply. An apparent slope = 0.9 was found on Si-face, while on C-face the slope

= 0.8. Since the apparent slope is equivalent to the exponent and closer to 1, it indicates

mostly the implication of single N-containing species, such as: HCN, HNC, SiNH to the doping

process. Among all the species containing single nitrogen, the thermodynamic calculations

predict that HCN, HNC, SiNH molecules are the dominating species since they have the highest

mole fractions [7].

Despite above mentioned arguments, some authors [34], [35] consider that N2 molecule

itself is involved to the [N] process by a dissociative adsorption to the grown surface.

Nevertheless, the role of all the different N-containing species to the nitrogen incorporation have

not been clearly established up to now. A similar slope value was reported by Forsberg et al. [7]

for both faces when the input N2 flow rate was higher than 0.9sccm, while others have reported

lower values [13] and even a direct proportionality [28].

In the case of Al incorporation, the study exposed for both SiC polarities a direct

proportionality between the TMA supply and the [Al], for the series of layers grown on R1

reactor. Since the = =1, it indicates the direct implication of the single Al-containing

molecules to the doping process. According to the thermodynamic calculation performed by

Forsberg et al. [14] the main Al-containing species are AlH2, Al, AlH, AlCH3, AlH3 due to their

high gaseous concentration. However, it is believed that the Al incorporation is governed by

single Al atoms as they have higher abundance in the gas phase. In literature, it was reported a

direct proportionality between the [Al] and the TMA supply for both SiC faces only by Kojima

et al [13]. Other research groups have reported stronger dependences for both Si- and C-face,


135

[14], [15], [27]. A direct proportionality =1 was obtained for the series of layers grown on

both Si- and C-face in R2 reactor but only for TMA supply higher than 0.01sccm. At TMA

supply lower than 0.01, a stronger dependence was seen especially on Si-face with 1.8,

while on C-face a 1.4 was found.

This difference in dependence is not clearly understood yet. Nevertheless, we cannot

point out to the different growth temperature used between the two reactors (1820°C instead of

1610°C) as the responsible factor for the dependence obtained in R2 reactor. This is mainly

because we found a similar over-proportionality on 3C-SiC layers grown at lower growth

temperature of 1350°C [36]. In literature [37] such an over-proportionality was reported only for

4H-SiC layers grown with high TMA flow rates (>0.01sccm), and attributed to an excess of C-

containing molecules that lead to more available sites on the grown surface for Al atoms to

incorporate to.

As a result, we may just speculate that the visible dependence could be related either to

the geometry/size of the reactor or to other factors that affect the density of available sites on the

grown surface.

3.3.3 Growth rate dependence


The influence of growth rate on N incorporation was studied on Si-face 2° off-axis 4H-

SiC layers. For this study, the layers were grown at a growth temperature of 1770°C, a reactor

pressure of 150mbar and a nitrogen flow rate of 0.5sccm. The silane and propane flow rates were

varied simultaneously while keeping constant the C/Si ratio=1.2. In this way, the silane flow rate

was varied from 5 to 40sccm which corresponds to growth rates ranging from 3.5 to 33µm/h (see

fig. 2.8). Figure 3.16 displays the N incorporation versus the growth rate. The results indicate the

presence of two incorporation regimes. Hence, the first incorporation regime includes the range

of growth rates Gr<10µm/h where we see the N incorporation independent of the growth rate.

For Gr≥10µm/h the second regime is detected, in which the N incorporation decreases with

increasing growth rate.

Forsberg et al. [7] and Danielsson et al. [34] have reported a similar behavior on the Si-

face from both experimental and simulations viewpoint. However, on C-face they showed an

opposite behavior, i.e. the N incorporation showed an increase with increasing growth rate. Since


136

our study was limited only to films grown on Si-face, we cannot say if such a tendency could

occur on C-face in our setup.

Fig. 3.16 Growth rate dependence of nitrogen concentration


The influence of the growth rate on aluminum incorporation was studied on Si-face

2°off-axis 4H-SiC layers, in R2 reactor. Two series of multilayer stacks were grown at various

TMA flow rates (i.e. 5.63x10-3

sccm and 8.74x10-4

sccm) for this investigation. The invariable

growth conditions were: growth temperature 1820°C, reactor pressure 150mbar and a C/Si ratio

equal to 1.8. As described above, the silane and propane were simultaneously varied for

maintaining fixed the C/Si ratio while varying the growth rate. Thus, the growth rate was

changed from 3 to 38µm/h which corresponds to silane flow rates ranging from 5 to 42sccm (see

fig. 2.8). Figure 3.17 represents the growth rate dependence of Al incorporation obtained at

different TMA flow rates. It can be noticed that the [Al] increases with the increase of growth

rate independent of the TMA flow rate. In literature some similar behavior for the growth rate

trend was reported by Forsberg et al. [14] and Ji et al. [38].

1 10

1015

1016

1017

1018

Si-face

[N

] (c

m-3

)

Growth rate (µm/h)


137

Fig. 3.17 Growth rate dependence of aluminum concentration for two TMA flow rates

3.3.3.3 Discussions

Under carbon rich conditions (C/Si>1), the growth rate is limited by the supply of Si and

thus, the growth rate is proportional to the silane flow rate (see fig. 2.8). In the present study, the

dopant flow rate was kept constant and the growth rate was adjusted by increasing the silane and

adapting the propane flow in order to keep a constant C/Si ratio at the reactor inlet.

The observed independence of N incorporation on growth rate for Gr<10µm/h suggests

that the incorporation process is limited by kinetics of surface reactions, too fast to incorporate

all available N atoms. At Gr≥10µm/h the incorporation regime changes and we notice that the

[N] decreases with increasing growth rate. This behavior is assigned to a mass transport limited

incorporation regime, in which all available dopant atoms are incorporated. In this regime, we

expect the dopant incorporation inversely proportional to the growth rate. The experimental

results show a stronger dependence, with an exponent ≈-1.7. Similar behavior was reported by

Forsberg et al. [7]. The explanation could be as follows: for C/Si>1, an increase of growth rate

results in an increase of partial pressure of excess C-containing molecules – local C/Si ratio

1 10

1015

1016

1017

1018

TMA=8.74x10-4 sccm

TMA=5.6x10-3 sccm

[Al]

(c

m-3

)

Growth rate (µm/h)

Si-face


138

increases along the reactor axis. It enhances the competition between N and C leading to further

reduction of N incorporation.

Compared to the trend observed for nitrogen, the growth rate dependence of Al

incorporation showed yet another behavior: for both TMA supply rates used during the study, we

observed an increase of Al incorporation with increasing growth rate. It was found the

dependence with an exponent ≈0.8 for TMA flow rate equal to 8.7x10-4

sccm and exponent

≈0.9 for TMA flow rate of 5.6x10-3

sccm.

To explain this behavior reported previously by Forsberg et al. [14] we’ll make appeal

this time to the ratio between the partial pressure of precursors and the one of the dopant: i.e.

( . As a consequence, an increase of growth rate leads to an increase of the

( ratio which leads not only to an enhanced competition with N atoms but

also to a reduced desorption of the Al-containing species at the grown surface especially on Si-

face.

3.3.4 Reactor pressure dependence


The growth pressure dependence of the N incorporation was investigated on Si-face 4°

off-axis 4H-SiC layers in the pressure range 50-800mbar. The growth conditions were

maintained constant: growth temperature 1700°C, C/Si ratio equal to 2 and a N2 flow rate 5sccm.

A variation of the growth rate was noticed while changing the reactor pressure: from

~8µm/h at 200mbar to ~5µm/h at 800mbar. The effect of reactor pressure on growth rate was

previously presented and discussed in Chapter2, section 2.2.3. The results are represented in fig.

3.18. It can be seen, that increasing the reactor pressure results in an increase of N incorporation.

Similar tendencies were reported for both SiC polarities previously in literature by

Forsberg et al. [7] and Kimoto et al. [4].


139

Fig. 3.18 Nitrogen concentration as a function of reactor pressure


The dependence of Al incorporation on the reactor pressure was studied for a series of

layers grown by varying the pressure range between 50-800mbar. The study was performed on

epilayers grown in both reactors.

In R1 reactor, the study was accomplished by growing multilayer stacks on Si-face 8°

off-axis 4H-SiC substrate. The conditions for the epitaxial process were: growth temperature

1650°C, TMA flow rate of 1x10-3

sccm and a C/Si ratio of 1. The growth rate was found to

decrease with increasing reactor pressure from ~4.5µm/h at 200mbar to ~2.5µm/h at 800mbar.

On the other hand, in R2 reactor the investigation was performed on both Si- and C-face

4H-SiC substrates. The growth was performed also by growing multilayer stacks during the same

growth run on both SiC polarities. The multilayer stacks were grown with a growth temperature

1820°C, TMA flow rate of 5.6x10-3

sccm and a C/Si ratio equal to 1.8. The growth rate was

slightly varied as the reactor pressure was raised (i.e. from ~15µm/h at 150mbar to ~8.5µm/h at

800mbar). The results obtained from the series of layers grown in both reactors are represented

in fig. 3.19. It can be seen, that the Al concentration increases with the increase of growth

102

103

1014

1015

1016

1017

1018

Si-face

[N]

(cm

-3)

pgrowth

(mbar)


140

pressure on both Si- and C-face, respectively. Moreover, a similar slope is observed for the Si-

face films grown in R1 and R2 reactor. The difference in [Al] between the two series of films

grown in both reactors is attributed to the adopted process conditions for the study.

In literature, previous studies made by Forsberg et al. [14] and Ji et al. [38] have reported

a similar dependence.

Fig. 3.19 Aluminum concentration as a function of reactor pressure for Si- and C-face

4H-SiC layers grown in the two available reactors: R1 and R2

3.3.4.3 Discussions

Considering Henry’s law (eq. 3.1), the dopant incorporation can be described as being

directly proportional to the reactor pressure as described by eq. 3.2. An increased growth

pressure leads to an increase partial pressure of dopant-containing species.

An exponent ≈1.4 was found from the [N] as a function of growth pressure. On the

other hand, for [Al] an exponent ≈0.6 was found on Si-face layers grown in both reactors and,

respectively on C-face. We may assume that the mechanism behind the resulted tendencies is

related to a progressive decrease of growth rate for growth pressures above 400mbar. This

102

103

1015

1016

1017

1018

1019

R1:Si-face

R2:C-face

[Al]

(c

m-3)

pGrowth

(mbar)

R2:Si-face


141

decrease of deposition rate impacts directly the dopant concentration as exposed in the previous

section- [N] increases while the [Al] decreases.

Another mechanism reported to play also a role in the resulted tendency was the partial

pressure of the dopant-related species. For nitrogen, the thermodynamic calculations in ref. [7]

and [14] indicate a progressive increase of the mole fraction of N-containing species in the gas

phase with increasing reactor pressure. An opposite tendency was calculated for Al-related

species – their mole fraction decreases at high pressure. Consequently, the partial pressure of the

N-containing species and single Al-containing species is not directly proportional with the

reactor pressure. Hence, the influence of both, growth rate and mole fractions of dopant

containing species on the [N] and respectively, [Al] cannot be excluded.

3.3.5 C/Si ratio dependence

The last parameter investigated was the C/Si ratio, which reflects the state of the

chemical environment during the growth process. As mentioned in the chapter 2 (see eq. 2.7), the

C/Si ratio is actually the ratio between the propane and silane flow rates set during the growth.

Moreover, by varying the ratio of silane and propane flow rates, the dopant incorporation can be

controlled during the SiC growth due to the “site competition” mechanism [2].

Multiple studies have been performed not only for nitrogen and aluminum, but also for

boron and phosphorus [2], [6]–[8], [11], [13], [14], [19], [27], [30]. However, most of the studies

involving the N and Al incorporation have been performed at C/Si ratio>1 (i.e. under C-rich

conditions), and only few at C/Si ratio<1 (i.e. under Si-rich conditions). The tendencies reported

by the groups are qualitatively different. Remarkably, the only observation where all the results

qualitatively converge is the reduction of nitrogen incorporation on Si-face under C-rich

conditions. The discrepancies between different studies indicate a potentially important influence

of the growth equipment on the tendencies observed.

Within this section, the study of the influence of C/Si ratio on voluntary incorporation of

N and Al in both, Si- and C-face 4H-SiC films will be presented and discussed. As for the other

tendencies, the influence of C/Si ratio on N incorporation was studied in R2 reactor, while the Al

incorporation was investigated in both in R1 and R2 reactors.

Sample series were grown under slightly different growth conditions, as summarized in

table 3.1. The N2 and TMA flow rates were also varied between the series (N2=0.08-1.85sccm,


142

TMA=1.12x10-3

-2.27x10-2

sccm). Hence, the results will be represented as the normalized [N]

and [Al] with respect to N2 and TMA flow rates used during the growth.

Sample

series#

Dopant Orientation C/Si ratio

(°C)

(mbar)

(µm/h)

N1

N in R2

Si-face 4° C/Si>1

~1750

250 const.

N2 Si-face 4° C/Si<1 250 var.

N3 C-face 8° C/Si>1 150 const.

N4 C-face 8° C/Si<1 150 const.

A1

Al in R1

Si-face 8° C/Si>1

~1650

200 const.

A2 Si-face 8° C/Si<1 200 var.

A3 C-face 8° C/Si>1 200 const.

A4 C-face 8° C/Si<1 120 const.

A5

Al in R2

Si-face 4° C/Si>1

~1820

150 const.

A6 Si-face 4° C/Si<1 150 const.

A7 C-face 4° C/Si>1 150 const.

A8 C-face 4° C/Si<1 150 const.

Table 3.2 Growth conditions of different samples series grown for the C/Si ratio dependence of

the N and respectively, Al incorporation in R1 and R2 reactor


The influence of C/Si ratio on N incorporation is shown in fig. 3.20 on Si- and C-face

4H-SiC layers, respectively. As mentioned previously the results are normalized with respect to

nitrogen flow rate used during the growth.


143

Fig. 3.20 Evolution of normalized N incorporation in Si- and C-face 4H-SiC films as a function

of C/Si ratio. Symbol refers to the samples for which C/Si ratio was modified by adjusting

the silane flow during high temperature process (HT). The refers to the series of films grown

at different growth parameters (LT).

N incorporation under C-rich conditions

Under C-rich conditions (C/Si ratio>1) the C/Si ratio was adjusted by keeping fixed the

silane flow while the propane flow rate was changed. In this way, the growth rate is maintained

constant over the range of C/Si ratio. It can be noticed, that a C/Si ratio increase leads to a

decrease of N incorporation on both faces. However, a much stronger dependence is observed on

Si-face. The [N] on C-face is more important than on Si-face as expected, due to the well

explained lattice polarity influence. Previous studies [2], [6], [8], [20], [34] have reported a

similar trend for both faces at high C/Si ratios. Other studies [7], [11] have shown that at

atmospheric pressure the [N] on C-face is insensitive to C/Si ratio changes. The insensibility to

C/Si ratio on C-face was attributed to a stronger C-coverage of the grown surface (see ref. [11],

[20]), which seems to prevent any dependence of dopant incorporation on C/Si ratio. Kojima et

0.1 1 1010

14

1015

1016

1017

1018

1019

C-face (LT)

C-face (HT)

Si-rich conditions

[N]/

N

2 (

cm

-3/s

cc

m)

C/Si ratio

C-rich conditions

Si-face (HT)

0.1 1 1010

14

1015

1016

1017

1018

1019

C face Si face C face Si face C face

C-face

Si-rich conditions

[N]/

N2 (

cm

-3/s

cc

m)

C/Si ratio

C-rich conditions

Si-face

0.1 1 1010

14

1015

1016

1017

1018

1019


C-face

Si-rich conditions

[N]/

N2 (

cm

-3/s

cc

m)

C/Si ratio

C-rich conditions

Si-face


144

al. [13] suggested that under low pressure growth conditions, the N incorporation on C-face may

become sensitive to the C/Si ratio modifications.

Based on these arguments, an additional series of films was grown on C-face under C-

rich conditions (empty triangles symbols on the graph at LT). For this series some of the growth

parameters were changed as: reducing the growth temperature up to 1700°C and increasing the

growth pressure to 400mbar. For this series we wanted to verify the impact of higher growth

pressure on the C/Si ratio dependence of [N]. As a result, it was found that the N incorporation is

almost constant as the C/Si ratio increases.

N incorporation under Si-rich conditions

On C-face, the C/Si ratio was tuned by varying the silane flow rate while keeping fixed

the propane flow rate. The growth rate was maintained constant. On Si-face, silane flow rate was

fixed and the propane flow rate was adjusted. Consequently, on Si-face we have to consider the

influence of growth rate that decreases at low C/Si ratios.

It can be seen from fig. 3.20, that decreasing the C/Si ratio results to slight increase of N

incorporation on both faces. Yet, on Si-face a slightly stronger variation of [N] is observed as

compared to C-face. Different research groups [39], [40] reported results consistent with ours.

On the other hand, Kojima et al. [6] reported no influence of C/Si ratio on N incorporation on

both faces.


The influence of C/Si ratio on Al incorporation was studied for both, Si- and C-face 4H-

SiC layers. The study was initially performed in the R1 reactor and further extended to the R2

reactor. Figure 3.21(a)-(b) shows the C/Si ratio dependence of Al incorporation on series of films

grown in both reactors. The results are normalized with respect to TMA flow rate used

during the growth. The epilayers were grown during individual runs and/or during single growth

runs by creating a multilayer stacks on both faces.


145

Fig. 3.21 Evolution of Al incorporation in Si- and C-face 4H-SiC films as a function of

C/Si ratio grown in: (a) R1 and (b) R2 reactor. Symbol and refer to the samples for which

C/Si ratio was modified by adjusting the silane flow rate.

Al incorporation under C-rich conditions

As for the study of [N], the C/Si ratio was tuned by varying the propane flow rate and

fixing the silane flow. From fig. 3.21(a) and (b), a stronger Al incorporation is observed on Si-

face compared to C-face which is related to the lattice polarity influence. For both studied

polarities, no influence of C/Si ratio on [Al] was observed in our R1 reactor (fig. 3.21a).

However, for R2 setup in which higher temperature is applied during the growth process it was

found that the [Al] is increasing with the C/Si ratio increase on both faces (fig. 3.21b). However,

a considerably stronger variation was noticed on Si-face compared to C-face.

Previous works [2], [13], [19] reported an increasing [Al] on Si-face with the increase of

C/Si ratio. On C-face most reports announced an increase of [Al], and only few indicate an

insensibility to C/Si ratio changes (only at atmospheric pressure) [11], [14], [20].

Al incorporation under Si-rich conditions

Under these conditions, the C/Si ratio was adjusted by changing the silane supply and

fixing the propane flow rate for C-face layers grown in R1 reactor. For both Si- and C-face films

0.1 1 1010

16

1017

1018

1019

1020

1021

Tg:1650°C

R1:C-face

Si-rich conditions

[Al]

/T

MA

(c

m-3

/sc

cm

)

C/Si ratio

C-rich conditions

R1:Si-face

a)

0.1 1 1010

16

1017

1018

1019

1020

1021

b)


R2:C-face

[Al]

/T

MA

(c

m-3

/sc

cm

)

C/Si ratio

R2:Si-face

Tg:1820°C

0.1 1 1010

14

1015

1016

1017

1018

1019


C-face

Si-rich conditions

[N]/

N2 (

cm

-3/s

cc

m)

C/Si ratio

C-rich conditions

Si-face

0.1 1 1010

14

1015

1016

1017

1018

1019

Si-face C-face Si face C-face

b)


R2:C-face

[Al]

(c

m-3)

C/Si ratio

R2:Si-face

Tg:1820°C


146

grown in R2 reactor, the same C/Si ratio tuning was used. Only on Si-face layers grown in R1

reactor, the propane was varied to change the C/Si ratio.

We observe from fig. 3.21(a)-(b) a decrease of [Al] on both faces as the Si-supply is

increased (i.e. C/Si ratio decreases). A stronger dependence is noticed on Si-face layers grown on

both reactors. However, the variation is clearly more important on Si-face films grown in R1

reactor, keeping in mind the impact of the decrease growth rate during this experiment. Up to

now, previous studies reported mostly the C/Si ratio dependence on the Al incorporation only for

high C/Si ratios (C/Si ratio>1) and only one study was extended towards C/Si<1 [13].

To extend further the study, an additional series of 1µm thick films was grown on Si-face

4H-SiC substrates in R1 reactor using a C/Si ratio≤1.25. The substrates were rotated during the

growth. The growth conditions set for this study were: growth temperature 1650°C, TMA flow

rate of 1x10-3

sccm and a reactor pressure 200mbar. The C/Si ratio was adjusted by varying the

propane supply while the silane flow rate was fixed. A decrease of the growth rate was observed

from C/Si ratio=0.25 and moreover, at C/Si ratio=0 (no propane flow!) a deposition rate ~2µm/h

was found. This effect was previously discussed in the section allocated to the in Chapter 2

and assigned to a C-supply source inside the reactor.

For this study the [Al] at different C/Si ratios was investigated along the substrate

diameter as determine by C-V measurements. The results regarding the C/Si ratio in the range

1.25 to 0.5 dependence of Al incorporation are displayed in fig. 3.22. It can be observed, a

decrease of [Al] with the increase in Si-supply (i.e. when C/Si ratio decreases).


147

Fig. 3.22 Net acceptor concentration [Al] at different C/Si ratio measured along the

epilayer diameter. The arrow indicates the measurement direction along the sample.

Furthermore, higher [Al] was found at the wafer border compared to the center. It was

found that a decrease of C/Si ratio leads to higher relative dispersion of [Al]. A further decrease

of the C/Si ratio below 0.25 leads to a change in conduction type from p-type to n-type

( = !). It is difficult to determine which effect is more dominate, i.e. a decrease of voluntary

[Al] related to high Si-supply (i.e. Al-Si competition) or an increase of residual [N] (i.e. N-C

competition).

3.3.5.3 Discussions

The decrease of [N] observed for both faces under C-rich conditions can be explained by

the presence of site competition between N and C atoms. From the experimental data an

exponent ≈ -4.7 was found for Si-face and ≈ -1.5 for C-face, respectively.

Under C-rich conditions, the gas phase presents an excess of C-related species that can

outcompete nitrogen atoms. If the nature of this site competition was purely statistical, we could

consider that the probability P of finding an N atom ready to incorporate on a C-site is equal to

the population of N-containing species reported to the total population of species that contribute

to the occupation of C-sites. The expression should be written as:


148

(3.3)

where and are the proportionality coefficients indicating the fraction of nitrogen and

carbon containing species, respectively participating to the film formation. If we estimate that

then the probability becomes:

(3.4)

Thus, purely statistical competition between N and C results in an exponent = -1. In

our case as previously related, a much stronger variation of [N] on Si-face was found (exponent

~ -4.7). In literature various groups (see e.g. ref. [6], [7], [11]) reported an exponent in the

range -3≤ ≤-2 on Si-face. On C-face, the variation of [N] was found to be smaller ( ~ -1.5).

A similar exponent was reported in ref. [6], while other groups [7], [11] exposed that the C-face

is almost insensitive to C/Si ratio in C-rich conditions.

The observation of exponent absolute values higher than -1 contradicts the usual

qualitative explanation in which the observed decrease of [N] is attributed to a stronger

probability of N to be outcompeted by C atoms at higher propane flow rates. Consequently, the

nature of the site competition is not purely statistical, at least not on Si-face.

On C-face where the observed exponent was close to -1, we may associate the [N]

variation with the statistical competition between the N and C atoms. The additional series of

layers grown on C-face under C-rich conditions with slightly lower growth temperature and

higher reactor pressure showed that [N] is quasi-independent on C/Si ratio.

The carbon coverage of the surface increases at low temperature and high pressure. Thus,

we may speculate that the increase of reactor pressure and reduction of growth temperature led to

saturated carbon coverage of the surface and, consequently, to the insensitivity on C/Si ratio

variation. This behavior is similar with the one reported on C-face at atmospheric pressure.

A model [7] was proposed wherein the C-face is C-terminated even under Si-rich

conditions, making the N incorporation in C-sites independent on C-supply. Another model [13]

suggested that applying low pressure conditions for the growth on C-face, results in a re-

evaporation of C-atoms from the grown surface. This will create more free C-sites for N to

incorporate and become dependent on the C-supply. Such a behavior can be qualitatively


149

described by a Langmuir adsorption type model [41] as presented in Annex III of this

manuscript. However, in literature the authors give qualitative explanations for the

experimentally obtained tendencies without relying explicitly on the support of a Langmuir

isotherm type equation.

Our results have been obtained by changing two growth parameters, i.e. growth pressure

and temperature. Unfortunately, it is impossible to evaluate which of two parameters is

dominant: both of them affect the adsorption/desorption rates (taken into account by adsorption

model) and both of them influence strongly the mole fraction of reactive species in the gas phase.

Under Si-rich conditions, for both SiC polarities we observed an increase of [N] when the

C/Si ratio decreases. On C-face the observed tendency ( ≈ -1) cannot be explained by the

competition between C and N since the C-supply is kept fixed. On the other hand, the excess of

Si in the gas phase results in higher Si coverage of the surface creating more C-sites available for

incorporation of N atoms [40]. Slightly stronger variation was observed on Si-face ( ≈-1.5). As

the propane flow rate decreases, the growth rate reduces and the capacity of carbon to

outcompete N atoms becomes lower. Also, the Si-coverage of the surface becomes more

important as in the experiments conducted on C-face. Thus, all this factors contribute to the

increase of N incorporation.

Other works [39], [40] reported similar tendencies as ours. No dependence of [N] on low

C/Si ratio for both faces was reported in ref. [6]. However, the experiments were conducted at

lower temperature and we may speculate that the C-coverage of the surface was almost

independent on C/Si ratio.

In the experiments conducted at low temperature (1650°C) in R1 reactor, a well-defined

decrease of [Al] was observed for both faces under Si-rich conditions, which can be clearly

associated to the competition between Al and Si atoms. Only one study reported a similar

tendency even though the range of C/Si ratio study was very narrow [13].

Stronger variation was found on Si-face with an exponent ≈2.6, compared to C-face

( ≈1.5) films grown in R1 reactor. This variation can be attributed to the progressive reduction

of at low propane flow rates. In the case of films grown in the R2 reactor at 1820°C, the data

for very low C/Si ratio are unfortunately missing. An exponent ≈1.6 was found on Si-face,

while on C-face the variation is almost negligible ( ≈0.3 for the C/Si in the range 0.7-3).

Therefore, we can use the same analysis as the one performed for [N] to determine the nature of


150

the site competition between Al and Si atoms. In R1 reactor as shown previously an exponent

close to 1 was found on C-face, suggesting the statistical nature of the site competition. A similar

conclusion can be draw for the epilayers grown in R2 reactor on Si-face. However, stronger

variation was obtained on Si-face in R1 reactor associated to the progressive decrease of growth

rate at low C/Si ratio. Thus, we cannot exclude a statistical nature of site competition with the

addition of growth rate influence. Unfortunately, on C-face layers grown in R2 reactor data are

missing which makes difficult the results interpretation.

Under C-rich conditions, the site competition between the Al and Si atoms is not changed

since the C/Si ratio was adjusted by a variation of the propane flow rate. In R2 reactor, a

variation of the [Al] with C/Si ratio was observed on both faces. Such an observation is

associated to a lower C-coverage on the growing surface due to stronger growth temperature

(1820°C). Therefore, a reduced C-coverage allows the increase of the C-containing species to

create more available bonds for Al atoms to incorporate. Previous works have reported an similar

behavior of [Al] on Si-face [11], [14], [20].

