King

SURFACE AND INTERFACE CHARACTERIZATION OF SiC

AND III-V NITRIDES

by

SEAN WESLEY KING

A dissertation submitted to the Graduate Faculty of

North Carolina State University in partial fulfillment of the

requirements for the Degree of Doctor of Philosophy

MATERIALS SCIENCE AND ENGINEERING

Raleigh

1997

APPROVED BY:

_____________________ _________________________

_____________________ _________________________

Chair of Advisory Committee

ABSTRACT KING, SEAN WESLEY. Surface and Interface Characterization of SiC and III-V

Nitrides. (Under the direction of Robert F. Davis.)

The effects of various wet chemical and chemical vapor processes on the surfaces

of (0001) silicon carbide (SiC), aluminum nitride (AlN), and gallium nitride (GaN)

have been investigated using x-ray photoelectron spectroscopy (XPS), Auger electron

spectroscopy (AES), low energy electron diffraction (LEED), electron energy loss

spectroscopy (EELS), temperature programmed desorption (TPD), and x-ray

photoelectron diffraction (XPD). XPS was also used to determine both the growth

mechanism of GaN on (0001) AlN as well as the heterojunction valence band

alignment between 2H-AlN and 6H-SiC at the (0001) interface. Polycrystalline

scandium nitride (ScN) films were also grown on 3C/6H-SiC (0001) substrates by gas

source molecular beam epitaxy (GSMBE).

Thermally oxidized surfaces of (0001)Si 6H-SiC were visually observed to be

hydrophilic after removal of the oxide with HF solutions. The hydrophilic nature of

this surface was correlated with a monolayer coverage of oxygen (hydroxide)

observed on the surface by AES and XPS. A completely dry ex situ cleaning process

based on UV/ozone (O3) oxidation for adventitious carbon removal, followed by HF

vapor exposure for oxide removal was demonstrated and observed to be equivalent to

conventional wet chemical processes. Removal of the monolayer coverage of oxygen

from the (0001)Si 6H-SiC surface via chemical vapor cleaning (CVC) in a flux of

silane (SiH4) was observed to produce higher purity surfaces compared to typical

thermal desorption techniques. Low energy electron diffraction showed that (0001)Si

6H-SiC surfaces prepared by chemical vapor cleaning exhibited a (3x3)

reconstruction. X-ray photoelectron spectroscopy indicated that the (3x3)

ii

reconstruction consisted of an incomplete bilayer of silicon terminating the SiC

surface. X-ray photoelectron diffraction indicated that the stacking sequence of the

silicon bilayer was similar to that of a faulted Si (111) stacking structure. Selective

removal/etching of silicon from (3x3) (0001)Si 6H-SiC surfaces by atomic H

processes was observed by both AES and XPS.

Wet chemical processes based on HF were found to produce AlN surfaces with

the lowest oxygen whereas HCl based wet chemical processes were to produce the

lowest oxygen coverages for GaN surfaces. This observation was correlated with

XPS and AES studies which showed fluorine (chlorine) termination of AlN (GaN)

surfaces after HF (HCl) wet chemical processes. Complete desorption of F from AlN

surfaces was not observed to occur until Tsub > 800°C. Desorption of Cl from GaN

surfaces was observed to be complete by 650°C. Annealing GaN surfaces in NH3 at

800°C was observed to remove both oxygen and carbon below the detection limits of

both AES and XPS.

Using XPS, it was determined that GaN growth on (0001) AlN by GSMBE

occurs via a three dimensional Stranski-Krastonov growth mechanism at Tsub < ≈

780°C. At higher temperatures (Tsub > 800°C), the growth mechanism switches to a

layer by layer/Frank van der Merwe two dimensional growth mechanism. The

change in growth mechanism with temperature was attributed to a change from a

fully hydrogenated growth surface to an incompletely hydrogenated growth surface at

higher temperatures.

Using x-ray photoelectron spectroscopy and published theoretical valence band

density of states, a type I heterojunction valence band discontinuity of 1.4 ± 0.3 eV

was determined for the (0001) 2H-AlN/6H-SiC interface.

iii

DEDICATION

In honor of my parents:

Carrol David King and Nanci Marie Curry King

In memory of my grandmother:

Hilda Bell

iv

BIOGRAPHY

Sean Wesley King, son of Carrol David and Nanci Marie King, was born in Seneca,

South Carolina November 12, 1968. Shortly thereafter, his family moved to Stokesdale,

North Carolina where he was blessed with a younger sister, Julie Virginia King. The King

family eventually settled in the small community of Green Spring, Virginia which is the

ancestral home of many generations of this family. Sean then went on to quietly graduate

from Abingdon High School and follow in the footsteps of his father by enrolling in

Engineering at Virginia Polytechnic Institute and State University. After a brief stint as an

Aerospace Engineering major, Sean eventually found his niche in Materials Engineering and

went on to graduate summa cum laude. Shortly after completion of his Bachelors, Sean

enrolled in the Materials Science and Engineering Ph.D. program at North Carolina State

University, with Professor Robert F. Davis as his advisor. The work herein represents partial

completion of this degree.

v

ACKNOWLEDGMENTS

I would like to first express my appreciation to my advisory committee chairman, Dr.

Robert F. Davis, for his continued patience and support during the course of this study. I

would also like to express my appreciation to Dr. Robert J. Nemanich for his endless

enthusiasm and guidance in this research. Appreciation is also extended Dr. S. M. Bedair

and Dr. D. Griffis for their contributions as members of my advisory committee.

The initial stages of this research involved an abundant amount of equipment and

experimental design which without the assistance and guidance of Drs. L. Rowland, J.

Sumakeris, and J. van der Weide would have never come to completion and to whom I am

indebted. Appreciation is also expressed to Dr. J. Yates Jr., and his students for the

numerous telephone conversations regarding various aspects of surface science equipment. I

would also like to thank Mr. A. Illingsworth and Mr. J. Emerick for helping me make my

equipment designs work despite their many flaws.

I would like to thank the Office of Naval Research for its financial support of this

research and the Department of Education for its assistance through a GAANN fellowship.

Thanks are also extended to Dr. W. Perry, Dr. A.D. Batchelor, Mr. M. Bremser, Mr. E.

Carlson, Mr. R. Therrien, and Mr. D. Bray for their help with various measurements and

analysis. Appreciation is also expressed to Dr. W.R.L. Lambrecht for his guidance and

collaboration in the heterojunction valence band discontinuity effort.

Perhaps the most rewarding aspect of this venture was the chance to work and

interact with many diversified people who comprise both the Davis Laboratories and the

NCSU Surface Science laboratory. During the course of this research, many of these people

vi

have assisted and encouraged me through the highs and lows and with whom special

friendships have been formed including: Dr. J. Barnak, Mr. M. Benjamin, Dr. L. Bergman,

Mr. R. Busby, Mr. R. Carter, Mr. J. Christman, Mr. H. Ham, Dr. R.S. Kern, Dr. Ja-Hum Ku,

Dr. T. Schnieder, Mr. A. Sowers, Dr. S. Tanaka, Mr. K. Tracy, Mr. B. Ward, Mr. W. Yang,

Mrs. H. Ying, and Mr. S. Wagoner.

Finally, I would like to recognize the unwavering support and patience of my parents

and family during the completion of this endeavor.

vii

TABLE OF CONTENTS

LIST OF TABLES .................................................................................................... xiv

LIST OF FIGURES .................................................................................................. xvi

1. Introduction.................................................................................................. 1

2. Wet Chemical Processing of (0001)Si 6H-SiC: Hydrophobic and

Hydrophilic Surfaces...................................................................... 5

2.1 Abstract............................................................................................. 6

2.2 Introduction....................................................................................... 7

2.3 Experimental Procedure..................................................................... 10

2.4 Results............................................................................................... 12

2.4.1. (0001)Si 6H-SiC, Oxidized.......................................... 12

2.4.2. (000-1)C, (11-20), & (10-10) 6H-SiC, Oxidized......... 20

2.4.3. (0001)Si 6H-SiC, As-polished..................................... 22

2.4.4. Si Passivation Layer.................................................... 26

2.5 Discussion.......................................................................................... 28

2.5.1. (0001)Si 6H-SiC, Oxidized ......................................... 28

2.5.2. (0001)Si 6H-SiC, As-polished..................................... 35

2.5.3. Si Passivation Layer.................................................... 37

2.6 Conclusions........................................................................................ 39

2.7 Acknowledgments.............................................................................. 39

2.8 References.......................................................................................... 40

3. Dry Ex Situ Cleaning Processes for (0001)Si

viii

6H-SiC Surfaces.............................................................................. 47

3.1 Abstract............................................................................................. 48

3.2 Introduction....................................................................................... 49


3.4 Results............................................................................................... 54

3.4.1. Solvents and UV/O3 Oxidation................................... 54

3.4.2. HF vapor..................................................................... 58

3.5 Discussion.......................................................................................... 59

3.5.1. UV/O3 oxidation......................................................... 59

3.5.2. HF Vapor.................................................................... 63

3.6 Conclusions........................................................................................ 66


3.8 References.......................................................................................... 67

4. Chemical Vapor Cleaning of (0001)Si, (000-1)C, (11-20), and (10-10)

6H-SiC surfaces............................................................................... 70

4.1 Abstract............................................................................................. 71

4.2 Introduction....................................................................................... 72


4.3.1. Integrated Surface Prep. and Analysis System........... 75

4.3.2. Substrate Preparation.................................................. 78

4.4 Results............................................................................................... 79

4.4.1. (0001)Si 6H-SiC.......................................................... 79

4.4.2. (000-1)C 6H-SiC......................................................... 86

ix

4.4.3. (11-20) and (10-10) 6H-SiC........................................ 91

4.4.4. Low vacuum CVC/LPCVD clean............................... 92

4.5 Discussion.......................................................................................... 93

4.5.1. (0001)Si 6H-SiC.......................................................... 93

4.5.2. (000-1)C 6H-SiC......................................................... 96

4.5.3. (11-20) and (10-10) 6H-SiC........................................ 98

4.6 Conclusions........................................................................................ 99


4.8 References.......................................................................................... 100

5. X-ray Photoelectron Diffraction of (3x3) and (√3x√3)R30°

6H-SiC (0001)Si surfaces............................................................... 105

5.1 Abstract............................................................................................. 106

5.2 Introduction....................................................................................... 107

5.3 Experimental...................................................................................... 116

5.4 Results............................................................................................... 118

5.4.1. (v3xv3)R30° (0001)Si 6H-SiC..................................... 127

5.4.2. (3x3) (0001)Si 6H-SiC.................................................. 130

5.5 Discussion.......................................................................................... 132

5.5.1. (v3xv3)R30° (0001)Si 6H-SiC..................................... 132

5.5.2. (3x3) (0001)Si 6H-SiC.................................................. 133

5.6 Conclusions........................................................................................ 136


5.8 References.......................................................................................... 137

x

6. Interaction of Atomic Hydrogen with (3x3) 6H-SiC (0001)Si surfaces. 140

6.1 Abstract............................................................................................. 141

6.2 Introduction....................................................................................... 142

6.3 Experimental...................................................................................... 146

6.4 Results............................................................................................... 149

6.4.1. Interaction with rf atomic H........................................ 149

6.4.2. Interaction with thermal atomic H............................... 155

6.5 Discussion.......................................................................................... 159

6.6 Conclusions........................................................................................ 173


6.8 References.......................................................................................... 174

7. Ex Situ and In Situ Methods for Oxide and Carbon Removal from

(0001) AlN and GaN Surfaces........................................................ 179

7.1 Abstract............................................................................................. 180

7.2 Introduction....................................................................................... 180


7.3.1. Integrated Surface Prep. and Analysis System............ 185

7.3.2. Samples and Ex Situ Preparation................................. 188

7.4 Results............................................................................................... 189

7.4.1. Ex Situ Cleaning of AlN.............................................. 189

7.4.2. In Situ Cleaning of AlN.............................................. 200

7.4.3. Ex Situ Cleaning of GaN............................................ 209

7.4.4. In Situ Cleaning of GaN............................................. 221

xi

7.4.5. Ex Situ Cleaning of AlxGa1-xN................................... 227

7.5 Discussion.......................................................................................... 229

7.5.1. As Recieved and UV/O3 Surfaces............................... 229

7.5.2. Wet Chemical and HF Vapor Processing................... 234

7.5.3. Thermal Desorption and Capping Layers.................. 240

7.5.4. Chemical Vapor Cleaning and H Plasma.................... 245

7.6 Conclusions........................................................................................ 249


7.8 References.......................................................................................... 250

8. X-ray Photoelectron Spectroscopy Analysis of GaN/AlN and AlN/GaN

Growth Mechanisms....................................................................... 258

8.1 Abstract............................................................................................. 259

8.2 Introduction....................................................................................... 260

8.3 Experimental...................................................................................... 262

8.3.1. Integrated Growth Analysis System........................... 262

8.3.2. Substrate and Thin Film Preparation........................... 264

8.3.3. Growth Mode Analysis............................................... 266

8.4 Results............................................................................................... 268

8.4.1. GaN Growth on (0001) AlN....................................... 268

8.4.2. AlN Growth on (0001) GaN....................................... 274

8.5 Discussion.......................................................................................... 276

8.5.1. GaN Growth Mechanism on (0001) AlN.................... 276

8.5.1.1 Strain Effects...................................... 276

xii

8.5.1.2. Hydrogen Desorption....................... 279

8.5.2. AlN Growth Mechanism on (0001) GaN.................... 287

8.5.3. Surface Reconstruction................................................ 287

8.6 Conclusions........................................................................................ 288


8.8 References.......................................................................................... 289

9. Interface Chemistry and Electronic Structure for the

(0001) 2H-AlN/6H-SiC Interface....................................................293

9.1 Abstract............................................................................................. 294

9.2 Introduction....................................................................................... 294

9.3 Experimental...................................................................................... 295

9.4 Theory............................................................................................... 296

9.5 Results............................................................................................... 297

9.6 Discussion.......................................................................................... 302

9.7 Conclusions........................................................................................ 303


9.9 Addendum.......................................................................................... 304

9.10 References.......................................................................................... 305

10. Dependence of (0001) GaN/AlN Valence Band Discontinuity on

Surface Reconstruction and Growth Temperature..................... 307

10.1 Abstract............................................................................................. 308

10.2 Introduction....................................................................................... 308

10.3 Experimental...................................................................................... 309

xiii

10.3.1. Thin Growth and Analysis........................................ 309

10.3.1. GaN/AlN ?Ev Analysis............................................. 311

10.4 Results............................................................................................... 313

10.5 Discussion.......................................................................................... 316

10.6 Conclusions........................................................................................ 320

10.7 Addendum.......................................................................................... 320


10.9 References.......................................................................................... 322

11. Gas-Source Molecular Beam Epitaxy growth of ScN on (111)/(0001)

3C and 6H-SiC Substrates............................................................. 324

11.1 Abstract............................................................................................. 325

11.2 Introduction....................................................................................... 325

11.3 Experimental Procedures.................................................................... 327

11.4 Results............................................................................................... 329

11.4.1. Thermodynamics....................................................... 329

11.4.2. Growth....................................................................... 333

11.5 Discussion.......................................................................................... 337

11.6 Conclusions........................................................................................ 338


11.8 References.......................................................................................... 338

xiv

List of Tables

Table 1.1. Selected materials properties of Si, GaAs, 6H-SiC, 2H-GaN and 2H-AlN...................................................................................... 3 Table 2.1. Binding energy (in eV) of core level positions from (0001)Si 6H-SiC as polished and oxidized surfaces......................................... 18 Table 2.2. Summary of XPS Si2p/O1s, Si2p/F1s, C/C, and C-C data for (0001)Si 6H-SiC surfaces............................................................. 19 Table 2.3. Summary of wetting characteristics of as polished and oxidized (0001)Si 6H-SiC and (111) Si.......................................................... 20 Table 2.4. Peak to peak height (pph) ratios for various 6H-SiC surfaces................ 22 Table 3.1. XPS core level binding energies (eV) from (0001)Si 6H-SiC surfaces after various exposures........................................................ 57 Table 3.2. The XPS core level intensity ratios from (0001)Si 6H-SiC after various treatments .................................................................... 58 Table 3.3. Summary of SiC-C1s/surface C1s and Si/O intensity ratios from XPS data.................................................................................... 62 Table 5.1. Expected forward scattering/focusing peaks from bulk terminated (111) Si, (111) 3C-SiC, and (0001)Si 6H-SiC................. 127 Table 5.2. Expected Si 2p and C 1s photoelectron diffraction peaks for adatom scattering in T4 and H3 positions......................................... 133 Table 5.3. Estimated forward scattering peaks for (3x3) reconstructed (111)/(0001) 3C/6H-SiC surfaces...................................................... 135 Table 6.1. Kinetic Parameters used to model hydrogen adsorption/ desorption from Si and C sites on SiC surfaces................................. 162

xv

Table 7.1. OKLL/NKLL and CKLL/NKLL AES pph ratios from OMVPE AlN surfaces given various wet chemical treatments following a UV/O3 oxidation............................................................. 195 Table 7.2. XPS core level positions and full width half maxima (Γ, FWHM) from a GSMBE AlN surface after dipping in 10:1 BHF and annealing at various temperatures...................................................... 197 Table 7.3. XPS core level positions from a 200Å GaN capping layer on AlN after annealing at various temperatures................................ 206 Table 7.4. Ratio of integrated intensity of Al 2p to Ga 3d and Ga 2p3/2.................. 206 Table 7.5. XPS core level positions for OMVPE GaN surfaces after various ex situ treatments (Γ= FHWM)............................................ 211 Table 7.6. AES pph ratios of UV/O3 and wet chemical processed OMVPE GaN surfaces...................................................................... 216 Table 7.7. XPS core levels from GSMBE GaN surfaces after various treatments... 222 Table 7.8. Ga and N core levels from GaN relative to the GaN VBM after various processes....................................................................... 226 Table 7.9. Bond energies of Cl, F, and H with Al, Ga, and N.................................. 235 Table 9.1. Valence-band maxima and core levels measured on the same Au 4f7/2 based reference scale............................................................ 301 Table 9.2. Various measurements of Si2p-Al2p from a AlN/6H-SiC...................... 305 Table 10.1. Published data for GaN/AlN ?Ev.......................................................... 309 Table 10.2. Al 2p and N 1s core levels referenced to AlN VBM............................. 315 Table 10.3. CL-VBM data for AlN and GaN reported by various investigators..... 317 Table 11.1. Properties of ScN................................................................................... 327 xvi