In R1 reactor, where a lower growth temperature (1650°C) was used, no [Al] variation

was observed with the changes of C/Si ratio on both faces. This behavior can be associated to the

lower temperature applied during the process, which results in stronger C-coverage that is not

sensitive to the C/Si ratio changes.

The importance of the effective C/Si ratio on the dopant incorporation was previously reported

[26], [42]. The effective C/Si ratio at the surface is a good indicator of the density of available

sites. Consequently, the dopant incorporation depends on the total number of available sites,

surface coverage, residence time of molecules impinging on the surface and their concentration

in gas phase.

3.4 Conclusions

In this chapter we have evaluated the dependence of growth parameters of N and Al

incorporation. We could assess initially that various SiC polytypes and/or substrates off-angles

does not influence considerably the N/ Al incorporation. Due to the site preferences of the

impurities atoms in the SiC matrix, a stronger incorporation was found as expected on C-face for

N incorporation [N] and respectively, on Si-face for Al incorporation [Al]. Furthermore, the


151

experimental study of the influence of growth conditions (growth temperature, dopant supply,

growth rate, reactor pressure, C/Si ratio) on voluntary [N] and [Al] in 4H-SiC heteroepitaxial

films reveal different incorporation trends. Growth temperature dependence on [N] put in

evidence an unexpected tendency for both polarities. In the case of Al, in both reactors a similar

reduction of [Al] with increasing growth temperature was observed on both polarities. We

proposed an qualitatively explanation for the experimentally obtained tendency, which considers

a direct implication of factors such as sticking coefficient of dopant-containing species, carbon

coverage of the surface and the dopant flow rate.

The dopant supply dependence on N and Al incorporation has showed that the dopant

incorporation increases almost proportionally with the increase of dopant flow rate. However, for

Aluminum, in R2 reactor a much stronger dependence was found at lower TMA flow rate

compared to R1 reactor. The mechanism behind this dependence was not fully understood but we

can correlate it to the growth conditions, resulting to more available incorporation sites.

Also, the growth rate dependence on [N] and [Al] lead to opposite experimentally

tendencies on Si-face. The [N] tends to decrease at higher growth rates (>10µm/h) and remains

insensitive at growth rates ≤10µm/h. On the contrary, the [Al] displayed an increase with

increase growth rate. The dependence of reactor pressure showed a similar tendency for both [N]

and [Al]. Moreover, the similar dependence of [Al] was observed in both reactors and for both

polarities.

The exhaustive study of the influence of C/Si ratio on voluntary [N] and [Al] put in

evidence various dopant incorporation tendencies: (i) site competition between N and C was

found under C-rich condition while site competition between Al and Si was visible under Si-rich

conditions; (ii) a surface coverage influence of [N] under Si-rich conditions and analogously of

[Al] which can be determined by an Langmuir isotherm model. Besides, under C-rich conditions,

the [N] on C-face was found independent on C/Si ratio when a lower growth temperature and

higher reactor pressure were applied. In the case of [Al], in R2 reactor a distinctive tendency was

obtained under C-rich conditions for both polarities compared to experimentally observed trends

in R1 reactor. This evidences the role of growth conditions on surface coverage.


152

3.5 Bibliography


superior silicon carbide electronics,” Applied Physics Letters, vol. 65, no. 13, p. 1659–1661,

1994

[2] D. J. Larkin, “SiC Dopant Incorporation Control Using Site-Competition CVD,” phys. stat.

sol. (b), vol. 202, no. 1, p. 305–320, 1997


[4] T. Kimoto, T. Hirao, S. Nakazawa, H. Shiomi, and H. Matsunami, “Homoepitaxial growth of

4H–SiC(0 3 3 8) and nitrogen doping by chemical vapor deposition,” Journal of Crystal

Growth, vol. 249, no. 1–2, p. 208–215, 2003

[5] T. Kimoto, K. Hashimoto, and H. Matsunami, “Effects of C/Si Ratio in Chemical Vapor

Deposition of 4H-SiC(1120) and (0338),” Jpn. J. Appl. Phys., vol. 42, no. 12R, p. 7294,

2003

[6] K. Kojima, S. Kuroda, H. Okumura, and K. Arai, “Nitrogen incorporation characteristics on

a 4H-SiC epitaxial layer,” Applied Physics Letters, vol. 88, no. 2, p. 21907, 2006

[7] U. Forsberg, Ö. Danielsson, A. Henry, M. K. Linnarsson, and E. Janzén, “Nitrogen doping of

epitaxial silicon carbide,” Journal of Crystal Growth, vol. 236, no. 1–3, p. 101–112, 2002.

[8] H. Pedersen, F. C. Beyer, J. Hassan, A. Henry, and E. Janzén, “Donor incorporation in SiC

epilayers grown at high growth rate with chloride-based CVD,” Journal of Crystal Growth,

vol. 311, no. 5, p. 1321–1327, 2009

[9] M. Yamanaka, H. Daimon, E. Sakuma, S. Misawa, K. Endo, and S. Yoshida, “Temperature

Dependence of Electrical Properties of Nitrogen-Doped 3C-SiC,” Jpn. J. Appl. Phys., vol.

26, no. 5A, p. L533, 1987

[10] M. Zielinski, M. Portail, T. Chassagne, S. Juillaguet, and H. Peyre, “Nitrogen doping of

3C-SiC thin films grown by CVD in a resistively heated horizontal hot-wall reactor,”

Journal of Crystal Growth, vol. 310, no. 13, p. 3174–3182, 2008

[11] T. Kimoto, A. Itoh, and H. Matsunami, “Incorporation mechanism of N, Al, and B

impurities in chemical vapor deposition of SiC,” Applied Physics Letters, vol. 67, no. 16, p.

2385, 1995

[12] K. Kojima, S. Kuroda, H. Okumura, and K. Arai, “Influence of lattice polarity of nitrogen

and aluminum doping on 4H-SiC epitaxial layer,” Microelectronic Engineering, vol. 83, no.

1, p. 79–81, 2006

[13] K. Kojima, T. Suzuki, S. Kuroda, J. Nishio, and K. Arai, “Epitaxial growth of high-

quality 4H-SiC carbon-face by low-pressure hot-wall chemical vapor deposition,” Japanese

journal of applied physics, vol. 42, no. 6B, p. L637, 2003

[14] U. Forsberg, O. Danielsson, A. Henry, M. Linnarsson, and E. Janzén, “Aluminum doping

of epitaxial silicon carbide,” Journal of Crystal Growth, vol. 253, no. 1–4, p. 340–350, 2003

[15] G. Wagner, W. Leitenberger, K. Irmscher, F. Schmid, M. Laube, and G. Pensl,

“Aluminum Incorporation into 4H-SiC Layers during Epitaxial Growth in a Hot-Wall CVD

System,” Materials Science Forum, vol. 389–393, p. 207–210, 2002

[16] D. J. Larkin, “Dopant Incorporation Efficiency in CVD Silicon Carbide Epilayers,” in

MRS Proceedings, vol. 410, p. 337, 1995

[17] A. Henry and E. Janzén, “Epitaxial Growth and Characterisation of Phosphorus Doped

SiC Using TBP as Precursor,” Materials Science Forum, vol. 483–485, p. 101–104, 2005

[18] A. Gali, P. Deák, P. R. Briddon, R. P. Devaty, and W. J. Choyke, “Phosphorus-related

deep donor in SiC,” Physical Review B, vol. 61, no. 19, p. 12602, 2000


153

[19] H. Pedersen, F. C. Beyer, A. Henry, and E. Janzén, “Acceptor incorporation in SiC

epilayers grown at high growth rate with chloride-based CVD,” Journal of Crystal Growth,

vol. 311, no. 13, p. 3364–3370, 2009

[20] T. Yamamoto, T. Kimoto, and H. Matsunami, “Impurity Incorporation Mechanism in

Step-Controlled Epitaxy Growth Temperature and Substrate Off-Angle Dependence,”

Materials Science Forum, vol. 264–268, pp. 111–114, 1998

[21] W. Chen and M. A. Capano, “Growth and characterization of 4H-SiC epilayers on

substrates with different off-cut angles,” Journal of Applied Physics, vol. 98, no. 11, p.

114907, 2005

[22] B. Kallinger, P. Berwian, J. Friedrich, and B. Thomas, “Step-controlled homoepitaxial

growth of 4H–SiC on vicinal substrates,” Journal of Crystal Growth, vol. 381, p. 127–133,

2013

[23] H. Matsunami and T. Kimoto, “Step-controlled epitaxial growth of SiC: High quality

homoepitaxy,” Materials Science and Engineering: R: Reports, vol. 20, no. 3, p. 125–166,

1997

[24] T. Kimoto, Z.Y. Chen, S. Tamura, S. Nakamura, N. Onojima, and H. Matsunami,

“Surface morphological structures of 4H-, 6H- and 15R-SiC (0001) epitaxial layers grown

by CVD,” Japanese Journal of Applied Physics, vol. 40, p. 3315, 2001


superior silicon carbide electronics,” Applied Physics Letters, vol. 65, no. 13, p. 1659, 1994

[26] S. Nishizawa, K. Kojima, S. Kuroda, K. Arai, and M. Pons, “Modeling of SiC-CVD on

Si-face/C-face in a horizontal hot-wall reactor,” Journal of Crystal Growth, vol. 275, no. 1–

2, pp. e515–e520, 2005

[27] C. Sartel, V. Soulière, M. Zielinski, Y. Monteil, J. Camassel, L. Smith, and S. Rushworth,

“Control of Al-doping in 4H-SiC homo-epitaxial layers grown with a HMDS/TMA/P

mixture,” physica status solidi (a), vol. 202, no. 4, p. 550–554, 2005

[28] T. Aigo, M. Kanaya, M. Katsuno, H. Yashiro, and N. Ohtani, “Nitrogen Incorporation

Mechanism and Dependence of Site-Competition Epitaxy on the Total Gas Flow Rate for

6H-SiC Epitaxial Layers Grown by Chemical Vapor Deposition,” Jpn. J. Appl. Phys., vol.

40, no. 4R, p. 2155, 2001

[29] K. Chindanon, H. D. Lin, G. Melnychuk, and Y. Koshka, “Nitrogen doping in low-

temperature Halo-carbon homoepitaxial growth of SiC,” Materials Science Forum, vol. 600–

603, p. 159–162, 2009

[30] T. Kimoto, A. Itoh, N. Inoue, O. Takemura, T. Yamamoto, T. Nakajima, and H.

Matsunami, “Conductivity Control of SiC by In-Situ Doping and Ion Implantation,”

Materials Science Forum, vol. 264–268, p. 675–680, 1998

[31] G. Wagner, M. Schmidbauer, K. Irmscher, P. Tanner, and R. Fornari, “P-Type Doping of

Epitaxial 3C-SiC Layers on Silicon (001),” Materials Science Forum, vol. 615–617, p. 165–

168, 2009

[32] C. Sartel, “Homoépitaxie du SiC-4H a partir de différents précurseurs. Realisation du

dopage p,” Universite Claude Bernard-Lyon1, 2003

[33] H. J. Kim and R. F. Davis, “Theoretical and Empirical Studies of Impurity Incorporation

into β-SiC Thin Films during Epitaxial Growth,” Journal of The Electrochemical Society,

vol. 133, p. 2351, 1986


154

[34] O. Danielsson, U. Forsberg, and E. Janzén, “Predicted nitrogen doping concentrations in

silicon carbide epitaxial layers grown by hot-wall chemical vapor deposition,” Journal of

Crystal Growth, vol. 250, no. 3–4, p. 471–478, 2003



Growth, vol. 267, no. 3–4, p. 436–451, 2004

[36] M. Zielinski, R. Arvinte, T. Chassagne, A. Michon, M. Portail, P. Kwasnicki,

Konczewicz L., S. Contreras, S. Juillaguet, and H. Peyre, “p-type doping of 4H-SiC and 3C-

SiC epitaxial layers with Aluminum,” Materials Science Forum, vol. 858, p. 137, 2016



[38] S. Ji, K. Kojima, Y. Ishida, S. Saito, T. Kato, H. Tsuchida, S. Yoshida, and H. Okumura,

“The growth of low resistivity, heavily Al-doped 4H–SiC thick epilayers by hot-wall

chemical vapor deposition,” Journal of Crystal Growth, vol. 380, p. 85–92, 2013

[39] R. Wang, I. B. Bhat, and T. P. Chow, “Epitaxial growth of n-type SiC using phosphine

and nitrogen as the precursors,” Journal of Applied Physics, vol. 92, no. 12, p. 7587, 2002

[40] J. Zhang, A. Ellison, M. K. Linnarsson, E. Janzén, and A. Henry, “Nitrogen incorporation

during 4H-SiC epitaxy in a chimney CVD reactor,” Journal of crystal growth, vol. 226, no.

2, p. 267–276, 2001

[41] https://en.wikipedia.org/wiki/Reactions_on_surfaces

[42] S. Nishizawa and M. Pons, “Growth and Doping Modelling of SiC-CVD in a horizontal

hot wall reactor,” Chemical Vapor Deposition, vol. 12, p. 516–522, 2006

Chapter 4 Material properties of Al doped films

155

Contents

Chapter 4 Material properties .................................................................................................. 156

4.1 Introduction .......................................................................................................................... 156

4.2. Samples preparation ........................................................................................................... 157

4.2.1 Doping estimation ............................................................................................................ 159

4.3. Surface morphology ............................................................................................................ 160

4.4. Structural characterization ................................................................................................ 164

4.4.1. X-ray rocking curve measurements ................................................................................ 164

4.4.2. Determination of lattice constants .................................................................................. 165

4.4.3 Reciprocal space mapping (RSM) ................................................................................... 169

4.4.4 Polytype analysis ............................................................................................................. 172

4.4.4.1 Discussions: 3C-SiC polytype .................................................................................. 174

4.5 Physical properties of Al doped epilayers .......................................................................... 179

4.5.1 Optical properties ............................................................................................................ 179

4.5.1.1. Low Temperature Photoluminescence (LTPL) measurements ................................ 179

4.5.2 Electrical properties ......................................................................................................... 182

4.5.2.1 Hall Effect measurements ......................................................................................... 182

4.6 Conclusions ........................................................................................................................... 190

4.7. Bibliography ........................................................................................................................ 191


156

Chapter 4 Material properties

4.1 Introduction

Considering the device fabrication, the doping of active layer can be obtained either by

epitaxial growth or ion implantation. The fabrication of vertical devices requires the presence of

low-resistivity layers or substrates. Nowadays, n-type 4H-SiC substrates doped in the range

5x1018

up to 1x1019

cm-3

show resistivity values of 15-28m .cm [1], [2].

The n-type 4H-SiC heavily doped layers have proven their development, reaching a

resistivity of 1.5m .cm for layers with nitrogen concentration [N] up to 1x1021

cm-3

[3].

Contrarily, the typical resistivity for p-type 4H-SiC substrates is around 2.5 .cm [4]. This causes

an impediment in the progress of p-type based electronic devices, as e.g. the case of insulated

gate bipolar transistors (IGBT’s) where n-channel is usually implemented [5]. In the case of p-

type doped epilayers, resistivity values as low as 16.5m .cm for layers with an Al concentration

[Al] of ~4x1020

cm-3

have been reported [6]. Moreover, high Al concentrations lead to surface

morphology deterioration. Therefore, the active layer of the SiC power devices requires optimal

doping density and thickness to keep high structural quality.

In our setups, as reported in the previous chapter, the nitrogen and aluminum

concentrations in the 4H-SiC films can be controlled in the range of 1x1015

up to 1x1019

cm-3

by

tuning various growth parameters. Considering, the strong consequence of dopant incorporation

on epitaxial growth and device production, it is essential to apply various characterization

techniques to obtain all the aspects of crystal quality and properties.

In this chapter, detailed studies of aluminum doped 4H-SiC layers properties will be

presented and discussed. After the examination of the surface morphology dependence of Al

concentration, the structural, optical and electrical properties were explored. Consequently, to

understand the impact of [Al] on material quality, various characterization techniques were

employed such as: XRD, LTPL, Raman spectroscopy and Hall effect analysis.


157

4.2. Samples preparation

Al-doped homoepitaxial layers were grown on commercially available two inch n+ 4H-

SiC wafers, 8° off-oriented towards [11 0] by using our 2’’ hot wall CVD system (R1 reactor).

The 4H-SiC substrates with (0001) Si-face were cut in quarters for the study. C-V measurements

were performed to detect the net donor concentration ( - ) of the 4H-SiC substrates. The -

concentration of the two substrates was found to be equal to 1.7x1018

cm-3

and 5x1018

cm-3

,

respectively.

A series of 8 samples were grown using a flow rate of SiH4=10sccm and C3H8=3.3sccm,

resulting in a C/Si ratio equal to 1 and growth rate of ~10µm/h. The flow rate of the H2 carrier gas

was 14slm. The growth temperature was 1650°C, and a reactor pressure of 120mbar was applied.

The TMA bubbler temperature was set to 20°C, and the bubbler pressure at 2bars. By adjusting

the H2 flow rates in the dilution line, the TMA mixture flow rate was varied between 5.6x10-4

to

1.43x10-1

sccm. The sample holder was rotated during the growth to obtain the thickness and

doping uniformity of the epilayers.

SIMS and C-V analysis were used to evaluate the dopant concentration of the grown

epilayers. The epilayers thickness was determined by FTIR and SEM techniques. Table 4.1

contains the information regarding the growth conditions of all the grown samples. In order to

obtain higher Al incorporation (>1x1019

cm-3

), a decrease in growth temperature was chosen for

sample 14-009. The rest of the growth parameters were kept constant. The applied condition was

based on the results obtained and described in chapter 3, i.e. at fixed TMA flow rate a lower

growth temperature results in higher [Al] [7], [8].

Due to intriguing material quality of the highly doped epilayer (sample 14-009) which

will be described in the next sections, two additional samples have been grown using similar

growth conditions as sample 14-007 and 14-009. For this study, 2-inch 4H-SiC 8° off-oriented

wafers were used and the information regarding the grown epilayers is also summarized in table

4.1. The slight thickness difference between the epilayers is attributed to the variation of the

epitaxial growth duration.


158

Sample # Substrate/

reactor

TMA flow

(sscm)

(°C)

C-V

-

(cm-3

)

SIMS

[Al]

(cm-3

)

FTIR

Thickness

(µm)

14-012

8° off-axis

4H-SiC/

R1

0 1650 2.21x1016

2.3x1016

14

14-015 5.62x10-4

1650 8.9x1016

6.2x1016

13.8

13-331 5x10-4

1650 1.9x1017

1.4x1017

10

14-003 4.41x10-3

1650 1.5x1018

9.1x1017

11.7

14-004 2.21x10-2

1650 5x1018

6.8x1018

10.4

14-005 4.25x10-2

1650 8x1018

3.1x1018

10.1

14-007 1.43x10-1

1650 ~1x1019

1.7x1019

12.4

14-009 1.43x10-1

1550 >5x1019

7.8x1019

13.8

14-377 1.43x10-1

1550 ~1x1020 8

14-395 1.43x10-1

1650 ~6.5x1019 8

16H4-010 4° off-axis

4H-SiC/

R2

1.1x10-1

1820 ~1.5x1019

8.2x1018

15

16H4-011 6.1x10-1

1820 ~5x1019 5.5x10

19 16

Table 4.1 Growth parameters of the as-grown layers and their reciprocal Al concentration and

thickness obtained by different analysis techniques

The study has been extended also in the 100mm hot wall CVD system (R2 reactor). A

series of highly Al doped layers have been grown on pieces of 3-in n+

4H-SiC wafers, 4° off-

oriented towards [11 0]. The series of 2 samples were grown using a SiH4 flow rate of 40sccm

and a C3H8 flow of 24sccm, resulting to a C/Si ratio equal to 1.8 and growth rate of ~30µm/h.

The flow rate of the H2 carrier gas was equal to 38slm. A growth temperature of 1820°C and a

reactor pressure of 100mbar were applied. The TMA flow rate was varied between 1.1x10-1

to

6.1x10-1

sccm, and a rotation of the substrate holder was applied during process. The Al

concentration was evaluated by C-V measurements. Details about the growth conditions and [Al]

were also added in table 4.1.


159

4.2.1 Doping estimation

The last series of samples (14-377 and 14-395) grown in R1 reactor, were characterized

by conventional analysis methods available in the laboratory. To determine the Al incorporation

of the grown layers, Hg-probe C-V measurement was performed. However, at high doping

concentration (≥1019

cm-3

) the C-V technique is not reliable since the Hg/SiC contact loses his

rectifying characteristic. This results in a very narrow space charge region width, that allows a

tunneling of the carriers across the barrier [9]. To acquire an indication of the [Al] in the grown

films, I-V measurements were performed.

Hg-probe I-V measurements cannot be considered as quantitative analysis of the sample,

since it combines the information about the contact and material resistivity. However, it can be

used as an indicator of C-V measurements reliability. Also, for strongly doped samples it lets a

qualitative comparison of doping level between samples. Figure 4.1 represents the I-V

characteristics obtained at room temperature (RT) by Hg-probe measurement on the samples with

various [Al]. As can be seen, the rectifying characteristics (sample 14-007) are reducing until

they are completely replaced by ohmic characteristics (sample 14-377). This obtained trend

corresponds to higher [Al] if we take into consideration as reference sample 14-007 on which the

[Al] was determined previously by SIMS analysis.

Fig. 4.1 I-V characteristics acquired from the grown epilayers with various Al concentrations

0 1 2 3 4 510

-7

10-6

10-5

10-4

10-3

10-2

10-1

100

101

-3 -2 -1 0 1 2 3

-0.20

-0.15

-0.10

-0.05

0.00

0.05

0.10

0.15

0.20

Cu

rre

nt

(A)

Voltage (V)

14-007

14-009

14-395

14-377


160

Furthermore, the I-V characteristics of the first series of samples (14-007 and 14-009) and

the second series (14-395 and 14-377) showed different behaviors even though they have been

grown under similar conditions. Considering that the two series of samples have been grown in

different periods of time, we may suggest that the effect is related to the actual growth conditions

(e.g. growth temperature, bubbler temperature/pressure, and/or susceptor position along reactor

axis). Assuming that the effect is related to the growth temperature, we may consider lower

growth temperatures (~1500-1600°C) for the second series of epilayers. This assumption can

explain the stronger [Al] obtained.

4.3. Surface morphology

The surface morphology of the Al-doped layers has been analyzed after growth by Optical

Microscopy (OM). For a more elaborated study, the layers were examined with SEM and AFM

techniques. It is important to mention that the sample series grown in R1 and R2 reactors, showed

specular surfaces even for very high [Al]. Let’s start our discussions with the first two series of

samples grown in R1 reactor on 8° off-axis 4H-SiC substrates.

From OM analysis, it was found that the layers surfaces are featureless up to an [Al] equal

to 1.7x1019

cm-3

(sample 14-007). Figure 4.2(a)-(b) represents the surface morphology images

obtained from OM and AFM analysis on sample 14-007.

Fig. 4.2 Example of surface morphology obtained by: (a) OM with 20x magnification and (b)

AFM 20x20µm2 scan image on sample 14-007


161

The surface morphology obtained and shown in fig. 4.2 for sample 14-007, describes

actually the whole series of epilayers with an [Al]≤1.7x1019

cm-3

. As can be observed from the

20x20µm2 scan image (fig. 4.2b), the AFM measurements confirm the featureless surfaces of the

film, with a roughness RMS~0.4nm.

Furthermore, the samples grown with an [Al]>1.7x1019

cm-3

(i.e. samples 14-009, 14-377

and 14-395) showed also a specular surface. The investigations performed by OM and AFM

techniques revealed a deterioration of the films surfaces, as they exhibit macroscopic step

bunching. Figure 4.3 depicts the SEM micrographs of the 4H-SiC layers grown with highly Al

concentrations. Apart from the formation of step bunching, we can observe the meandering of the

surface and a sort of macrosteps parallel to the wafer flat (i.e. parallel to [11 0] direction). As

shown in fig. 4.3(a) and (c), the macrosteps parallel to the wafer flat appear to be non-uniform in

size, which may suggest that they may be domain boundaries on the surface.

Fig. 4.3 SEM micrographs of highly Al doped epilayers grown in R1 reactor: (a) 14-009, (b) 14-

377 and (c) 14-395

The AFM investigation confirmed the observed surface modification for the series of

highly Al doped layers. The AFM images of the highly doped samples are displayed in fig. 4.4.

As shown by the 10x10µm2 scan images, the surface morphologies of sample 14-009 (fig. 4.4a)

and 14-395 (fig. 4.4c) are quite similar. It was found for the two samples a high surface

roughness with an RMS value of ~22nm and respectively, ~15.6nm. On the other hand, for

sample 14-377, i.e. with the highest [Al] (fig. 4.4b), a lower surface roughness with a RMS value

of ~7nm was obtained.


162

Fig.4.4 AFM images of the highly Al doped epilayers grown in R1 reactor: (a) 14-009

(b) 14-377 and (c) 14-395

This surface effect seems comparable with the one noticed for non-intentionally doped

samples when lower growth temperatures are applied. Figure 4.5 shows the AFM images of two

8° off-axis 4H-SiC films grown under a growth temperature of 1450°C and 1350°C, respectively.

As can be seen from the 5x5µm2 scan images, the surface morphology deteriorates and becomes

rougher with the decrease in process temperature. The roughness values exhibited by the surfaces

were found to be: a RMS~2.5nm for the layer grown at 1450°C (fig. 4.5a) and a RMS~5nm at

1350°C (fig. 4.5b), respectively.

Consequently, factors such as (low) temperature and/or (high) dopant supply can impact

on some degree the surface morphology.


163

Fig. 4.5 AFM images of non-intentionally doped epilayers grown in R1 reactor at various growth

temperatures: (a) 1450°C and (b) 1350°C

The series of layers grown on 4° off-axis 4H-SiC with various TMA flow rates in the R2

reactor (i.e. 16H4-010 and 16H4-011 samples), also displayed specular surfaces. The

investigation performed by AFM technique indicates a surface covered only by highly-oriented

step-bunching that are perpendicular to wafer flat, as displayed in fig. 4.6. However, it was

observed that increasing the [Al] from 1.5x1019

cm-3

(fig. 4.6a) to 1x1020

cm-3

(fig. 4.6b) leads to

surface roughening with RMS values from 3nm to respectively 11nm, for a 20x20µm2 scan.

Fig.4.6 AFM images of highly Al doped films grown in R2 reactor with various TMA flow rates:

(a) 16H4-010 and (b) 16H4-011


164

In this study, we observed that at high Al concentrations the surface morphology is

disturbed by step-bunching and wavy features, resulting in higher surface roughness.

Consequently, this effect can be attributed to high dopant supply introduced during the growth.

4.4. Structural characterization

The crystal structure of the grown Al-doped 4H-SiC epilayers was investigated by high-

resolution X-ray diffraction (HR-XRD) measurements. The analysis was carried out on a

PANalytical X’Pert PRO MRD diffractometer in triple axis X-ray mode. The diffractometer was

equipped with four-reflection asymmetric Ge(220) incident beam monochromator, selecting the

Cu K 1 wavelength.

4.4.1. X-ray rocking curve measurements

The HR-XRD analysis of the grown films was performed by measuring the double axis

rocking curves (RC) of the symmetric 4H-SiC(0004) reflection. For the highly Al-doped films

([Al]≥1x1019

cm-3

), the RC of the asymmetric 4H-SiC(10 8) reflection was measured in a skew

diffraction geometry [10]. The full width at half maximum (FWHM) values of the x-ray rocking

curves ( -scans) of both 4H-SiC(0004) and (10 8) reflections as a function of Al concentration

are represented in fig. 4.7. It was found from the symmetrical 4H-SiC(0004) reflection an

increasing in FWHM values from 15 to 34arc sec, with the increase of [Al] from 1.5x1018

to

1.7x1019

cm-3

. Such FWHM values indicate a good crystalline quality of the epilayers [11], [12].

Furthermore, for epilayers with an [Al]>2x1019

cm-3

a faster increase rate is observed, as the

FWHM values are expanding from 73-127arc sec. These high values points out to a degradation

of the crystal quality which seems to coincide with the surface morphology deterioration.

For the RC case of the asymmetrical 4H-SiC(10 8) reflection, the measurement was

performed only on the highly Al-doped epilayers. As can be seen from fig. 4.7 (red symbols), the

FWHM values are found to increase at a similar rate as the one seen for symmetrical reflection.

However, somewhat lower FWHM values have been obtained, i.e. between 40-91arc sec with the

increase of Al concentration.


165

Fig. 4.7 FWHM values of x-ray rocking curves as a function of Al incorporation of the

symmetric 4H-SiC(0004) and asymmetric 4H-SiC(10 8) reflection. The green symbols represent

the FWHM values of the symmetric 4H-SiC(0008) and asymmetric 4H-SiC(10 8) reflection of

the epilayer grown in R2 reactor

The RC analysis was performed also for the sample with the highest [Al]=5x1019

cm-3

(sample 16H4-011) grown in R2 reactor. The obtained results are represented also in fig. 4.7 by

green half closed symbols. For this study, the 4H-SiC(0008) symmetrical reflection was chosen

since it gives more information (e.g. peak separation) due to a deeper exploration of the structure.

Hence, the FWHM value obtained from the RC analysis on symmetrical 4H-SiC(0008)

reflection was found to be equal to 38arc sec which indicates high crystalline quality. On the

asymmetrical 4H-SiC(10 8) reflection, a FWHM value of 67arc sec was found.

These results may be an indication that stronger [Al] can result in a degradation of the structural

quality as previously reported in ref. [6].

4.4.2. Determination of lattice constants

The lattice constants of the grown films had been evaluated from 2 - scans of

symmetrical 4H-SiC(0004), (0008) and (00012) reflections, and asymmetrical 4H-SiC(10 2),

1018

1019

1020

0

20

40

60

80

100

120

140 (0004)

(10-18)

R2

FW

HM

of

XR

C (

arc

se

c.)

[Al] (cm-3)


166

(10 3), (10 5) reflections. For the epilayers with an [Al]≥5x1019

cm-3

grown in R1 reactor, the

corresponding 2 - scans of the symmetrical reflections revealed the presence of peak splitting,

i.e. a second peak emerged. Accordingly, the -lattice parameter of the first peak ( ) and of the

second peak ( ) were derived from the 2 - positions of symmetrical reflections scan by using

Bragg’s law formula for hexagonal symmetry [10]:

(4.1)

(4.2)

The -lattice parameter of the epilayers were estimated from 2 - positions of

asymmetrical 4H-SiC(10 2), (10 3), (10 5) reflections using the mean value of -lattice

parameters determined for each grown sample.

A theoretical model was proposed by Jacobson et al. [13] to predict the change in lattice

constants when the dopant incorporation is varied. This model assumes that the dopant atoms are

substituting at the host atoms sites. Since the size, i.e. the covalent radius of the dopant and

host atoms are different, a change of the volume is induced. This change is translated to a

relative change of lattice constants that are dependent on dopant concentration and the

substitution position in the SiC lattice. In the SiC matrix, the Al atoms with a covalent radius of

0.125nm are substituting at Si atoms sites, which have a covalent radius of 0.11nm [14]. Hence,

an increase of [Al] should result to an expansion of the 4H-SiC lattice. In contrast, the N atoms

with a covalent radius of 0.065nm are substituting at the C atoms sites with a covalent radius of

0.07nm [14] and thus, should result to a compression of the 4H-SiC lattice.

The relative change of lattice constants due to dopant incorporation can be calculated as

shown by Jacobson et al. [13] as follows:

(4.3)


167

where rs and rh are the covalent radii of the substitutional (Al) and host (Si) atoms, is the

substitutional dopant concentration, =225GPa [13] represents the bulk modulus, and =167GPa

[13] is the shear modulus, respectively.

The model prediction regarding the lattice parameters change and the experimental results

are shown in fig. 4.8 as a function of Al concentration. The Jacobson’s model predicts an increase

of lattice constants for [Al] slightly above 1x1019

cm-3

compared to the lattice constants of the n+

4H-SiC substrates. It’s important to point out that the films were grown on four n+ 8° off-axis

4H-SiC substrates with slightly different [N] which results to a slight variation of -lattice

parameter.

In fig. 4.8(a) the experimental results represents the -lattice parameters values of the

epilayer ( ) and peak splitting ( ) obtained from the two series of samples 14-003 to 14-395

grown in R1 reactor. The value of relaxed -lattice parameter of the n+ 8° off-axis 4H-SiC

substrate is displayed as a dotted black line in the graph. It was found that epilayers with [Al] in

the range 9x1017

-1.7x1019

cm-3

show no peak splitting. Accordingly, a constant -lattice

parameter with a mean value =1.0083nm was calculated. The single peak indicates an identical

-lattice parameter of the substrate and the strained epilayer. For the epilayers with

[Al]≥5x1019

cm-3

where the peak splitting occurred, two -lattice parameter values were

calculated, i.e. an attributed to the substrate and a ascribed to the strain epilayer. It was

found a constant -lattice parameter value assigned to the substrate ( ) with an increase of [Al].

Contrary to the theoretical prediction of Jacobson’s model, the -lattice parameter of the

strain epilayers ( ) was found to decrease with the increase in [Al]. Therefore, our results are in

contradiction to Jacobson’s model and previous results reported in literature [6], [11], [15], which

are all indicating larger lattice constants with stronger [Al].

The calculated -lattice parameters of the epilayers are plotted in fig. 4.8(b). For the films

grown with [Al] in the range 9x1017

-1x1020

cm-3

, the 2 - scans of the asymmetrical reflections

showed single peaks. It can be seen that the calculated -lattice parameters of the epilayers ( )

are independent of the [Al] of the epilayers. For comparison, the -lattice parameter value of

relaxed n+ 8° off-axis 4H-SiC substrate is displayed as a dotted line in the graph. Moreover, the

-lattice parameter of the highly Al doped epilayer grown in R2 reactor ([Al]~5x1019

cm-3

) was

also calculated from the peak recorded for symmetrical 4H-SiC(0008) reflection using eq. 4.1 and

4.2.


168

Fig. 4.8 Experimental and the Jacobson’s model prediction of lattice constants change as

a function of Al concentration. (a) -lattice parameter of epilayer ( ) and peak splitting of the

epilayer ( ). The -lattice constant of the highly doped epilayer grown in R2 reactor is

displayed as green closed symbol. (b) -lattice parameter of epilayer ( ). The value of relaxed

- and -lattice constants of n+ 4H-SiC substrates is displayed as black dotted line.

1018

1019

1020

1.004

1.006

1.008

1.010

1.012

1.014

c-l

att

ice

pa

ram

ter

(nm

)

[Al] (cm-3)

Jacobson's modelepilayer c

e

peak split ce1

a)

R2

35.72

35.68

35.64

35.60

35.56

35.52

35.48

35.44

35.40

2 (°)

1018

1019

1020

0.304

0.306

0.308

0.310

0.312

0.314

b)

a-l

att

ice

pa

ram

ter

(nm

)

[Al] (cm-3)

Jacobson's model epilayer a

e

35.72

35.68

35.64

35.60

35.56

35.52

35.48

35.44

35.40

2 (°)


169

The obtained value of the -lattice parameter is represented in fig. 4.8(a) by the green

triangle symbols. No peak splitting was observed from the 2 - scan, suggesting an equal -

lattice parameter between substrate and epilayer. Consequently, the corresponding

[Al]~5x1019

cm-3

doesn’t induce any lattice constant change as predicted by Jacobson’s model.

This suggests that a much stronger [Al] in the epilayer than the one obtained in our reactors is

necessary to induce a lattice parameter change.

4.4.3 Reciprocal space mapping (RSM)

To explore further the observed peak splitting for highly doped films, a more precise

measurement has been performed, i.e. the Reciprocal Space Mapping (RSM). The RSMs have

been measured on the samples grown with an [Al]≥1.7x1019

cm-3

in R1 reactor (i.e. sample 14-

007, 14-009, 14-377 and 14-395). To determine any lattice deformation and lattice constant

change, RSMs of symmetrical 4H-SiC(0008) and asymmetric (10 8) reflections were measured

by monitoring a series of -2 scan at consecutive values in triple axis geometry. As

mentioned previously in section 2.3.3 of Chapter 2, a compensation of the 8° off-angle substrate

is required to obtain the diffraction. In consequence, the compensation was adjusted as a function

of the azimuth angle by adapting the omega ( ) and chi ( ) angles (see chapter 2).

The RSMs were measured at angle of 0° and 90° as a function of peak intensities.

Therefore, for 4H-SiC(0008) reflection the diffraction was acquired at set angles 2 ~75.4° and

~46° at an azimuth angle =0°. In the case of an azimuth angle =90°, the 8° off-cut

compensation was made by angle, with a diffraction acquired at angles 2 ~75.3° and ~37.6°,

respectively. In the case of asymmetric (10 8) reflection, the measurements were performed at

=90° at corresponding diffraction angles 2 ~85° and respectively, ~17.2°.

The RSMs of symmetrical (0008) and asymmetric (10 8) reflections registered from the

various Al-doped epilayers ([Al]=1.7x1019

-1x1020

cm-3

) are displayed in fig. 4.9. The 2 and

coordinates of the RSMs were converted to Cartesian and coordinate, where they are

parallel to the direction [11 0] and [0001], respectively. The conversion of 2 and into

reciprocal lattice units having the wave vector = , and can be expressed as [10]:

(4.4)


170

(4.5)

where the is the x-ray wavelength Cu K 1=1.54056

Fig. 4.9 Reciprocal space maps measured on symmetrical (0008) from (a)-(d) and asymmetric

(10 8) reflections from (a’)-(d’) of epilayers with Al concentration of (a, a’) 1.7x1019

cm-3

, (b, b’)

7.8x1019

cm-3

, (c) ~6.5x1019

cm-3

and (d, d’) ~1x1020

cm-3

. The dotted line marks the n+ 4H-SiC

substrate diffraction peaks. Red: high intensity. Blue: low intensity


171

It can be observed from fig. 4.9(b-d), that 4H-SiC(0008) reflection reveals the presence of

two separated diffraction peaks along axis. Of the two peaks, the position of the diffraction

peak from n+ 4H-SiC substrate was identified at smaller 2 angles (marked by dotted line)

compared to the epilayer position. The substrate and epilayer peaks were identified based on the

relative peak intensities measured from symmetrical 4H-SiC(0004), (0008), and (00012)

reflections.

Contrarily, a single well-shaped peak was found from the asymmetric (10 8) reflection

regardless of the [Al] in the epilayer, as illustrated by fig. 4.9(a’-d’). Considering that the

diameter of Al atom is larger than that of host Si atom, an increase of lattice parameter should

have been observed for high Al concentrations. This was previously described by Jacobson’s

model [13]. Therefore, a shift of the epilayer peak towards smaller 2 angles was expected. Yet,

our results are in opposition with the theoretical prediction of Jacobson’s model [13] and the

results reported by various research groups [6], [11], [15] which may suggest the presence e.g. of

cubic polytype inclusions.

On the other hand, RSM of symmetrical 4H-SiC(0008) reflection was also carried out on

sample 16H4-011 grown R2 reactor. The measured RSM is illustrated in fig. 4.10. As can be

seen, only a single well-shaped peak was observed.

Fig. 4.10 Reciprocal space map of symmetrical 4H-SiC(0008) reflection measured on the

sample grown in R2 reactor with an [Al]~5x1019

cm-3

. Dotted line marks the on-axis 4H-SiC

substrate diffraction peak. Red: high intensity. Blue: low intensity.


172

As mentioned previously, the [Al] in the epilayer is not sufficient to induce a change in

the lattice parameters between the substrate and the epilayer since the peak positions are

identical.

4.4.4 Polytype analysis

Since the epilayer peak position observed from the RSMs of 4H-SiC(0008) reflection

doesn’t correspond to the one calculated from Jacobson’s model [13], the study was further

extended to understand the origin of this second peak assigned to the epilayer.

In the SiC growth, it is known that the 3C-SiC can nucleate spontaneously on hexagonal

SiC substrates with various off-cut angles when the growth conditions are not properly chosen,

i.e. lower growth temperature and/or high growth rate [16]. Therefore, to determine whether the

epilayer peak can be associated to the 3C-SiC polytype presence, RSMs measurements were

performed on a heteroepitaxial structure. The structures is formed from an n-type 3C-SiC

epilayer grown on an on-axis n+ 4H-SiC substrate.

Since no compensation of the off-cut is necessary, the RSM of the symmetrical 4H-SiC(0008)

reflection was performed at diffraction angles 2 ~75.23° and ~37.61°, respectively. The

measured RSM of 4H-SiC(0008) reflection is displayed in fig. 4.11. It can be noticed the

presence of two distinct well-shaped peaks, where the peak position of the epilayer is located at

higher 2 angle and assigned to the 3C-SiC(222) reflection. It was found that the 4H-SiC(0008)

and 3C-SiC(222) reflections are separated along the 2 angle by ~440arc sec. The shift of the

epilayer peak towards higher angles corresponds to a decrease in -lattice parameter, which is in

agreement with the results obtained for highly Al doped layers. Converting the 2 angular

positions to interplanar spacing based on eq. 4.1, it was found an out of plane mismatch c/c~ -

1.382x10-3

(i.e. > ). Similar results were reported in ref. [17], [18], where it was suggested

a partially relaxed 3C-SiC epilayer.

Therefore, we may consider that our highly Al doped epilayers grown in R1 reactor are

formed at least partially from 3C-SiC(111) polytype.


173

Fig. 4.11 Reciprocal space map of symmetrical (0008) reflection measured from 3C-

SiC/4H-SiC heterostructure. Dotted line marks the on-axis 4H-SiC substrate diffraction peak.

Red: high intensity. Blue: low intensity.

Since the crystal structure of the substrate and epilayer are different, it is necessary to

consider them independently in order to calculate the -lattice constant. For comparison

purposes, the cubic lattice can be converted to hexagonal lattice [19], i.e. the 3C-SiC(222) cubic

index is converted to 3C-SiC(0006) hexagonal index in order to calculate lattice constants.

Dudley et al. [17] proposed a multiple order reflection method in order to determine the -lattice

constants. However, in our study we decide to compare the 4H-SiC substrate and 3C-SiC epilayer

at the level of bilayer thickness i.e. at internal planes with ratio where is the number of

bilayers per unit cell. The calculated bilayer thickness or derived -lattice parameter of both

substrate and epilayer as a function of [Al] are represented in fig. 4.12. It can be seen that the

bilayer thickness of the substrates is slightly larger than the one of the epilayer .

Moreover, the values of remains constant with the increase of [Al], while a slight variation is

noticed for . Considering that the nucleation of the 3C-SiC polytype is related to the growth

conditions, while the lattice parameter changes due to the dopant atom size, we suggest that the

observed trend for 3C-SiC epilayers is not entirely influenced by the Al concentration. In the case

of 4H-SiC substrate the constant behavior was expected, since the [Al] should not influence

its lattice structure.


174

Fig. 4.12 Bilayer thickness of the 4H-SiC substrate and 3C-SiC epilayer calculated from

the measured RSMs on symmetrical 4H-SiC(0008) reflections

4.4.4.1 Discussions: 3C-SiC polytype

If we suppose a sporadic presence of 3C-SiC polytype (i.e. local domains or inclusions) in

the highly Al doped films, it is important to identify its exact presence. As discussed in chapter 2,

section 2.3.3, the 3C-SiC113 reflections were chosen as candidates to quantify the 3C-SiC

polytype. Since the 4H-SiC planes are not parallel to 3C-SiC planes (except for 4H-SiC(0001)

and 3C-SiC(111)), the diffraction of the specific peaks will take place at different , and

angles. Thus, the presence/absence of the 3C-SiC(113) peak is an unequivocal indication of the

presence/ absence of cubic polytype in the layer. Because the Al doped layers were grown on

vicinal 4H-SiC substrates, the 8°/4° off-cut was compensated by adjusting the and angles.

Furthermore, to avoid any misinterpretation due to any incorrect settings, the compensation

adjustment was validated by a detection of the 4H-SiC108 substrate related peak at similar

azimuth angle.

1019

1020

1021

2.517

2.518

2.519

2.520

2.521

2.522

3C-SiC

De

riv

ed

c-l

att

ice

pa

ram

ete

r (Å

)

[Al] (cm-3)

4H-SiC


175

Most of the 3C-SiC films grown on -SiC substrates present crystallographic defects such

as stacking faults (SFs) and/or the so called double positioning boundaries (DPBs) owing to a 60°

rotation of the 3C-SiC stacking sequence [20], [21].

In order to evaluate the structural properties of the 3C-SiC polytype, 2 - scans of the

asymmetrical 3C-SiC(113) reflection were performed on four highly Al doped films grown in R1

reactor (i.e. sample 14-007, 14-009, 14-377 and 14-395). The diffraction angle was set at

2 ~71.8 for asymmetrical 3C-SiC(113) reflection, while the value was adjusted to compensate

the 8° off-cut angle. Because the diffractometer used in the study allows the compensation by an

adequate disorientation of the sample holder only up to 4° off-cut angles, the analysis of

diffraction pole figure to determine the DPBs presence was not possible. Nevertheless, to

determine the presence/absence of DPBs in the grown epilayers, 2 - scans were performed at

different angle, i.e. the sample was rotated by 60° around its symmetry axis. Figure 4.13

displays a representative example of 2 - scan of 3C-SiC(113) reflection at various angle

obtained on sample 14-377 (with [Al]~1x1020

cm-3

).

Fig. 4.13 2 - scan of the asymmetrical 3C-SiC (113) reflection measured at various angle on

sample 14-377 (with an [Al]~1x1020

cm-3

)

69.5 70.0 70.5 71.0 71.5 72.0 72.5 73.0 73.50

40000

80000

120000

160000

3C-II

Inte

ns

ity

(c

ou

nts

)

2 scan (°)

3C(113) =30°

3C(113) =90°

3C-I


176

It can be seen that the presence of 3C-SiC polytype in the epilayer is formed by twin 3C-

SiC domains labeled as 3C-I (at =30°) and 3C-II (at =90°). The presence of twin 3C-SiC

domains were observed on all four investigated epilayers with an [Al]≥1.7x1019

cm-3

.

For each sample, the FWHM values and integrated peak area from both 3C-SiC(113)

peak registered at angle 30° and respectively 90°, were obtained using a Pseudo-Voigt function

for the fit. This analysis was performed in order to check the existence of any correlation between

the presence of 3C-SiC domains in the epilayer and the Al concentration. The FWHM values and

integrated areas obtained from the twin 3C-SiC(113) domains as a function of epilayer Al

concentration are represented in fig. 4.14(a) and (b), respectively.

From fig. 4.14(a) it can be observed that for the epilayer with an [Al]=1.7x1019

cm-3

(sample 14-007), a similar FWHM value was found for both 3C-SiC domains. In the case of

epilayers with stronger [Al], a slightly higher FWHM value was found for 3C-II domain

compared with 3C-I domain. Moreover, the FWHM values of both 3C-SiC domains tend to

decrease with the increase in [Al]. The observed behavior may indicate an improvement of the

3C-SiC crystal properties with the increase in [Al]. However, we cannot exclude any influence

coming from wafer curvature or growth conditions on the obtained FWHM values of both 3C-

SiC domains.

Apart from the FWHM, the integrated peak areas of both 3C-SiC domains were analyzed

as a function of [Al] as displayed in fig. 4.14(b). As it can be seen, we found identical peak areas

of both 3C-SiC domains on the epilayer with an [Al] =1.7x1019

cm-3

(sample 14-007). However,

when the [Al] is increased, the area of 3C-II domain is also increasing, while the area of 3C-I

domain remains almost constant. The detected behavior cannot be assigned entirely to [Al]

dependence, but at least partially to the 3C-SiC nucleation which depends strongly on growth

conditions. It was reported [21], [22] that during growth, the 3C-I domain outcompetes the other

3C-II domain, since the 3C-SiC nuclei with one orientation expands faster than the one rotated by

60°.


177

Fig. 4.14 Analysis of the 3C-SiC(113) peaks: (a) FWHM and (b) integrated peak area as a

function of Al concentration

1019

1020

0.0

0.2

0.4

0.6

0.8

1.0

1.2

1.4

1.6

FW

HM

(°)

[Al] (cm-3)

3C-I =30°

3C-II =90°

a)

1019

1020

0

20000

40000

60000

80000

100000

120000

140000

Pe

ak

Are

a (

a.u

)

[Al] (cm-3)

3C-I =30°

3C-II =90°

b)


178

Considering that the highly Al doped ([Al] ≥1x1019

cm-3

) samples grown in R1 reactor are

partially formed by 3C-SiC polytype, a similar investigation was performed for the highly Al

doped sample (16H4-011) grown in R2 reactor. As mentioned, to avoid any misinterpretation, the

compensation adjustment of the 4° off-cut angle was validated by the detection of the 4H-

SiC(10 8) peak at an azimuth angle =90°. The scan of the asymmetrical 3C-SiC(113)

reflection recorded on sample 16H4-011 (with [Al]~5x1019

cm-3

) is shown in fig. 4.15.

Fig. 4.15 scan of the asymmetrical 3C-SiC(113) reflection measured at =90° on

sample16H4-011 (with an [Al]~5x1019

cm-3

) grown in R2 reactor

As can be seen, no 3C-SiC(113) peak was found which indicates the absence of any cubic

polytype inclusions in the grown epilayer. Therefore, the study revealed that highly Al doped

films grown in R1 reactor ([Al]≥1.7x1019

cm-3

) are composed, at least partially, from double 3C-

SiC domains. Contrarily, the presence of 3C-SiC was not detected for the epilayer grown in R2

reactor with an [Al]~5x1019

cm-3

. We suggest that the detected differences are associated to the

applied growth conditions during the process of high Al doped layers. Moreover, the growth

conditions in R2 reactor appear to be more appropriate to grow single crystalline 4H-SiC layers

with high [Al]. On the down side, we have a limitation in [Al] up to ~5x1019

cm-3

of epilayers

when such growth conditions are applied. However, the maximum obtained [Al] in the layer

4 5 6 7 80

4

8

12

16

20

24

Inte

ns

ity

(c

ou

nts

)

scan (°)

3C(113) =90°


179

doesn’t induce any visible lattice constant change between the substrate and layer as already

detected by the XRD study.

4.5 Physical properties of Al doped epilayers

4.5.1 Optical properties

4.5.1.1. Low Temperature Photoluminescence (LTPL) measurements

LTPL measurements were carried out to evaluate the influence of Al concentration on

optical properties on the first series of 4H-SiC epilayers grown in R1 reactor (see table 4.1). The

experimental set-up for LTPL technique used for collecting the data was described previously in

Chapter 2. The LTPL measurements were performed at temperature of ~4K in the range 380-

750nm at Laboratoire Charles Coulomb in Montpellier. A detailed investigation of the evaluation

of the Al concentration in the grown epilayers by LTPL and Raman techniques, are described in

detail in the PhD manuscript of Pawel Kwasnicki [23].

In fig. 4.16, the LTPL spectra collected on the samples with different Al concentrations

([Al]=3x1016

to 1.7x1019

cm-3

) are illustrated. In the Near Band Edge (NBE) region of the LTPL

spectra of 4H-SiC shown in fig. 4.16(b), we notice the presence of the Al-bound exciton ( )

peak. With the increase of [Al], it was found that the peak intensity changes, also is

broadening and shifts towards lower energies.

In the range 390-450nm on the LTPL spectra (see fig. 4.16a) the presence of a broad band

assign to N-Al Donor Acceptor Pair (DAP) can be identified. The intensity of the N-Al DAP

band is increasing when the [Al]≥6.8x1018

cm-3

, i.e. the epilayer compensation becomes stronger.

These effects have been previously reported for both aluminum and nitrogen (for N-bound

excitons) doped layers [24]–[27].


180

Fig. 4.16 Comparison of LTPL spectra collected on 4H-SiC samples grown with various

Al concentrations in the range: (a) 380-480nm and (b) 380-396nm of the NBE region

380 400 420 440 460 480

Inte

ns

ity

(c

ou

nts

)

Wavelength (nm)a)

3.2 3.1 3.0 2.9 2.8 2.7 2.6

[Al]=1.7x1019

cm-3

[Al]=2.3x1016

cm-3

[Al]=1.4x1017

cm-3

[Al]=9.1x1017

cm-3

[Al]=6.8x1018

cm-3

Energy (eV)

NBEN-Al DAP

Reference

380 384 388 392 396

Reference

[Al]=1.7x1019

cm-3

[Al]=2.3x1016

cm-3

[Al]=1.4x1017

cm-3

[Al]=9.1x1017

cm-3

[Al]=6.8x1018

cm-3

Inte

ns

ity

(c

ou

nts

)

Wavelength (nm)b)

3.26 3.24 3.22 3.20 3.18 3.16 3.14

Al0

Al0

Al0

P0Q0

Energy (eV)


181

The LTPL spectrum collected on the strongly Al doped sample 14-009 (with

[Al]=7.8x1019

cm-3

) is shown in fig. 4.17. It can be seen that the collected LTPL spectra differs

significantly compared with the previous ones. A single broad band situated at around 600nm can

be noticed.

Fig. 4.17 LTPL spectrum collected on strongly Al doped 4H-SiC sample with the insert of an n-

type 3C-SiC sample as reference

By comparing the collected spectrum with the one of an n-type 3C-SiC sample, we find at

similar position the presence of a broad band assign to N-Al DAP that indicates a fairly

compensated layer. Therefore, we may associate the broad band of the strongly Al doped epilayer

to an N-Al DAP feature of the 3C-SiC polytype. This observation is in agreement with our XRD

measurements which exposed the presence of double 3C-SiC domains in the epilayer.

It has been reported that strong [Al] results to saturation of the N-Al DAP band due to the

consumption of the available nitrogen atoms [27]. Consequently, we suggest that the band

distortion of the N-Al DAP feature originates from the high Al content. However, we cannot

exclude the influence of growth conditions and crystallographic defects. Furthermore, in the

350 400 450 500 550 600 650 700 750

Inte

ns

ity

(a

.u)

Wavelength (nm)

N-Al DAP[Al]=7.8x10

19cm

-3

Reference 3C-SiC

3.4 3.2 3.0 2.8 2.6 2.4 2.2 2.0 1.8

Energy (eV)


182

range 380-480nm, no clear signature of the 4H-SiC polytype is visible though, we may

distinguish the presence of some features corresponding to 4H-SiC DAP.

4.5.2 Electrical properties

4.5.2.1 Hall Effect measurements

The electrical properties of Al-doped 4H-SiC films were characterized by Hall effect

measurements, to determine the dependence of carrier concentrations and mobility on

temperature and Al concentration. In this study, the ionization energy and resistivity of the Al-

doped layers have been also determined.