List of Figures

Figure 2.1. Particle accumulation on removal from solution for a hydrophilic surface and a hydrophobic surface..................................................... 9 Figure 2.2. AES survey spectra of (a) (0001)Si 6H-SiC after thermal oxidation and removal of the oxide with 10:1 HF, (b) (0001)Si 6H-SiC as polished surfaces after solvent cleaning........................................ 13 Figure 2.3. XPS spectra of the F 1s core level from an oxidized (0001)Si 6H-SiC surface after (a) oxide removal with 10:1 HF, followed by (b) rigorous rinsing in running DI water............................................ 14 Figure 2.4. XPS spectra of the O 1s core level from (a) (0001)Si 6H-SiC after removal of a 750Å thermal oxide with 10:1 HF, and (b) (111) Si after a dip in 10:1 HF..................................................... 15 Figure 2.5. Typical XPS spectra of the C 1s core level from (a) oxidized (0001)Si 6H-SiC after oxide removal with 10:1 HF, (b) as-polished (0001)Si 6H-SiC after solvent cleaning, and (c) as-polished (0001)Si 6H-SiC after RCA SC1.............................. 17 Figure 2.6. EELS spectra of (0001)Si 6H-SiC (a) oxidized followed by 10:1 HF, (b) as-polished with only solvent cleaning......................... 18 Figure 2.7. The AES survey spectra of various 6H-SiC surface orientations after removal of a thermal oxide using 10:1 HF............ 21 Figure 2.8. XPS spectra of the F 1s core level from as-polished (0001)Si 6H-SiC, (a) after solvent cleaning, (b) Piranha etch, and (c) RCA SC1............................................................................... 25 Figure 2.9. AES survey spectra of as-polished (0001)Si 6H-SiC after (a) solvent cleaning, (b) a 10:1 HF dip, (c) 5 min. Piranha etch, and (d) 30 min. RCA SC1 clean........................................................ 25 Figure 2.10. (a) AES spectrum from 20Å a-Si/(0001)Si 6H-SiC after a 10:1 HF dip. (b) after thermal desorption of Si passivation layer at 1100°C.................................................................................. 27 Figure 2.11. XPS spectra of Si 2p core level from silicon passivated

xvii (0001)Si 6H-SiC (a) before thermal desorption and

(b) after thermal desorption at 1100°C.............................................. 28 Figure 2.12. (a) Schematic of (0001)Si 6H-SiC surface after thermal oxide removal with 10:1 HF. (b) Schematic of as-polished (0001)Si 6H-SiC................................................................................ 29 Figure 2.13. (a) Schematic illustrating mechanism of hydrogen termination of silicon in HF solutions. (b) Schematic illustrating stability of F- or OH- termination of SiC in HF solutions rather than H termination. (c) Schematic illustration of crystal potential in SiC.... 33 Figure 2.14. Energy diagram of (a) Si and (b) SiC in aqueous solutions................. 34 Figure 3.1. (a) Schematic of UV/ O3 oxidation system. (b) Schematic of HF vapor procedure...................................................................... 53 Figure 3.2. XPS of the C 1s core level from (0001)Si 6H-SiC after (a) solvent cleaning, sequentially followed by (b) UV/O3, and (c) HF vapor exposures............................................................... 56 Figure 3.3. XPS of the F1s core level from (0001)Si 6H-SiC after (a) solvent cleaning, sequentially followed by (b) UV/O3, and (c) HF vapor exposures............................................................... 56 Figure 3.4. XPS of the Si 2p core level from (0001)Si 6H-SiC after (a) solvent cleaning, followed by (b) UV/O3, and (c) HF vapor exposures..................................................................... 57 Figure 3.5. Schematic illustrating the surface termination from UV/O3 and HF vapor exposures on (0001)Si 6H-SiC................................... 60 Figure 3.6. Schematic illustrating the mechanism for F- and OH- termination of (0001)Si 6H-SiC........................................................ 65 Figure 4.1. AES of (0001)Si 6H-SiC surfaces after (a) 200£ SiH4 at 750°C, (b) 200£ SiH4 at 820°C, and (c) 200£ SiH4 at 880°C....................... 80 Figure 4.2. LEED patterns from (a) HF dipped (0001)Si 6H-SiC, (b) (1x1) (0001)Si 6H-SiC, (c) (3x3) (0001)Si 6H-SiC, (d) (v3xv3)R30° (0001)Si 6H-SiC, (e) (1x1) (000-1)C 6H-SiC, (f) (11-20) 6H-SiC, (g) (10-10) 6H-SiC............................................ 83 xviii

Figure 4.3. XPS spectra of the Si 2p core level from a (3x3) reconstructed (0001)Si 6H-SiC surface.................................................................... 84 Figure 4.4. XPS spectra of the Si 2p core level from (3x3), (1x1) and (v3xv3)R30° reconstructed (0001)Si 6H-SiC surfaces...................... 84 Figure 4.5. AES survey spectra from (v3xv3)R30° reconstructed (0001)Si 6H-SiC surfaces prepared by (a) SiH4 CVC, and (b) thermal desorption....................................................................... 85 Figure 4.6. XPS spectra of the C 1s core level from (v3xv3)R30° reconstructed (0001)Si 6H-SiC surfaces prepared by (a) SiH4 CVC, and (b) thermal desorption........................................ 85 Figure 4.7. AES survey spectra from (0001)C 6H-SiC surfaces prepared by (a) thermal desorption, (b) SiH4 CVC, and (c) SiH4/C2H4 CVC..... 87 Figure 4.8. EELS spectra from (000-1)C 6H-SiC surfaces after (a) annealing in UHV at 1050°C, (b) annealing in 2000£ SiH4 at 1050°C followed by (c) annealing in 2000£ C2H4 at 950°C.......................... 88 Figure 4.9. XPS spectra of the C 1s core level from (000-1)C 6H-SiC surfaces after (a) 2000£ SiH4 at 1050°C, (b) thermal desorption in UHV at 1050°C.............................................................................. 89 Figure 4.10. EELS spectrum from (v3xv3)R30° reconstructed (0001)Si 6H-SiC surface prepared by annealing in UHV at 1050°C.. 89 Figure 4.11. EELS spectra of (3x3), (1x1), and (v3xv3)R30° reconstructed (0001)Si 6H-SiC surfaces prepared via SiH4 CVC............................. 90 Figure 4.12. EELS spectra from (000-1)C 6H-SiC surfaces after (a) 2000£ SiH4 at 1000°C, (b) 400£ C2H4 at 850°C, and (c) 800£ C2H4 at 850°C............................................................................................. 90 Figure 4.13. EELS spectra of (10-10) 6H-SiC (a) after annealing in SiH4 at 1000°C, and (b) then annealing in C2H4 at 850°C.......................... 92 Figure 5.1. Schematic illustrating forward focusing/scattering effects xix

in x-ray photoelectron diffraction experiments.................................. 108 Figure 5.2. Schematics illustrating various adatom adsorption sites for (v3xv3)R30° reconstructions on (111)/(0001) surfaces. (a) Top down view along [000-1], (b) Side view along [11-20]........ 111 Figure 5.3. Model proposed by Kulalov et al for the (3x3) reconstructed (0001)Si 6H-SiC surface. (a) Top down view along [000-1], (b) side view along [11-20], and (c) [10-10]...................................... 113 Figure 5.4 Model proposed by Kaplan for the (3x3) reconstructed (0001)Si 6H-SiC surface. (a) Top down view along [000-1], (b) side view along [11-20]................................................................ 114 Figure 5.5. Top down view of model proposed by Li and Tsong for the (3x3) reconstructed (0001)Si 6H-SiC surface......................... 115 Figure 5.6. XPD pattern from (2x1) Si (100) along the [110] azimuth.................... 118 Figure 5.7. Si 2p x-ray photoelectron diffraction pattern along [1-21]/[10-10] azimuths from (a) (7x7) Si (111), (b) (3x3) 6H-SiC (0001)Si, and (c) (v3xv3)R30° (0001)Si 6H-SiC............................................... 119 Figure 5.8. Si 2p x-ray photoelectron diffraction patterns from (a) (3x3) 6H-SiC (0001)Si along [01-10], (b) (3x3) 6H-SiC (0001)Si along [10-10], and (c) (v3xv3)R30° 6H-SiC (0001)Si along [01-10]........................................................................ 120 Figure 5.9. Si 2p x-ray photoelectron diffraction patterns along [-110]/[11-20] from (a) (7x7) Si (111), (b) (3x3) 6H-SiC (0001)Si, and (c) (v3xv3)R30° 6H-SiC (0001)Si...................................................... 121 Figure 5.10. C 1s XPD patterns from (v3xv3)R30° 6H-SiC (0001)Si along (a) [01-10], and (b) [11-20]................................................................ 122 Figure 5.11. C 1s XPD patterns from (3x3) 6H-SiC (0001)Si along (a) [10-10], (b) [11-20], and (c) [10-10] azimuths................................................ 123 Figure 5.12. Schematic illustrating the differences in stacking along the [111]/[0001] direction for 3C and 6H-SiC........................................ 125 xx

Figure 5.13. Schematic illustrating expected forward scattering/focusing peaks in XPD along the [11-20] azimuth of 3C/6H-SiC............................. 126 Figure 5.14. Schematic illustrating expected forward scattering/focusing peaks in XPD along the [10-10] azimuth for 3C/6H-SiC............................ 126 Figure 6.1. AES of (3x3) reconstructed (0001)Si 6H-SiC (a) before remote H plasma, and (b) after remote H plasma.......................................... 152 Figure 6.2. XPS of the Si 2p core level from (3x3) reconstructed (0001)Si 6H-SiC (a) before remote H plasma, and (b) after remote H plasma............................................................................................ 152 Figure 6.3. TPD of (1x1) 6H-SiC (0001)Si after remote H plasma exposure (1 min., 20 W, 15 mTorr, and 450°C), (ß = 1°C/sec.)........................ 153 Figure 6.4. TPD of (a) sample heating stage after outgassing, and (b) molybdenum plate after remote H plasma exposure.................... 153 Figure 6.5. XPS of the C1s core level from (0001)Si 6H-SiC (a) before remote H plasma, (b) after remote H, and (c) after annealing at 1000°C....... 154 Figure 6.6. TPD of Si (111) after room temperature exposure to 2000£ H2 with rhenium filament at > 1700°C (ß=1°C/sec.).............................. 156 Figure 6.7. AES of (3x3) 6H-SiC (0001)Si (a) before atomic H exposure and (b) after room temperature exposure to 2000£ H2 with hot filament at > 1700°C.......................................................................... 157 Figure 6.8. TPD of (0001)Si 6H-SiC after room temperature exposure to 2000£ H2 with rhenium filament at > 1700°C (ß=1°C/sec.).............. 158 Figure 6.9. TPD of (3x3) 6H-SiC (0001)Si after cooling to 300°C in 10-6 Torr SiH4 (ß=1°C/sec.)............................................................... 159 Figure 6.10. Mono-hydride surface coverage on silicon sites of SiC....................... 164 Figure 6.11. Di-hydride surface coverage on silicon sites of SiC............................. 165 xxi

Figure 6.12. Hydrogen surface coverage on carbon sites of SiC based on kinetic data of Hamza....................................................................................165 Figure 6.13. Hydrogen surface coverage on carbon sites of SiC based on kinetic data of Thomas et al........................................................................... 166 Figure 6.14. Percent dissociation of H2 into H as a function of temperature........... 168 Figure 7.1. AES survey spectra of OMVPE AlN: (a) as received, (b) solvent cleaned and 20 min. UV/O3 exposure, and (c) 3 min. dip in 10:1 buffered HF (BHF)............................................................................ 192 Figure 7.2. XPS of O 1s core level from bulk AlN wafer (a) as recieved, (b) UV/O3 exposure, and (c) 10:1 BHF............................................. 192 Figure 7.3. XPS of Al 2p core level from bulk AlN wafer after a 10:1 BHF dip..... 193 Figure 7.4. XPS of N 1s core level from bulk AlN wafer after a 10:1 BHF dip....... 193 Figure 7.5. AES of (0001) OMVPE AlN after UV/O3 oxidation and oxide removal with (a) 1:1 NH3OH:H2O2, (b) 1:1 HCl:DI, (c) 10:1 BHF, (d) RCA SC1, and (e) RCA SC2.............................................. 195 Figure 7.6. Close up of AlLVV from AES of OMVPE AlN after (a) UV/O3 and (b) 10:1 BHF...................................................................................... 196 Figure 7.7. XPS of the F 1s core level from a 30Å AlN GSMBE film on (0001) 6H-SiC after (a) dipping in 10:1 BHF, and annealing for 15 min. at: (b) 400°C, (c) 600°C, (d) 800°C, and (e) 950°C............................ 198 Figure 7.8. XPS of the C 1s core level from a 30Å AlN GSMBE film on (0001) 6H-SiC after (a) dipping in 10:1 BHF, and annealing for 15 min. at: (b) 400°C, (c) 600°C, (d) 800°C, and (e) 950°C............................ 198 Figure 7.9. XPS of the O 1s core level from a 30Å AlN GSMBE film on (0001) 6H-SiC after (a) dipping in 10:1 BHF, and annealing for 15 min. at: (b) 400°C, (c) 600°C, (d) 800°C, and (e) 950°C............................ 199 Figure 7.10. XPS of Na 2p from polycrystalline AlN surface after (a) etching in NaOH and (b) an RCA clean............................................................. 200 xxii

Figure 7.11. TPD of m/e- (a) 18, (b) 20 and (c) 38 from 10:1 BHF dipped AlN (ß = 20°C/min.)......................................................................... 202 Figure 7.12. XPS of F 1s core level from polycrystalline AlN wafer cleaned in 10:1 BHF (a) before and (b) after remote H plasma exposure at 450°C (15 mTorr, 20W)................................................................. 204 Figure 7.13. AES survey spectra of OMVPE AlN after: (a) UV/O3 and 10:1 BHF dip and (b) remote H plasma at 450°C..................................... 204 Figure 7.14. XPS of O 1s from (0001) GSMBE AlN after (a) annealing at 1000°C and (b) annealing in a 0.1 ML/sec flux of Al at 1000°C....... 205 Figure 7.15. The XPS O 1s core level from an AlN surface (a) before and (b) after annealing in a SiH4 flux....................................................... 206 Figure 7.16. XPS of O 1s from In capping layer on OMVPE AlN (a) as recieved, (b) after annealing at 600°C, and (c) after annealing at 750°C............................................................................................. 207 Figure 7.17. XPS of O 1s core level from 200Å GaN capping layer on (0001) AlN buffer layer, (a) as recieved, (b) after annealing at 500°C, (c) 750°C, (d) 950°C, and (e) > 1000°C................................. 208 Figure 7.18. XPS of C 1s core level from 200Å GaN capping layer on (0001) AlN buffer layer, (a) as recieved, (b) after annealing at 500°C, (c) 750°C, (d) 950°C, and (e) >1000°C.................................. 208 Figure 7.19. XPS of Ga 2p3/2 core level from 200Å GaN capping layer on (0001) AlN buffer layer, (a) as recieved, (b) after annealing at 500°C, (c) 750°C, (d) 950°C, and (e) >1000°C.................................. 209 Figure 7.20. AES survey spectra from GSMBE GaN after (a) 1 day in air on laminar flow bench, (b) UV/O3 oxidation, and (c) 5 min. etch in 1:1 HCl:DI.............................................................................. 210 Figure 7.21. AES survey spectra from (001) GaAs (a) before and (b) after UV/O3. 211 Figure 7.22. XPS of C 1s core level from (0001) OMVPE GaN xxiii

after (a) ultrasonification in trichloroethylene, acetone, and methanol, and (b) UV/O3 exposure.................................................... 212 Figure 7.23. AES survey scan of OMVPE GaN after a 24 hr UV/O3 exposure with 1L/sec. flowing O2..................................................................... 213 Figure 7.24. XPS of O 1s core level from (0001) OMVPE GaN (a) after solvent cleaning, (b) after UV/O3 oxidation for 25 min., and (c) after UV/O3 oxidation for 24 hr. with 1000 L/sec flowing O2..... 214 Figure 7.25. AES of (0001) OMVPE GaN after UV/O3 oxidation and oxide removal with (a) 1:1 HCl:DI, (b) 10:1 BHF, and (c) 1:1 NH3OH:H2O2 (spectra normalized to NKLL)......................... 217 Figure 7.26. AES of (0001) OMVPE GaN after UV/O3 oxidation followed by (a) RCA SC1 and (b) RCA SC2 (spectra normalized to NKLL)......... 217 Figure 7.27. AES of (0001) OMVPE GaN (a) HCl:DI dip, (b) UV/O3 oxidation, and (c) HF vapor cleaning.................................................................. 219 Figure 7.28. XPS of F1s core level from (0001) GaN after UV/O3 oxidation followed by (a) 10:1 BHF vapor clean, and (b) a DI rinse................ 219 Figure 7.29. XPS of Ga2p core level from (0001) OMVPE GaN after UV/O3 oxidation and (a) BHF vapor clean and (b) DI rinse and NH3OH:H2O2 clean........................................................................... 220 Figure 7.30. XPS of N 1s core level from (0001) OMVPE GaN after UV/O3 oxidation and (a) BHF vapor clean and (b) DI rinse and NH3OH:H2O2 clean........................................................................... 220 Figure 7.31. m/e- 69 (Ga) signal from GSMBE GaN as a function of surface temperature............................................................................ 223 Figure 7.32. AES of (0001) OMVPE GaN after (a) HCl vapor clean, and annealing at (b) 300°C, (c) 450°C, and (d) 600°C............................. 223 Figure 7.33. AES survey spectra from GSMBE GaN (a) exposed to UV/O3, and after annealing in: (b) UHV at 650°C, 20 min., (c) NH3 at 650°C, 20 min, and (d) NH3 at 800°C, 25 min................................... 224 xxiv

Figure 7.34. AES survey spectra from GSMBE GaN after various sequential scans (a) 1, (b) 3, (c) 8, and (d) 9 scans............................................. 224 Figure 7.35. UPS of OMVPE GaN after (a) rinsing in methanol, (b) 1:1 HCl:DI, and (c) annealing in NH3 at 800°C, (d) as grown (2x2) GSMBE GaN........................................................................... 225 Figure 7.36. AES survey spectra from OMVPE GaN (a) after oxide removal with 1:1 HCl:DI, (b) 100°C H plasma exposure, and (c) 450°C H plasma exposure............................................................ 227 Figure 7.37. AES of AlGaN (a) as recieved, (b) after solvent cleaning in trichloroethylene, acetone, methanol, and isopropanol, (c) 1:1 NH3OH:H2O2, (d) 1:1 HCl:DI, and (e) 10:1 BHF...................................................................................... 228 Figure 7.38. AES of AlGaN after (a) UV/O3 exposure and (b) HF vapor oxide removal.................................................................................... 228 Figure 7.39. XPS of Al 2p core level after UV/O3 oxidation and (a) HF vapor oxide removal, and (b) DI rinse and NH3OH:H2O2..... 229 Figure 7.40. Schematic illustrating alignment of Cl- and F- ions with VBM of GaN and AlN in 1:1 HCl:DI and 10:1 HF respectively................... 237 Figure 8.1. Attenuation of Al 2p core level from AlN buffer layer as a function of overlying GaN film thickness for Tsub=650°C................ 270 Figure 8.2. LEED diffraction patterns from (a) (1x1) (0001) GaN, and (b) (2x2) (0001) GaN......................................................................... 270 Figure 8.3. SEM micrographs at 10 kX from GaN films grown in GSMBE at (a) 650°C, (b) 750°C, and (c) 800°C.................................................. 271 Figure 8.4. Photoluminesence (PL) at 4K of NH3-GSMBE GaN grown at (a) 650°C and (b) 800°C..................................................................... 272 Figure 8.5. Attenuation of Al 2p core level from AlN buffer layer as a function of overlying GaN film thickness for Tsub=800°C................ 273 xxv

Figure 8.6. Attenuation of Ga 3d and 3p core levels from OMVPE GaN as a function of overlying GaN film thickness for Tsub = 800°C.............. 275 Figure 8.7. Hydrogen surface coverage of Ga sites as a function of temperature and flux.............................................................................................. 283 Figure 8.8. Hydrogen surface coverage of N sites as a function of temperature and flux.............................................................................................. 284 Figure 8.9. Hydrogen surface coverage of Al sites as a function of temperature and flux.............................................................................................. 284 Figure 9.1. XPS spectra (arbitrary units) and theoretical densities of states (in states per unit cell per eV) of 6H-SiC.......................................... 298 Figure 9.2. XPS spectra (arbitrary units) and densities of states (in states per unit cell per eV) of 2H-AlN......................................................... 300 Figure 10.1. UPS Spectra of (2x2) (0001) AlN surface............................................ 313 Figure 10.2. UPS spectra of (2x2) (0001) GaN surface............................................ 315 Figure 11.1. ScN equilibrium diagram computed using HSC.................................. 330 Figure 11.2. AlN equilibrium diagram computed using HSC.................................. 330 Figure 11.3. GaN equilibrium diagram computed using HSC.................................. 331 Figure 11.4. GaN equilibrium diagram computed using experimental data of Thurmond and Logan, and Munir and Searcy............................... 331 Figure 11.5. InN equilibrium diagram computed using HSC and experimental data of Jones and Rose....................................................................... 332 Figure 11.6. AES survey scan of ScN film grown at 800°C on (111) 3C-SiC......... 334 Figure 11.7. XPS spectrum of N 1s and Sc 2p3/2,1/2 2p3/2,1/2 core levels xxvi

from a ScN film grown at 800°C...................................................... 335 Figure 11.8. UPS spectrum of ScN film grown at 800°C on (111) 3C-SiC............. 335 Figure 11.9. TEM of ScN film grown at 800°C on (111) 3C-SiC............................ 336 Figure 11.10. TEM of 800°C ScN on 5000Å (0001) GaN/AlN/6H-SiC.................. 336

xxvii

1. Introduction

Surfaces are the starting points for all semiconductor processes and hence represent

the foundation on which all subsequent microelectronic device fabrication processes are built.

Accordingly, proper surface preparation is an absolutely critical first step in all

semiconductor processes in order to achieve optimum results. The consequences of improper

surface preparation/cleaning are numerous leading to less than optimized results through

increased epitaxial defect densities (dislocations and stacking faults), lower dielectric

breakdown voltages, higher densities of interface states, variations in threshold voltage,

increased Ohmic contact resistance, and variations in Schottky barrier heights. All of these

effects manifest themselves in the bottom line through decreased device performance and

yield. In turn, these observations, have led to numerous studies of the surfaces of

technologically important semiconductors surfaces such as Si, Ge, GaAs, and InP. These

studies have shown that the entire chemical, electrical, and structural state of a semiconductor

surface must be considered when developing surface processes [1-4].

For the above reasons, the author has chosen to examine the surfaces and interfaces

of SiC, GaN, and AlN. These compounds are wide band gap semiconductors and are of

significant technological importance due to their extreme materials properties (see Table 1.1)

[5,6]. As documented in Table 1.1 [7], AlN and SiC are extremely strong materials with high

melting points and have accordingly long been of interest as high temperature structural

ceramics [5]. However, SiC exhibits both n and p type conductivity which with its large

band gap and high temperature stability has gained it significant consideration as a possible

semiconductor for high temperature microelectronic devices. The high thermal conductivity,

large breakdown voltage, and moderate electron and hole mobilities of SiC has also made it

of interest for possible high power and high frequency device applications. Due to the

insulating properties of AlN and moderately close lattice matching to SiC (?a/ao = 0.9%),

AlN/SiC quantum well structures and MISFET device applications have also been recently

proposed and/or fabricated [5].

However many of the same extreme materials properties exhibited by SiC are also

exhibited by GaN (see Table 1.1). Accordingly, GaN is also of interest for use in high

temperature, high power, and high frequency device applications [5,6]. In contrast to SiC,

however, GaN possesses a direct band gap which can be varied from the visible region into

the deep UV through alloying with InN and AlN. Accordingly, GaN and III-V nitride

materials are of considerable interest for numerous blue/UV optoelectronic device

applications [5,6]. The recent demonstration of a InGaN quantum well blue laser diode by

Shuji Nakamura of Nichia Chemical Co. is a highlight of the recent advances achieved in the

III-V nitride field [8].

2

Table 1.1. Selected materials properties of Si, GaAs, 6H-SiC, 2H-GaN and 2H-AlN [7].

Si GaAs 6H-SiC GaN AlN Eg (eV) 1.1 (i.) 1.43 (d) 3.01 (i) 3.40 (d) 6.2(i) Tmelt (°C) 1415 1238 2830 2200 2850 µn (cm2/V sec) 1350 8500 600 2000 µp (cm2/V sec) 480 400 40 150 14 E (GPa) 130 85.5 480 300 400 σT (W/cm K) 1.5 0.5 4.9 1.3 2.0 Vsat (cm/sec) 1x107 3x107 2x107

2.5x107 EB (MV/cm) 0.25 0.3 2.5 2.0 a (Å) 5.43 5.65 3.08 3.19 3.11 c (Å) 15.12 5.185 4.982 (α 106/K) 6.0 3.59 4.2 5.59 4.2 FJ (V/sec) 5x1011 13x1011 125x1011 159x1011

(EB vsat/2p)

3

In the following chapters of this text, a detailed examination of both ex situ and in situ

surface cleaning processes for (0001) surfaces of 6H-SiC, AlN, and GaN is presented. The

presented findings are based on results obtained from a variety of surface analytical

techniques used to characterize these surfaces including: Auger electron spectroscopy (AES),

x-ray photoelectron spectroscopy (XPS), low energy electron diffraction (LEED), electron

energy loss spectroscopy (EELS), temperature programmed desorption (TPD), and x-ray

photoelectron diffraction (XPD). Emphasis was first placed on understanding the nature of

the native contaminants present on these surfaces during ambient exposure. Processes were

accordingly then developed for the removal of these contaminants and the effects of these

processes on these surfaces were investigated. In addition, the surface processes involved in

the growth of GaN and AlN in gas source molecular beam epitaxy were also examined and

modeled. Finally, the heterojunction valence band discontinuity between AlN and SiC at the

(0001) interface was also examined.

1. W. Kern, RCA Review, 39, 278 (1978). 2. W. Kern, J. Electrochem. Soc., 137, 1887 (1990). 3. T. Ohmi, J. Electrochem. Soc., 143, 1957 (1996). 4. D.E. Aspnes and A.A. Studna, Appl. Phys. Lett., 39, 316 (1981). 5. R.F. Davis, Proc. of IEEE, 79, 702 (1992). 6. S. Strite and H. Morkoc, J. Vac. Sci. Technol. B, 10 1237 (1992). 7. J.H. Edgar, Ed., Properties of Group III Nitrides, Inspec, London (1994). 8. S. Nakamura, S. Masayuki, and S. Yasunobo, Appl. Phys. Lett., 68, 3269 (1996).

4

2. Wet Chemical Processing of (0001)Si 6H-SiC:

Hydrophobic and Hydrophilic Surfaces

To Be Submitted for Consideration for Publication

to the

Journal of the Electrochemical Society.

by

Sean W. King, Robert J. Nemanich, and Robert F. Davis,

Department of Materials Science and Engineering

North Carolina State University,

Raleigh NC 27695

5

2.1 Abstract

The wetting characteristics of oxidized and as-polished (0001)Si 6H-SiC (on and off

axis) surfaces in various acids and bases were compared to that of (111) Si. Auger electron

spectroscopy, x-ray photoelectron spectroscopy, and low energy electron diffraction were

used to characterize the chemical state and order of these surfaces. It was observed that the

oxidized SiC surfaces were hydrophilic after oxide removal with a 10:1 HF solution and were

terminated with approximately a monolayer of OH, CO, CH, and F species. In contrast, as-

polished SiC surfaces, were observed to be hydrophobic and covered with a thin (5-10Å)

contamination layer composed primarily of C-C, C-F, and Si-F bonded species. Removal of

this contamination layer using an RCA SC1 etch or Piranha clean resulted in a disordered

hydrophilic SiC surface. Similar effects were observed for other orientations of 6H-SiC (i.e.

(000-1)C, (11-20), and (10-10)). A 20Å amorphous Si capping layer was identified as an

alternative passivation layer for producing hydrophobic SiC surfaces.

6

2.2. Introduction

Preparation of clean, structurally well ordered surfaces is an important first step in all

semiconductor microelectronic fabrication processes [1-3]. Experience gained in silicon

technology has shown that the criteria for surface cleanliness must include removal of not

only native oxides and organic contaminants but also metallic impurities, particulate

contaminants, adsorbed molecules, and residual species left by previous processes [1-3]. The

effects of incomplete removal of all of these various contaminants by non-optimized surface

cleaning processes has been observed to consequently result in decreased device performance

and yield [4,5]. This is primarily through the generation of increased densities of electrical

defects (higher interface state densities, lower breakdown fields, increased leakage currents)

[6-13] and structural defects (dislocations and stacking faults) in epitaxial layers [14-25]. For

these reasons, an intensive effort has been made to understand the nature and source of

surface contaminants accumulated during silicon microelectronic processing [26-69].

Accordingly, these studies have developed a wide range of wet and dry (ex situ and in situ)

surface cleaning processes specifically optimized for silicon surfaces

[1,2,26,28,29,34,40,51,54-58].

In the case of SiC, numerous studies have been concerned with the nature and

removal of native oxides and other contaminants on silicon carbide surfaces [70-97].

However, few of these studies [87-92] have actually considered and/or investigated the effect

of wet chemical processes on SiC surfaces and any differences that might exist between

silicon and silicon carbide wet chemical processing. Therefore due to a lack of such detailed

studies for silicon carbide surfaces, many of the wet chemical processes optimized for silicon

7

have been implemented in SiC device processing [70-97]. This is in despite of the fact that

differences should be expected in the surface chemistry of these two semiconductors based

on the extreme chemical inertness of SiC compared to Si [98,99]. Thus given the plethora of

examples of the deleterious effects of non-optimized surface cleaning processes, it is evident

that further optimization of silicon carbide surface cleaning processes should result in a

reduced level of electrical and structural defects in silicon carbide device structures. Further

reduction in epitaxial and electrical defects should assist the development of SiC into the

semiconductor of choice for high speed, high frequency, and high temperature electronic

devices as well as the substrate of choice for III-V nitride heteroepitaxy [99,100].

In this paper, we investigate the effect on silicon carbide surfaces of various wet

chemical processes common to silicon. Emphasis is placed on HF processes which typically

serve as one of the last steps in silicon wet chemical processing and is primarily used for

surface oxide removal [40-69]. This process has become common due to the fact that it

produces a hydrogen terminated/hydrophobic silicon surface which is stable against

oxidation in air for several hours/days [56,63-69]. In this paper, we closely examine the

effect of HF based chemical processes on (0001)Si 6H-SiC surfaces. In this study, we have

observed that HF processing of oxidized (0001)Si 6H-SiC and other surface orientations (i.e.

(000-1)C, (11-20), and (10-10)) always leaves a hydrophilic SiC surface. Based on x-ray

photoelectron and Auger electron spectroscopy, we have additionally determined that HF

processing leaves a SiC surface terminated predominantly with O (or OH) rather than H.

This is in contrast with HF processing of silicon surfaces which is known to leave a

hydrophobic hydrogen terminated silicon surface [56,63-69].