The measurements were performed on four samples (i.e. samples 13-331, 14-003, 14-004

and 14-007) from the first series grown in R1 reactor with an [Al] in the range 1.4x1017

to

1.7x1019

cm-3

(see table 4.1). For the three lowly doped epilayers (i.e. with [Al] in the range

1.4x1017

to 6.8x1018

cm-3

) an additional ~200nm thick p+-layer with a [Al]~5x10

19cm

-3 was grown

on top of the epilayers to obtain ohmic contacts. For creating the ohmic contacts, Ti/Al was

evaporated on the sample surfaces in a van der Pauw configuration, followed by a flash annealing

at 700°C under N2 atmosphere. For correct measurements, the p+-layer was removed by Reactive

Ion Etching (RIE) as shown in fig. 4.18, which illustrates the schematic representation of the

samples.

a)

b)

Fig. 4.18 Schematic representation of the sample for Hall effect measurements in van der

Pauw configuration as: (a) cross-section and (b) top view

As mentioned in Chapter 2, the Hall effect measurements were performed at Laboratoire

Charles Coulomb (L2C) in Montpellier under the guidance of Dr. Sylvie Contreras and Dr.


183

Leszek Konczewicz, and at the Institute of Microelectronics and Microsystems (CNR-IMM) in

Catania under the guidance of Dr. Fabrizio Roccaforte and Dr. Marilena Vivona. These

measurements were realized during various shared trainings done in the framework of NetFiSiC

project. The experimental set-up for Hall Effect measurement available in Montpellier was

described previously in Chapter 2. The measurements were performed in a van der Pauw

configuration in the temperature range 297 up to 900K and a magnetic field B up to 1T. At CNR-

IMM in Catania, the resistivity and Hall carrier concentrations were obtained in van der Pauw

configuration in the temperature rage 300-450K under an magnetic field B=0.1T. For this study,

also Ti/Al contacts were evaporated on the sample surfaces (see fig.4.18) and followed similar

steps as described above to obtain the ohmic characteristic.

Temperature dependence of Hall carrier concentration

Figure 4.19 displays the Hall carrier concentration ( ) as a function of the reciprocal

temperature for the four studied samples, both at Montpellier (closed symbols) and Catania (open

symbols). The dashed lines are represented only as guidance for the eyes.

Fig. 4.19 Hall carrier concentration vs. reciprocal temperature for samples with different

Al concentrations. Closed symbols: Montpellier work, open symbols: Catania work

1 2 3 410

16

1017

1018

1019

1020

1021

Ha

ll c

on

ce

ntr

ati

on

(c

m-3)

1000/T (K-1)

13-331

14-003

14-004

14-007


184

By plotting Hall resistivity as a function of magnetic field B as shown in fig. 4.20, the

slope can be determined and used to calculate the Hall carrier concentrations (see eq. 2.28 in

Chapter 2). The Hall scattering factor ( ) was assumed to be equal to unity, although some

studies have reported a variation of the Hall scattering factor with the temperature and doping

concentration [28], [29].

Fig. 4.20 Experimental Hall resistivity as a function of magnetic field measured at various

temperatures on sample 14-003. Empty symbols display the values at 297K (left scale) and filled

symbols the values at 600K (right scale)

It can be seen from fig. 4.19 that the Hall concentration is increasing with the increase in

temperature. In the case of sample 14-007 ([Al]=1.7x1019

cm-3

), a saturation of the is observed

for higher temperatures. This behavior can be associated to an excitation of all the hole carriers

resulting to a filled acceptor level in the valence band . Regarding the Hall measurements

performed at Catania, it was found that the investigated samples (13-331 and 14-007) have

similar dependence as the ones in Montpellier, i.e. the increases with temperature. Lower

were obtained though, for sample 13-331 ([Al]=1.4x1017

cm-3

) investigated at Catania. The visible

difference between the two Hall measurements may be associated with factors such as

measurement equipment and/or deposition of the contacts between the two laboratories.

-1.0 -0.5 0.0 0.5 1.0-4

-2

0

2

4

T=297K

H

all

re

sis

tiv

ity

(

)

Magnetic field (Tesla)

T=600K

-0.2

-0.1

0.0

0.1

0.2


185

Regarding the Hall measurements performed at Montpellier, the three lowly doped

samples (i.e. 13-331, 14-003 and 14-004) show higher above 350K than the [Al] determined

by SIMS. The magnitude of the values at higher temperature measurements seems to decrease

when the [Al] in the epilayer is stronger. We can make a first interpretation of this Hall carrier

concentration overestimation as being affiliated to an inability of the PN junction isolation, which

gives rise to current leakage through the substrate [30]. Therefore, the resistance of the epilayer

needs to be compared with the one of the substrate at different temperatures when a PN junction

is used for isolation [31]. Schöner et al. [32] showed that applying a two layer Hall effect model,

which takes into consideration the electrical properties of both the epilayer and substrate, an

overestimation of the free carrier density values can be avoided. Consequently, we need to

consider the role of the n-type substrate in the Hall measurements when a temperature variation is

applied. Another explanation for such a behavior can be the related to the Hall scattering factor

which is strongly temperature dependent (i.e. not equal 1) and hard to determine for the

holes in SiC material [33]. Since we dealt with the overestimation, another series of films with

various Al concentrations have been grown simultaneously on both n-type and semi-insulating

4H-SiC substrates. These series were grown in order to understand more clearly the influence of

the substrate and the on the Hall concentration. The results of the electrical transport obtained

from this study have been presented at ECSCRM 2016 conference and also will be published in

the Phys. Status Solidi A journal.

Ionization energy of acceptors

From the Arrhenius plot described by fig. 4.19, ionization energies can be estimated using

the slope of the curves. The estimated mean ionization energies of the samples as a function of

Al concentration is represented in fig. 4.21. A decrease of ionization energy with the increase

of [Al] can be seen. In consequence, the decreases from 211meV (for [Al]=1.4x1017

cm-3

) to

180meV (for [Al]=1.7x1019

cm-3

). The values obtained in this study are in agreement with the

values reported by different groups [34]–[36]. Such values of indicates the presence of a

shallow acceptor level. In the case of sample 13-331 measured at Catania, it was estimated at

slightly lower of ~125meV ([Al]=1.4x1017

cm-3

) compared to ~211meV obtained from the

Montpellier measurements in the temperature range 300-450K. On the other hand, for sample 14-

007 ([Al]=1.7x1019

cm-3

) comparable ~180meV were found for both measurements.


186

The ionization energy dependence on [Al] can be associated to the average distance

between the Al acceptors and the various scattering mechanisms. Consequently, the dependence

can be described by the following relation [34], [37]:

(4.6)

where the is the ionization energy for infinite dilution, is a proportion constant and is

the acceptor concentration. Fitting the data with eq. 4.6, we determine the parameters as being

equal to =230 10meV and =2.8 0.3x10-5

meV.cm-1

. In the graph, this dependence is

represented by the solid line.

Fig. 4.21 Estimated mean ionization energies as a function of Al concentration in 4H-SiC

samples. Solid line was obtained using eq. 4.6. Experimental values of Koizumi et al. [34],

Matsuura et al. [35] and Pernot et al. [36] are also shown.

Furthermore, the ionization ratio / at room temperature (~300K) was determined and

plotted as a function of Al concentration in fig. 4.22. It can be seen, that increasing the Al

concentration of the epilayers leads to a reduction of the ionization ratio from 0.4 to 0.1. The

1014

1015

1016

1017

1018

1019

1020

1021

0

50

100

150

200

250

300

Ion

iza

tio

n e

ne

rgy

(m

eV

)

[Al] (cm-3)

Montpellier work

Catania work

Koizumi et al.

Matsuura et al.

Pernot et al.


187

obtained values of / indicates that most part of the incorporated Al acceptors correspond to

neutral impurities in the layer. The obtained dependence is in agreement with previous results

reported in literature [34] where the effect was associated with the displacement of Fermi level

towards the maximum of the valence band as the neutral impurities are increased. Moreover,

the ionization energy is dependent on the compensation donor density and the activation energy.

Consequently, slightly lower ionization energies were found for the samples measured at Catania

which translated also in smaller ionization ratio in the range 0.16-0.075.

Fig. 4.22 Ionization ratio / as a function of Al concentration at ~300K. Dotted line is

only a guide for the eyes

Resistivity of Al doped epilayers

The resistivity of Al doped samples has been determined during the Hall measurements

as described by eq. 2.29 in Chapter 2. Figure 4.23 shows the dependence of: (a) the resistivity at

room temperature on Al concentration and (b) the resistivity as a function of temperature. As it is

shown in fig. 4.23(a), the resistivity of the sample measured at RT decreases with the increase of

[Al]. For the epilayer with an [Al]=1.4x1017

cm-3

(sample 13-331), a resistivity value of 2.3 cm

was found from both measurements performed at Montpellier and Catania. In the case of the

epilayer with an [Al]=1.7x1019

cm-3

(sample 14-007), a resistivity as low as 0.25 cm was found.

1016

1017

1018

1019

1020

0.1

1

Ion

iza

tio

n r

ati

o (

p N

a-1

)

Na (cm

-3)

Montpellier work

Catania work


188

Form the measurements performed at Catania on a similar sample, slightly lower

resistivity ~0.16 cm was found. Ji et al. [6] reported similar dependence of resistivity on Al

concentration, with a resistivity as low as ~0.017 cm for an epilayer with a [Al]~3.5x1020

cm-3

.

On the other hand, the resistivity of the Al implanted layers reported in ref. [38], [39] are higher

for [Al]>1020

cm-3

and lower for [Al]>1015

cm-3

, indicating a dependence on implantation

temperature and annealing.

From fig. 4.23(b), it can be noticed that the resistivity is decreasing with the increase in

temperature in the range 300-800K. Similar temperature dependence was previously reported by

various research groups [37], [38].

Fig. 4.23 (a) Resistivity obtained at room temperature as a function of Al concentration

and (b) resistivity vs. temperature. Data from literature [6], [37], [38] are shown for both epitaxial

and implantation layers are also displayed (open symbols).

Hall mobility

The Hall mobility was determined from the measured Hall hole density and

resistivity (see eq. 2.30 in Chapter 2). Figure 4.24 shows the Hall mobility as a function of (a)

temperature and (b) Al concentration for different temperatures (in the range 300-500K). A rapid

1015

1016

1017

1018

1019

1020

1021

10-2

10-1

100

101

102

Re

sis

tiv

ity

(

cm

)

[Al] (cm-3)

Montpellier work

Catania work

Ji et al.

Heera et al.

Negoro et al.

a)

300 400 500 600 700 80010

-2

10-1

100

101

b)

Re

sis

tiv

ity

(

cm

)

Temperature (K)

13-331

14-003

14-004

14-007


189

decrease in Hall mobility with the increase of the temperature from 300K to 800K can be

observed from fig. 4.24(a) for the three lightly doped samples (13-331, 14-003 and 14-004). Yet,

for the highly doped sample 14-007, the Hall mobility shows a lower dependence on temperature.

It was found a temperature dependence of ~T-3.6

on mobility for the lightly doped samples

measured both, at Montpellier and Catania. A comparable dependence was reported in ref. [34]

for temperatures above 250K and epilayers with an [Al]≥1.8x1017

cm-3

. Consequently, the rapid

decrease in can be assigned to an increase of the phonon-lattice scattering rate with

temperature and Al concentration [29], [36] such as, acoustic phonon scattering and non-polar

optical phonon scattering [34].

Fig. 4.24 Hall mobility as a function of (a) the temperature and (b) Al concentration in the

temperature range 300-500K. Closed symbols: Montpellier work. Open symbols: Catania work.

Figure 4.24(b) presents the Hall mobility as a function of [Al] for more restricted range of

temperatures. It has been found that at RT the Hall mobility decreases from ~50cm2/Vs to

~14cm2/Vs as the [Al] increases. Moreover, when the temperature is increased up to 500K the

reaches the minimum values of ~6.4 cm2/Vs for [Al] of 1.4x10

17cm

-3 (sample 13-331) and ~1.5

400 600 800 10000.1

1

10

100

Ha

ll m

ob

ilit

y (

cm

2/V

s)

Temperature (K)

13-331

14-003

14-004

14-007

a)

1017

1018

1019

1

10

100

500 K

450 K

400 K

350 K

300 K

Ha

ll m

ob

ilit

y (

cm

2/V

s)

[Al] (cm-3)

Montpellier

Catania

b)


190

cm2/Vs for [Al] of 1.7x10

19cm

-3 (sample 14-007). Comparable results were reported by Kimoto et

al. [16] for the Al-doped 6H-SiC epilayers.

For the two samples measured at Catania, higher values were observed. Hence, at RT

the Hall mobility was found to be ~90cm2/Vs for sample 13-331 and ~30cm

2/Vs for sample

14-007. In literature Koizumi et al. [34] have reported comparable values of the Hall mobility for

4H-SiC epitaxial layers.

4.6 Conclusions

In this chapter the structural, optical and electrical properties of the Al doped 4H-SiC

films were investigated by different characterization techniques. For the study, two series of

samples were grown in R1 reactor with an Al concentration in the range 2.3x1016

to ~1x1020

cm-3

.

Additionally, a series of two samples were grown in R2 reactor with an Al concentration in the

range ~1x1019

to ~5x1019

cm-3

. All the samples presented specular surfaces after growth though, it

was found a meandering of the surface with step-bunching and macrosteps parallel to wafer flat

for highly doped layers ([Al]>1.7x1019

cm-3

) grown in R1 reactor. Contrarily, the highly doped

films grown in R2 reactor showed a surface covered by highly oriented step-bunching and a slight

increase of roughness with [Al].

The structural characterization of the Al doped samples revealed high crystalline quality

of the lowly Al doped layers, with FWHM values of symmetrical 4H-SiC(0004) reflection in the

range 15-34arc sec. Concerning the highly doped layers ([Al]>1.7x1019

cm-3

), a degradation of the

crystal quality was found. To determine the lattice constants, 2 - scans and RSMs were

recorded on both symmetric 4H-SiC(0008) and asymmetric (10 8) reflections. For highly doped

samples a peak splitting was observed, with the epilayer peak situated at higher angular position.

The experimental values of the lattice constants were compared with the theoretical ones

predicted by Jacobson’s model [13]. In spite of the lattice change, our results were found to be in

opposition to the theoretical predictions. Further structural investigations, reported that all the

samples with an [Al]≥1.7x1019

cm-3

are formed at least partially by 3C-SiC polytype with double

domain boundaries. However, no 3C-SiC polytype inclusions were found on the sample with an

[Al]~5x1019

cm-3

grown in R2 reactor. Consequently, the growth conditions used in R2 reactor


191

such as higher temperature of 1820°C, and higher growth rate of 30µm/h seem to be conducive

for obtaining highly Al doped layers without 3C-SiC inclusions.

The optical properties characterized by LTPL method showed the presence of an Al-

bound exciton peak ( ) in the NBE region and that of N-Al DAP feature, which changes with

the increase of [Al]. For the highly doped sample the LTPL spectrum displayed the presence of

only one broad feature allocated to the 3C-SiC N-Al DAP, attesting the existence of 3C-SiC

polytype in the layer.

The electrical characterization of the Al doped layers revealed an overestimation of the

Hall hole density with temperature which could be attributed to the Hall scattering factor

which is temperature dependent or to a poor isolation of the junction. Moreover, a large

ionization energy was estimated which leads the most Al acceptors to behave as neutral

impurities. Resistivity as low as 2.3 cm for layers with [Al] of 1.4x1017

cm-3

and 0.25 cm for the

heavily Al doped layer at 1.7x1019

cm-3

were estimated. In the range of [Al] study, the Hall

mobility determined at room temperature was found to decrease from ~90cm2/Vs to ~14cm

2/Vs.

4.7. Bibliography

[1]http://www.cree.com/~/media/Files/Cree/Chips%20and%20Material/Data%20Sheets%20Mate

rial/MATCATALOG.pdf

[2] http://norstel.com/wp-content/uploads/2016/01/Standard-100mm-off-axis-n-type.pdf

[3] K. Kojima, T. Kato, S. Ito, J. Kojima, F. Hirose, Y. Kito, S. Yamauchi, K. Nishikawa, and A.

Adachi, “Reducing stacking faults with highly doped n-type 4H-SiC crystal,” Materials


[4]http://www.cree.com/~/media/Files/Cree/Chips%20and%20Material/Data%20Sheets%20Mate

rial/MATCATALOG.pdf

[5] T. P. Chow, V. Khemka, J. Fedison, N. Ramungul, K. Matocha, Y. Tang, and R. J. Gutmann,

“SiC and GaN bipolar power devices,” Solid-State Electronics, vol. 44, no. 2, p. 277–301,

2000




[7] R. Arvinte, M. Zielinski, T. Chassagne, M. Portail, A. Michon, P. Kwasnicki, S. Juillaguet,

and H. Peyre, “Investigation of Aluminum incorporation in 4H-SiC epitaxial layers,”

Materials Science Forum, vol. 806, p. 45, 2015

[8] K. Kojima, S. Kuroda, H. Okumura, and K. Arai, “Influence of lattice polarity of nitrogen

and aluminum doping on 4H-SiC epitaxial layer,” Microelectronic Engineering, vol. 83, no.

1, pp. 79–81, 2006


192

[9] Schroder D.K., Semiconductor material and device characterization, 3d ed. John Wiley &

Sons, Inc, 2006

[10] M. A. Moram and M. E. Vickers, “X-ray diffraction of III-nitrides,” Rep. Prog. Phys.,

vol. 72, no. 3, p. 36502, 2009

[11] S. W. Huh, H. J. Chung, M. Benamara, M. Skowronski, J. J. Sumakeris, and M. J.

Paisley, “Doping-induced strain and relaxation of Al-doped 4H-SiC homoepitaxial layers,”

Journal of Applied Physics, vol. 96, no. 8, p. 4637, 2004




[13] H. Jacobson, J. Birch, C. Hallin, A. Henry, R. Yakimova, T. Tuomi, E. Janzén, and U.

Lindefelt, “Doping-induced strain in N-doped 4H–SiC crystals,” Applied Physics Letters, vol.

82, no. 21, p. 3689–3691, 2003


[15] B. Kallinger, P. Berwian, J. Friedrich, G. Müller, A.-D. Weber, E. Volz, G. Trachta, E.

Spiecker, and B. Thomas, “Doping induced lattice misfit in 4H–SiC homoepitaxy,” Journal

of Crystal Growth, vol. 349, no. 1, p. 43–49, 2012

[16] T. Kimoto, A. Itoh, and H. Matsunami, “Step-Controlled Epitaxial Growth of High-

Quality SiC Layers,” phys. stat. sol. (b), vol. 202, no. 1, pp. 247–262, 1997

[17] Dudley M., Huang X., and Vetter W. M., Synchrotron White Beam X-ray Topography

and High Resolution Triple Axis X-ray Diffraction Studies of Defects in SiC Substrates,

Epilayers and Devices. W.J. Choyke, H.Matsunami, G. Pensl, 2004

[18] H. J. Chung and M. Skowronski, “High-resolution X-ray diffraction and optical

absorption study of heavily nitrogen-doped 4H–SiC crystals,” Journal of Crystal Growth,

vol. 259, no. 1–2, p. 52–60, 2003

[19] M. O’Keefe and B.G. Hyde, Crystal Structures: I Patterns and symmetry. Mineralogical

Society of America, 1996

[20] X. Li, H. Jacobson, A. Boulle, D. Chaussende, and A. Henry, “Double-Position-

Boundaries Free 3C-SiC Epitaxial Layers Grown on On-Axis 4H-SiC,” ECS J. Solid State

Sci. Technol., vol. 3, no. 4, p. P75–P81, 2014

[21] M. Soueidan, G. Ferro, B. Nsouli, F. Cauwet, J. Dazord, G. Younes, and Y. Monteil,

“Effect of growth parameters on the heteroepitaxy of 3C-SiC on 6H-SiC substrate by

chemical vapor deposition,” Materials Science and Engineering: B, vol. 130, no. 1–3, p. 66–

72, 2006


the C-face of 6H–SiC (0001) substrates,” Journal of Crystal Growth, vol. 348, no. 1, pp. 91–

96, 2012



[24] A. Henry, U. Forsberg, M. K. Linnarsson, and E. Janzén, “Determination of Nitrogen

Doping Concentration in Doped 4H-SiC Epilayers by Low Temperature

Photoluminescence,” Phys. Scr., vol. 72, no. 2–3, p. 254, 2005

[25] I. G. Ivanov, C. Hallin, A. Henry, O. Kordina, and E. Janzén, “Nitrogen doping

concentration as determined by photoluminescence in 4H–and 6H–SiC,” Journal of applied

physics, vol. 80, no. 6, p. 3504–3508, 1996


193

[26] J. Camassel, S. Juillaguet, M. Zielinski, and C. Balloud, “Application of LTPL

Investigation Methods to CVD-Grown SiC,” Chemical Vapor Deposition, vol. 12, no. 8–9, p.

549–556, 2006

[27] Balloud Carole, “Méthodes de caractérisation optique de SiC,” Université Montpellier II

Sciences et Techniques du Languedoc, 2005

[28] F. Schmid, M. Krieger, Laube, M., G. Pensl, and G. Wagner, “Hall scattering factor for

electrons and holes in SiC,” Silicon Carbide Recent Major Advances edited by W.J. Choyke,

H. Matsunami, and G. Pensl, Springer Berlin, p. 517, 2004

[29] A. Parisini and R. Nipoti, “Analysis of the hole transport through valence band states in

heavy Al doped 4H-SiC by ion implantation,” Journal of Applied Physics, vol. 114, no. 24, p.

243703,. 2013

[30] S. Contreras, L. Konczewicz, P. Kwasnicki, R. Arvinte, H. Peyre, T. Chassagne, M.

Zielinski, M. Kayambaki, S. Juillaguet, and K. Zekentes, "Electrical transport properties of

highly aluminum doped p-type 4H-SiC", Materials Science Forum, vol. 858, p.249, 2016



[32] A. Schöner, S. Karlsson, T. Schmitt, N. Nordell, M. Linnarsson, and K. Rottner, “Hall

effect investigations of 4H–SiC epitaxial layers grown on semi-insulating and conducting

substrates,” Materials Science and Engineering: B, vol. 61–62, p. 389–394, 1999

[33] G. Pensl, F. Schmid, F. Ciobanu, M. Laube, S.A Reshanov, N. Schulze, K. Semmelroth,

H. Nagasawa, A. Schoner, and G. Wagner, “Electrical and Optical characterization of SiC,”


[34] A. Koizumi, J. Suda, and T. Kimoto, “Temperature and doping dependencies of electrical

properties in Al-doped 4H-SiC epitaxial layers,” Journal of Applied Physics, vol. 106, no. 1,

p. 13716, 2009

[35] H. Matsuura, M. Komeda, S. Kagamihara, H. Iwata, R. Ishihara, T. Hatakeyama, T.

Watanabe, K. Kojima, T. Shinohe, and K. Arai, “Dependence of acceptor levels and hole

mobility on acceptor density and temperature in Al-doped p-type 4H-SiC epilayers,” Journal

of Applied Physics, vol. 96, no. 5, p. 2708–2715, 2004

[36] J. Pernot, S. Contreras, and J. Camassel, “Electrical transport properties of aluminum-

implanted 4H–SiC,” Journal of Applied Physics, vol. 98, no. 2, p. 23706,. 2005

[37] G. L. Pearson and J. Bardeen, “Electrical Properties of Pure Silicon and Silicon Alloys

Containing Boron and Phosphorus,” Phys. Rev., vol. 75, no. 5, p. 865–883, 1949

[38] V. Heera, K. N. Madhusoodanan, A. Mu cklich, D. Panknin, and W. Skorupa, “Low-

resistivity, p-type SiC layers produced by Al implantation and ion-beam-induced

crystallization,” Applied Physics Letters, vol. 81, no. 1, p. 70, 2002

[39] Y. Negoro, T. Kimoto, H. Matsunami, F. Schmid, and G. Pensl, “Electrical activation of

high-concentration aluminum implanted in 4H-SiC,” Journal of Applied Physics, vol. 96, no.

9, p. 4916–4922, 2004

Chapter 5. Fabrication of PN junction power devices

194

Contents

Chapter 5 Fabrication of PN junction power diodes .............................................................. 195

5.1. Introduction ......................................................................................................................... 195

5.2. Fundamentals of PN junction ............................................................................................ 196

5.2.1. PN junction at thermodynamic equilibrium .................................................................. 197

5.2.2. P-N junction under forward bias ................................................................................... 198

5.2.3. P-N junction under reverse bias .................................................................................... 199

5.2.4 Current-Voltage (I-V) characteristics of P-N junction ................................................... 200

5.3. Fabrication of SiC P+-N junction diodes ........................................................................... 203

5.3.1. Epitaxial process ........................................................................................................... 203

5.3.2 Sample cleaning ............................................................................................................. 204

5.3.3. Diode fabrication procedure .......................................................................................... 205

5.3.3.1 Photolithography ..................................................................................................... 205

5.3.3.2 Etching ..................................................................................................................... 206

5.3.3.3 Ohmic contact deposition ........................................................................................ 208

5.4 Electrical analysis of the diodes .......................................................................................... 211

5.4.1 Linear TLM structures ................................................................................................... 214

5.4.2. I-V measurements of the P+-N and PiN junction diodes ............................................... 216

5.5 Conclusions ........................................................................................................................... 219

5.6. Bibliography ........................................................................................................................ 220


195

Chapter 5 Fabrication of PN junction power diodes

5.1. Introduction

The power semiconductor diodes are divided in two main categories: unipolar metal-

semiconductor diodes (also named Schottky diodes) and bipolar junction diodes. The basic

structure of these devices corresponds to a junction between two different materials or between

similar materials with different doping type.

The major requirements for efficient power devices include: low on-state voltage, low

reverse leakage current and fast switching with minimum losses [1]. Therefore, to minimize the

total power dissipation and to enhance the safe operating area during device turn-on or turn-off, a

particular attention should be given to parameters such as reverse blocking voltage, forward

current capability, reverse leakage current, operating temperature and switching frequency [2].

For example, Schottky diodes provide high switching speed, although usually suffer larger

reverse leakage current and on-resistance compared to bipolar diodes. In unipolar devices, the

forward current conduction is governed mostly by the majority carriers. On the other hand, in PN

junction diodes, the current conduction is governed by both majority (in the contact regions) and

minority carriers (in the junction region).

Accordingly, a slow reverse recovery characteristic attributed to minority carrier storage

is displayed together with a low leakage current and high voltage operation. The current-voltage

characteristics of a power rectifier device are represented in figure 5.1, where the ideal and real

characteristics are compared.


196

Fig. 5.1 Schematic representation of ideal and real current-voltage characteristics of a diode

As can be observed, in ideal power devices a “zero” voltage drop in the on-state and zero

leakage current in the off-state are expected. Moreover, no breakdown voltage seems to be

achieved. Although, in a real device, things change and the device exhibits finite resistance,

leakage current due to electrical defects and maximum breakdown voltage due to ionization or

dielectric breakdown phenomena in the semiconductor.

This chapter focuses on the study of 4H-SiC and 3C-SiC P-N and PiN diodes fabricated

in the clean-room of our laboratory. In this work, the P-N structure was epitaxially grown. The

diodes fabrication process and electrical characterizations will be reported and discussed.