8

The importance of controlling the wetting characteristics of surfaces in silicon

technology has been shown to not only be important from a surface termination point of view

but also important in reducing particulate and metallic contamination during wet chemical

processing [40-51]. In principle, wet chemical processing of hydrophobic surfaces should

lead to lower levels of particulate contamination due to a "snow plow" effect as the surface

traverses the liquid/air interface (see Figure 2.1). However in reality, particles have been

empirically observed to be more attracted to hydrophobic surfaces then hydrophilic surfaces

during wet chemical processing. This observation may be related to the fact that hydrophilic

surfaces tend to acquire a negative charge in solution which electrically repels particles in

solution which also have a negative charge [3, 41-43]. Hydrophobic surfaces on the other

hand are more passivated and develop less of a charge (or positive charge) in ionic solutions

and therefore negatively charged particles in the solution are attracted to the hydrophobic

surface [3, 41-43].

(a) (b)

9

Figure 2.1. Particle accumulation on removal from solution for (a) a hydrophilic surface and (b) a hydrophobic surface.

In the case of SiC, we have previously noted that the hydrophilic nature of SiC

surfaces in wet chemical processing leads to trapping of wet chemicals in the micro pipes of

SiC wafers and the unwanted removal/outgassing and incorporation of these chemicals in

following processes [101]. For this reason and those mentioned above, we have additionally

investigated other alternative processes and passivation layers which could lead to a

hydrophobic (0001)Si 6H-SiC surface. In particular, we have identified the as as-polished

(0001)Si 6H-SiC surface as being hydrophobic. However, XPS, AES and low energy

electron diffraction (LEED) analysis indicated this surface was composed of a thin (≈ 5-10

Å) contamination layer composed mainly of CFx and SiFx species. Removal of this

contamination layer using RCA SC1 (1:1:5 NH3OH:H2O2:H2O) or Piranha (7:3

H2SO4:H2O2) cleaning results in a defective and disordered surface. In contrast, we have

found that a hydrophobic SiC surface can be obtained using an amorphous 20Å Si capping

layer. The excess silicon may be removed by thermal desorption resulting in a structurally

well ordered SiC surface as evidenced by a (3x3) LEED pattern.

2.3. Experimental Procedure

On axis and vicinal n-type (typically Nd=1018/cm3) (0001)Si 6H-SiC wafers were

used in these experiments. These wafers were supplied by Cree Research, Inc. with or

without a 500-1000Å thermally grown oxide. Both the oxidized and as-polished, as-received

10

6H-SiC substrates were ultrasonically cleaned/degreased in trichloroethylene, acetone, and

methanol each for 10 min. prior to any other wet chemical treatments. The oxide was

removed from the thermally oxidized wafers using a 10 min. dip in a 10:1 HF solution,

followed by rinsing in de-ionized (DI) water and N2 blow drying. The wetting

characteristics of this surface and the as-polished surfaces were then investigated by

immersion in other acid/base solutions. The wet chemistries examined included 100:1 HF,

10:1 HF, 1:1 HF, 10:1 buffered HF (7:1 NH4F:HF), 30:1 buffered HF, 40% NH4F, 38%

HCl, 70% HNO3, Piranha Etch (7:3 H2SO4:H2O2@120°C), RCA SC1 & SC2 (1:1:5

NH3OH:H2O2:H2O@85°C) and 1:1:5 HCl:H2O2:H2O@85°C), 100% HC2H3O2, and 40%

KOH. These chemistries were chosen for examination primarily due to their extensive use in

the silicon microelectronics industry. Except where noted, after all wet chemical treatments

the samples were rinsed in DI water (18 M?) and blown dry with N2. At this point, the wafer

was visually inspected to determine whether a hydrophobic or hydrophilic surface had been

retained or obtained. To assist in the comparison to silicon, (100) and (111) Si (on and off

axis) wafers were dipped in the same acid/base solution immediately after the SiC wafer and

the wetting characteristics of these surfaces were compared. All wet chemicals were of

CMOS grade purity (J.T. Baker).

SiC wafers with a 200Å amorphous Si capping layer were prepared in the following

manner. First, an oxidized (0001)Si 6H-SiC wafer was given a 10 min. dip in 10:1 HF, DI

rinsed followed by N2 drying, and loaded into a Si-Ge MBE where it was then degassed at

450°C and annealed at 1000°C for 20 min. in a 10-6 Torr SiH4 flux. This produced an

oxygen free, Si rich (1x1) SiC surface. Next, an electron beam was used to evaporate/deposit

≈ 200Å of amorphous Si onto the SiC surface at room temperature. The a-Si/SiC sample was 11

then removed from vacuum for subsequent wet chemical processing as described above for

oxidized and as-polished SiC surfaces.

Surfaces prepared as described above were further subjected to surface analysis in an

integrated ultra-high vacuum (UHV) system incorporating the following analytical

techniques: x-ray photoelectron spectroscopy (XPS), Auger electron spectroscopy (AES),

electron energy loss spectroscopy (EELS), and low energy electron diffraction (LEED). The

details of these systems and the transfer line have been previously described [102]. After

each wet chemical treatment, the SiC wafer was mounted to a molybdenum sample holder

and loaded into a load lock for subsequent analysis by AES, XPS, EELS, and LEED. The

XPS analysis was performed using a VG CLAM II electron energy analyzer and an Al anode

(hν=1486.6 eV) at 20 mA and 12 kV. The AES spectra were obtained using a Perkin Elmer

CMA, a beam voltage of 3 keV and an emission current of 1 mA. The EELS spectra were

also obtained using the CMA but with a 100 eV electron beam and an emission current of 1

mA. The LEED was performed using Princeton Scientific rear view optics, a beam voltage

of approximately 115 eV, and an emission current of 1 mA. Calibration of the XPS binding

energy scale was performed by measuring the position of the Au 4f7/2 and shifting the

spectra such that the peak position occurred at 83.98 eV.

2.4. Results

2.4.1. (0001)Si 6H-SiC, Oxidized

12

In contrast to (100) and (111) Si, a 10:1 HF solution was visually observed to wet the

surface of (0001)Si 6H-SiC as this wafer was withdrawn from the HF solution after a 10 min.

dip (Note: this is sufficiently long to remove the 750Å thermal oxide since the SiO2 etch rate

with 10:1 HF is ˜ 10 Å/sec [60-62]). The 10:1 HF solution was noticed to slowly pool up on

the SiC surface over time. However after rinsing in DI water, the SiC surface was observed

to be clearly wetted by H2O (i.e. no pooling up of the H2O). Figure 2.2(a) displays an AES

survey spectrum taken from an off axis (0001)Si 6H-SiC wafer after removal of the 750Å

thermal oxide with a 10 min. 10:1 HF dip (Note: similar results were obtained from on axis,

(0001)Si 6H-SiC as well). Si, C, and significant amounts of oxygen were detected by AES.

Using XPS, further analysis of the (0001)Si 6H-SiC surface after a 10:1 HF dip, also

revealed the presence of significant amounts of fluorine, ≈ 1/4 ML (see Figure 2.3(a)), on the

surface. The fluorine was not detected by AES due to a lower sensitivity to fluorine and

possible electron beam stimulated desorption effects [103-107]. By analogy to intentionally

fluorinated silicon surfaces where the F 1s peak was located at 685.9-686.2 eV [103-107], the

F 1s peak at 686.8 eV observed from the oxidized (0001)Si 6H-SiC surface after the HF dip

was likewise attributed to Si-F bonded fluorine. The fluorine surface coverage was initially

observed to vary from 0-1/4 ML, but further investigation revealed that the fluorine surface

coverage was highly dependent on the DI rinsing procedure. In fact, fluorine was not

detected by XPS for HF dipped surfaces rigorously rinsed in DI water (see Figure 2.3(b)).

13

100 200 300 400 500 600 700

dN(E

)/dE

Electron Energy (eV)

(a)

(b)

Si

C

NO

F

Figure 2.2. AES survey spectra of (a) (0001)Si 6H-SiC after thermal oxidation and removal of the oxide with 10:1 HF, (b) (0001)Si 6H-SiC as polished surfaces after solvent cleaning.

680 682 684 686 688 690 692

(a)

(b)

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

Figure 2.3. XPS spectra of the F 1s core level from an oxidized (0001)Si 6H-SiC surface after (a) oxide removal with 10:1 HF, followed by (b) rigorous rinsing in running DI water.

14

Comparison of the relative intensities of the O KLL AES transition and the XPS O 1s

core level from oxidized (0001)Si 6H-SiC and (100) and (111) Si surfaces after oxide

removal with 10:1 HF indicated a significantly larger surface concentration of oxygen on SiC

surfaces in comparison to Si surfaces (see Figure 2.4). The intensity of these two

peaks/transitions were estimated to correspond to oxygen surface coverage's of ≈ 3/4±1/4

ML for (0001)Si 6H-SiC and < 1/10 ML for (100) and (111) Si. This is in agreement with

the observed hydrophilic and hydrophobic nature of these surfaces respectively. More

detailed analysis of the (0001)Si 6H-SiC surface using photoemission from the O 1s core

level indicated the presence of two peaks. The larger peak at 532.1 eV is indicative of Si-O

or Si-OH bonded oxygen, and the smaller peak at 533.5-533.9 eV is indicative of C-O

bonding [30,31,56,57]. The intensity of the C-O O 1s bonding peak like the F 1s Si-F

bonding peak was observed to depend on the DI rinsing procedure with no DI rinsing

resulting in the observance of the most C-O bonded oxygen at the SiC surface. The presence

of C-O bonded oxygen on the (0001)Si surface was further supported by photoemission from

the C 1s core level which detected two C 1s peaks, one at 282.8 eV indicative of C-Si bonds,

and one at 284.7 eV indicative of a mixture of C-H and C-O bonds [19,20] (see Figure

2.5(a)). In contrast, EELS did not detect the presence of any p-p* transitions (≈ 3-6 eV)

typically observed from organic molecules and contamination (see Figure 2.6(a)). As

mentioned by Mizokawa et al [89], the failure of EELS to detect adventitious surface carbon

could be due to electron stimulated desorption.

15

526 528 530 532 534 536 538 540

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

Si-O

C-O

Figure 2.4. XPS spectra of the O 1s core level from (a) (0001)Si 6H-SiC after removal of a 750Å thermal oxide with 10:1 HF, and (b) (111) Si after a dip in 10:1 HF.

Due to the inherent asymmetry of the XPS Si 2p core level arising from the

unresolved Si 2p3/2,1/2 doublet, it was difficult to accurately determine if a Si-O Si 2p

bonding peak actually existed. Assuming a FWHM of 1.45 eV for the main Si-C Si 2p peak

at 100.7 eV, a second peak at ˜ 102.2 eV could be fitted to the spectrum, though the line

width of this peak was quite narrow < 1.4 eV. However, the presence of significant amounts

of Si-O bonded oxygen on the (0001)Si surface was supported by the Si KLL line shape in

AES (see Figure 2.2) which as noted by Mizokawa is similar to that of oxidized silicon

surfaces [89]. The intensity of the Si-O O 1s or Si 2p bonding peaks were not observed to

depend on the DI rinsing procedure with the O 1s (Si-O)/Si 2p (Si-C) intensity ratio

remaining essentially constant and independent of wafer cut (i.e. off or on axis). Finally,

these surfaces displayed intense (1x1) LEED patterns with broad dots which were clearly

16

visible at beam energies (Ep) as low as 60-100 eV. The AES, XPS, and EELS results from

the oxidized (0001)Si 6H-SiC surface are summarized in Tables 2.1 and 2.2.

280 282 284 286 288 290 292

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(c)

(b)

(a)

Figure 2.5. Typical XPS spectra of the C 1s core level from (a) oxidized (0001)Si 6H-SiC after oxide removal with 10:1 HF, (b) as-polished (0001)Si 6H-SiC after solvent cleaning, and (c) as-polished (0001)Si 6H-SiC after RCA SC1.

17

-30 -25 -20 -15 -10 -5 0 5

Cou

nts (

arb.

uni

ts)

Electron Loss Energy (eV)

(a)

(b)

š-š *

Figure 2.6. EELS spectra of (0001)Si 6H-SiC (a) oxidized followed by 10:1 HF, (b) as-polished with only solvent cleaning.

Table 2.1. Binding Energy (in eV) of core level positions from (0001)Si 6H-SiC as polished and oxidized surfaces after various treatments. The full width half maxima (Γ) of the peaks are also indicated.

Treatment Si 2p, Γ O1s, Γ C1s, Γ F1s, Γ N1s, Γ Oxidized 10:1 HF 100.7, 1.4 532.1, 1.9 282.8, 1.1 284.7, 2.1 As polished Solvents 100.5, 1.4 531.9, 2.3 282.6, 1.1 685.8, 1.9 398.2, 3.0 283.7, 2.7 687.5, 2.5 286.0, 4.5 10:1 HF 100.5, 1.4 531.6, 2.4 282.5, 1.1 685.6, 1.9 398.2, 3.0 283.6, 2.7 687.3, 2.5 286.0, 4.4 H2SO4:H2O2 100.4, 1.4 531.5, 2.5 283.5, 1.1 686.0, 3.0 398.2, 3.0 283.6, 2.7 RCA SC1 100.3, 1.4 531.3, 2.8 282.4, 1.1

18 283.9, 2.4

Aqua Regia 100.5, 1.4 531.6, 2.3 282.6, 1.1 685.5, 1.7 398.2, 3.0 283.7, 2.7 686.9, 3.2

able 2.2. Summary of XPS Si2p/O1s, Si2p/F1s, C/C, and C-C data (uncorrected for

Treatment Si2p/O1s Si2p/F1s C/C C-C

Tsensitivity factors) for (0001)Si 6H-SiC surfaces.

Oxidized

1.4 ∞ 6.65 1.9 eV ed

.2 1.6 1.1 1.3 eV

10:1 HF As Polish Solvents 110:1 HF 1.8 1.8 0.9 1.1 eV H2SO4:H2O2 0.9 4.0 1.1 1.1 eV RCA SC1 1.0 ∞ 2.2 1.5 eV Aqua Regia 1.2 3.8 1.2 1.1 eV

19

Following the 10:1 HF dip, the thermally oxidized treated (0001)Si 6H-SiC surfaces

were dipped in a variety of different acids and bases and the surface wetting characteristics

visually noted. The results are summarized in Table 2.3 and illustrate that this surface was

found to be hydrophilic in all acids and bases investigated. In all cases, the SiC surface was

observed to retain a monolayer coverage of oxygen. For the oxidized (0001)Si 6H-SiC

surface, the surface coverage of oxygen was not found to change appreciably with dipping

time (1-24 hr.), HF concentration (1:1 - 1000:1), composition (HF-NH4F), or pH (1-10).

Table 2.3. Summary of wetting characteristics of as polished and oxidized (0001)Si 6H-SiC nd (111) Si.

(0001)Si 6H-SiC (0001)Si 6H-SiC (111) Si &

a

Treatment As Polished Thermally Oxidized a-Si Passivated 6H-SiC

None Hydrophobic Hydrophilic Hydrophilic10:1 HF Phobic Philic Phobic 38% HCl Phobic Philic Phobic 70% HNO3 Phobic Philic Philic RCA SC1 Philic Philic Philic RCA SC2 Phobic Philic Philic Piranha Philic Philic Philic Aqua Regia Phobic Philic Philic Acetic Phobic Philic PhobicNH4F Phobic Philic Phobic KOH Phobic Philic Philic

2.4.2. (000-1)C, (11-20), and (10-10) 6H-SiC surfaces

Thermally oxidized surfaces of other orientations of 6H-SiC such as (000-1)C, (11-

20), and (10-10) were also investigated. After removal of the thermal oxide with 10:1 HF,

these surfaces were also observed to be hydrophilic in all acids and bases investigated and

very little difference was observed between these surfaces and those with the (0001)Si

orientation. Figure 2.7 displays a series of AES spectra obtained from all four orientations

investigated. The spectra were obtained after removal of the thermal oxide with 10:1 HF

(note that all spectra are normalized to the C KLL Auger transition). Table 2.4 lists the Si/C,

O/Si, and O/C pph ratios (uncorrected for differences in sensitivity) calculated for each

surface. Table 2.4 indicates that the O/C ratio for all the different orientations of 6H-SiC

surfaces is centered around 0.3. Given that two of these surfaces are polar ((0001)Si and

20

(000-1)C) and the others non polar ((11-20) and (10-10)), it is surprising that they would

exhibit this similarity. Further, the Si/C pph ratio for the (0001)Si, (11-20), and (10-10)

surfaces are all centered around 0.6 which is surprising in that the ideal (0001)Si surface

would be terminated exclusively with Si, whereas the (10-10) and (11-20) are non polar

surfaces ideally with equal numbers of carbon and silicon atoms at the outermost surface.

The Si/C ratio as shown in Table 2.4 for the (000-1)C surface, however, is half that found for

the (0001)Si surface. The lower Si/C pph ratio for the (000-1)C face is expected based on the

differences in polarity for these two surfaces.

30 130 230 330 430 530 630 730

Si LVV

O KLL

C KLL

N KLL

dN(E

)/dE

(b)


(c)

(d)

2.7. The AES survey spectra of various 6H-SiC surface orientations after reml oxide using 10:1 HF. (a) (0001)Si, (b) (000-1)C, (11-20), and (d) (10-10).

(a)

Figure oval of thermaa

21

Table 2.4. Peak to peak height (pph) ratios for Various 6H-SiC surfaces (uncorrected for ensitivity factors).

(0001)

s

Si (000-1)C (11-20) (10-10) Si/C 0.6 0.3 0.6 0.6

.4 O/Si 0.5 1.0 0.4 0 O/C 0.3 0.3 0.2 0.3

2.4.3. (0001)Si, As-Polished

As-polished (0001)Si 6H-SiC on and off axis surfaces which were not oxidized were

observed to be hydrophobic as received without any further processing. Figure 2.2(b)

displays an AES survey spectrum obtained from an as-polished, off axis (0001)Si 6H-SiC

surface after ultrasonic degreasing in trichloroethylene, acetone, and methanol for 10 min.

each. In contrast to hydrophilic SiC surfaces in which a thermally grown oxide had been

removed with 10:1 HF, traces of N and F were detected by AES from the as-polished SiC

surfaces, as shown in Figure 2.2(b). The presence of F and N on the surfaces of these as-

polished SiC wafers was also subsequently confirmed by XPS analysis. A more detailed

analysis of the F 1s spectrum from the as-polished SiC surface revealed the presence of two F

1s peaks, one at 685.6 eV indicative of Si-F bonding [103-107] and a second peak at 687.3

eV indicative of C-F, N-F, or SiFx bonding [106, 107] (see Figure 2.8(a)). The N 1s peak

was observed at 398.2 eV and is indicative of Si-N bonding [108,109]. The surface

concentration of F and N based on the XPS data was estimated to be ≈ 1ML and < 1/10 ML

22

respectively. These surface concentrations were not observed to change appreciably with

subsequent HF processing.

Like the SiC surfaces which had undergone a thermal oxidation treatment, XPS

Figures 2.2(a) and 2.2(b), the oxygen surface coverage is

cleaning.

detected the presence of significant amounts of surface carbon from the solvent cleaned

surface as polished/unoxidized (0001)Si 6H-SiC surface. Figure 2.5(b) shows an XPS

spectrum of the C 1s core level obtained from this surface. This spectrum was fitted to three

peaks centered at 282.6, 283.7, and 286.0 eV indicative of C-Si, C-C, and C-F bonding

respectively [106,107] (see Table 2.1 for Si2p, O1s, C1s, N1s, and F1s core level positions).

Additionally, a loss peak at 5-6 eV indicative of the p-p* transition of graphite like carbon

was also detected in EELS (see Figure 2.6(a)). Further, a (1x1) diffraction pattern was barely

discernible in LEED at Ep=100eV and only clearly discernible at Ep ≈ 180 eV. This is

indicative of either a thin contamination layer as suggested by the non-SiC C 1s peaks

observed or due to a disordered surface from subsurface defects or damage produced by the

polishing treatment [96].

As can be determined from

lower for the solvent cleaned, as-polished (0001)Si 6H-SiC surface than the HF dipped

oxidized SiC surface. However, the oxygen surface coverage based on the AES O KLL and

XPS O 1s intensities still is ˜ 1/2 ML. The binding energy of the O 1s core level from the as

polished SiC surface is slightly larger than that from a SiC surface which had undergone an

oxidation treatment 531.1 (ox) vs. 531.6 eV (no-ox). The differences in binding energy are

most likely due to band bending. Small amounts of a C-O O 1s peak was also detected from

the as polished SiC surfaces after wet chemical processing, but was not detected after solvent

23

After dipping in 10:1 HF for 10 min., the as-polished surface was still observed to be

hydrophobic when rinsed in DI water. F, N, and surface carbon were also still detected by

acids and bases, but would become hydrophilic in

XPS, AES, and EELS. As indicated in Figure 2.9(a), the HF dip did remove some oxygen

from the surface and the XPS Si2p/O1s ratio was found to increase after the HF dip (see

Table 2.2), but the oxygen surface coverage was still ≈ 1/2 ML. The C1s(Si-C)/C1s(C-C)

ratio was found to decrease after the HF dip indicating that the HF dip left more surface

carbon on the surface. Additionally, the adventitious C1s peak was observed to shift from

283.9 eV to 283.6 eV (see Table 2.1).

After the 10:1 HF dip, the as-polished SiC surface was observed to remain

hydrophobic when dipped in various

extended dip/etches in RCA SC1 or H2SO4:H2O2 (Piranha etch). In some instances, this

hydrophilic surface could be made hydrophobic again by boiling in Aqua Regia (3:1

HCl:HNO3) for 5-10 min. Subsequent XPS analysis revealed complete removal of fluorine

from hydrophobic (0001)Si 6H-SiC surfaces which had been permanently converted to

hydrophilic surfaces by prolonged immersion in RCA SC1 (see Figure 2.8(c)). In cases

where hydrophilic (0001)Si 6H-SiC surfaces were observed to be converted back to

hydrophobic surfaces by a boiling Aqua Regia treatment, XPS revealed incomplete removal

of fluorine from the surface by the RCA SC1 or Piranha etch treatment (see Figure 2.8(b)).

In addition, the third C1s peak observed at 286.0 eV was observed to track with the fluorine

coverage possibly suggesting that this surface is terminated with a contamination layer of

fluorocarbons. As a further note, none of the wet chemical processes employed here were

successful in converting hydrophilic SiC surfaces to hydrophobic surfaces for surfaces which

had undergone a thermal oxidation treatment followed by oxide removal with HF. A

24

summary of the wetting characteristics observed for the as received SiC surface after various

wet chemical treatments is provided in Table 2.3.

680 682 684 686 688 690 692

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

C-FSi-F

Figure 2.8. XPS spectra of the F 1s core level from as-polished (0001)Si 6H-SiC, (a) after solvent cleaning, (b) Piranha etch, and (c) RCA SC1.

25

100 200 300 400 500 600 700

dN(E

)/dE


(c)

Si

C

NO

F

(d)

(b)

(a)

Figure 2.9. AES survey spectra of as-polished (0001)Si 6H-SiC after (a) solvent cleaning, (b) a 10:1 HF dip, (c) 5 min. Piranha etch, and (d) 30 min. RCA SC1 clean.

2.4.4. Si Passivating Layer

As the hydrophobic nature of as received/polished SiC surfaces seemed to be related

to a fluorocarbon contamination layer, a Si capping layer was investigated as a more

controllable alternative hydrophobic passivation layer for (0001)Si 6H-SiC surfaces. SiC

wafers terminated with an amorphous 200Å Si passivation layer were observed to be

hydrophobic after dipping in various HF and NH4F solutions. To investigate thinner Si

passivation layers, the 200Å Si passivation layer was thinned by repeated UV/O3 oxidation

and 10:1 HF dips. The Si passivation layer was observed to maintain a hydrophobic SiC

surface down to a Si thickness of ≈ 20Å, below which the SiC surface became increasingly

hydrophilic. However, this effect could be partially related to non-uniformity's in the

thickness of the Si passivation layer. The wetting characteristics of this Si passivated 26

(0001)Si 6H-SiC surface in other acids and bases after an HF dip were additionally observed

to be similar to those of silicon (see Table 2.3).

Figure 2.10(a) displays an AES spectrum from a 20Å Si/(0001)Si 6H-SiC surface

after a 10:1 HF dip and as can be seen lower oxygen and non-SiC carbon levels were

observed. This passivation layer was easily removed in vacuum by annealing at 1100°C for

5 min. prior to epitaxy. Figure 2.10(b) shows that a silicon rich, oxygen free, well ordered

surface characterized by a (3x3) LEED pattern was obtained. Figure 2.11 shows an XPS

spectrum of the Si 2p core level from the same surface and further illustrates the loss of the

silicon passivation layer by the reduction in the Si-Si bonding peak at 99.5 eV after the

1100°C anneal. As previously noted [101], the silicon passivation layer also resulted in

lower outgassing rates in vacuum due to lower levels of wet chemicals trapped in micro pipes

in the SiC wafer. Additional advantages and applications of the silicon capping layer for SiC

surfaces will be more fully discussed in the following section.

27

100 200 300 400 500 600 700

dN(E

)/dE


(a)

(b)

Si

C

O

Figure 2.10. (a) AES spectrum from 20Å a-Si/(0001)Si 6H-SiC after a 10:1 HF dip. (b) after thermal desorption of Si passivation layer at 1100°C.

95 97 99 101 103 105 107

(a)

(b)

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

Si-Si

Si-C

Figure 2.11. XPS spectra of Si 2p core level from silicon passivated (0001)Si 6H-SiC (a) before thermal desorption and (b) after thermal desorption at 1100°C. 28

2.5. Discussion

2.5.1. (0001)Si 6H-SiC, Oxidize

The data presented above illustrates that a significant fraction of the (0001)Si 6H-SiC

surface is covered with oxygen and adventitious carbon (≈1-2 ML) after oxide removal using

10:1 HF or other HF/NH4F compositions (see Figure 2.12(a)). As essentially all ex situ

cleaning practices are limited by the ambient in which they are conducted, the presence of

small quantities of contaminants such as oxygen and carbon may be expected. However, the

amounts detected on (0001)Si 6H-SiC after removal of a thermal oxide with HF are several

times (5-10X) larger than those detected on silicon in our laboratory under the same

processing conditions. Additionally, the observation of a hydrophilic surface as opposed to a

hydrophobic surface further illustrates that the chemistry and interactions occurring at Si and

SiC surfaces in HF are clearly different.

CO

CH

x

FOH

OH

OH

OH

C=O

OH

OH

OH

OH

OH

OH

OH

OH

OH

OH

OH

OH

OH

OH

OHF F CO

C=O

C= O

(a)

(b)

10Å C-C, C-F, Si-F Contamination Layer

29

Figure 2.12. (a) Schematic of (0001)Si 6H-SiC surface after thermal oxide removal with 10:1 HF. (b) Schematic of as-polished (0001)Si 6H-SiC.

The relatively large levels of oxygen remaining on the surface of SiC after an HF dip

are particularly surprising given that concentrated HF is known to etch silicon oxide at rates

as high as 1000Å/sec [2]. Clearly something is prohibiting the HF from removing the last

monolayer of oxygen or hydroxide from the SiC surface. One possible explanation for the

differences in oxygen coverage between Si and SiC is that the extra oxygen observed from

SiC surfaces is due to more strongly bound oxygen located at the steps of the SiC surface and

which are bonded to both Si and C. If this scenario were true, one would expect to see in

XPS and AES a difference in the O/Si ratio between on axis and off axis (0001)Si 6H-SiC

surfaces. However, this was not observed and it was found that the O/Si ratios were

essentially the same for both on and off axis wafers. Additionally, excess oxygen trapped at

SiC steps would simply not be large enough of an effect to explain the observed 1 ML

surface coverage.

Another possible explanation for the observed differences between Si and SiC is that

the oxygen is bonded to only carbon. Evidence of C-O bonding at the SiC surface is seen in

the C 1s and O 1s XPS spectra with peaks at 284.7 eV and 533.9 eV respectively. However,

XPS shows most of the oxygen to occur predominantly at 532.1 eV which is clearly

indicative of Si-O or Si-OH bonding. This is further supported by the Si LVV line shape in

Figure 2.1 which is also indicative of Si-O bonding.

In order to explain the apparent oxygen or OH termination of oxidized (0001)Si 6H-

SiC surfaces after an HF dip, it is necessary to first consider why hydrogen termination of Si

is achieved with HF and secondly to account for the polar and ionic nature of (0001) SiC 30

surfaces and HF solutions respectively. In the case of silicon, it was originally suggested that

HF processes produced Si-F terminated surfaces due to the relatively large bond strength of

Si-F compared to Si-H bonds ( 6 eV vs. 3.5 eV respectively) [59]. However, subsequent IR,

TPD, and HREELS analysis showed that HF processed silicon surfaces were in fact

terminated largely with hydrogen with < 1/10 ML fluorine coverage [66-69]. Trucks et al

[63] explained the hydrogen termination as being a result of the instability of Si-F bonds due

to the polarization of Si-Si backbonds by the strongly heteropolar Si-F bond (see Figure

13(a)). Polarization of the Si backbonds leaves these bonds susceptible to the strongly

polarized H+F- molecule which can then attack the backbond and fluorinate the silicon

surface atom and hydrogen terminate the nearby atom. This continued scenario eventually

leads to removal of the Si surface atom by complete fluorination (i.e. SiF4) leaving behind

only Si-H species. The stability of the hydrogen terminated silicon surface in HF can be

explained by comparison of the electronegativites of Si, H, and F which are 1.9, 2.2, and 4.0

respectively [110]. Due to the similarities in electronegativities of Si and H, the Si-H bond is

non polar (as compared to Si-F) and is therefore not attacked by HF. Trucks et al [63] has

additionally shown that the reverse reaction Si-H + HF => Si-F + H2 is energetically

unfavorable.

In the case of (0001)Si SiC surfaces, the underlying Si-C bonds are already polarized

(see Figure 2.13(b)) similar to Si-F bonds on Si due to the differences in electronegativity

between Si and C (1.9 and 2.6 respectively). Due to the stacking sequence along the [0001]

direction in SiC, the polarization of the Si-C bonds leads to the build up of a large crystal

potential/field which must be canceled in order to stabilize the crystal (see Figure 2.13(c))

[111]. In vacuum, cancellation of this field is achieved by desorption of surface atoms which

31

produces a compensating charge that cancels the internal field [111]. However in an ionic

solution such as HF, cancellation of this field can be achieved by simply adsorbing ionic

species of opposite polarity/charge. For the (0001)Si SiC surface, this would require the

adsorption of negatively charged ions such OH- or F- instead of hydrogen which is exactly

what is observed. Termination of the SiC surface with OH- is also supported by the recent

high resolutions EELS (HREELS) studies of Starke et al [90] which were able to identify the

O-H stretch from HF processed (0001) 6H-SiC substrates.

By analogy to silicon [56,65-69], the observation that the fluorine coverage on SiC

surfaces after an HF dip depended on the DI rinsing procedure suggests that the fluorine is

located at defects sites on the SiC wafer. During the DI rinse, the surface fluorine is removed

from the surface by saturation of these defect sites with OH- instead. However, we did not

observe a change in the O/Si ratio with or without DI rinsing. Therefore, this suggests that F

either substitutes for H in OH (i.e. Si-OF) instead of bonding directly with Si at the surface

(i.e. Si-F) or the fluorine coverage is << 1/4 ML such that a change in the O/Si ratio can not

be detected.

Another way at looking at this is to consider the quasi band diagram shown in Figure

2.14. In this figure, we consider the liquid (HF) to have a quasi Fermi level which is aligned

with the Fermi level of Si and/or SiC [112]. All of the ionic species in the HF solution have

an energy somewhere below this quasi Fermi level and which is dependent on the

electronegativity of the specie [3, 112]. For Si it has been postulated that the energy position

of the H+ ion in the HF solution is located very closely to the Si VBM where effective charge

transfer can occur and hence is the reason for hydrogen termination of silicon by HF [3,112].

SiC, however, has a larger bandgap and its VBM should lie below that of Si and hence

32

termination of the SiC surface with species which lie at lower energy levels such as OH- and

F- is expected.

F

Si

Si Si Si

Si

Si Si Si

F F

H

F

F

F FSi

Si

Si Si Si

FF

F

H H

Si Si Si

H H H

δ−

δ−

δ+ H - Fδ−δ+

δ−

H - Fδ−δ+H - Fδ−δ+

δ−

(a)

Si

C C Cδ−

δ+

F

Si

C C C

δ−

δ−

δ+OH-

Si

C C Cδ−

δ+

(b)

(c)

σ+σ−

σ+σ−

σ+σ−

σ+σ−

[0001]

Ε

ΟΗ−

ΟΗ

−

ΟΗ

−

ΟΗ

−

ΟΗ

−

ΟΗ

−

ΟΗ

−

ΟΗ

−

Figure 2.13. (a) Schematic illustrating mechanism of hydrogen termination of silicon in HF solutions. (b) Schematic illustrating stability of F- or OH- termination of SiC in HF solutions rather than H termination. (c) Schematic illustration of crystal potential in SiC.

33

Si10:1 HF

H+

OH-

F-

K+

SiC 10:1 HF

OH-

F-

SO42-

O3-

(a) (b)

EF

EF

Figure 2.14. Energy diagram of (a) Si and (b) SiC in aqueous solutions.

The OH- termination of SiC wafers is important not only from a chemical purity point

of view but also from the viewpoint of particle contamination during wet chemical

processing. In silicon microelectronics, it has been found that the zeta potential of silicon

surfaces in acids is negative and that particles in the solution are charged positively [3,41-43].