5.2. Fundamentals of PN junction

The P-N junction is formed by two regions with different doping concentration of the

same semiconductor material. The fabrication of the P-N diode can be accomplished by epitaxial

growth or by implantation in which the layers are doped n- or p-type. The obtained P-N diodes

have rectifying current-voltage (I-V) characteristics as illustrated in fig. 5.1.

Usually, the P-N diodes (rectifiers) are the basic block of many electronic and

optoelectronic devices such as junction field effect transistors [3], [4] light emitting diodes

(LEDs) [5], solar cells [6] or laser diodes [7].


197

5.2.1. PN junction at thermodynamic equilibrium

We can assume that the P and N layers are initially separated as shown in figure 5.2(a)

with their energy band diagram. Hence, on the P region the Fermi level ( ) is closer to the

valence band ( ), while on the N region the is closer to the conduction band ( ). The

difference between and or depends on the concentration and on the energy position of

the dopant in the material, and usually one region is more doped that the other. In our study, the

p-type region is more doped compared with the n-type region.

When the two doped regions are brought into contact as represented in figure 5.2(b), the

concentration changes abruptly (energy diagram bending), from where the name abrupt p-n

junction. Therefore, at contact the holes (majority carriers) from the p-type region tend to diffuse

towards the n-type side due to concentration gradient, leaving behind negatively ionized

acceptors charges ( ).

Fig. 5.2 Schematic representation of a P-N junction and their energy band diagram:

(a) before contact and (b) after contact of the two semiconductor regions

Similarly, the electrons from the n-type side diffuse towards p-type side exposing the

positively ionized donors charge ( ). In the vicinity of the metallurgical junction, a depletion

layer is created (space charge region) of a certain width ( ) between the two neutral p-n

sides. The generated electric field in the space charge region drifts the minority carriers in the

opposite direction across the junction. The thermodynamic equilibrium is established when the

diffusion and drift currents are balanced.

However, as a result of the depletion region created at the junction, the energy band is

bended resulting in voltage potential differential or built-in potential ( ) determinate as follows

[8]:


198

(5.1)

where the and are the net ionized donors and acceptor density, respectively, the

intrinsic concentration, the elementary charge ( =1.6022x10-19

C), is the Boltzmann constant

( =1.3806x10-23

J.K-1

) and is the temperature (in K).

Under the influence of the built-in potential, the majority carriers should acquire at least

an energy of to overcome the barrier height and diffuse into either n or p region. Once the

potential barrier is known, other parameters such as electric field profile, depletion layer width

and/or depletion capacitance can be evaluated.

The depletion layer width formed at the junction can be obtained as:

(5.2)

where is the dielectric constant of the semiconductor ( =9.7) and the is the vacuum

permittivity ( =8.85x10-12

F/m).

When one side of the junction is heavily doped compared to the other, e.g. P+-N or N

+-P

junction (P+ or N

+ denote the heavily doped side), the depletion takes place mostly in the lightly

doped region. Therefore, the heavily doped side can be neglected and the depletion width

can be evaluated from [8]:

(5.3)

5.2.2. P-N junction under forward bias

Figure 5.3 represents the P+-N junction and the energy band diagram under a forward

voltage (positive bias applied to the P+ contact side). Initially, we can observe that under a

forward bias applied to the junction, the potential barrier height is reduced from at

equilibrium to .


199

Fig. 5.3 P-N junction with the energy band diagram under a forward-bias

The reduced barrier height grants an increase of the diffusion current resulting in

enhanced minority carrier density, as the electrons from n-region are injected to the p-region and

the holes from the p-region to the n-region.

Accordingly, the width of the depletion region is reduced with the applied voltage as:

(5.4)

5.2.3. P-N junction under reverse bias

A reverse bias is applied to the P+-N junction when the negative voltage is applied to p-

region relative to the n-region. Figure 5.4 shows the schematic representation of the P+-N

junction and the energy band diagram when a reverse bias (negative bias applied to P+

contact) is applied to the diode. Under reverse voltage, the barrier height is increased from to

and the depletion region gets widen. Thus, the depletion region width can be

calculated by substituting the barrier height in eq. 5.4.


200

Fig. 5.4 P-N junction with the energy band diagram under a reversed-bias

The current conduction under a reverse bias application is realized by the minority

carriers since the barrier height is too high and the depletion region to thick for any injection of

the majority carriers. A moderate reverse bias allows only little current to pass through by

accelerating the minority carrier across the depletion region. This results in the so called leakage

current. However, increasing the reverse bias over a critical value will make the structure to

breakdown and the current to over-flow [9]. If a limitation of the current is adapted, the P-N

junction may be operated even under breakdown voltage.

5.2.4 Current-Voltage (I-V) characteristics of P-N junction

Varying the voltage at the P-N junction, the exhibited I-V characteristics are rectifying as

shown in fig. 5.1 which are comparable to Schottky diodes. However, different processes are

involved in the generation of current conduction in the P-N junction diode.

The current crossing an ideal P-N or Schottky diode is usually described as a function of applied

voltage as [10]:

(5.5)

where is the saturation current (in A) and is the ideality factor. The ideality factor is used to

determine the interface quality and usually has a value in the range 1 to 2.


201

The saturation current in p-n junctions consists from the electron-hole pairs generation in the

quasi-neutral region as [8]:

(5.6)

where is the diode area (in cm2), and are the diffusion constant for holes and electrons

(

where is the mobility of the carriers), and are the diffusion length of

holes/electrons ( where is the recombination lifetime of carriers) [11].

When the junction diode has an asymmetrical doping concentration e.g. P+-N junction,

the highly doped region can be ignored and the equation 5.6 simplifies, since only the lightly

doped region is limiting the current.

Under forward bias, the current is present due to the passage of electrons and holes carriers

across the junction and towards the quasi-neutral region. The quasi-neutral regions (qnr) are

actually the regions adjacent to the depletion region in which no electrical field is present and the

free carriers density are close to the net doping concentration. Moreover, the depletion region or

space charge region (scr) is the region around the metallurgical junction in which an electrical

field is created that sweeps away the free carriers and leaves behind ionized acceptor and donors.

However, an additional current produced by different generation-recombination of the

carrier in the depletion region has to be added to the total current. In the SiC material; multiple

numbers of deep and shallow impurity levels exist which act as recombination sites, therefore

causing the appearance of additional current in the depletion region [12], [13]. The generation-

recombination current created in depletion region is added to the total current and the expression

can be written:

(5.7)

where the second component in the eq. is the generation-recombination current, is the diode

area (cm2) and is the effective carrier lifetime.


202

By plotting the measured characteristics of the current as a function of forward bias, in

semi-logarithmic scale, the mechanism of current conduction can be determined as illustrated by

fig. 5.5.

Fig. 5.5 Schematic representation of the forward I-V characteristics of a P-N diode with the four

different conduction regimes [10]

The illustrated curve reveals the presence of four different current regimes.

Consequently, the junction diode current is composed by: (i) recombination current in the

depletion region (the first two regimes on the curve), (ii) diffusion current forming the third

regime and (iii) series resistance regime that limits the current [2], [10]. Determining the slope of

the linear regions located on the I-V curve allows the calculation of the ideality factor by:

(5.8)


203

5.3. Fabrication of SiC P+-N junction diodes

5.3.1. Epitaxial process

The structures of the P-N diodes were grown epitaxially by the CVD technique. The Al

and N were the dopants adopted to create the P-region and N-region of the structure,

respectively. The FTIR and C-V techniques were used to characterize the thickness and the net

dopant concentrations of the epilayers.

The study attributed to the development of P-N junction was an exploratory study in

which the fabrication process for the SiC material in the laboratory was tested. For this study, it

was assigned three different substrates and diode structures. Therefore, the substrates used to

form the epitaxially grown P-N structures were: (i) an 8° off-axis Si-face 4H-SiC commercially

available n-type substrate, (ii) an 8° off-axis Si-face 4H-SiC semi-insulating (SI) substrate and

(iii) a template of 3C-SiC layer grown on n-type Si(100) oriented substrate. Figure 5.6 illustrates

the schematic representation of the grown multilayer stack on the different substrates under

study.

Fig. 5.6 Schematic representation of grown epilayers structure for p-n junction diodes fabrication

on different substrates: (a) standard n-type 4H-SiC, (b) semi-insulating 4H-SiC

and (c) 3C-SiC/Si(100) template

On the Si-face of the 8°off-axis 4H-SiC n+-type substrate ( - ~5x1018

cm-3

) was

grown, as represented in fig. 5.6(a), a P+-N epilayers stack. The first layer that will form the n-


204

region of the structure was grown with a thickness of ~6µm and a net donor concentration -

~2x1016

cm-3

. On top, a second layer that forms the p-region was deposited with a thickness of

~1.4µm and a net acceptor concentration - estimated at~5x1019

cm-3

.

Figure 5.6(b) represents the multilayer stack grown on the Si-face of 8°off-axis 4H-SiC

semi-insulating substrate on which a PiN structure was fabricated. The stacking includes a

~1.9µm thick n-type doped layer at a concentration of - ~3.5x1018

cm-3

, followed by ~1.2µm

thick non-intentionally doped epilayer and completed with a ~0.7µm thick p-type epilayer with a

net acceptor concentration - estimated at 5x1019

cm-3

.

The last multilayer stacking was grown on a 3C-SiC/Si(100) substrate template to form a

PN diode as shown in fig. 5.6(c). The structure is formed by a ~3.4µm thick n-type doped layer

at - ~2x1017

cm-3

and a ~1.6µm thick p-type doped layer with a net acceptor concentration

- estimated at 1x1020

cm-3

.

Table 5.1 summarizes the grown multilayer stacking on the different substrates that will

be used to fabricate the three p-n diodes.

ID Substrate Grown structure

D1 n+-type 4H-SiC p+/n-

D2 SI 4H-SiC p+/n-/n+

D3 n-type 3C-SiC/n+-type Si(100) p+/n+

Table 5.1 Epitaxial grown multilayer stacking for diode fabrication

5.3.2 Sample cleaning

Before the fabrication process of the power devices, a thorough cleaning of the sample

surfaces was performed. The cleaning procedure of the SiC surfaces is usually based on the well-

established cleaning procedure of the silicon process technology [14], [15].

For this study the surface cleaning involved two step operations. The first step consisted

in elimination of any particles or organic compounds from the surface by dipping the samples

into a heated acetone bath under ultrasound for 15min. The second step included the elimination

of any residues from the first treatment by dipping the samples into an isopropanol solution bath


205

under ultrasound for another 15min. After the cleaning procedure, the samples were dried and

ready for the fabrication steps.

5.3.3. Diode fabrication procedure

To create the structure of the P-N diode and uncover the buried n+ layer to form the

cathode, an etching step is necessary. However, an appropriate etching process has to be selected

in order to obtain the smooth surface required for high quality contacts. As a result of the

photoresist small selectivity regarding the etching by Reactive Ion Etching (RIE) [16], a Nickel

(Ni) hard mask was used to etch away the p+ SiC layer.

5.3.3.1 Photolithography

The photolithography step process was used to expose and prepare the areas for contacts

fabrication on both n- and p-type layers. Figure 5.7 displays the cross-sectional schematic

representation of the PN diodes structure that we tried to fabricate on the substrates described

above.

Fig. 5.7 Section view representation of the diode structure on the different substrates

An identical photolithography process was applied for all the samples. On the cleaned

surfaces of the samples, a coating of 500-600nm thick reversible AZ514E photoresist (of

negative-type) was deposited. The photoresist coating was performed on a spinner with 6000

rotation per minute (rpm) to expand uniformly the resist on the entire sample surface. The excess


206

of solvent was removed by placing the samples on a heated plate at a temperature of 114-115°C

for ~2min.

The diodes were designed following the pattern mask represented in figure 5.8. The

patterns were formed by introducing the coated sample by photoresist into a mask aligner

apparatus, for the photomask alignment.

Fig. 5.8 P-N diode patterns used for the study

After the adjustments, the photoresist get exposed to ultraviolet (UV) light for 1sec in

which the selected areas exposed to the UV light are polymerized. Afterwards, the samples were

baked again before the whole surface was exposed to UV for about 15sec. To reveal the

photomask patterns, the samples were immersed in developer solution for 30sec, followed by

water dipping to remove the developer solution from the surface.

5.3.3.2 Etching

Prior to the etching, the samples were dipped in Buffer Oxide Etch (BOE) commercial

solvent to remove any undesired oxide, and then were introduced in an electron-beam evaporator

where ~40nm thick Ni layer was evaporated as etching mask onto each of the samples. The

deposition rate of the Ni metal was 0.25nm/s and the chamber pressure was ~1x10-7

mbar. A lift-

off process in acetone was used to reveal the exposed surfaces and remove the photoresist.

The etching was carried out in a RIE reactor under a gas mixture of Sulfur Hexafluoride

(SF6) and Oxygen (O2) with the flow rates of 6.5sccm and 1.2sccm, respectively. It has been


207

shown by Lazar et al. [16] that using dry etching, and particularly the RIE method to etch the

SiC material, deep tranches can be achieved.

The plasma etching conditions involved a reactor pressure of 0.27Pa with a fixed

radiofrequency (RF) power of 20W. The etching rate for 4H-SiC and 3C-SiC under the given

conditions was ~30nm/min. Investigation of the etched thickness was done by two techniques:

the Dektak profilometer and SEM to be sure that the n+ layer has been reached.

Table 5.2 shows the data regarding the etched thickness obtained from the

characterization techniques.

ID Etched thickness (µm)

Dektak SEM

D1 2.5 2.3

D2 2.6 2.8

D3 2.8 2.2

Table 5.2 Estimated thicknesses after the RIE etching on all three structures by Dektak

profilometer and SEM

To remove the Ni hard mask, aqua regia (a mixture of water, hydrochloric and nitric acid)

H2O:HCl:HNO3 (1:1:2) was used as etchant. Although, the removal of the Ni mask proved to be

difficult and time consuming due to a small etching rate (i.e. strong adherence to SiC surface).

Figure 5.9 shows the SEM images of all three samples after the etching process and

removal of the Ni layer. As can be seen, the side-walls present a smooth surface with well-

defined features indicating appropriate plasma etching process. Yet, it has been found on the

etched surfaces some regions in which the etching process was faulty (rough areas). In addition,

on the surface of samples D1 and D2, the presence of a very thin Ni layer was found. Therefore,

the samples were dipped again in aqua regia solution to remove the remaining Ni layer.


208

Fig. 5.9 SEM images x7000 of etched layers on D1 and D2 samples (conductive and semi-

insulating 4H-SiC substrates), and x10000 etch trench of the D3 samples (3C-SiC substrate)

5.3.3.3 Ohmic contact deposition

When a metal comes in intimate contact with a semiconductor, two different behaviors

can be expected: i.e. the contact can behave as a rectifier (Schottky) or like an ohmic contact.

The character of the contacts between a metal and a semiconductor is explained by the metal

work function and the electron affinity of the semiconductor [10]. For ohmic contacts, it is

necessary that the current is not limited, i.e. to allow the transfer of majority carriers from one

material to the other. To achieve such a behavior, it is essential to use a heavily doped

semiconductor in which the tunneling (field emission) is the dominant carrier transport

mechanism. A good ohmic contact is obtained on a n-type doped SiC epilayer if the doping level

is larger than 1x1017

cm-3

, and on a p-type doped SiC epilayer if the doping concentration is equal

or larger than 1x1019

cm-3

[17], [18].

The selected metals to fabricate the ohmic contacts in our study were: (i) Ni to form the

cathode on the n+ SiC layer and (ii) a bilayer of Al/Au to form the anode on the p

+ SiC layer. The

schematic representation of the contact deposition steps is displayed in fig. 5.10. Initially, a


209

cleaning of the surfaces using the method described in section 5.3.2 was performed to prepare

the samples for the fabrication process chain.

The sample surfaces obtained after the etching (fig. 510a) were coated with a photoresist

(fig. 5.10b), exposed to UV and treated as previously described to open the space on the surface

for the metal deposition as shown in fig. 5.10(c). A 100nm thick Ni layer was evaporated on the

entire surface (fig. 5.10d) by electron beam evaporation method. The last step in the fabrication

process involved the removal of the photoresist by lift-off process, i.e. dipping the samples in

acetone and exposing the deposited Ni contacts onto the open areas (fig. 5.10e).

Fig. 5.10 Schematic representation of the process chain for metal deposition of ohmic

contacts


210

Annealing of the deposited metal is an important step since it enhances the ohmic

characteristics of the contact by forming different interfacial layers such as silicide or carbide at

the metal/SiC interface. The interfacial layers formed at the metal/SiC interface after annealing

induce a reduction of the Schottky barrier height (lower specific contact resistance) which

determines the contact properties of the devices. It was reported that an annealing process to

form the Ni/SiC ohmic contacts requires temperatures in the range 900-1000°C [19]–[21]. In our

study, the deposited Ni metal was annealed at a temperature of ~1000°C for 2min under argon

atmosphere in a Rapid Thermal Annealing (RTA) furnace.

On the other hand, the formation of good ohmic contacts on p-type SiC layers presents

challenges since metals with sufficient low Schottky barrier height values are hard to find.

However, it was found that contacts based on Al metallization and more specifically Al/Ti based

metal schemes give the best ohmic characteristics and low specific contact resistance after an

annealing process above 800°C [22]–[24]. Like in the case of ohmic contacts on n-type SiC

layers, the ohmic behavior is associated with the formation of interfacial phases that assist in

lowering the barrier height [25].

The deposition process of the contacts on the p+ type layers were performed by the lift-

off method also because it offers the possibility to deposit different metals stacking on the same

region of diode structure. Hence, a lift-off resist mask was patterned by photolithography (see

fig. 5.10b) on diode surface, allowing access only to the p+ region of the diode. A 200nm thick

Al layer was deposited by electron-beam evaporation to create the p-type ohmic contact, while

the unwanted metal and resin was discarded from the structure by lift-off in acetone. The

deposited Al metal contacts have a high predisposition to oxidation at high annealing

temperature. Based on the A-E. Bazin [15] work , the Al contacts were annealed at 460°C for a

period of 2min under argon in the RTA furnace to obtain the ohmic behavior. To reduce the

series resistance and to avoid any oxidation of Al contacts, a 200nm thick Au capping layer was

deposited on top without any thermal treatment.

Figure 5.11 displays the SEM images of the deposited and annealed contacts surface of

the p-n junctions. As can be seen from fig. 5.11(a), the annealing at 1000°C of the Ni contacts

deposited on the n-type SiC layer reveals a homogeneous surface with slight roughness.


211

Fig. 5.11 SEM images of the deposited contacts: (a) Ni after annealing (x8000) and

(b) Al/Au stack without annealing of the Au (x5000)

On the other hand, the Au/Al contact stacks formed on p-type SiC layer represented in

fig. 5.11(b) display a different morphology which seems to indicate the presence of some clusters

on the surface. It must be pointed out that it is hard to identify whether the observed surface

morphology is specific only for Al or Au metal in the deposited metal stacks since no

investigation was performed after the deposition and annealing of Al contacts. In literature,

similar morphology was observed for the Al-based contacts, and attributed to an excess of Al

formed during the cooling down process [22]. However, the morphology could be related also to

the Au since no thermal treatment was performed which could result in lower adhesion to the Al

surface.

5.4 Electrical analysis of the diodes

The electrical behavior of the deposited Ni and Au/Al contacts was studied only after the

annealing treatment on all three samples. The current-voltage (I-V) characteristics were acquired

at room temperature from the deposited linear Transmission Line Model (TLM) patterns.

The TLM patterns were fabricated by the process described above (photolithography, lift-off,

and metal deposition) during the same process run as the diodes. The ohmic contact patterns

have dimensions of 60x130µm2. An electrical isolation surrounds the TLM structures obtained

by RIE etching only on p+ layers and not n

+ layers (see fig. 5.13).

The electrical measurement on TLM structure to determine the contact behavior was

carried out by two-point probe using a KARL SUSS probing station and a 2400 KEITHLEY


212

source meter. Figure 5.12 reports the I-V characteristics acquired from two adjacent contacts of

TLM structures separated by a distance of 5µm: (a) for Ni contacts formed on n-type layers

without isolation and (b) for Au-Al contacts formed on the p-type layers with isolation of the

junctions grown on 4H-SiC, SI 4H-SiC and 3C-SiC substrates.

Fig. 5.12 I-V characteristics collected from two adjacent TLM patterns of TLM structures with a

distance of 5 µm: (a) in Ni contacts formed on n-type and (b) in Al-Au contacts formed on the p-

type layers

-10 -5 0 5 10

-8x10-2

-4x10-2

0

4x10-2

8x10-2 a)

Cu

rre

nt

(A)

Voltage (V)

4H-SiC(D1)

SI 4H-SiC(D2)

3C-SiC(D3)

Ni contacts

-10 -5 0 5 10-4x10

-2

-2x10-2

0

2x10-2

4x10-2

b)

Cu

rre

nt

(A)

Voltage (V)

4H-SiC(D1)

SI 4H-SiC(D2)

3C-SiC(D3)

Al/Au contacts


213

As can be seen from fig. 5.12(a), only the Ni contacts deposited on the n-type layer of

samples D2 and D3 present an ohmic behavior. The rectifying behavior of the contact deposited

on sample D1 can be explained by an inappropriate contact on the n-type epilayer: i.e. the Ni

contact was deposited on the layer with a donor concentration of 1.5x1016

cm-3

. Therefore, a high

barrier height is present at the interface and the conduction of electrons carriers by tunneling

(field) effect is not possible since the doping incorporation is not sufficient to reduce the width of

the depletion region. The Ni contacts formed on D3 sample seem to present a smaller linear

slope (from Ohm’s law I-V slope equal to inverse of the resistance) compared with the Ni

contact formed on D2 sample. We may consider that the contact deposited on D3 sample or the

layer underneath the contact has a higher resistance than the contact/layer on D2 sample.

On the other hand, from fig. 5.12(b) we observe that the Au/Al contacts deposited on the

p+ type layers on the three substrates present rectifying behavior. We may explain this behavior

by assuming some hypothesis such as: annealing temperature, the time exposed to the thermal

treatment, the material itself or the appearance of an oxidation layer at the contact that may block

the electrical measurements. A-E. Bazin [15] however, reported in her PhD thesis an ohmic

character of the Al contacts deposited on p-type 3C-SiC layers when a thermal treatment was

performed in the range 300-500°C.

Moreover, in literature [25]–[27] it was shown that pure Al contacts on p-type 4H-SiC

layers after thermal treatment give rise to ohmic behavior. However, a higher temperature seems

to be applied (T>800°C) which contributes to the formation of Al4C3 alloy by an out-diffusion of

silicon and/or formation of lattice vacancies which apparently gives rise to desired ohmic

behavior of the p-type contacts [17], [19], [25].

It is hard to draw a conclusion and to explain the direct effect which caused the rectifying

character since no other experimental studies were performed. Therefore, we have to consider

the necessity of more studies. Nonetheless, we may disclaim the role of epilayer surface as one

of the causes since an ohmic contact was previously obtained on our p-type doped 4H-SiC

layers. The p-type contacts were formed by Ti/Al followed by a flash annealing at 700°C at

Laboratoire Charles Coulomb in Montpellier and at CNR-IMM Catania for the Hall

measurements described in Chapter 4. In consequence, to understand the correct fabrication of

the p+-n junction diodes in the frame of our study, we need further investigations. Unfortunately,

no more samples or pieces of the grown samples are available to extend such investigation.


214

5.4.1 Linear TLM structures

The Ni metal-semiconductor specific contact resistance was estimated from the two

point TLM measurements [10] of D2 and D3 sample which presented the ohmic behavior. The

TLM structure used in our work is represented in fig. 5.13. The structure is composed of 8

adjacent rectangular TLM pads characterized by a contact spacing that varies from 5µm to

160µm.

Fig. 5.13 TLM structure deposited on the surface of the layers with and without isolation

for Al/Au and Ni contacts, respectively

As mentioned, the Ni TLM structure was fabricated without any electrical isolation near

the isolated Al/Au TLM structure, which allows the current the possibility to pass through the

SiC volume. Therefore, the measurements have been limited to the pads of the TLM structure

with low distance (first three pads on fig. 5.13).

The total resistance collected from I-V measurements at different contacts spacing can

be plotted as a function of the different contacts spacing distance as shown in fig 5.14. From

this plotting, one can estimate the values of semiconductor sheet resistance and the transfer

length .


215

Fig. 5.14 Total resistance RT as a function of the contact spacing distance from the Ni TLM

patterns on D2 sample (SI 4H-SiC substrate)

The equation describing the total resistance depending on the contacts spacing

distance is presented below [19]:

(5.10)

where the total resistance depends on the contact resistance , semiconductor sheet resistance

and TLM pads width

From eq. 5.10, the specific contact resistance can be deduced as follows:

(5.11)

where and are the TLM pads length and the transfer length can be estimated from the

function of plotting and extrapolating the intercept on axis.

0 10 20 300

4

8

12

16

Slope=RSH

/w

To

tal

res

ista

nc

e R

T (

)

Contact spacing (m)


216

The specific contact resistance was extracted for both D2 and D3 samples (Ni ohmic

contacts as shown in fig. 5.12a) from the linear fit at values of 3.6x10-5 cm

2 and respectively,

1.3x10-4 cm

2. A similar value of 2x10

-4 cm2

was reported by A-E. Bazin [15] for 3C-SiC

layers with a donor concentration ~1x1017

cm-3

, annealed at 950°C under N2 for 1min.

Additionally, an increase in dopant concentration of the 3C-SiC layer seems to lead to a decrease

of the values up to 10-5 cm

2 [15], [28]. In the case of Ni contacts formed on -SiC layers

with an ~7x1018

cm-3

, it was reported values as ~4-5x10-5 cm

2 [29], [30], which are

comparable with the values obtained in our study on the D2 sample. A stronger dopant

concentration was reported to show lower values up to 3x10-6 cm

2 [28], [31]–[33].

Also, from the slope of the linear fit, the sheet resistance of the n-type region was evaluated

for both samples: at ~488 / for D3 and respectively, ~38 / for D2. The observed

difference between the two samples can be related to the presence of extended defects in the

grown layer or at the metal/SiC interface. For a Ni contact on 3C-SiC, it was reported a

value as low as 22 / [34], while for 4H-SiC a 173 / [35], or even lower ~0.29 /

when a multilayer contact is used [33].

5.4.2. I-V measurements of the P+-N and PiN junction diodes

Electrical measurements of the fabricated junction diodes with different structures were

performed. However, we have to keep in mind that the deposited n- and p-type contacts are

ohmic and Schottky, respectively, meaning that the I-V characteristics may account for a P-N

junction, a resistor and a Schottky junction connected in series. The acquired forward and reverse

I-V characteristics of the fabricated P-N diodes structures with an area of 0.25mm2 are displayed

in figure 5.15. The I-V measurements performed under forward bias for the SiC diodes are

plotted in fig. 5.15(a) under logarithmic and linear scale. From a first view, it can be seen that the

D1 diode fabricated provides smaller current compared with the other two diodes (D2 and D3).

The response may be associated with the presence of the inappropriate contacts created onto the

surface (i.e. both contacts have a rectifying behavior).


217

Fig. 5.15 I-V characteristics acquired on a p-n diode with an area of 0.25mm2

fabricated on the

three structures: (a) under forward bias and (b) under reverse bias. The logarithmic and linear

responses are plotted.