As a result, particles are electrostatically attracted to silicon surfaces in acid processing and

consequently acid processes generally result in higher levels of particle accumulation [3, 41-

43]. To account for this, silicon surfaces have to be cleaned in basic (high pH) solutions for

particle removal where both Si surfaces and particles acquire the same negative charge and

electrostatically repel one another [3, 41-43]. Due to the OH termination of SiC, one would

expect SiC surfaces however to exhibit a zeta potential pH dependence more similar to SiO2.

In fact, this is exactly what has been observed in zeta potential measurements on SiC

powders [113-115]. The importance of this is that the zeta potential pH dependence of SiO2

34

and SiC particles are very similar [3,41-43]. Therefore, particles should be repelled from SiC

surfaces in both acids and bases.

2.5.2. (0001)Si 6H-SiC, As-Polished

The above results have shown that the (0001)Si 6H-SiC surfaces, which are not

oxidized after polishing, are hydrophobic before and after dipping in 10:1 HF. However,

only slightly smaller amounts of oxygen were detected from the as-polished surfaces after

even an HF dip. As previously shown in Figure 2.4, the oxygen surface coverage for both as-

polished and oxidized SiC surfaces is 5-10 times higher than that for hydrophobic hydrogen

terminated Si surfaces. So clearly, the hydrophobic, as-polished SiC surface observed in this

study is not related to the hydrophobic, hydrogen terminated silicon surface. This is evident

simply from the observation that the as-polished (0001)Si 6H-SiC surface is hydrophobic in

HNO3 and H2SO4 whereas the hydrogen terminated (111) Si surface is not. As the most

significant difference observed between the oxidized and unoxidized SiC surfaces is the

identification of large amounts of fluorine (≈ 1-2 ML) on the as-polished surface, the

hydrophobic nature of this surface is most likely related to fluorine termination (C-F or Si-F)

instead of hydrogen termination.

As previously noted, two fluorine peaks were detected by XPS from the as-

polished/unoxidized (0001)Si 6H-SiC surface after solvent cleaning (see Figure 2.8). The

first peak was located at 685.8 eV and attributed to Si-F bonding based on previous

examinations of fluorinated silicon surfaces [103-107]. The second F 1s peak was detected at

687.5 eV and attributed to either C-Fx, N-Fx, or SiFx bonding [106,107]. Based on the

observation of a third broad C 1s peak at 286.0 eV (FWHM=4.1eV), we suggest that the

35

second F 1s peak is due to C-Fx bonding. Any SiFx species present on the SiC surface after

polishing would have been removed by the DI rinse or HF dip. It is possible that the nitrogen

detected by XPS and AES is simply nitrogen dopants which have segregated or which have

been preferentially left at the SiC surface during polishing. Hence, we propose that the

nitrogen is primarily bonded to Si (i.e. Si-N) at the SiC surface

As a LEED pattern was only observable from as-polished SiC surfaces at high beam

energies of ≈ 200 eV, we suggest that the as-polished SiC surface is terminated with a thin (≈

5-10Å) contamination or disordered/defective layer. This contamination layer is composed

primarily of a mixture of C-C, C-F, Si-F bonded species and is directly responsible for the

hydrophobic nature of the as-polished SiC surfaces. Once this contamination layer has been

completely removed by oxidation of the C-C and C-F bonds, the F 1s and C 1s peaks at 685-

687 and 286.0 eV respectively disappear and the C-C C 1s peak at 283.6 eV shifts to 283.9

eV indicative of more C-O bonding. After oxidation, this surface is irreversibly hydrophilic.

As far as why some hydrophilic as-polished SiC surfaces can be reverted back to being

hydrophobic by a boiling Aqua Regia treatment, the authors speculate that the Aqua Regia

treatment somehow replaces C-O bonds with C-H bonds.

As for the source of the fluorine detected on the as-polished SiC surfaces, the authors

offer two explanations. One possible source of the fluorine are the Fluoroware containers in

which the SiC wafers are shipped from Cree. The containers are made of natural

polypropylene which has a rather high outgassing rate for particles [116]. It has been

previously noted by others working on Si [117], that SiO2 films stored in these containers for

sufficiently long times can become apparently hydrophobic due the large concentrations of

hydrophobic particles found on them. However, we do not think that the wafers investigated

36

here were so contaminated. Therefore we think that the only other source of fluorine should

be Cree's polishing or etching procedure. Unfortunately, Cree has not been willing to divulge

or share any information concerning this. We do note, however, that we have observed

similar contamination layers from Si surfaces etched in CF4 RIE systems [118].

At this point, it is worth considering which surface would be better to work with: the

hydrophilic SiC surface or the as-polished hydrophobic SiC surface. The hydrophobic as

polished SiC surface has the advantage of minimizing trapping of wet chemicals in micro

pipes and reduced particulate contamination from wet chemical processing. However, the

results indicate that a thin contamination layer is responsible for the hydrophobic nature of

as-polished SiC surfaces. The hydrophilic SiC surface produced by thermal oxidation

followed by oxide removal with HF is most likely the more appropriate surface for devices

due to the better crystallinity and surface order. Thermal oxidation of as-polished SiC

surfaces not only removes the thin contamination layer but also oxidizes and removes much

of the subsurface damage present after polishing [96].

2.5.3. Si Passivation Layer

37

We now consider a more controllable passivation layer based on a 20-200Å

amorphous Si capping layer. As described in the results section, an amorphous Si capping

layer behaves similarly in acids and bases to Si (111) surfaces and may be easily removed in

situ by annealing in UHV at 1100°C. Hence, the a-Si passivation allows the application of

knowledge developed from ex situ processing of silicon. The use of this capping layer was

also found to result in lower outgassing rates in vacuum due to the production of a

hydrophobic surface preventing liquids from being trapped in micro pipes. The authors

would like to point out that the a-Si capping layer has the added advantage that it can

potentially be incorporated into already existing processing routes. Potentially, the Si

capping layer could be deposited during cooling from SiC thin film CVD epitaxy. Rupp et al

[119], have already currently demonstrated the ability to control the surface stoichiometry of

SiC epitaxial films by controlling the gas phase composition in their LPCVD system during

cooling. Therefore, the deposition of a 20-200Å amorphous or polycrystalline layer of Si

during cooling after SiC epitaxial growth is possible. Another advantage of the Si capping

layer is that it may be oxidized or nitrided to form the oxide/insulator for MOSFET/MISFET

structures. The advantage here is that the Si capping layer would protect the SiC surface

from various metallic and other contaminants during processing which could effect the

quality of the SiC/SiO2 interface. In instances where a hydrophilic surface is needed the Si

capping layer can be easily made hydrophilic by immersion in HNO3 or H2SO4.

38

2.6. Conclusions

In conclusion, it has been found that removal of a thermal oxide from (0001)Si 6H-

SiC surfaces using 10:1 HF results in a hydrophilic surface terminated primarily with Si-OH

and C-O species. Other crystallographic orientations ((000-1)C, (11-20), and (10-10)) were

also observed to be hydrophilic after oxide removal with 10:1 HF. In contrast, non-

oxidized/as polished (0001)Si 6H-SiC surfaces were observed to be hydrophobic as received.

This surface was observed be to terminated by a thin (5-10Å) disordered contamination layer

composed mainly of C-C, C-F, and Si-F bonded species. Removal of this contamination

layer using RCA SC1 or Piranha etch converted the surface to hydrophilic. As an alternative

passivation layer, a 20-200Å a-Si capping layer was demonstrated to produce a hydrophobic

SiC surface. This a-Si passivation layer was easily removed in situ via thermal desorption at

1100°C.

2.7. Acknowledgments

The authors would like to thank Cree Research Inc., for supplying the SiC wafers

used in these studies. This research was sponsored by the Office of Naval Research and

through the Department of Education via a an Electronic Materials/GAANN fellowship.

39

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46

3. Dry Ex Situ Cleaning Processes for (0001)Si 6H-SiC Surfaces

To be Submitted for Consideration for Publication

to the

Journal of the Electrochemical Society

by

Sean W. King, Robert J. Nemanich, and Robert F. Davis.


North Carolina State University

Raleigh, NC 27695

3.1. Abstract

47

A completely dry ex situ cleaning process based on UV/O3 oxidation for surface

carbon contamination removal and HF vapor exposure for oxide removal has been

demonstrated for (0001)Si 6H-SiC surfaces. Using x-ray photoelectron spectroscopy (XPS)

analysis for comparison, this cleaning procedure has been demonstrated to reduce residual

surface carbon and oxide contamination to levels equivalent to or better than conventional

wet chemical ex situ processing. Specifically, XPS showed that (0001)Si 6H-SiC surfaces

exposed to UV generated ozone were oxidized resulting in the formation of both carbon and

silicon oxides. This was clearly illustrated in XPS by the formation of a broad Si-O Si 2p

peak at 102.4 eV (FWHM=2.1 eV) and a shift in the surface C1s peak from 283.6 to 284.2

eV. A reduction in the amount of surface carbon was evidenced by an increase in the ratio of

the SiC C1s peak to the surface C1s from 0.8 to 2.7 after the UV/O3 treatment. Removal of

the UV/O3 silicon oxide via exposure to the equilibrium vapor from a 10:1 buffered HF

solution was deduced from the absence (below the XPS detection limit) of the Si-O Si 2p

peak at 102.4 eV. Significant amounts of fluorine remained on the surface after the HF vapor

exposure suggesting that the process results in a fluorine terminated surface.

48

3.2. Introduction

For SiC to succeed as the semiconductor/substrate of choice for high frequency, high

temperature, high power devices and as a substrate for III-N heteroepitaxy, a considerable

reduction in defects (line, planar, point, etc.) must be achieved [1,2]. Following Si

technology, where surface cleaning and preparation are critical first steps in all processes [3-

5], a continued reduction in defects in SiC should be expected as a result of improved SiC

wafer surface cleaning techniques. In Si technology for example, improper removal of

surface contamination and oxides prior to Si homoepitaxy has been shown to result in an

increase in the density of line and planar defects in epitaxial films from < 104/cm2 to >

1010/cm2 [6-8]. The increased defect densities were in turn found to correspond with a

reduction of device yield [6]. In the case of heteroepitaxy, studies on SixGe1-x alloy growth

on Si (100) have additionally shown that surface defects produced in the Si substrate by

residual organic/carbon contamination act as the preferred sites for misfit dislocation

generation [9]. These examples clearly illustrate that surface preparation and cleaning should

be equally important to the control of defects in both homo and heteroepitaxial growth of SiC

and III-V nitrides on (0001) 6H-SiC.

Due to a limited number of studies concerned with ex situ SiC cleaning practices [10-

13], most SiC ex situ wet chemical processing has been based on processes specifically

developed for and employed in Si technology [13,14]. SiC ex situ cleaning/surface

preparation has typically consisted of some variation of solvent degreasing, organic

contaminant removal using RCA or Piranha cleans, and finishing with oxide removal using

HF based solutions [10-14]. An important assumption underlying the use of these procedures

49

is that the SiC surfaces should behave similar to silicon surfaces in these wet chemicals. In a

previous study we have provided examples of where this assumption fails, specifically with

regard to oxide removal from SiC surfaces using an HF dip process. In Si technology, oxide

removal with a dilute HF etch is known to generate a hydrophobic, hydrogen terminated

surface, stable against oxidation in air for several hours [15-20]. However, we have

previously shown that SiC surfaces are inherently hydrophilic after oxide removal with HF

due to a preference for OH termination [21]. The hydrophilic surface allows water and HF to

become trapped in micropipes in the SiC wafer which can lead to large concentrations of

oxygen and fluorine at the SiC-dielectric interface if not properly outgassed. In order to

produce a hydrophobic surface, passivation layers based on silicon or fluorocarbons were

required [21].

An alternative to the use of passivation layers to form hydrophobic SiC surfaces

would be to develop a completely dry cleaning process and thus remove the need for wet

chemical processing. In Si and GaAs technology, dry removal of carbon contaminants from

surfaces using UV/O3 oxidation has become an alternative to wet chemical processing [22-

30]. In the UV/O3 oxidation process, UV radiation from a Hg lamp (specifically the 184.9

nm line) is used to photo excite molecular oxygen (O2) and generate ozone (O3) [22].

Additionally, the 253.7 nm line of Hg assists in removal of carbon contaminants as the light

is adsorbed by most hydrocarbons and excites C-H and C-C bonds [22]. Removal of the

UV/O3 generated oxide is typically achieved by wet chemical processing and/or in situ

thermal desorption [25,26,28,30]. However, Iyer et al [32] have shown that the equilibrium

vapor from an HF solution can be alternatively used to remove the oxide from a silicon wafer

via a dry process. Thus the combined use of UV/O3 oxidation for removal of carbon

50

contaminants and HF vapor exposure for oxide removal represents a completely dry cleaning

process. Use of completely dry processing techniques may eliminate the need for large

quantities of expensive high purity chemicals while simultaneously reducing the costs for

disposal of these toxic materials [32-33].

In this paper, we demonstrate for the first time a completely dry cleaning process for

(0001)Si 6H-SiC surfaces which is based on the combined use of UV/O3 oxidation and HF

vapor cleaning. This clean has been found to be equivalent to or better than typical wet

chemical processes in terms of residual surface carbon and oxide contamination levels as

measured by XPS. The combined UV/O3-HF vapor treatment eliminates the need for a

hydrophobic SiC surface and avoids the use of exotic passivation layers.


As polished on axis, n-type (typically Nd=1018/cm3) (0001)Si 6H-SiC wafers

supplied by Cree Research, Inc. were used in these experiments. Selection of the as polished

wafers for examination was based on previous investigations which showed these surfaces to

be terminated with a thin (˜5-10Å) contamination layer of C-C, C-F, and Si-F bonded species

[21]. Prior to UV/O3 oxidation, each wafer was first ultrasonically degreased in

trichloroethylene, acetone, and methanol for 10 min. each. The experimental system

employed for the UV/O3 exposures described in this study employed a high intensity Hg

lamp positioned in close proximity (≈ 1 cm) to the SiC wafer (see Figure 1). In order to

increase the concentration of generated O3 (i.e. to increase the oxidation rate), the UV/O3

box was purged with 1 L/sec O2 during the UV exposure. Further details of this process

51

have been described previously [22,26]. The HF vapor exposures were achieved by simply

positioning the SiC wafer in ambient air within approximately 5 mm of a 10:1 buffered HF

solution for times ranging from 5-30 min. (see Figure 1). Condensation of HF on the SiC

surface was not observed for exposures of this length.

Surfaces prepared in the above manner were subjected to surface analysis in an

integrated ultra-high vacuum system incorporating the following analytical techniques: x-ray

photoelectron spectroscopy (XPS), Auger electron spectroscopy (AES), electron energy loss

spectroscopy (EELS), and low energy electron diffraction (LEED). Details of this system are

given elsewhere [34]. After each treatment above, the SiC wafer was mounted to a

molybdenum sample holder and loaded into the load lock for subsequent analysis by AES,

XPS, EELS, and LEED. The XPS analysis was performed using the Al anode (hν=1486.6

eV) at 20 mA and 12 kV. The AES spectra were obtained using a beam voltage of 3 keV and

an emission current of 1 mA. The EELS spectra were obtained using a 100 eV electron beam

and an emission current of 1 mA. The LEED was performed using rear view optics, a beam

voltage of approximately 115 eV, and an emission current of 1 mA. Calibration of the XPS

binding energy scale was performed by measuring the position of the Au 4f7/2 core level

(from a ~1 µm thick Au film) and shifting the spectra such that the peak position occurred at

the accepted value of 83.98 eV [15].

52

b.)

HF

Hg UV Lamp

hν=185nm 254 nm

O3

SiC

a.)

Air or O2

Figure 3.1. (a) Schematic of UV/ O3 oxidation system. (b) Schematic of HF vapor procedure.

53

3.4. Results

3.4.1. Solvents and UV/O3:

Figures 3.2(a) and (b) show XPS spectra of the C 1s core level obtained from a as-

polished, solvent cleaned (0001)Si 6H-SiC surface followed by a UV/O3 oxidation treatment.

As is indicated in Figure 3.2(a), a broad C 1s feature was obtained. Previous analysis of this

spectrum [21] revealed the presence of three C 1s peaks centered at 282.5, 283.6, and 286.0

eV. The most intense peak at 282.5 is associated with carbon bonded to silicon in SiC.

Based on the large FWHM (full width half maximum) of 2.6 and 4.5 eV, the latter two peaks

were respectively assigned to a mixture of C-C and C-Hx (283.6 eV), and C-Fx (286.0)

bonded carbon [35-37]. The presence of C-Fx species was further supported by the XPS

spectra of the F 1s core level (see Figure 3.3(a)). The spectra from the as polished (0001)Si

6H-SiC surface after solvent cleaning shows two F 1s peaks located at 685.4 and 687.2 eV.

These two peaks have been assigned to Si-F [40-42] and C-F [35-37] bonding respectively.

The presence of a thin "fluorocarbon" contamination layer is also indicated by the inability to

obtain a LEED pattern from these surfaces at beam energies (Ep) < 200 eV.

Following a 2 hour UV/O3 exposure, the C-Fx C 1s peak was observed to disappear

and the C1s peak associated with C-C and C-Hx was observed to shift from 283.6 to 284.2

eV (see Figure 3.2(b) and Table 3.1). In contrast, the SiC C1s peak showed an increase in

intensity and shifted by only 0.1 eV to 282.6 eV. In addition, the ratio of the C1s peak

intensities associated with the SiC and the surface carbon (uncorrected for sensitivity factors)

was observed to increase from 0.8 to 2.7 (see Table 3.2). This result indicates that the

UV/O3 process removes the contamination layer via oxidation. The shift and reduction in

the C-C and CHx C1s peak is consistent with the formation of C-O bonding at the surface

and removal of some surface carbon via desorption of CO and CO2 [22,24,37]. Removal of

54

the contamination layer was also supported by the complete disappearance (below the XPS

detection limit) of the F 1s peak at 687.2 eV after the UV/O3 treatment (see Figure 3.3b).

Only a slight trace of the lower binding energy F 1s peak was detected, and it was observed

to shift by 0.5 eV to 685.9 eV (see Figure 3.3(b)).

The formation of silicon oxides on the SiC surface is displayed in the XPS of the Si

2p core level from the (0001)Si 6H-SiC surface before and after UV/O3 treatment (see

Figures 3.4(a) and (b)). As shown in Figure 3.4(a), a single Si 2p peak was detected before

UV/O3 oxidation. The line shape of this Si 2p peak is asymmetric suggesting the possibility

of a Si-O bonding peak on the higher BE (binding energy) side. Unfortunately,

deconvolution of this peak is complicated by the fact that the Si 2p peak is an unresolved

doublet (i.e. Si 2p3/2,1/2 ) and fitting a second peak to this spectrum showed only a small

peak at 102.2 eV with a FWHM narrower than the substrate peak (i.e. 1.1 vs. 1.4 eV). As

such, it was not possible to conclusively detect a Si-O peak prior to the UV/O3 exposure.

However after the UV/O3 exposure, a broad Si-O peak centered at 102.4 eV (FWHM=2.1

eV) was clearly observed (see Figure 3.4(b)). The width of the Si 2p peak at 102.4 eV

indicates that silicon in +2,+3, and +4 oxidation states is bonded to the oxygen (i.e. SiOx)

[16,38,39]. Based on the attenuation of the Si-C Si 2p peak, the thickness of the SiOx layer

was estimated to be < 20Å.

55

280 282 284 286 288 290

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

Figure 3.2. XPS of the C 1s core level from (0001)Si 6H-SiC after (a) solvent cleaning, sequentially followed by (b) UV/O3, and (c) HF vapor exposures.

680 682 684 686 688 690 692

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

Si-F C-F

56

Figure 3.3. XPS of the F1s core level from (0001)Si 6H-SiC after (a) solvent cleaning, sequentially followed by (b) UV/O3, and (c) HF vapor exposures.

96 98 100 102 104 106 108

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

Figure 3.4. XPS of the Si 2p core level from (0001)Si 6H-SiC after (a) solvent cleaning, followed by (b) UV/O3, and (c) HF vapor exposures.

Table 3.1. XPS core level binding energies (eV) from (0001)Si 6H-SiC surfaces after various exposures. The full width half maxima (Γ) are also included.

Si 2p, Γ C 1s, Γ O 1s, Γ F 1s, Γ Solvents 100.4, 1.5 282.5, 1.1 531.6, 2.3 685.4, 1.8 283.6, 2.6 687.2, 2.7 286.0, 4.5 UV/O3 100.5, 1.4 282.6, 1.1 532.1, 2.4 685.9, 1.9 102.4, 2.1 284.2, 2.1 HF Vapor 100.5, 1.5 282.6, 1.1 531.8, 2.3 685.8, 1.9 283.8, 2.8

57

Table 3.2. The XPS core level intensity ratios from (0001)Si 6H-SiC after various treatments (uncorrected for differences in sensitivity factors). The first ratio represents the C1s peaks attributed to carbon in SiC and to adventitious or surface carbon.

SiC C1s/surface C1s Si2p/O1s Si2p/F1s Solvents 0.8 1.1 5.4 UV/O3 2.7 0.3 10.6 HF Vapor 1.7 1.3 0.5

3.4.2. HF vapor

After oxidation of the (0001)Si 6H-SiC surface using a UV/O3 exposure, removal of

the thin silicon oxide layer was achieved by a 30 min. exposure of the SiC surface to a vapor

from a 10:1 buffered HF solution. As shown in Figure 3.4(c), the higher binding energy Si

2p peak centered at 102.4 eV was not detectable after the HF vapor exposure. Shorter HF

vapor exposures were observed to result in an observable Si-O Si2p peak but at larger

binding energies (≈104 eV). After the HF vapor treatment, the amount of surface carbon was

monitored from the intensity ratio of the SiC C1s to the surface carbon C1s peaks. The ratio

was observed to decrease from 2.7 to 1.7 after the HF vapor exposure indicating an increase

in the amount of surface carbon contamination. The surface C1s peak was likewise observed

to shift back to 283.8 eV and displayed a large increase in FWHM from 2.1 to 2.8 eV.

However, the SiC to surface carbon intensity ratio of 1.7 after the HF vapor treatment is still

much larger (i.e. less surface carbon) than the 0.8 value that was found after the solvent

cleaning process. In addition to an increase in surface carbon after the HF vapor treatment,

the amount of fluorine on the surface was also observed to increase significantly (see Figure

3.3(a)). Prior to the HF vapor treatment, the Si/F ratio was equal to 10.6. However, after the

HF vapor treatment the Si/F ratio decreased to 0.56. Further the binding energy of the F 1s

58

was not observed to shift but remained centered at 685.9 eV which suggests that fluorine is

bonded only to silicon atoms at the surface.

3.5. Discussion

3.5.1. UV/O3 Oxidation

In the above section, it was demonstrated that exposure of (0001)Si 6H-SiC surfaces

to ozone generated by a Hg UV lamp oxidized and removed adventitious and CFx bonded

carbon from the SiC surface (see Figure 3.5(a) and (b)). This resulted in an increase of the

SiC/non-SiC carbon ratio from 0.8 to 2.7. This result is in agreement with previous studies

of UV/O3 oxidation of Si and GaAs surfaces which have also shown a reduction of carbon

contaminants [22-30]. Despite the long exposure (2 hr.), some adventitious carbon or surface

carbon, however, was observed to be present on the SiC surface after the UV/O3 treatment.

Some of this surface carbon is likely due to recontamination of the SiC surface during sample

transfer and mounting in a laboratory ambient prior to insertion into vacuum. However,

contamination levels of this magnitude are usually not observed from silicon wafers cleaned

in the same environment. Alternatively, the remaining surface carbon could be due to carbon

trapped in the SiOx layer and/or carbon bonded to both silicon and oxygen at the SiC/SiOx

interface. We also note that the studies of Fominski et al [24] and Baunack and Zehe [30]

report incomplete removal of carbon contaminants from Si surfaces using O3 generated from

a Hg lamp. In fact Fominski et al [24] found that it was necessary to employ deeper UV

radiation from a D2 lamp and immerse the wafer in an O2/NF3/H2 gas mixture.

59

F F F F F F FOH

OH

OH

OH

F F F F F F FOH

OH

F

5-10Å C-C, CHx, and CFxcomtamination layer

- 20Å SiOx layer

UV/O3

HF Vapor

Figure 3.5. Schematic illustrating the surface termination from UV/O3 and HF vapor exposures on (0001)Si 6H-SiC.

The shift in the position of the surface C 1s peak from 283.6 to 284.2 eV with UV/O3

oxidation is consistent with the oxidation of C-C, CHx, and CFx bonds to form CO. For HF

dipped Si wafers, it has been previously determined that residual carbon contaminants with C

1s peaks positions of 284.6, 286.3, and 288.4 eV are composed mostly of C-H2, C-O, and O-

C=O bonded carbon respectively which illustrates the trend to higher binding energies for C-

60

O bonds [43]. Unfortunately, a direct comparison between the surface C1s peak position and

bonding configuration for both Si and SiC surfaces is complicated by the probable existence

of 0.5-1.0 eV of band bending at the SiC surface due to surface Fermi level pinning.

Although the UV/O3 oxidation treatment was not completely successful in removing

all of the non-carbidic carbon from the SiC surface, a comparison of this technique with wet

chemical processes does show the utility of the technique. In a previous study [21], we

examined the effect that standard wet chemical treatments such as RCA SC1 and Piranha

etch have on the removal of the same carbon surface contamination observed in this study.

Table 3.3 provides a direct comparison of the SiC C1s/surface C1s and Si2p/O1s intensity

ratios for each treatment. As can be seen, the UV/O3 treatment provides the highest SiC

C1s/surface C1s ratio of all the treatments examined. We also note that Afanas'ev et al [33]

has recently found UV/O3 oxidation to be a useful cleaning or pre-oxidation procedure prior

to thermal oxide growth for p and n-SiC/SiO2 MOS structures. In comparison to RCA

cleaned SiC samples, they observed that the UV/O3 pre oxidation treatment resulted in a

reduction of defects (fast interface states) and a decrease in positive charge at the p-SiC/SiO2

interface from 2x1012/cm2 to 6-8x1011/cm2. They suggested that the reduction in positive

charge by the UV/O3 treatment was due to the removal of carbon clusters (i.e. C-C bonding)

that remain on the SiC surface after the growth of epitaxial layers and which are not removed

by RCA cleaning or the thermal oxidation process itself. We note that this suggestion is

supported by our observation that UV/O3 oxidation removes the non-SiC carbon (C-C, CHx,

and CFx) from SiC surfaces.

Finally, we would like to emphasize the ability of a room temperature UV/O3

treatment to grow or form thin SiOx layers ( < 20Å) on SiC. In a separate study [21], we

investigated the ability to oxidize SiC surfaces using other wet chemical treatments

commonly employed to form passivating oxides on silicon. As shown in Table 3.3, the

Si2p(Si-C)/O1s intensity ratio of 0.3 resulting after a UV/O3 exposure is much lower than the

61

≈ 1.0 obtained after wet chemical treatments such as boiling Aqua Regia or RCA SC1. In

fact, the Si2p(Si-C)/O1s intensity ratio after the RCA SC1 clean is not significantly different

from Si2p(Si-C)/O1s ratio of 1.1 observed from solvent cleaned SiC surfaces. However, this

observation is consistent with the known inability of any of these acids to etch SiC.

Therefore, the ability of UV/O3 to grow a thin (10-20Å) passivating oxide is an added

benefit over conventional wet chemical processing. Also, the ability to form the passivating

oxide at room temperature is an additional bonus compared to thermal oxidation of SiC

which typically requires temperatures of 1000-1200°C [45]. Finally, it is noted that the Si 2p

spectrum obtained from SiC treated by UV/O3 oxidation (see Figure 4b) bares a striking

resemblance to the Si 2p spectra obtained from a SiC sample previously exposed to a total

fluence of 9x1021 oxygen atoms (i.e. O instead of O2) during low earth orbit on a satellite

[46]. This suggests that UV/O3 oxidation could be used to simulate the operational

conditions of SiC devices in outer space and other harsh oxidizing environments.

Table 3.3. Summary of SiC-C1s/surface C1s and Si/O intensity ratios from XPS data (uncorrected for differences in sensitivity factors).

Treatment SiC C1s/surface C1s Si2p/O1s Solvents 0.8 1.1 Piranha 1.1 0.9 RCA SC1 2.2 1.0 Aqua Regia 1.2 1.2 UV/O3 2.7 0.3 HF vapor 1.3 1.2

3.5.2. HF Vapor

62

As Figures 3.4(b) and 3.4(c) show, the equilibrium vapor from an HF solution alone

can be used to effectively remove thin silicon oxide layers from SiC surfaces. Though Figure

3.4(c) shows the complete removal (below detection limits) of the higher BE Si 2p peak at

102.4 eV, some oxygen was observed to remain on the SiC surface (probably in the form of

suboxides or hydroxides of silicon and carbon i.e.C3-Si-O(H) and Si3-C-O). Table 3.3

shows that the resulting Si/O XPS intensity ratio after the vapor treatment was observed to

increase from 0.3 to 1.3. The Si/O intensity ratio of 1.3 compares well with the value of 1.4

obtained from a SiC surface after removal of a thermal oxide using a 10:1 HF dip [21].

These results indicate that HF vapor exposure is equally as effective as an HF dip in

removing surface silicon oxides from SiC surfaces. It should be noted, however, that the

silicon oxide etch rates for HF vapor and an HF dip are substantially different. In the HF

vapor case, a 30 minute exposure was required to remove only 10-20Å of surface oxide

resulting from a UV/O3 treatment, whereas in the HF dip case, only 10 minutes were

required to remove 1000Å of thermal oxide.

Unfortunately, the SiC C1s/surface C1s intensity ratio was observed to decrease

from 2.7 to 1.7 after the HF vapor treatment. Some of this increased surface carbon may be

attributable to the ambient exposure during and after the vapor treatment. An in situ HF

vapor exposure, however, could eliminate this recontamination. Takayuki et al [25] has

previously demonstrated the removal of native oxides on Si (001) using photoexcited fluorine

gas. Our results presented here suggest that an in situ HF vapor exposure should work as

well.

The authors also note that the HF vapor exposure results in significant amounts of

fluorine on the (0001)Si 6H-SiC surface (see Figure 3.5c). The observed fluorine coverage

following the HF vapor treatment was 3-4 times larger than that previously observed from

(0001)Si 6H-SiC wafers dipped in 10:1 HF and blown dry (without a de-ionized water rinse).

The fluorine surface coverage approaches that of 1/2 to a full monolayer. As the peak

63

position of the F 1s core level after the HF vapor treatment remains essentially unchanged at

685.9 eV (i.e. Si-F bonding), this suggests that the HF vapor treatment leaves a Si-F

terminated SiC surface [40-42]. These results are in contrast with those of Iyer et al [31] for

Si (100) in which no fluorine was detected by XPS and hydrogen termination was confirmed

by TPD. However, in a previous study on HF wet chemical processing of SiC [21], it was

argued that OH- termination would be preferred for (0001)Si 6H-SiC as opposed to hydrogen

due to the polarity of the Si-C bond. Termination of the (0001)Si SiC surface with OH- tends

to cancel the dipole created by the Si-C bond whereas termination with H does not (see

Figure 3.6). Similarly, F- ions (from HF vapor) could also cancel this dipole and in the HF

vapor exposure the F- ions are more readily available than OH- ions. Therefore, fluorine

termination of SiC surfaces should be expected after HF vapor processes as opposed to either

H+ or OH- termination.

64

Si

C C Cδ−

δ+

F

Si

C C C

δ−

δ−

δ+OH-

Si

C C Cδ−

δ+

σ+σ−

σ+σ−

σ+σ−

σ+σ−

[0001]

Ε

(a)

(b) ΟΗ

−

ΟΗ−

ΟΗ−

ΟΗ−

ΟΗ−

ΟΗ−

ΟΗ−

ΟΗ−

Figure 3.6. Schematic illustrating the mechanism for F- and OH- termination of (0001)Si 6H-SiC.

3.6. Conclusions

65

In conclusion, a completely dry process which removes carbon contamination from

(0001)Si 6H-SiC surfaces via UV/O3 oxidation and removes surface oxides via HF vapor

exposure has been demonstrated. Based on the levels of surface carbon and oxide

contaminants, this dry cleaning procedure has been found to be equivalent to or better than

other standard wet chemical processes. In contrast to silicon, the HF vapor exposure was

observed to result in a fluorine terminated SiC surface as opposed to a hydrogen terminated

SiC surface. This process was observed to leave some residual adventitious carbon. The

residual carbon was largely attributed to recontamination in the laboratory ambient.

However, this effect may be eliminated by either in situ vapor phase cleaning or well

controlled mini environments.

3.7. Acknowledgments:

The authors would like to thank Cree Research, Inc., for supplying the SiC wafers

used in these studies. This research was sponsored by the Office of Naval Research and

through the Department of Education via an Electronic Materials/GAANN fellowship.