Three regions can be identified on the logarithmic scale plot: a low voltages region up to

~1.2V, a linear region up to ~3.6V pursued by a series resistance dominated region. A similar

behavior with three different regions is observed from the registered I-V curves for the D2

0 1 2 3 4 510

-10

10-8

10-6

10-4

10-2

100

Cu

rre

nt

(A)

Voltage(V)

4H-SiC(D1)

SI 4H-SiC(D2)

3C-SiC(D3)

a)

-200 -150 -100 -50 0

-0.02

-0.01

0.00

Cu

rre

nt

(A)

Voltage (V)

4H-SiC(D1)

SI4H-SiC(D2)

3C-SiC(D3)

b)


218

sample. On the other hand, on the D3 sample, the I-V curve under forward bias displays the

presence of only two regions under semi-logarithmic scale. The first region formed under low

voltage values up to ~1.2V, followed by a second regime in which the series resistance limits the

current.

Using the measured I-V characteristics, the barrier height can be determinate as follows [10]:

(5.12)

where is the diode area and is the Richardson’s constant that can be calculated as:

(5.13)

where is the effective mass of electron in the material ( =0.677 and the

=0.42 [36], [37]) and is the Planck constant ( =6.626x10

-34J.s).

The extracted barrier heights for the diodes fabricated on D1 and D3 samples with an

area of 0.25mm2 are 0.75eV and respectively, 0.56eV. On D2 sample, a more particular forward

characteristic was observed, i.e. a double potential barrier was extracted from the linear regions.

This behavior is typical associated to an “inhomogeneous” Schottky barrier at the metal/SiC

interface and reported previously in literature for Ni/Ti contacts [38], [39]. Considering that in

our case the PiN diode is formed from a rectifier and an ohmic contact, this can give easily rise

to a difference in forward voltage conduction [40]. Therefore, the barrier height of the P-N

and Schottky region is 0.57eV and respectively, 0.81eV. The value of the barrier height obtained

for Ni on 3C-SiC (D3 sample) corresponds with results previously reported [15], [34]. In the

case of 4H-SiC diodes, values in the range 1.6-1.01 [41], [42] or even lower to about 0.5eV

[43] have been reported as a function of annealing temperature and surface preparation.

Due to the state of the contacts formed on P-N regions, the ideality factor cannot be taken into

consideration, as it indicates a strong degradation of metal/SiC interface.

Figure 5.15(b) displays the reverse I-V characteristics of the diodes under linear plot. It

can be seen for D3 sample a fast increase of the reverse leakage current from 2.85x10-5

A at -

2.5V to 2.64x10-2

A at -20V. This rapid leakage current can be associated with the presence of


219

high density of structural defects in the growing 3C-SiC layers. On the contrary, samples D1 and

D2 are blocking higher voltages, i.e. up to -50V and -120V, respectively. Consequently, the D1

sample presents a leakage current that increases from 6x10-5

A at -50Vup to 4x10-2

A at -200V.

For the D2 sample, a leakage current of 4x10-5

A at -120V was observed. During the I-V

measurements on sample D2, it was observed that applying a high reverse current of -50V results

in the luminescence of the whole edge of the p-n junction and the melting of the contact. A

similar behavior was observed by Negoro et al. [44] and attributed to an avalanche multiplication

due to an increase electric field crowding at the p-n junction edges.

The extracted values of the main electrical parameters of the diodes and TLM structures

from the I-V measurements are summarized in table 5.3.

Diode# I-V measurements TLM structures

theoretical (eV) (V) (Ωcm2) (Ω/sq)

D1 1.5 eV 0.75 50

D2 1.5 eV 0.81 120 3.6x10-5

~38

D3 1.1 eV 0.56 5 1.3x10-4

~488

Table 5.3 Main electrical parameters of Schottky diodes and TLM structures from I-V

characteristics

The purpose of this study was to fabricate p-n diodes with different epitaxial grown

structures on SiC substrates based on predetermined fabrication process. Unfortunately, the

predetermined fabrication mechanism that we followed led us to an unsuccessful result as

presented by the electrical parameters represented in table 5.3.

5.5 Conclusions

During this chapter, the fabrication process of P-N diodes with different epitaxial grown

structure and their electrical characterization was discussed. The P-N diodes were fabricated

using a predetermined fabrication recipe on three different substrates that include 8° off-axis

conductive and semi-insulating 4H-SiC and 3C-SiC/Si(100) templates. The process steps


220

involved in the diode fabrication are the: (i) photolithography in which a photoresist coating is

deposited on the sample surface, the patterns are created by a mask alignment and exposure to

UV light; (ii) etching to reach the buried n region; (iii) metal deposition to fabricate the contacts

and (iv) a lift-off to remove the unnecessary metal and photoresist. In this study, the n- and p-

type contacts were formed by Ni and Al/Au metal, respectively. The Ni and Al contacts were

exposed to a thermal treatment, in which a temperature of 1000°C and 460°C was applied under

an Ar atmosphere into an RTA furnace for 2 min. However, it was found out from the electrical

I-V analysis that the p-type contacts have a rectifying behavior.

On the other hand, the Ni contacts formed on 3C-SiC (D3 sample) and SI 4H-SiC (D2

sample) substrates display ohmic character, and the specific contact resistance was calculated

at 1.3x10-4 and 3.6x10

-5 , respectively from the TLM measurements.

Given the rectifying behavior of the p-type contacts, the electrical characterization of the

diodes reveal that we are far from ideality and improvements of the fabrication processes are

necessary. A more direct attention needs to be given to every step in the process of diode

fabrication and nonetheless, to the material chosen for contacts and their thermal treatment.

5.6. Bibliography

[1] T. Kimoto, “Material science and device physics in SiC technology for high-voltage power

devices,” Jpn. J. Appl. Phys., vol. 54, no. 4, p. 40103, 2015

[2] T. P. Chow, V. Khemka, J. Fedison, N. Ramungul, K. Matocha, Y. Tang, and R. J. Gutmann,

“SiC and GaN bipolar power devices,” Solid-State Electronics, vol. 44, no. 2, p. 277–301,

2000

[3] N. G. Wright, A. B. Horsfall, and K. Vassilevski, “Prospects for SiC electronics and

sensors,” Materials Today, vol. 11, no. 1–2, p. 16–21, 2008

[4] http://www.mitsubishielectric.com/company/rd/advance/pdf/vol149/149_TR6.pdf

[5] J. Brault, D. Rosales, B. Damilano, M. Leroux, A. Courville, M. Korytov, S. Chenot, P.

Vennéguès, B. Vinter, P. D. Mierry, A. Kahouli, J. Massies, T. Bretagnon, and B. Gil, “Polar

and semipolar GaN/Al 0.5 Ga 0.5 N nanostructures for UV light emitters,” Semicond. Sci.

Technol., vol. 29, no. 8, p. 84001, 2014

[6] B. S. Richards, A. Lambertz, R. P. Corkish, C. A. Zorman, M. Mehregany, M. Ionescu, and

M. A. Green, “3C-SiC as a future photovoltaic material,” in Proceedings of 3rd World

Conference on Photovoltaic Energy Conversion, vol. 3, p. 2738–2741, 2003

[7] A. Kuramata, K. Domen, R. Soejima, K. Horino, S. Kubota, and T. Tanahashi, “InGaN laser

diode grown on 6H-SiC substrate using low-pressure metalorganic vapor phase epitaxy,”

Journal of Crystal Growth, vol. 189–190, pp. 826–830, 1998


221

[8] C.C. Hu, Modern Semiconductor Devices for Integrated Circuits. New Jersey:

Pearson/Prentice Hall, 2010

[9] B.J. Baliga, Fundamentals of Power Semiconductor Devices. Springer, Berlin, 2008

[10] Schroder D.K., Semiconductor material and device characterization, 3d ed. John Wiley

& Sons, Inc, 2006

[11] E. Rosencher and B. Vinter, Optoelectronics. Cambridge Univ Press, 2004

[12] A. O. Evwaraye, S. R. Smith, and W. C. Mitchel, “Shallow and deep levels in n‐type

4H‐SiC,” Journal of Applied Physics, vol. 79, no. 10, p. 7726–7730, 1996

[13] K. Kawahara, H. Watanabe, N. Miura, S. Nakata, and S. Yamaka, “Shallow and Deep

Levels in Al+ implanted p-type 4H-SiC measured by Thermal Admittance Spectroscopy,”


[14] K. A. Reinhardt and W. Kern, Handbook of Silicon wafer cleaning technology, Second.

Wiliam Andrew Inc., 2008

[15] A. E. Bazin, “Conception de diodes Schottky sur 3C-SiC épitaxie sur silicium,”

Université François-Rabelais de Tours, 2009

[16] M. Lazar, H. Vang, P. Brosselard, C. Raynaud, P. Cremillieu, J.-L. Leclercq, A.

Descamps, S. Scharnholz, and D. Planson, “Deep SiC etching by RIE,” Superlattices and

Microstructures, vol. 40, p. 388, 2006

[17] J. Crofton, L. M. Porter, and J. R. Williams, “The Physics of Ohmic Contacts to SiC,”

phys. stat. sol. (b), vol. 202, no. 1, p. 581–603, 1997

[18] H. Vang, M. Lazar, P. Brosselard, C. Raynaud, P. Cremillieu, J.-L. Leclercq, J.-M. Bluet,

S. Scharnholz, and D. Planson, “Ni–Al ohmic contact to p-type 4H-SiC,” Superlattices and

Microstructures, vol. 40, no. 4–6, p. 626–631, 2006

[19] F. Roccaforte, F. La Via, and V. Raineri, “Ohmic contacts to SiC,” International Journal

of High Speed Electronics and Systems, vol. 15, no. 4, p. 781, 2005

[20] F. La Via, F. Roccaforte, A. Makhtari, V. Raineri, P. Musumeci, and L. Calcagno,

“Structural and electrical characterisation of titanium and nickel silicide contacts on silicon

carbide,” Microelectronic Engineering, vol. 60, no. 1–2, p. 269–282, 2002

[21] K. V. Vassilevski, I. P. Nikitina, N. G. Wright, A. B. Horsfall, A. G. O’Neill, and C. M.

Johnson, “Device processing and characterisation of high temperature silicon carbide

Schottky diodes,” Microelectronic Engineering, vol. 83, no. 1, p. 150–154, 2006

[22] F. Roccaforte, A. Frazzetto, G. Greco, F. Giannazzo, P. Fiorenza, R. L. Nigro, M. Saggio,

M. Leszczyński, P. Pristawko, and V. Raineri, “Critical issues for interfaces to p-type SiC

and GaN in power devices,” Applied Surface Science, vol. 258, no. 21, p. 8324–8333, 2012

[23] N. Thierry-Jebali, A. Vo-Ha, D. Carole, M. Lazar, G. Ferro, D. Planson, A. Henry, and P.

Brosselard, “Very low specific contact resistance measurements made on a highly p-type

doped 4H-SiC layer selectively grown by vapor-liquid-solid transport,” Applied Physics

Letters, vol. 102, no. 21, p. 212108, 2013

[24] M. Vivona, G. Greco, R. L. Nigro, C. Bongiorno, and F. Roccaforte, “Ti/Al/W Ohmic

contacts to p-type implanted 4H-SiC,” Journal of Applied Physics, vol. 118, no. 3, p. 35705,

2015

[25] B. J. Johnson and M. A. Capano, “Mechanism of ohmic behavior of Al/Ti contacts to p-

type 4H-SiC after annealing,” Journal of Applied Physics, vol. 95, no. 10, pp. 5616–5620,

2004

[26] B. J. Johnson and M. A. Capano, “The effect of titanium on Al–Ti contacts to p-type 4H-

SiC,” Solid-State Electronics, vol. 47, no. 9, p. 1437–1441, 2003


222

[27] R. Kakanakov, L. Kassamakova, I. Kassamakov, K. Zekentes, and N. Kuznetsov,

“Improved Al/Si ohmic contacts to p-type 4H-SiC,” Materials Science and Engineering: B,

vol. 80, no. 1–3, pp. 374–377, Mar. 2001

[28] R. Estève, “Fabrication and Characterization of 3C- and 4H-SiC MOSFETs,” KTH Royal

Institute of Technology, Stockholm, Sweden, 2011

[29] F. La Via, F. Roccaforte, V. Raineri, M. Mauceri, A. Ruggiero, P. Musumeci, L.

Calcagno, A. Castaldini, and A. Cavallini, “Schottky–ohmic transition in nickel silicide/SiC-

4H system: is it really a solved problem?,” Microelectronic Engineering, vol. 70, no. 2–4, p.

519–523, 2003

[30] F. Roccaforte, F. La Via, V. Raineri, L. Calcagno, and P. Musumeci, “Improvement of

high temperature stability of nickel contacts on n-type 6H–SiC,” Applied Surface Science,

vol. 184, no. 1–4, p. 295–298, 2001

[31] T. Marinova, A. Kakanakova-Georgieva, V. Krastev, R. Kakanakov, M. Neshev, L.

Kassamakova, O. Noblanc, C. Arnodo, S. Cassette, C. Brylinski, B. Pecz, G. Radnoczi, and

G. Vincze, “E-MRS 1996 Spring Meeting, Symposium A: High Temperature Electronics:

Materials, Devices and ApplicationsNickel based ohmic contacts on SiC,” Materials Science

and Engineering: B, vol. 46, no. 1, p. 223–226, 1997

[32] A. Kakanakova-Georgieva, T. Marinova, O. Noblanc, C. Arnodo, S. Cassette, and C.

Brylinski, “Characterization of ohmic and Schottky contacts on SiC,” Thin Solid Films, vol.

343, p. 637–641, 1999

[33] A. V. Kuchuk, P. Borowicz, M. Wzorek, M. Borysiewicz, R. Ratajczak, K. Golaszewska,

E. Kaminska, V. Kladko, A. Piotrowska, “Ni-Based Ohmic Contacts to n-Type 4H-SiC: The

Formation Mechanism and Thermal Stability,” Advances in Condensed Matter Physics, vol.

2016, p. e9273702, 2016

[34] B. C. Kim, “Electrical characterization of Schottky barrier diodes on heteroepitaxial 3C-

SiC grown on Si substrates,” Purdue University, Indiana, 2008

[35] H. Na, H. Kim, K. Adachi, N. Kiritani, S. Tanimoto, H. Okushi, and K. Arai, “High

quality Schottky and ohmic contacts in planar 4H-SiC metal semiconductor field effect

transistor and device performance,” Journal of Electronic Materials, vol. 33, no. 2, p. 89,

2004

[36] W. M. Chen, N. T. Son, E. Janzén, D. M. Hofmann, and B. K. Meyer, “Effective Masses

in SiC Determined by Cyclotron Resonance Experiments,” phys. stat. sol. (a), vol. 162, no.

1, p. 79–93, 1997

[37] N. T. Son, W. M. Chen, O. Kordina, A. O. Konstantinov, B. Monemar, E. Janzén, D. M.

Hofman, D. Volm, M. Drechsler, and B. K. Meyer, “Electron effective masses in 4H SiC,”

Applied Physics Letters, vol. 66, no. 9, p. 1074–1076, 1995

[38] D. Defives, O. Noblanc, C. Dua, C. Brylinski, M. Barthula, V. Aubry-Fortuna, and F.

Meyer, “Barrier Inhomogeneities and electrical Characteristics of Ti/4H-SiC Schottky

recitfiers,” IEEE Transactions on Electron Devices, vol. 46, no. 3, p. 449, 1999

[39] F. Roccaforte, F. La Via, A. Baeri, and V. Raineri, “Structural and electrical properties of

Ni/Ti Schottky contacts on SiC upon thermal annealing,” Journal of Applied Physics, vol.

96, no. 8, p. 4313, 2004

[40] H. Norde, “A modified forward I‐V plot for Schottky diodes with high series resistance,”

Journal of Applied Physics, vol. 50, no. 7, p. 5052–5053, 1979


223

[41] L. Calcagno, A. Ruggiero, F. Roccaforte, and F. L. Via, “Effects of annealing

temperature on the degree of inhomogeneity of nickel-silicide/SiC Schottky barrier,” Journal

of Applied Physics, vol. 98, no. 2, p. 23713, 2005

[42] F. Roccaforte, F. L. Via, V. Raineri, R. Pierobon, and E. Zanoni, “Richardson’s constant

in inhomogeneous silicon carbide Schottky contacts,” Journal of Applied Physics, vol. 93,

no. 11, p. 9137–9144,. 2003

[43] S. Tanimoto, M. Miyabe, T. Shiiyama, T. Suzuki, H. Yamaguchi, S. Nakashima, and H.

Okumura, “Toward a better understanding of Ni-based ohmic contacts on SiC,” Materials


[44] Y. Negoro, N. Miyamoto, T. Kimoto, and H. Matsunami, “Avalanche phenomena in 4H-

SiC p-n diodes fabricated by aluminum or boron implantation,” IEEE Transactions on

Electron Devices, vol. 49, no. 9, p. 1505, 2002

General conclusions

224

General conclusions and perspectives

The major objective of this thesis was to explore the influence of process conditions on

Nitrogen and Aluminum incorporation in 4H-SiC homoepitaxial films. Although, the dopant

incorporation characteristics on 4H-SiC and 3C-SiC epitaxial layers were reported previously in

literature, it was found almost always a dependence of reactor geometry and growth conditions.

Thus, we set our interest in studying the dopant incorporation characteristics in our own CVD

system and understand the aspects behind the doping mechanism. Moreover, the Aluminum

doping was investigated for the first time in our setup and the obtained results allowed the

transfer from the smaller reactor R1 (2 inch capacity) towards the bigger reactor R2 (4 inch

capacity). All these aspects are important for increasing the layers potential in electronic

applications, especially nowadays, when 4H-SiC wafers with 150mm diameter and excellent

quality (zero micropipes) are commercially available. The study was performed in the framework

of NetFISiC project (2012-2014) in strong collaboration with some of the partners in the project.

In the first part of the manuscript (chapter 1 and 2), I have presented in a general manner

the silicon carbide material, from the attractive properties for microelectronics to the techniques

used for the material growth. Besides, the current status of the material from wafer to device

fabrication was considered, which disclosed an industrial interest and maturity of the SiC.

Furthermore, the deposition kinetics of SiC have been revised in our CVD setup. The boundary

layer model proposed by Eversteyn et al. was found very useful to describe in a qualitative

manner the growth mechanism of SiC inside the reactor chamber. The experimental study

revealed the strong influence of growth parameters on deposition rate. Additionally, the variation

of the input dopant (N2 or TMA) mixture flow rate during the growth leads to the tuning of p-

type doping in the range 1016

cm-3

to 1020

cm-3

and of n-type in the range 1014

cm-3

to 1019

cm-3

on

Si-face 4H-SiC epilayers.

The third chapter is dedicated to the main study involving the N and Al incorporation

particularly in 4H-SiC homoepitaxial films. An exhaustive experimental study of the influence of

process conditions (growth temperature, dopant supply, growth rate, reactor pressure, C/Si ratio)

on voluntary N and Al incorporation in 4H-SiC films was presented and discussed. The

investigation was performed consecutively on Si- and C-face 4H-SiC substrates in both available

General conclusions

225

CVD reactors. Due to site preference of the dopants atoms in the SiC matrix, a systematically

higher N incorporation was obtained on C-face, and opposite for Al incorporation, i.e. stronger

on Si-face. The influence of process conditions on dopant incorporation resulted to different

behaviors that depend on the dopant nature, crystal orientation and chemical environment. On the

other hand, the SiC polytypes and substrates off-angles did not display any significant influence

on N or Al incorporation. An unexpected experimental tendency was obtained from the growth

temperature dependence on N incorporation [N] compared to Al incorporation [Al]. The growth

temperature dependence of the [N] showed two opposite incorporation regimes on C-face and Si-

face in the temperature range 1400-1600°C. Moreover, their dependence are shifting when the

temperature is increased between 1600-1800°C, but the trends remain opposite between the two

polarities. This indicated a more complicated incorporation mechanism of nitrogen compared to

aluminum, as the [Al] was found to decrease with increasing growth temperature on both

polarities and reactors. We suggested that the mechanism behind the experimentally observed

tendencies of N and Al incorporation is related with temperature dependent factors such as:

sticking coefficient of dopant-containing species, carbon coverage and the dopant-containing

species flux towards the grown surface. Besides, the opposite temperature dependence between

the two polarities should be related to desorption/adsorption probability of dopant-containing

species.

The dependence of [N] and [Al] on dopant supply showed on both polarities an increase

almost/ directly proportional to dopant flow rate. However, in R2 reactor (1x100mm) at TMA

flow rate ≤0.01sccm, a much stronger dependence of [Al] was found compared to R1 reactor

(1x2”). The effect behind this dependence remained unfortunately uncertain. The influence of

growth rate on dopant incorporation have shown that the [N] is limited by kinetics of surface

reactions for growth rates ≤10µm/h, and mass transport limited at high growth rates. On the other

hand, the influence of growth rate on [Al] was found to be related to an enhanced C-coverage

given by the increase growth rate. As expected, the incorporation of both dopants increases with

increasing reactor pressure.

The influence of C/Si ratio on [N] and [Al] put in evidence various dopant incorporation

tendencies as: (i) site competition between N and C was found under C-rich condition while site

competition between Al and Si was visible under Si-rich conditions; (ii) a surface coverage

influence of [N] under Si-rich conditions and analogously of [Al] under C-rich conditions.

General conclusions

226

Qualitatively, similar trends were recorded for Si-face and C-face. We proposed the Langmuir

isotherm model to describe the C-coverage at the surface which we found to play an important

role especially on C-face since depends strongly on growth temperature.

In the fourth chapter, the evolution of structural, optical and electrical properties of the Al

doped 4H-SiC films were examined. It was found that the surface morphology of the highly Al

doped 8° off-axis 4H-SiC epilayers grown in R1 reactor deteriorates with the increase of [Al] in

the range 5x1019

cm-3

to ~1x1020

cm-3

, due the low growth temperatures applied. The high

resolution XRD measurements performed on these highly doped off-angle epilayers, highlighted

the presence of double domains 3C-SiC polytype. The evolution of 3C-SiC domains with [Al]

was studied, leading to conclude that both, growth conditions and Al concentration played a role

in the observed outcome. On the other hand, on the highly Al doped 4° off-axis 4H-SiC layers

grown in R2 reactor, the absence of any cubic polytype inclusions was found. However, with the

given growth conditions in R2 reactor, the [Al] in 4H-SiC is limited to ~5x1019

cm-3

. The LTPL

analysis have confirmed the presence of 3C-SiC polytype in the strongly Al doped epilayer and

the good quality of 4H-SiC epilayers at lower [Al]. The electrical characterization by Hall effect

technique of Al doped 4H-SiC layers were investigated for the first time, giving access to

valuable information about the Hall concentration, Hall mobility and resistivity.

The fifth chapter concerns the fabrication process of PN diodes on three different

substrates: standard n-type 4H-SiC, semi-insulating 4H-SiC and a 3C-SiC/Si(100) template

substrate. We found out that the fabrication process and the quality of the ohmic contacts need a

better optimization in order to obtain diodes with superior electrical characteristics.

The present thesis tackled issues from material growth (characteristics of Nitrogen and

Aluminum incorporation) to device fabrication. For a better comprehension regarding the dopant

incorporation mechanism, modelling and simulation of the surface processes are necessary.

Additional experimental data on C-face are required for a complete study of the growth rate

dependence of dopant incorporation. Regarding the device fabrication, a follow-up study of the

technological steps is imperative to reach better quality of the metal/semiconductor interface

which translates into superior electrical characteristics.

Annex I: Growth of cubic SiC on α-substrates

227

Annex I Growth of 3C-SiC epilayers on -SiC substrates

Introduction

The cubic SiC polytype (or 3C-SiC) presents some interest for device fabrication due to

the owned characteristics, although the quality limitation of the grown epilayers sets a drawback

for their industrial use. The quality of 3C-SiC epilayers heteroepitaxial grown on -substrates is

constrained by the formation of twin defects (DPBs) and/or staking faults (SFs). Despite this

challenging issue few groups [1]–[4] have proven that paying special attention to substrate

surface preparation and growth conditions, 3C-SiC free of DPBs can be obtained.

In this annex, we will discuss the growth of 3C-SiC performed on -substrates (i.e. on-

axis 6H and 4H-SiC). The goal of this study was the elaboration of large 3C-SiC domains. The

growth of all layers was performed in R2 reactor. The surface morphology of the epilayers was

characterized by Nomarski interference contrast microscopy (OM) and AFM techniques. The

polytype identification was realized by both LTPL and Raman spectroscopy analysis in

collaboration with Laboratoire Charles Coulomb in Montpellier.

Experimental details about the 3C-SiC growth

A series of samples were deposited in a first attempt to understand and determine the

process conditions necessary for qualitative 3C-SiC epilayers. The growth conditions were set at:

reactor pressure 200mbar, H2 gas flow 30slm and a C/Si ratio equal to 1.5. The SiH4 and C3H8

flow rates were 6sccm and 3sccm, respectively, resulting in a growth rate of ~5µm/h. The growth

was performed on Si- and C-face of nominally on-axis 6H-SiC substrates. For this study the two

inches wafers were cut in quarters of wafers. In order to find the optimal process conditions the

growth temperature was varied in the range 1550-1750°C while keeping all the other parameters

constant. To prepare the surface for the growth, the propane flow rate was introduced during the

temperature ramp-up at 1300°C.

Temperature effect: The variation of growth temperature leads to surface morphology

modifications of the epilayers. The recorded surface morphologies by Nomarski OM of the

epilayers are represented in fig. I.1. Initially, the 3C-SiC presence was evaluated from the color

appearance (i.e. yellow for 3C-SiC and transparent-bluish for 6H-SiC polytype). It was found

from the recorded OM images represented in fig I.1(a)-(c), an enlargement of the 3C-SiC


228

domains with increase temperature from 1550-1650°C on Si-face. A higher density of 3C-SiC

domains can be observed in fig I.1(c), which can be associated to the growth conditions-higher

growth rate (~8µm/h) leading to a epilayer thickness of ~16µm. However, when the growth

temperature was increased up to 1750°C the surface morphology changed as seen in fig I.1(d),

resulting in some domains assign to 6H-SiC polytype. On the other hand, on C-face, the growth

temperature variations just lead to the formation of polycrystalline material as displayed in fig

I.1(e). Yet, when the growth temperature was raised to 1750°C it was observed that the

polycrystalline material formation is inhibited and a specular surface appears.

Fig.I.1 Optical microscopy images with a 50x magnification of the on-axis Si-face 6H-

SiC epilayers with growth temperature variation: (a)1550°C, (b) 1600°C, (c) 1650°C, (d)

1750°C, and (e) on C-face at1600°C, respectively. Epilayers thickness was ~5µm except for (c)

~16µm thick epilayer

Photoluminescence characterization

To confirm the presence of 3C-SiC polytype on both polarities of on-axis 6H-SiC layers,

LTPL measurements were performed at ~4K. Figure I.2 represents the examples of LTPL spectra

of three samples grown on Si-face substrates with different growth temperatures. The collected

spectra are shown only for the near band edge range of 3C-SiC polytype. As can be observed, the

energetic position and the number of the peaks reveal the undeniable 3C-SiC polytype presence.


229

Moreover, the increase in growth temperatures leads to an apparent increasing in intensity of the

peaks. However, since sample 12C1-103 was grown with higher growth rate besides higher

temperature, i.e. the thickness of the epilayer may play a role in the observed peaks intensity.

Also, for this sample no peaks related to the 6H-SiC polytype were recorded compared to the

other two samples. This can suggest that the grown epilayer is formed entirely by 3C-SiC

polytype.

When the sample grown on both Si- and C-face at a growth temperature of 1750°C were

analyzed, the recorded peaks in the NBE correspond mostly to the 6H-SiC polytype. Therefore,

such high temperature plus the growth conditions applied are detrimental for the growth of 3C-

SiC material on both polarities of on-axis 6H-SiC substrates. This first series of samples showed

that the suitable temperatures for the growth of 3C-SiC layers are in the range 1600-1650°C.

Additionally, higher growth rate seems to improve the quality of the obtained 3C-SiC layer.

Fig.I.2. LTPL spectrum collected at 4K on three samples grown on Si-face on-axis 6H-

SiC substrates with various growth temperatures

The effect of C/Si ratio was also analyzed on both Si- and C-face. For this study, the

growth temperature was set at 1650°C while the C/Si ratio was varied between 0.8-1 by changing

520 530 540 550

Inte

ns

ity

(a

.u)

Wavelength (nm)

12C1-099 12C1-102 12C1-103

T~4K

2.38 2.36 2.34 2.32 2.30 2.28 2.26

LOTO

LA

TA

1550°C

1600°C

Energy (eV)

1650°C

ZPL


230

the propane flow rate and keeping constant the silane flow rate at 6sccm. All the other process

parameters were kept the same as described above. The epilayers have a thickness of ~5µm.

On Si-face it was observed that when the C/Si ratio was decreased to 0.8, the surface

morphology deteriorates. The domains associated with 3C-SiC material are replaced by

inclusions of polycrystalline material. Besides, on C-face, the surface morphology remains

simply deteriorated by polycrystalline material indifferent of the applied C/Si ratio.

Consequentially, this study displayed that a C/Si ratio of 1.5 remains desirable since larger

domains correlated to 3C-SiC material are obtained.

Moreover, a slight study of pressure effect was performed again on both polarities of 6H-

SiC substrate. For this study the growth conditions were: growth temperature 1650°C, a C/Si

ratio=1.5, a growth rate of ~5µm/h and a reactor pressure varied in the range 100-50mbars. The

first analysis of the surface morphology revealed polycrystalline material on C-face for both

applied reactor pressures. On Si-face, a reduce growth pressure lead to larger domains assign to

3C-SiC polytype. This first investigation gives actually a first indication of the process conditions

necessary for the 3C-SiC growth.

Pressure effect

Leaving from the first study, an additional investigation of the pressure effect on surface

morphology was performed on both Si- and C-face of on-axis 6H-SiC substrates. For this series

the growth conditions were: growth temperature 1780°, C/Si ratio equal to 1.5 given by a SiH4

flow rate of 15sccm and a C3H8 flow rate of 7.5sccm. A growth rate of ~8µm/h was obtained.

These conditions are the standard process conditions for the homoepitaxial growth of 4H-SiC.

For the study the reactor pressure was varied in the range 100-400mbars. Additionally, the

doping of the layers was accomplished by a variation of the N2 flow rate in the range 1-200sccm.

The doping of the layers was necessary for another study regarding the investigation of nitrogen

doped layers by optical techniques described in the manuscript of P. Kwasnicki [5]. The epilayers

thickness was determined from FTIR measurements at ~12µm.

It was observed that varying the reactor pressure the surface morphology changes-the

domains characteristics for 3C-SiC can be enlarged or diminished. In the reactor pressure range

150-250mbar the surface morphology showed on large areas of 3C-SiC domains, especially on

C-face. However, at high pressure 400mbar and reduce pressure 100mbar the cubic domains on


231

both polarities were much reduced with only small nuclei visible. To determine the 3C-SiC

polytype presence, Raman measurements were performed at room temperature in the range

spectra 120-1100cm-1

at different points on the sample. For the measurements a He-Ne laser with

a 532nm wavelength was used.

Fig. I.3 Raman spectra collected at room temperature on three samples grown with

various reactor pressures

Figure I.3 display the Raman spectra in the range 700-1100cm-1

collected on three

samples grown on Si-face (sample 13CH-007) and C-face (samples 13CH-006/13CH-012) with

various growth pressures. We can see that for the two samples grown on Si- and C-face at reactor

pressure of respectively, 150mbar and 250mbar, the frequencies of the folded modes of

transverse and longitudinal phonons are associated with both 6H-SiC and 3C-SiC polytypes. In

the case of the sample grown at high growth pressure (13CH-012) only the peaks of the

transverse and longitudinal phonons of 6H-SiC polytype are present on the collected spectrum.

Consequentially, the Raman measurements has determined that indeed growth pressures

in the range 150-250mbar are suitable with the applied process conditions for the growth 3C-SiC

domains especially on C-face. As could be observed, the growth of 3C-SiC on both polarities of

6H-SiC substrate, slight difference in growth conditions seems to be necessary for stronger

700 800 900 1000

C-face

LOFTOx(1)

FTOx(1/3)

Si-facep=150 mbar

p=400 mbar

In

ten

sit

y (

a.u

)

Raman shift (cm-1)

13CH-012 13CH-006 13CH-007

p=250 mbar C-face

FLOx(0)

TO


232

coverage in 3C-SiC as previously reported in literature [6], [7]. This difference between the Si-

and C-face is associated to the surface free energy (6H-SiC C-face surface free energy is 718

erg/cm² and on Si-face is 1767 erg/cm² [8]) which leads to distinctive nucleation mechanisms [9].

Another series of epilayers have been grown on 2” nominally on-axis 6H-SiC substrates,

but this time only on C-face. The growth conditions were set as: reactor pressure 70mbar, C/Si

ratio equal to 1.5 given by the flow rate of SiH4=50sccm and C3H8=25sccm. A N2 flow rate of

1.6sccm was used to dope the epilayers in the range 4-8x1016

cm-3

, as determined by C-V

measurements. For this study the growth temperature was varied in the rage 1720-1820°C while

all the other parameters were kept constant. An immediate analysis of the surface morphology

(full wafer mapping) was performed by a differential interference contrast (DIC) optical

microscopy. The DIC images with full wafer scan of the grown epilayers are shown in fig. I.4.

Besides, the insert images display the scan version of the grown surfaces, where the yellow areas

were associated to 3C-SiC polytype.

Fig. I.4. Full mapping by DIC microscopy at 2.5x magnification of the epitaxial layers

grown with various temperatures: (a) 1820°C, (b) 1770°C and (c) 1720°C

It was found that reducing the growth temperature the area covered by domains associated

to 3C-SiC is increased as shown in fig. I.4(a)-(b). Moreover, the 3C-SiC domains seem to be

more defined and also enlarged when a growth temperature of 1770°C was applied (fig. I.4b).


233

When the growth temperature was decreased to 1720°C, the density of 3C-SiC domains is

presented mostly at the edge of the sample as shown by fig. I.4(c). Also, we can see that the

formed 3C-SiC domains are much smaller in length compared with ones grown on previous

samples. The dark area on the wafer is the result of the formation of polycrystalline material.

We concluded that for the formation of 3C-SiC with large domains on C-face, the best

growth temperature is 1770°C under the applied process conditions. Thus, keeping the 1770°C

temperature and growth conditions described above, another investigation was performed by

changing the C/Si ratio. Initially, the C/Si ratio was reduced to 1.2 by decreasing the propane

flow rate and maintaining the same growth rate as previous experiments. The C/Si ratio reduction

had a direct effect on the surface morphology that showed to be characteristics to 6H-SiC with

only small 3C-SiC nuclei. In order to increase the C/Si ratio up to 1.88 a decrease of growth rate

was necessary (SiH4 flow rate set at 40sccm) due to the limitation of the propane flowmeter

(max. 25sccm). These modifications lead to a surface morphology typical for 6H-SiC, with some

small 3C-SiC inclusions. Since the two parameters can have a different effect on the grown

surface an additional sample was grown. This time the C/Si ratio was set to 1.8 by reducing the

silane and propane flow rate at 30sccm and 18sccm. Also, the growth temperature was decreased

to 1720°C.

Fig. I.5 Full mapping by DIC microscopy at 2.5x magnification of the epitaxial layer

grown at a temperature of 1720°C and a C/Si=1.8

Figure I.5 display the full mapping of the grown epilayer with the modifications of the

growth parameters described above. As can be observed, the epilayer surface showed the


234

presence of 3C-SiC domains which are larger in size. However, the surface area covered by these

domains is somewhat limited.

The study was extended further, to compare different on-axis substrates for the growth of

cubic material. Therefore, a comparison was performed using similar growth conditions for all

samples as described above: growth temperature 1820°C, growth pressure 70mbar, C/Si=1.5 and

a growth rate of ~40µm/h. The substrates in the study were: Si- and C-face on-axis 6H-SiC and

Si-face on-axis 4H-SiC substrates. Their surface morphology was investigated by AFM

technique.

Fig. I.6 AFM images obtained on:(a) 20x20µm² scan of a Si-face on-axis 6H-SiC

substrate, (b) 30x30µm² scan of a C-face on-axis 6H-SiC substrate and (c) 50x50µm² scan of a

Si-face on-axis 4H-SiC substrate

Figure I.6 represents the AFM scans of the three grown samples at same growth

conditions. The surface morphology recorded on Si-face and C-face 6H-SiC substrates are shown

in fig. I.6(a)-(b). On Si-face (fig. I.6a), the 20x20µm2 AFM scan indicates a spiral growth which

is the morphology characteristic for 6H-SiC growth. On C-face (fig. I.6b), the 30x30µm2 scan

display the presence of stacking faults characteristic for 3C-SiC growth. However, on Si-face 4H-

SiC substrate (see fig. I.6c) the triangular features and SFs characteristics of 3C-SiC polytype are

much more visible. The surface roughness RSM with a value of 2.32nm was found for a

50x50µm2 scan.

Consequently, this study exposed that the applied growth conditions are suitable for the

growth of 3C-SiC polytype only on C-face 6H-SiC and Si-face 4H-SiC substrates. Although, to


235

determine the quality of the grown 3C-SiC, more analysis are necessary since the AFM can give

only an indication about the surface morphology.

Bibliography

[1] A. Henry, X. Li, H. Jacobson, S. Andersson, A. Boulle, D. Chaussende, and E. Janzén, “3C-

SiC Heteroepitaxy on Hexagonal SiC Substrates,” Materials Science Forum, vol. 740–742, p.

257–262, 2013.

[2] S. Leone, F. C. Beyer, A. Henry, O. Kordina, and E. Janzén, “Chloride-based CVD of 3C-

SiC Epitaxial Layers on On-axis 6H(0001) SiC Substrates,” AIP Conference Proceedings,

vol. 1292, no. 1, p. 7–10, 2010

[3] X. Li, H. Jacobson, A. Boulle, D. Chaussende, and A. Henry, “Double-positioning

boundaries free 3C-SiC layers grown on-axis 4H-SiC,” Journal of Solid Science and

Technology, vol. 3, p. P75, 2014

[4] K. Nishino, T. Kimoto, and H. Matsunami, “Reduction of Double Positioning Twinning in

3C-SiC Grown on α-SiC Substrates,” Jpn. J. Appl. Phys., vol. 36, no. 8R, p. 5202, Aug. 1997.



[6] L. Latu-Romain, D. Chaussende, and M. Pons, “High-Temperature Nucleation of Cubic

Silicon Carbide on (0001) Hexagonal-SiC Nominal Surfaces,” Crystal Growth & Design,

vol. 6, no. 12, p. 2788–2794, 2006

[7] R. Vasiliauskas, S. Juillaguet, M. Syväjärvi, and R. Yakimova, “Cubic SiC formation on the

C-face of 6H–SiC (0001) substrates,” Journal of Crystal Growth, vol. 348, no. 1, p. 91–96,

2012

[8] R. Yakimova, M. Syvajarvi, and E. Janzén, Materials Science Forum, vol. 264–268, p. 159,

1998

[9] R. Vasiliauskas, M. Marinova, M. Syväjärvi, R. Liljedahl, G. Zoulis, J. Lorenzzi, G. Ferro, S.

Juillaguet, J. Camassel, E. K. Polychroniadis, and R. Yakimova, “Effect of initial substrate

conditions on growth of cubic silicon carbide,” Journal of Crystal Growth, vol. 324, no. 1,

pp. 7–14, 2011

Annex II: Growth temperature dependence of dopant incorporation

236

Annex II Discussions: growth temperature dependence of Al and N

incorporation

The relationship between the environmental conditions (growth temperature, growth

pressure, gas phase composition, environment geometry) and growth characteristics

(deposition rate and dopant incorporation) is complex. Detailed description of volume and

surface processes (precursor decomposition, molecule adsorption, diffusion, incorporation or

desorption) requires advanced modeling and simulation which are beyond the scope of this

work. The discussion below set the goal to point out the principal mechanisms that could be

responsible for the experimentally observed tendencies of dopant incorporation.

In SiC, the dopant incorporation sites are well defined, as already mentioned: nitrogen bonds

to silicon and occupies carbon sites, while aluminum bonds to carbon and occupies silicon

sites. During the incorporation process, the dopant atoms (N and Al) compete with the host

atoms (C and Si, respectively) to occupy the free lattice sites created by the presence of

complementary element atoms (Si and C, respectively).

We have to keep in mind that depending on growth conditions (pressure, temperature),

the proportions between the dopant-containing species and their sticking coefficients may

vary. The same statement concerns also the carbon and silicon containing species.

Consequently, several dopant-containing species may contribute simultaneously, each in

different way to dopant incorporation. A thorough quantitative description of dopant

incorporation is almost impossible if this multitude of species and chemical reactions are

neglected. However, we can use a simplified approach that disregards the particularities of

precursor decomposition at high temperature, and still get a coherent qualitative description of

the observed phenomena. By the way, such an approach is in fact applied in a majority of

literature reports.

To understand the mechanism behind the dopant incorporation as a function of growth

temperature, we consider the exposed crystal surface to the controlled flux of Si-, C- and

dopant-containing species during growth. Actually, only a part of the dopant atoms that

collide with the surface will stick to it – we call sticking coefficient the ratio of the number of

adsorbed atoms to the total number of atoms that struck the surface during the same period of

time . If all the stuck dopant atoms are definitively incorporated into the crystal, then their

volume density [D] (dopant incorporation) should be:


237

- Proportional to the flux of dopant-containing molecules towards the surface. is

in turn proportional to the partial pressure of dopant precursor ;

- Proportional to the sticking coefficient of the respective dopant-containing species;

- Inversely proportional to the growth rate .

(II.1)

The sticking coefficient can be written as [1]:

(II.2)

where is the probability that an adatom adsorbs to a free surface site, and are the

probabilities of desorption from above a free and occupied site respectively, is the surface

coverage (ratio between occupied surface sites and total surface sites densities). Under some

conditions (desorption probabilities are stronger than adsorption probability), the formula can

be further simplified and, represented as the product of sticking coefficient at low surface

coverage , and the fraction of free surface sites for dopant incorporation:

(II.3)

The dopant incorporation becomes then:

(II.4)

Usually, the dopant incorporation is lower than 11020

cm-3

. By comparing it with the

bulk density of host element atoms in SiC (~51022

cm-3

) we can conclude that for the most

highly doped layers the surface coverage by dopant atoms is below 1.6%. However, since

dopant and host element atoms incorporate on the same sites, the surface coverage

involved in eq. II.2 to II.4, is dominated by the surface coverage of host element atoms:


238

, . We suppose that the evolution of and with the growth conditions

can be described by a standard adsorption model, such as Langmuir-Freundlich approach [2]:

(II.5)

where stands for C or Si, is the partial pressure of precursor species and

exp is the temperature dependent equilibrium constant. Exponent indicates the

level of surface coverage dependence on precursor concentration. The description is further

simplified by a common assumption, that the sticking coefficient for silicon =1,

independently on surface coverage and growth conditions [3]. This assumption allows us to

consider only the influence of growth conditions on carbon coverage and thus, calculate

the silicon coverage directly as =1- .

After these preliminary considerations, let´s focus on the analysis of temperature

dependence of aluminum and nitrogen incorporation. By expressing in terms of carbon

coverage, the equation II.4 can be explicitly rewritten for both N and Al:

(II.4.1)

(II.4.2)

In the above equations, all factors rely on growth temperature and can potentially influence

the temperature dependence of dopant incorporation. Therefore, we will analyze one by one

their possible contributions.

The growth rate is slightly reduced at high temperature, however this variation is

small compared to the experimentally observed variations of dopant incorporation.

Furthermore, the term would act in the same way in both formulas. In other words, the

term is not responsible for the qualitative differences between N and Al incorporation

tendencies, or for the extent of observed effect. Consequently, we can and will neglect its role

in the further considerations.

For both dopants, the sticking coefficient at low surface coverage and ,

decreases with increasing temperature due to an increased desorption probability .


239

Our experimental study presented in section 3.3.1 was conducted under carbon-rich

conditions (C/Si1). According to equation II.5, the carbon coverage of the surface reduces as

the growth temperature increases. As can be seen from eq. II.4.1 and II.4.2, this reduction will

have a different impact on Al/ N incorporation. For [Al], carbon is the complementary

element – i.e. a reduction of carbon coverage will result in a decrease of the density of

available incorporation sites, leading to the reduction of Al incorporation. For [N], reduced

carbon coverage can contributed to more available sites for nitrogen atoms and, potentially, to

an increase of N incorporation.

The last factor we need to discuss is the flux of dopant containing species towards the

grown surface. In the fig. 2.11 the evolution of N and Al incorporation along the reactor axis

was presented. We found that for nitrogen, stronger incorporation occurs in downstream part

of the reactor. This effect is usually associated with the difficult decomposition of the strong

NN bonding in N2 molecule [4]. For aluminum, the decomposition of precursor molecule is

much faster, as the TMA starts to decompose at temperatures slightly above 100°C.

Consequently, the Al-containing species are available in the upstream part of the reactor and

progressively exhaust along the axis. For both dopants, an increase of temperature in the

deposition zone will accelerate the dopant precursor decomposition and shift the dopant

containing species towards the reactor inlet. For [Al], such shift will reduce further the mole

fraction of available Al-containing species in the deposition zone. For [N], an opposite result

will be observed, i.e. more species will be available for incorporation.

To resume, for [Al], all the three important temperature dependent factors (i.e. , ,

) present in equation II.4, decrease with the increase in growth temperature. This

explains qualitatively the reduction of aluminum incorporation at high temperature

experimentally observed.

For nitrogen, the situation is more complex. We deal with a competition between the

nitrogen flux which increases with temperature, and the nitrogen sticking coefficient ,

which decreases at high temperature. The overall tendency for [N] (increasing or decreasing)

will depend on which of the factors would dominate. A very interesting point is that the

thermal behavior of N incorporation is different on Si- and C-face (Fig. 3.11). Unfortunately,

we can only present some speculations on the origins of the observed behavior.

Nitrogen bonding on Si-face is weaker than on C-face. This fact should have two

consequences: (i) the N desorption probability on Si-face should become more important at

lower temperatures than on C-face; (ii) on C-face, the activation energy of desorption


240

probability should be higher than on Si-face. Taking into account the analytical form of ,

we may speculate that on Si-face, the nitrogen sticking coefficient continuously decreases

with increasing temperature. On C-face two regimes can be observed: in “low temperature”

(LT) regime the nitrogen sticking coefficient is almost constant; while in “high temperature”

(HT) regime it strongly decreases with the increase in temperature, as schematically

represented by dot lines in fig. II.1.

Fig. II.1: Arrhenius plot of possible thermal dependence of nitrogen sticking coefficient (dot

lines), nitrogen flux (dash line) and nitrogen incorporation on C- and Si-face (continuous

lines). Sticking coefficient was calculated using the complete formula described by eq.II.2.

We also assume that the flux of N-containing species (independent on surface

polarity) should increase with increasing temperature. These assumptions are sufficient to

explain qualitatively the tendency observed on C-face. In fact, in LT regime, the sticking

coefficient is being close to 1 – all nitrogen atoms available are incorporated in the layer, and

their quantity increases with temperature. In HT regime, the sticking coefficient reduction is

stronger than the increase of nitrogen flux, leading to a drop down of [N]. In parallel, this

approach explains also the behavior observed in LT regime on Si-face: the increase of

nitrogen flux can be compensated by a decrease of sticking coefficient, which results in

experimentally observed decreasing tendency of [N].

However, if the nitrogen flux presented the same activation energy in all range of

temperatures, [N] on Si-face would continue to decrease in HT regime. To account for the


241

experimentally observed increase of [N] in HT regime, we have to make another assumption:

suppose that the nitrogen flux in HT regime grows faster than in LT regime.

Bibliography:

[1] P. Kisliuk, “The sticking probabilities of gases chemisorbed on the surfaces of solids,” J.

Phys. Chem. Solids, vol. 3, p. 95, 1957

[2] G. P. Jeppu and T. P. Clement, “A modified Langmuir-Freundlich isotherm model for

simulating pH-dependent adsorption effects,” Journal of Contaminant Hydrology, vol.

129–130, p. 46–53, 2012

[3] O. Danielsson, O. Kordina, and E. Janzén, “Simulations of SiC CVD-Perspectives on the

need for surface reaction model improvements,” Materials Science Forum, vol. 778–780,

p. 218, 2014

[4] D. S. Dandy, “Influence of the gas phase on doping in diamond chemical vapor

deposition,” Thin Solid Films, vol. 381, no. 1, p. 1–5, 2001

Annex III: Surface coverage

242

Annex III Surface coverage

We supposed that the evolution of carbon coverage at surface with C/Si ratio can be

described by a Langmuir isotherm adsorption type equation:

(III.1)

where stands for carbon-containing species, is the partial pressure of precursor species and

exp is the temperature dependent equilibrium constant.

We have to keep in mind that the surface coverage is strongly dependent on the substrate

polarity, i.e. a much stronger C-coverage is found on C-face compared to Si-face. Figure III. 1

shows the relationship between the C-coverage and the input C/Si ratio at various growth

temperatures.

Fig. III.1 Surface coverage as a function of input C/Si ratio described by a Langmuir adsorption

model at different growth temperatures

It was found that the surface C-coverage increases almost linearly with the input C/Si

ratio at a given growth temperature. However, under high C/Si ratios (somewhere higher than 1)

Annex III: Surface coverage

243

the C-coverage approaches nearly unity, i.e. the surface reaches the saturation level. The growth

temperature impacts the C-coverage as can be seen. A lower growth temperature leads to a much

faster saturation in carbon of the surface ( ~1) at smaller C/Si ratios. While a high temperature

results to lower C-coverage at the surface which can give rise depending of the dopant nature to

more available sites for incorporation.

Besides the growth temperature, other process parameters such as growth rate and reactor

pressure have a direct impact on the coverage of surface in carbon. In the case of high reactor

pressures or growth rates during the epitaxial process, the at the surface can approach unity

especially on C-face which is already C-terminated.

Therefore, the experimentally observed tendencies of dopant incorporation can be explained by

accounting the coverage of the surface at different growth conditions.

Introduction générale

244


Nous vivons dans l’ère des nouvelles technologies où les composants électroniques

nous entourent. Avec l’avancée des technologies, de plus en plus d’énergie est consommée ce

qui nous amène au problème le plus important de ce siècle : la consommation énergétique.

Afin de diminuer de manière efficace la consommation et la dissipation de l’énergie, nous

devons développer des composants plus performants et plus efficaces. Depuis des décennies,

le marché de l’électronique est dominé par l’utilisation du matériau semiconducteur Silicium

(Si) pour la fabrication de composant électroniques, filière aujourd’hui d’une grande maturité

technologique. Cependant, les demandes ne cessent de croître pour des applications

photovoltaïques, éoliennes, pour les contrôleurs des moteurs des véhicules électriques entre

autres, et le Si semble avoir atteint ses limites pour améliorer davantage performances et

efficacité énergétique.

Pour aller plus avant, des nouveaux composants ont été développés et les

semiconducteurs à grande largeur de bande interdite (grand gap) ont fait leur entrée dans

l’industrie de l’électronique de puissance. Le carbure de silicium (SiC) est un semiconducteur

à grand gap possédant des propriétés physique intéressantes pour atteindre de plus grandes

densités de puissance, des températures de jonction plus importantes et la tenue à haute

tension.

Le SiC a été abondamment étudié durant des décennies dans les laboratoires et au niveau

industriel. De récentes avancées sur la croissance de boules de SiC (des substrats 6’’ de 4H-

SiC avec une faible densité de défauts sont aujourd’hui commercialisés) ont permis de

démontrer des réelles perspectives industrielles pour ce matériau qui présente des avantages

notables au vu des gammes de puissance accessibles et de l’efficacité des dispositifs.

Bien que des percées technologiques aient été mises en avant, l’épitaxie de SiC est encore

confrontée à de nombreux problèmes qui limitent les performances des dispositifs. Pour

dépasser ces limites, il est nécessaire de comprendre les mécanismes de croissance ainsi que

l’incorporation des impuretés nécessaires pour obtenir un dopage maîtrisé. Le dopage du SiC

est réalisé la plupart du temps in-situ pendant l’épitaxie des couches de SiC afin de pouvoir

ajuster les propriétés du matériau épitaxié à l’application visée.

Ces travaux réalisés dans cette thèse sont destinés à apporter de nouvelles

connaissances sur les mécanismes d’incorporation des dopants dans des films de SiC épitaxiés

par dépôt chimique en phase vapeur. La majeure partie des travaux ont porté sur l’influence


245

des paramètres de croissance sur le dopage de type n et p (obtenu avec l’Azote et

l’Aluminium respectivement) de films 4H-SiC. Cette thèse a été réalisé dans le cadre d’un

réseau européen ITN (Marie Curie) intitulé NetFiSiC (Network on Functional Interfaces for

Silicon Carbide). Par ailleurs, les travaux présentés dans cette thèse ont été menées en

collaborations avec 2 partenaires du projet européen NetFiSiC: le laboratoire Charles

Coulomb (L2C) à Montpellier et l’Institut de Microélectronique et de Microsystèmes de

Catane (CNR-INM).

Ce manuscrit se structure en cinq chapitres :

Le premier chapitre constitue une présentation générale du matériau SiC. Il inclut des

détails sur les propriétés chimiques, physiques et électriques des polytypes de SiC les plus

couramment utilisés dans la microélectronique. Les méthodes de croissance pour la

fabrication de substrats (wafers) ou de couches de SiC sont présentées en se focalisant sur les

difficultés de la croissance 4H-SiC et 3C-SiC et les avancées significatives réalisées dans ce

domaine. Finalement, nous faisons un état des lieux de l’épitaxie de SiC et de la fabrication de

composants.

La première partie du second chapitre est consacrée à la description du réacteur

d’épitaxie par dépôt chimique en phase vapeur (CVD) et à celle du procédé de croissance des

couches épitaxiées. Un model qualitatif a été étudié pour comprendre la cinétique de

croissance de SiC dans la chambre de croissance, en portant une attention particulière à

l’influence de la cinétique de croissance. La seconde partie de ce chapitre est dédiée aux

techniques de caractérisation utilisées pour analyser les différentes propriétés de nos couches

épitaxiées.

Le troisième chapitre concerne l’étude du dopage intentionnel avec l’Azote (N) et

l’’Aluminium (Al) durant la croissance de couches de SiC. Le début de ce chapitre est dédié à

l’influence du polytype, de la désorientation du substrat et de l’orientation cristalline sur

l’incorporation de dopant. Par la suite, l’influence de différents paramètres tels que la

température de croissance, le flux de dopant, la vitesse de croissance, la pression et le ratio

C/Si sur l’incorporation de N et Al est décrite et discutée pour de couches de 4H-SiC sur les

faces Si et C de substrats 4H-SiC. Cette étude a pour but d’avoir une vision globale sur les

mécanismes d’incorporation des dopants pendant la croissance CVD. Dans cette étude, nous

tentons d’expliquer les tendances observées expérimentalement en nous basant sur les

mécanismes d’adsorption/désorption sur la surface durant la croissance.

Le chapitre quatre décrit les propriétés structurales, optiques et électriques de couches

4H-SiC dopé Al. L’impact de l’incorporation d’Al sur le paramètre de maille et sur la


246

morphologie de surface des couches épitaxiées a été étudié. Par ailleurs, les propriétés

optiques et électriques des couches 4H-SiC dopé Al ont également été évaluées.

Dans le cinquième chapitre de ce manuscrit, la fabrication et les caractérisations

électriques de diodes à jonction p-n fabriquées sur différents substrats ont été étudiées.

Finalement, le manuscrit se termine par une conclusion générale sur les principaux résultats

obtenus lors de cette étude et des perspectives sont proposées.

Résume détaillées

247


Le SiC est un semiconducteur à large bande interdite (grand gap) important pour les

applications en électronique de puissance en raison de ses propriétés électroniques et physiques

uniques. La structure cristalline du SiC est particulière en ce sens qu’elle est polymorphe et

permet la formation de multiples variétés de mailles de SiC dénommées polytypes. Parmi ces

différents polytypes du SiC seuls le 4H-SiC, le 6H-SiC et le 3C-SiC sont utilisés de nos jours en

électronique. D’une manière générale, les propriétés électriques dépendent du polytype de SiC et

incluent en comparaison avec Si [1] une plus grande énergie de bande interdite (gap), un fort

champ de claquage, une grande conductivité thermique et une grande vitesse de saturation des

électrons. Par ailleurs, la forte liaison chimique Si-C (faible longueur 1.89Å) donne à ce matériau

une grande dureté, une grande stabilité chimique et une importante conductivité thermique. Grâce

à ces propriétés, le SiC suscite l’intérêt pour le développement de composants de puissance

notamment pour les applications de hautes puissances et hautes températures. Les propriétés

physico-chimiques du SiC conduisent également à des difficultés pour la synthèse des cristaux ce

qui dans un premier temps a ralenti le développement des composants. Cependant, de

nombreuses études ont permis l’avancée de la technologie de croissance du SiC résultant dans la

fabrication de substrats de bonne qualité ce qui a permis d’améliorer les performances des

composants à base de SiC. De nos jours, des substrats de 6’’ 4H-SiC désorienté de 4° avec une

densité de défaut faible (absence de ‘micropipes’) sont commercialisés [2]. Il est à noter que des

substrats de SiC 8’’ [3] sont en cours de développement. Parmi les différents polytypes de SiC, le

4H-SiC est le plus couramment utilisé pour les composants de puissance du fait de la

disponibilité de substrats et de couches épitaxiées de bonne qualité.

Par ailleurs, des couches épitaxiées de bonne qualité, d’épaisseur et de dopage uniformes

sont primordiales pour la fabrication de n’importe quel composants à base de SiC. La réelle

avancée pour l’épitaxie de SiC fut ‘la croissance par avancée de marche’ ou ‘step flow growth’

découverte par Kuroda et al. [4]. Ce mécanisme a permis l’homoépitaxie de couches 4H-SiC/6H-

SiC monocristallines sur substrats 4H-SiC/6H-SiC désorientés. La croissance des couches de SiC

se fait généralement avec la technique de dépôt chimique en phase vapeur (CVD) qui a été

largement améliorée au cours des années.


248

Pour un large déploiement des composants de puissance à base de SiC, des couches épitaxiées sur

de grandes surfaces avec une bonne uniformité d’épaisseur et de dopage sont nécessaires. De

plus, connaître les caractéristiques d’incorporation des dopants et les propriétés de cinétique de

croissance du matériau est essentiel pour contrôler le dopage sur une large gamme de

concentrations. Le dopage de SiC est généralement réalisé in-situ pendant la croissance par

introduction de sources gazeuses à base de dopants de type n ou p. Pour le SiC, le dopage de type

n peut être réalisé en utilisant l’Azote (N) ou le Phosphore (P) alors que les éléments tels que

l’Aluminium (Al) ou le Bore (B) sont utilisés pour le dopage de type p.

Dans cette étude, nous nous sommes focalisés sur l’influence des différentes conditions

d’incorporation des dopants dans des couches de 4H-SiC épitaxiées par CVD. De plus,

l’évolution des propriétés structurales, optiques et électriques de couches de 4H-SiC dopé Al a

été examinée. Cette étude a été menée sur des couches de SiC dopé n et p respectivement avec de

l’Azote et de l’Aluminium. Par ailleurs, l’incorporation de N et Al a été étudiée pour des couches

4H-SiC épitaxiées sur les faces C et Si de substrats 4H-SiC. Afin d’obtenir des informations

exhaustives sur les tendances de l’incorporation de N et Al dans SiC, les croissances ont été

réalisées avec différentes configurations expérimentales (taille de réacteur, conditions de

croissance, nature et orientation du substrat). Toutes les couches de SiC ont été épitaxiées dans un

réacteur de croissance CVD horizontal à basse pression avec un système de rotation du substrat.

Notons que l’étude de l’incorporation de N a été réalisée uniquement dans un réacteur de capacité

1x4’’ (réacteur R2), tandis que l’incorporation d’Al a tout d’abord été étudiée dans un réacteur de

capacité 1x2’’ (réacteur R1), avant transfert du procédé dans le réacteur R2. Une chimie standard

a été utilisée pour l’épitaxie des couches de SiC: hydrogène (H2) purifié comme gaz vecteur,

silane (SiH4) et propane (C3H8) comme précurseurs, N2 et triméthylaluminium (TMA) dilué dans

l’hydrogène comme sources de N et d’Al pour les dopage n et p. L’activation des dopants a été

déterminée par des mesures capacité-tension (C-V) avec une sonde à bille de mercure sur des

couches de 4H-SiC dopé n et p. De plus, des mesures de spectrométrie de masse des ions

secondaire (SIMS) ont été conduites sur des empilements de couches afin de déterminer le profil

d’incorporation d’Al.

Il a déjà été établi que l’incorporation de dopants dans la maille de SiC dépend de la taille

des atomes. Ainsi, les atomes de N et d’Al vont préférentiellement s’incorporer respectivement

sur les sites C et Si [5]. La préférence du site du dopant dans la maille SiC a une conséquence


249

directe sur: (i) l’effet de compétition de site entre le dopant et les atomes de la maille et (ii) sur

l’effet de l’orientation cristalline du substrat puisque l’incorporation des dopants dépend

fortement de sa polarité (face Si ou C). Le rôle des principaux paramètres de croissance tels la

température de croissance, le flux de dopant, la vitesse de croissance, la pression dans le réacteur

et le rapport C/Si a été étudié en détail.

Une incorporation d’azote [N] supérieure d’environ un ordre de grandeur a été trouvée sur

la face C [NC] par rapport à la face Si [NSi]. L’incorporation d’Al sur la face Si [AlSi] est

d’environ deux ordres de grandeur plus important que sur la face C [AlSi]. Des résultats similaires

ont été reportés dans la littérature [6]–[8]. Cette dépendance avec la polarité est attribuée aux

différentes liaisons entre les atomes à la surface. En prenant en compte le fait que les atomes N et

Al s’incorporent respectivement sur les sites carbone et silicium dans la maille de SiC, chaque

atome N sur la face C (et Al sur la face Si) présente trois liaisons covalentes alors que N sur la

face Si (et Al sur la face C) n’en présente qu’une seule. Ainsi, une plus forte liaison réduit la

probabilité de désorption et induit une plus grande incorporation des dopants.

L’influence de la température de croissance sur [N] et [Al] a été menée sur les faces Si et

C de substrats de 4H-SiC désorienté de 8° dans les réacteurs R2 et R1. Différentes croissances et

empilements ont été réalisés pour cette étude où la température a varié entre 1400 et 1850°C. Une

tendance inattendue a été trouvée pour [N] avec la variation de la température de croissance sur

les deux polarités. Par ailleurs, la dépendance de la [N] a montré un effet opposé sur la face Si et

C. Dans le cas d’Al, pour les deux réacteurs, une réduction similaire de l’incorporation [Al] avec

l’augmentation de la température de croissance a été observée pour les deux polarités. Nous

avons proposé une explication qualitative sur les tendances observées pour [N] et [Al] sur les

deux polarités comprenant une implication directe de différents facteurs tels que les coefficients

de collage des dopants, la couverture en carbone à la surface et le débit des dopants. Il est

important de noter que la dépendance avec la température de croissance, les proportions entre les

espèces contenant des dopants et leurs coefficients de collage sur la surface de croissance peuvent

varier. Le même comportement concerne également le carbone et le silicium.

La tendance générale pour l’incorporation [N] (augmentation ou diminution) dépend du

facteur dominant précédemment exposé. Il est intéressant de noter que le comportement en

température de [N] est différent selon la face C ou Si. Malheureusement, nous ne pouvons à ce

jour proposer que des spéculations sur les origines des mécanismes observés. Les liaisons azote


250

sont moins nombreuses sur la face Si que sur la face C. Cette observation aurait deux

conséquences: (i) la probabilité de désorption de N sur la face Si serait plus importante que sur la

face C, (ii) en conséquence, sur la face C, l’énergie d’activation de la probabilité de désorption

serait plus importante que sur la face Si. Bien sûr, pour pouvoir comprendre la totalité du

mécanisme d’incorporation des dopants, des simulations numériques peuvent être recommandées

de par la complexité de ces processus.

L’influence du flux de dopant sur l’incorporation de N et Al a montré que l’incorporation

de dopant augmente quasiment proportionnellement avec l’augmentation du débit de N2/TMA

respectivement sur les faces Si et C. Ceci met en évidence l’implication d’une seule espèce

contenant des dopants sur le mécanisme de dopage. Cependant, dans le réacteur R2, en

comparaison avec le réacteur R1, une plus forte dépendance a été trouvée pour l’incorporation de

Al sur les deux faces Si et C et ce pour des débits de précurseur inférieurs à 0.01sccm. Le

mécanisme expliquant cette dépendance n’a pas été entièrement compris mais nous supposons

que ce dernier correspond aux conditions de croissance menant à une disponibilité plus

importante des sites pour l’incorporation de dopant.

L’influence de la vitesse de croissance sur l’incorporation de N et Al a été étudiée

uniquement sur la face Si de couches de 4H-SiC. Deux comportements opposés ont été observés.

Ainsi, une diminution de [N] et une augmentation de [Al] ont été observées lorsque la vitesse de

croissance augmente. Ces comportements ont été attribués à une augmentation de la pression

partielle des espèces contenant du C (TMA et propane) résultant en une augmentation locale du

rapport C/Si. Ainsi, l’augmentation de la couverture en carbone augmente la compétition entre N

et C menant à une réduction de [N] et à une augmentation de [Al].

L’influence de la pression dans le réacteur montre une tendance similaire pour [N] et [Al]

avec une augmentation des concentrations lorsque la pression du réacteur augmente. Par ailleurs,

une tendance similaire concernant [Al] a été observée dans les réacteurs pour chaque polarité du

4H-SiC. De telles tendances expérimentales sont liées à la pression partielle des espèces

contenant des dopants.

De plus, nous avons réalisé une étude exhaustive de l’influence du rapport C/Si sur

l’incorporation volontaire de N et Al dans des couches 4H-SiC, étude couvrant les variations

d’incorporation sur les faces Si et C dans des conditions riche carbone (C/Si>1) et riche silicium

(C/Si<1). Cette étude a mis en évidence différentes tendances d’incorporation des dopants: (i) la


251

compétition de sites entre N et C a été observée pour des conditions riche carbone tandis que la

compétition de site entre Al et Si a été observé pour des conditions riche Si; (ii) l’influence de la

couverture en carbone de la surface sur [N] pour des conditions riche Si et de façon analogue sur

[Al] a été mise en évidence. Le modèle de Langmuir a été utilisé pour montrer l’évolution de la

couverture de C avec l’environnement chimique pendant la croissance. De plus, dans des

conditions riche C, l’incorporation [N] sur la face C se trouve être indépendante du rapport C/Si

lorsqu’une faible température de croissance et une plus forte pression dans le réacteur sont

utilisées. Dans le cas de [Al], une tendance distincte a été obtenue pour des conditions riches

carbones sur les deux polarités dans le réacteur R2 en comparaison avec le réacteur R1. Ceci met

en évidence le rôle des conditions de croissance, tout particulièrement la température, sur la

couverture de surface et les espèces contenant des dopants.

Par ailleurs, les propriétés structurales, optiques et électriques des couches 4H-SiC dopé

Al dans la gamme ~2x1016

cm-3

à ~1x1020

cm-3

épitaxiées dans le réacteur R1 ont été étudiées. Des

mesures de diffraction à haute résolution (HRXRD), de photoluminescence à basse température

(LTPL) et des mesures d’effet Hall ont été réalisées sur ces couches afin d’en estimer les

propriétés. Tous les échantillons présentent une surface spéculaire après croissance, avec pour les

couches fortement dopées ([Al]>1.7x1019

cm-3

) un effet de ‘step-bunching’ des marches parallèles

au méplat des substrats. La caractérisation structurale des couches de SiC dopé Al révèle une

bonne qualité cristalline pour les couches faiblement dopées, avec des valeurs de FWHM dans la

gamme 15-34arcsec pour les réflexions symétriques 4H-SiC(0004). Les couches fortement

dopées ([Al]>1.7x1019

cm-3

) présentent quant à elles, une dégradation de la qualité cristalline.

Pour pouvoir déterminer le paramètre de maille, des scans en 2θ-ω et des mesures RSMs ont été

réalisés sur les réflexions symétriques 4H-SiC(0008) et asymétriques 4H-SiC(10 8). Malgré le

changement de paramètre de maille, nos résultats se sont avérés être en opposition avec les

prédictions théoriques [9]. De plus tous les échantillons ayant un dopage [Al]>1.7x1019

cm-3

sont

partiellement formés de domaines 3C-SiC avec des orientations différent. Cependant, aucune

inclusion de 3C-SiC n’a été trouvée pour l’échantillon avec une concentration [Al]~5x1019

cm-3

épitaxié dans le réacteur R2. Ainsi, les conditions de croissance ont une influence essentielle pour

l’épitaxie des couches fortement dopées Al sans inclusion de 3C-SiC.

Les propriétés optiques obtenues par LTPL ont montré la présence de raies typiques de

l’aluminium ( ) dans la région du bord de bande du 4H-SiC et la présence caractéristique d’une


252

paire donneur-accepteur (N-Al DAP) changeant avec l’augmentation de [Al]. Pour les

échantillons fortement dopés, les spectres de LTPL ont révélé la présence d’un seul pic

caractéristique identifié comme une recombinaison N-Al DAP dans des inclusions de 3C-SiC.

Les caractéristiques électriques des couches dopées Al ont révélé un courant de fuite à

travers le substrat due à la pauvre isolation de la jonction pn résultant d’une surestimation de la

concentration de trous libre avec la température. Par ailleurs, une forte énergie d’ionisation a été

mesurée ce qui implique que la plupart des accepteur Al se comportent comme des impuretés

neutres. Des résistivités de 2.3 cm et 0.25 cm ont été évaluées pour des couches dopées [Al] à

1.4x1017

cm-3

et respectivement 1.7x1019

cm-3

. Dans la gamme de concentration [Al] étudiée, la

mobilité des trous, déterminée à température ambiante, s’est avérée diminuer de ~90cm2/Vs à

~14cm2/Vs.

La fabrication de diodes PN a été menée sur trois différentes couches épitaxiées en

utilisant un procédé de fabrication précédemment. Dans cette étude, les contacts de type n et p

ont été réalisés respectivement avec du Ni et avec l’empilement Al/Au. Les caractéristiques

électriques ont révélé que les contacts de type p avaient un comportement redresseur sur les trois

différentes couches. Par ailleurs, les contacts Ni formés sur les couches dopées n sur 3C-SiC et

sur le substrat 4H-SiC semi-isolant ont montré un caractère ohmique et la résistance spécifique de

contact a pu être calculée à 1.3x10-4 et respectivement 3.6x10

-5 avec des mesures

TLM. Etant donné le comportement redresseur des contacts de type p, les caractéristiques

électriques des diodes ont révélé que nous sommes encore loin d’un système idéal et que des

améliorations sur les procédés de fabrication étaient nécessaires.

Bibliographie

[1] P. G. Neudeck, The VLSI handbook/ Chapter 5 SiC Technology, CRC Press. 2006.

[2] http://www.cree.com/News-and-Events/Cree-News/Press-Releases/2012/August/150mm-

wafers

[3] http://globenewswire.com/news-release/2015/07/16/752880/10142036/en/II-VI-Advanced-

Materials-Demonstrates-World-s-First-200mm-Diameter-SiC-Wafer.html.

[4] N. Kuroda, K. Shibahara, W. Yoo, S. Nishino, and H. Matsunami, “Step-controlled VPE

growth of SiC single crystals at low temperatures,” Ext. Abstr. 19th Conf. on Solid State

Devices and Materials, Tokyo, p. 227, 1987.


253


superior silicon carbide electronics,” Applied Physics Letters, vol. 65, no. 13, pp. 1659–1661,

1994.

[6] K. Kojima, S. Kuroda, H. Okumura, and K. Arai, “Nitrogen incorporation characteristics on a

4H-SiC epitaxial layer,” Applied Physics Letters, vol. 88, no. 2, p. 21907, 2006.

[7] U. Forsberg, O. Danielsson, A. Henry, M. . Linnarsson, and E. Janzén, “Aluminum doping of

epitaxial silicon carbide,” Journal of Crystal Growth, vol. 253, no. 1–4, pp. 340–350, Jun.

2003.

[8] U. Forsberg, Ö. Danielsson, A. Henry, M. K. Linnarsson, and E. Janzén, “Nitrogen doping of

epitaxial silicon carbide,” Journal of Crystal Growth, vol. 236, no. 1–3, pp. 101–112, Mar.

2002.

[9] H. Jacobson, J. Birch, C. Hallin, A. Henry, R. Yakimova, T. Tuomi, E. Janzén, and U.

Lindefelt, “Doping-induced strain in N-doped 4H–SiC crystals,” Applied Physics Letters, vol.

82, no. 21, pp. 3689–3691, May 2003.

Conclusions générales et perspectives

254


L’objectif principal de cette thèse était d’explorer l’influence des conditions de

croissance sur l’incorporation d’Azote et d’Aluminium dans des couches homoépitaxiées de

4H-SiC. Bien que les caractéristiques de l’incorporation des dopants dans des couches de 4H-

SiC et 3C-SiC aient été rapportées dans la littérature, il a toujours été montré une dépendance

avec la géométrie du réacteur de croissance et/ou avec les conditions de croissance. Ainsi,

nous nous sommes intéressés aux différentes caractéristiques de l’incorporation dans nos

propres réacteurs de croissance CVD en essayant de comprendre les différents facteurs

opérant dans les mécanismes de dopage. Par ailleurs, pour la première fois au laboratoire, le

dopage Aluminium a été étudié et les résultats obtenus ont permis de transférer le procédé de

croissance d’un réacteur de capacité 1x2’’ (R1) à un réacteur de capacité 1x4’’ (R2). Tous ces

aspects sont importants pour améliorer les propriétés des couches épitaxiées pour des

applications électroniques notamment depuis que des substrats (wafers) de 150 mm de

diamètre ayant une excellente qualité structurale (zéro ‘micropipes’) sont commerciés. Cette

étude a été menée dans le cadre du projet européen NetFISiC (pendant 2012-2014).

Dans la première partie de ce manuscrit (chapitre 1 et 2), nous avons présenté le

carbure de silicium et ses propriétés attractives pour la microélectronique ainsi que les

techniques utilisées pour la croissance de SiC. D’autre part, l’état de l’art du SiC, de la

fabrication des wafers jusqu’à la fabrication de composants, a été présenté tout en montrant

l’intérêt industriel et la maturité d’une partie de la filière SiC. Le cinétique de dépôt du SiC a

été étudiée. Le modèle de la couche limite proposé par Eversteyn et al. a été utilisé pour

décrire de manière qualitative les mécanismes de croissance du SiC. Les études

expérimentales ont révélé la forte influence des paramètres de croissance sur la vitesse de

dépôt. De plus, la variation du débit de dopant (TMA ou N2) pendant la croissance permet de

contrôler respectivement le dopage de type p dans la gamme 1016

cm-3

à 5x1019

cm-3

, et le

dopage de type n entre 1014

cm-3

à 1019

cm-3

dans des couches 4H-SiC épitaxiées avec la

polarité Si.

Le troisième chapitre décrit l’étude principale de cette thèse concernant l’incorporation

de N et d’Al dans des couches homoépitaxiées de 4H-SiC. Une étude exhaustive de

l’influence des différentes conditions de croissance (température de croissance, flux de

dopant, vitesse de croissance, pression dans le réacteur et rapport C/Si) sur l’incorporation des

dopants N et Al dans des couches de 4H-SiC a été présentée et discutée. Cette étude a été


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menée sur les faces Si et C de substrats 4H-SiC dans les deux réacteurs de croissance CVD

disponibles. La prise en compte de la préférence de site des atomes dopants dans le réseau

cristallin SiC permet d’expliquer une incorporation plus importante d’azote obtenue sur la

face C et inversement pour l’aluminium sur la face Si. De plus, l’étude de l’influence des

conditions de croissance sur l’incorporation des dopants (N and Al) montre des

comportements différents selon la nature du dopant, l’orientation cristalline et

l’environnement chimique. D’autre part, les polytypes SiC et la désorientation du substrat

n’ont montré aucune influence significative sur l’incorporation de N ou d’Al.

Un résultat inattendu concerne l’influence de la température de croissance sur

l’incorporation de N [N] par rapport à l’incorporation d’Al [Al]. L’incorporation d’azote

montre deux régimes opposés sur la face C et Si dans la gamme de température 1400-1600°C.

Par ailleurs, les courbes se décalent lorsque la température est augmentée entre 1600 et

1800°C, mais la tendance reste opposée entre les deux polarités. Ceci indique un mécanisme

d’incorporation plus complexe pour l’Azote en comparaison avec l’Aluminium puisque la

concentration [Al] diminue lorsque la température augmente quels que soient la polarité et/ou

le réacteur utilisé. Nous supposons que le mécanisme expliquant les tendances observées pour

l’incorporation de N et d’Al est lié à des facteurs dépendant de la température tels que le

coefficient de collage pour les espèces contenant les dopants, la couverture en carbone à la

surface et le flux des espèces contenant les dopants. D’autre part, la dépendance opposée en

température entre les deux polarités du 4H-SiC serait reliée à la probabilité

d’adsorption/désorption des espèces contenant les dopants.

L’étude sur l’incorporation a également montré sur chaque polarité une augmentation

quasiment proportionnelle au flux de dopant. Cependant, dans le réacteur de croissance R2

(1x4’’) une forte dépendance de la [Al] a été trouvée pour un débit de TMA ≤0.01sccm par

rapport au réacteur de croissance R1 (1x2’’). La raison de cette dépendance reste cependant

incertaine. L’influence de la vitesse de croissance sur l’incorporation de dopant a montré que

l’incorporation d’azote est limitée par la cinétique de réaction à la surface pour des vitesses de

croissance ≤10μm/h, et par le transport de masse pour des vitesses de croissance plus grandes.

D’autre part, la dépendance de l’incorporation [Al] avec la vitesse de croissance s’est avérée

être reliée à une augmentation de la couverture en carbone à la surface. Comme attendu,

l’incorporation de chacun des dopants augmente avec l’augmentation de la pression dans le

réacteur.

L’influence du rapport C/Si sur [N] et [Al] a mis en évidence différentes tendances

d’incorporation telles que: (i) la compétition de site entre les N et C pour des conditions riche


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carbone tandis que le compétition de site entre Al et Si a été observé pour des conditions riche

Si; (ii) l’ influence de la couverture en carbone sur l’incorporation [N] pour des conditions

riche Si et de façon analogue pour l’incorporation [Al], dans des conditions riche carbone.

Ont été observées des tendances qualitativement similaires pour l’effet du rapport C/Si

sur les faces Si et C. Le modèle isotherme de Langmuir a été proposé pour décrire la

couverture en carbone qui se trouve jouer un rôle important sur la face carbone au vu de sa

grande dépendance avec la température de croissance.

L’évolution des propriétés structurales, optiques et électriques de couches 4H-SiC

dopées Al a été évaluée dans le chapitre quatre. Il a été constaté que la morphologie des

couches 4H-SiC fortement dopées Al épitaxiées dans le réacteur R1 sur des substrats 4H-SiC

désorientés de 8°, se détériore avec la augmentation de la concentration en Al dans la gamme

5x1019

cm-3

- ~1x1020

cm-3

. Cette détérioration de la morphologie a été attribuée à la plus faible

température de croissance utilisée pour la croissance des couches. Des mesures XRD à haute

résolution réalisées sur des couches fortement dopées Al ont montré la présence de domaines

des 3C-SiC avec des doubles orientations. L’évolution de ces domaines avec la concentration

[Al] a été étudiée. D’autre part, aucune inclusion 3C-SiC n’a été observée sur les couches 4H-

SiC fortement dopées Al sur des substrats désorientés de 4° et épitaxiés dans le réacteur R2.

Cependant, les conditions de croissance utilisées dans le réacteur R2 mènent à une

incorporation d’Al limitée à ~5x1019

cm-3

sur la face Si du couches 4H-SiC, ce qui ne nous

permet pas de conclure définitivement sur l’influence des conditions de croissance. L’analyse

LTPL a confirmé la présence du polytype 3C-SiC dans la couche fortement dopée Al dans

réacteur R1 et des bonnes qualités cristallines pour les couches 4H-SiC avec une

concentration plus faible. Des caractérisations électriques réalisées par effet Hall sur des

couches de 4H-SiC dopées Al ont été étudiées pour la première fois au laboratoire et

fournissent des informations intéressantes sur la concentration, la mobilité des porteurs et en

conséquence la résistivité des couches.

Le cinquième chapitre concerne la fabrication et l’étude de diodes PN sur trois

substrats différents: sur substrat 4H-SiC de type n, sur substrat 4H-SiC semi-isolant et sur un

tremplin 3C-SiC/Si(100). Nous avons observé que le processus de fabrication et en particulier

les contacts ohmiques sur les couches dopées p doivent être optimisés pour obtenir des diodes

présentant de meilleures caractéristiques électriques.

Cette thèse a abordé les problèmes de la croissance du matériau SiC (caractéristiques de

l’incorporation de l’Azote et de l’Aluminium) et de la fabrication de composants. Cette

approche expérimentale nous a permis d’établir un cadre phénoménologique intégrant les


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principaux mécanismes en jeu. Pour pouvoir comprendre plus en détail les mécanismes

d’incorporation des dopants, des simulations numériques des phénomènes en jeu, en

particulier à la surface seront nécessaires. De nouvelles expériences doivent être également

menées sur la face C afin d’obtenir une image plus complète de là l’influence de la vitesse de

croissance sur l’incorporation des dopants. En ce qui concerne la fabrication de composants,

un suivi étape par étape des procédés technologiques est impératif pour optimiser les

interfaces métal/semiconducteur et obtenir de meilleures caractéristiques électriques.


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investigation of dopant incorporation in silicon carbide

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