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22. J.R. Vig, J. Vac. Sci. Techonl. A, 3, 1027 (1985). 23. M. Tabe, Appl. Phys. Lett., 45, 1073 (1984). 24. V.Y. Fominski, O.I. Naoumenko, V.N. Nevolin, A.P. Alekhin, A.M. Markeev, and L.A. Vyukov, Appl. Phys. Lett., 68, 2243 (1996). 25. T. Takahagi, I. Nagai, A. Ishitani, H. Kuroda, and Y. Nagasawa, J. Appl. Phys., 64, 3516 (1988). 26. J.A. McClintock, R.A. Wilson, and N.E. Byer, J. Vac. Sci. and Technol., 20, 241 (1982). 27. R.F. Kopf, A.P. Kinsella, and C.W. Ebert, J. Vac. Sci. Technol. B, 9, 132 (1991). 28. M. Suemitsu, T. Kaneko, and M. Miyamoto, Jap. J. Appl. Phys., 28, 2421 (1989). 29. S.J. Pearton, F. Ren, C.R. Abernathy, W.S. Hobson, and H.S. Luftman, Appl. Phys. Lett., 58, 1416 (1991). 30. S. Baunack and A. Zehe, Phys. Stat. Solid A, 115, 223 (1989). 31. S.S. Iyer, M. Arienzo, and E. de Fresart, Appl. Phys. Lett., 57, 893 (1990). 32. R. Iscoff, Semiconductor International, 7, 58 (1993). 33. W.A. Cady and M. Varadarajan, J. Electrochem. Soc., 143, 2054 (1996). 34. J. van der Weide, Ph.D Dissertation, NCSU. 35. T.J. Chuang, H.F. Winters, and J.W. Coburn, Surface Science, 2, 514 (1978). 36. J.W. Coburn, H.F. Winters, and T.J. Chuang, J. Appl. Phys., 48, 3532 (1977). 37. V.S. Smentkowski, J.T. Yates, Jr., X. Chen, and W.A. Goddard, III, Surface Science, 370, 209 (1997). 38. G. Hollinger and F.J. Himpsel, Appl. Phys. Lett., 44, 93 (1984). 39. G. Hollinger and F.J. Himpsel, J. Vac. Sci. Technol. A, 1, 640 (1983). 40. T.J. Chuang, J. Appl. Phys., 51, 2614 (1980). 41. C.D. Stinespring and A. Freedman, Appl. Phys. Lett., 48, 718 (1986). 42. J.H. Thomas III, and L.H. Hammer, J. Vac. Sci. Technol. B, 5, 1617 (1987).

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69

4. Chemical Vapor Cleaning of (0001)Si, (000-1)C, (10-10) &

(11-20) 6H-SiC Surfaces


to the:

Journal of Applied Physics

by

Sean W. King, R. Scott Kern, *Mark C. Benjamin, John P. Barnak,

*Robert J. Nemanich, and Robert F. Davis,

Department of Materials Science and Engineering and

*Department of Physics


Raleigh, NC 27695

70

4.1. Abstract

A chemical vapor cleaning (CVC) procedure based on annealing in fluxes of SiH4

and C2H4 has been demonstrated for (0001)Si, (000-1)C, (11-20), and (10-10) 6H-SiC

surfaces. In comparison to SiC surfaces prepared by thermal desorption techniques, SiH4

CVC prepared surfaces were found to be of higher purity, free of both oxides and C-C

bonded carbon/graphite. For the (0001)Si orientation, the SiH4 CVC procedure was found to

produce (3x3) reconstructed surfaces which consisted of an incomplete bilayer of silicon on

top of the SiC surface. Reconstructed (√3x√3)R30° (0001)Si 6H-SiC surfaces could be

prepared by annealing the (3x3) SiH4 CVC surface in UHV at 1050°C. In contrast, no

reconstructions were observed for SiH4 CVC prepared (000-1)C, (11-20), and (10-10) 6H-

SiC surfaces. The SiH4 CVC procedure was found to be particularly effective in preventing

and removing graphite formation from (000-1)C surfaces. The stoichiometry of CVC

prepared (000-1)C and (11-20) and (10-10) surfaces was easily controlled via a second

exposure to C2H4.

71

4.2. Introduction

Preparation of clean, structurally well ordered surfaces is an important first step in all

semiconductor microelectronic fabrication processes [1-3]. Surface cleaning prior to epitaxy

is particularly important as improper removal of surface oxides and organic contaminants has

been shown in Si homoepitaxy to result in an increased density of stacking faults and

dislocations from < 104/cm2 to > 1010/cm2 [4-15]. In fact, studies on the heteroepitaxial

growth of SixGe1-x alloys on Si have shown that residual oxide and organic contaminants

act as the preferred sites for nucleation of misfit dislocations [15]. Epitaxial defects from

improper surface cleaning are important as they have been shown to cause a decrease in

device performance and yield [16-22]. Accordingly, the control of defects in epitaxial films

is clearly related to surface cleaning and preparation. These observations are of paramount

importance to SiC as a range of structural and electrical defects are currently prohibiting it

from becoming the semiconductor/substrate of choice for high power, high frequency, and

high temperature electronic devices and III-V nitride heteroepitaxy [23-25]. Thus based on

analogy to silicon technology, it is clear that improved surface cleaning procedures should

result in decreased epitaxial defect densities. In fact, Powell et al [26] and Burk and Rowland

[27] have previously shown surface pretreatments to be instrumental in the control of

polytype and reduction of interfacial Al in atmospheric CVD and VPE growth of SiC.

72

In a previous set of studies [27,28], we have shown that typical wet chemical

cleaning processes leave a monolayer coverage of oxygen and adventitious surface carbon on

(0001)Si 6H-SiC surfaces which must be removed in situ prior to epitaxial growth. Removal

of the oxide and adventitious carbon in situ has been achieved via a variety of techniques

including thermal desorption [30-46], ion bombardment/sputtering [47-50], ECR H2 plasma

cleaning [51], and annealing in a Si flux [55-62]. As others have shown [34,35,55], thermal

desorption of the monolayer of oxide from SiC occurs at ≈ 1000°C which is ≈ 200°C higher

than that required for silicon [63]. Additionally, thermal desorption of oxygen results in the

loss of silicon from the SiC surface and the formation of C-C bonding which may lead to

graphite bonded structures on the surface [33-36,38,39]. This is primarily due to the fact that

surface oxides on Si and SiC desorb as SiO instead of O2. This desorption depletes the SiC

surface of silicon leaving behind excess carbon which may form graphite bonded structures

[38,39]. In the case of plasma cleaning, an ECR H2 plasma clean has been previously shown

to be useful for removing C-C, C-F, and C-O bonded contaminants, but the technique is

inefficient or incomplete with regard to removing Si-O [51]. As we [52] and others [53,54]

have shown in separate studies, atomic H also selectively removes Si from the SiC surface

producing a carbon rich surface. Ion bombardment or sputtering inherently induces surface

damage and disorder which must be removed via high temperature annealing. This will also

lead to a loss of silicon and the possibility graphitic bonding at the surface [49].

An alternative to the above techniques is to anneal a SiC wafer in a flux of material

whose oxide or suboxide is more volatile than SiO2. The advantage of this technique is that

the oxide can be chemically removed at temperatures 100-200°C lower than simple thermal

desorption and without physically bombarding or damaging the surface with high energy

73

ions. In the case of silicon, fluxes of Ga [63], Ge [64], GeH4 [65], SiH4 [66,67], and Si2H6

[68] have been successfully used to chemically reduce and remove oxides on silicon surfaces.

In the case of SiC, Kaplan and Parril [55-57] have demonstrated that evaporated Ga or Si can

also be used to reduce and remove surface oxides from SiC surfaces at ≈ 850°C which is ≈

100-200°C lower than the temperature necessary to thermally desorb the surface oxide from

SiC surfaces [34,35]. The technique of Kaplan [55] has been recently implemented by Fissel

et al [62] in solid source MBE growth of SiC on (0001)Si 6H-SiC. . Unfortunately however,

the approach used by Kaplan [55] and Fissel [62] is most compatible with MBE growth

techniques. An alternative to this approach would be to use a gas source of silicon such as

SiH4 or Si2H6 which would allow this technique to be extended to lower vacuum with

techniques.

In this paper, we demonstrate that for preparation of silicon terminated (0001)Si 6H-

SiC surfaces, essentially similar results to evaporated silicon can be achieved using a gas

source of silicon such as SiH4. This chemical vapor cleaning (CVC) cleaning procedure

was found to be compatible with both GSMBE and LPCVD growth processes as well. We

additionally demonstrate that the combined sequential use of SiH4 and C2H4 can be used for

the preparation of carbon terminated (000-1)C 6H-SiC surfaces and non-polar (11-20) and

(10-10) 6H-SiC surfaces of varying stoichiometry. Direct comparison to thermal desorption,

shows the CVC technique to produce the purest SiC surfaces free of oxide and carbon-carbon

bonded contaminants.

74


4.3.1. Integrated Surface Preparation and Analysis System.

All experiments described below were conducted using a unique ultra high vacuum

(UHV) configuration which integrates several completely independent UHV surface

preparation, thin film growth and surface analysis systems via a 36 ft. long transfer line

having a base pressure of 9x10-10 Torr (additional details of the transfer line, and many of

the associated systems are provided in Refs. 69-71). The experiments described in this paper

employed the SiC atomic layer epitaxy (ALE), Auger electron spectroscopy (AES), electron

energy loss spectroscopy (EELS), low energy electron diffraction (LEED), x-ray

photoelectron spectroscopy (XPS), and remote H2/SiH4 plasma CVD systems. A brief

description of these systems is provided below.

The SiC ALE system consisted of a UHV chamber with a base pressure of 3x10-10

Torr and was equipped with a residual gas analyzer (RGA) and a variety of gas dosers. The

RGA (a 0-200 amu quadrapole gas analyzer from Hiden Analytical Ltd.) was housed in a

separate differentially pumped cylindrical chamber (similar in design to that of Smentkowski

and Yates [72] ) which had a 0.5 cm diameter orifice at the head of the RGA for TPD

experiments and an approximately 50 cm2 "sunroof" for monitoring residual gases in the

75

system. The sample heating stage in the ALE system consisted of a wound tungsten heating

filament positioned close to the back of the sample and mounted on a boron nitride disk [69].

A W/6%Re-W/26%Re thermocouple was employed to measure the temperature of the

backside of the wafer. Surface temperatures and heating profiles to 1100°C were easily

achieved using a programmable microprocessor and 20 amp SCR power supply. Actual

surface/sample temperatures (i.e. those reported herein) were recorded using an infra-red

pyrometer with a spectral response of 0.8 to 1.1 µm and a emissivity setting of 0.5. The

estimated experimental accuracy for the latter temperatures was estimated to be ± 25°C.

Gas sources in the ALE system included SiH4 (99.995%), C2H4 (99.99%), and H2

(99.995%). Mass spectroscopic analysis of the as received silane using the RGA revealed

that the primary resolvable impurities were H2O, CO2, and Si2H6 in concentrations of 180

ppm, 31 ppm, and 170 ppm respectively. Impurities such as CO and O2 were expected but

difficult to resolve due to overlap with the silane cracking patterns. However, we estimate

that the level of these impurities were at least below 600 ppm or better. No further

purification was deemed necessary and the silane was used in this purity. Similar purity was

found for the ethylene and hydrogen. Sample exposure to SiH4 and/or C2H4 was obtained

using "molecular beam" dosers similar to the design of Bozack et al [73]. Collimation of

SiH4 or C2H4 into a molecular beam focused onto the sample was achieved with this doser

using a 13 mm diameter x 2 mm thick glass capillary array with a ten micrometer pore size

(Galileo Electro Optics Inc.). Thought, the doser to sample distance was fixed at 2", no

attempts were made to accurately measure the flux of SiH4 or C2H4 and hence all exposures

are quoted as Langmuirs (£ = 10-6 Torr sec).

76

XPS experiments were performed in a stainless steel UHV chamber (base pressure =

2x10-10 Torr) equipped with a dual anode (Mg/Al) x-ray source and a 100 mm

hemispherical electron energy analyzer (VG CLAM II). All XPS spectra reported herein

were obtained using Al Kα radiation (hν = 1486.6 eV) at 12 kV and 20 mA emission

current. XPS analysis typically required less than 1 hour during which time the pressure

never increased above 9x10-10 Torr. Calibration of the binding energy scale for all scans

was achieved by periodically taking scans of the Au 4f7/2 and Cu 2p3/2 peaks from

standards and correcting for the discrepancies in the measured and known values of these two

peaks (83.98 and 932.67 eV, respectively [74]). Curve fitting of most data was performed

using the software package GRAMS 386. A combination Gaussian-Lorentzian curve shape

with a linear background was found to best represent the data. The Auger electron

spectrometer and the low energy electron diffraction optics were mounted on a six way cross

off the transfer line and pumped through the transfer line. In the AES analysis, a 3 keV, 1

mA beam was used. Each Auger electron spectrum was collected in the undifferentiated

mode and numerically differentiated. In LEED an 80 eV, 1 mA beam was used.

The remote plasma CVD system consisted of a metal seal stainless steel vacuum

chamber pumped by a 330 l/s turbomolecular pump. The base pressure of this system was

4x10-9 Torr. The process gases flowed through a quartz tube mounted at the top of the

chamber, and the plasma is excited by rf (13.56 MHz) applied through a copper coil wrapped

around the quartz tube. The sample was located 40 cm below the center of the rf coil. An

inline Nanochem purifier and filter was used for point of use purification of hydrogen and

silane. Sample heating in the plasma system was conducted using a sample heater similar in

design to the one previously described in the ALE system. The plasma system was also

77

equipped with a differentially pumped 0-100 amu RGA which allowed direct analysis of the

purity of the process gases. RGA analysis of the hydrogen and silane (both 99.999% purity)

used in these experiments after in situ purification revealed that the impurity level of these

gases were below the baseline of the system (<1 ppm)

4.3.2. Substrate Preparation:

Various orientations of 6H-SiC substrates supplied by Cree Research Inc. were

examined in this research including: on axis and vicinal (4° off axis toward (11-20))

(0001)Si, on axis and vicinal (000-1)C, (11-20), and (10-10). The size of these substrates

ranged from ≈ 1.5 cm2 for (0001)Si and (000-1)C orientations to 5 mm2 for the (11-20) and

(10-10) samples. All substrates were nitrogen doped n-type with a carrier concentration, Nd-

Na, equal to ≈ 1018/cm3. The off axis (0001)Si and (000-1)C substrates were provided with

an ≈ 1 µm n-type epitaxial layer (Nd=5x1017/cm3) by Cree Research Inc. Cree additionally

oxidized 500-1000Å of the surface of all SiC wafers by dry oxidation. This oxide was

removed via a 10 min. dip in 10:1 HF.

After removal of the thermal oxide, the unpolished back side of each wafer was

coated with tungsten via rf sputtering to increase the heating efficiency of the SiC, as the

latter is partially transparent to the infrared radiation emitted from our tungsten filament

heaters. Originally platinum was used as the refractory metal but was later found to be 78

unsuitable. Typically, platinum was found to react with the back of the wafer forming Pt-

silicides which evaporated at the temperatures used in these experiments. Further Pt was

observed to diffuse through micropipes in the wafer to the polished side of the wafer also

forming platinum silicides which were detected by AES and XPS. Platinum contaminated

surfaces were observed to exhibit (2x2), (4x4), and (5x5) LEED patterns which are

uncharacteristic of pristine SiC surfaces [55].

After coating the backside of the SiC wafer with tungsten, the SiC wafers were

ultrasonically rinsed in trichloroethylene, acetone and methanol each for 5 min. and then

exposed to the vapor from a 10:1 buffered HF solution for 10 min. The wafers were then

mounted to a 1" diameter Mo disk using Ta wire. Each wafer/Mo assembly was then

fastened to a ring shaped Mo sample holder using Ta wire and inserted into the transfer line

load lock for further experimentation.

4.4. Results

4.4.1. (0001)Si 6H-SiC

Figure 4.1 shows a series of AES spectra acquired from a vicinal (0001)Si 6H-SiC

surface exposed to 200£ SiH4 at various different temperatures in the SiC ALE system

Similar results were obtained from on axis (0001)Si 6H-SiC surfaces. Comparison between

Figures 4.1(a) and 4.1(b) shows that the SiH4 exposure at a substrate temperature of 750°C

results in very little change in the amount of surface oxide on SiC. A 200£ exposure at

820°C (Fig. 4.1(c).) results in an incomplete reduction of the surface oxide. Removal of the

79

oxygen to below the detection limits of both AES and XPS was achieved at temperatures of

880°C and greater. This temperature is 100-200°C lower than the temperature needed to

simply thermally desorb the oxide in UHV. Using these conditions, sharp (1x1) LEED

patterns were obtained whereas prior to the CVC clean a (1x1) pattern with broad diffraction

spots were obtained (see Figure 4.2(a),(b)). The changes in the Si and C concentrations on

the surface were examined from the ratio of the peak-to-peak heights (pph) of the AES

signals or the ratio of the XPS peak heights. The SiC surfaces prepared in this manner were

slightly silicon rich as observed by the increase in the AES Si LVV/C KLL pph ratio from <

1 to ≈ 3 after the CVC treatment. A similar increase from < 1 to ˜ 1.1 was observed for the Si

2p/C 1s ratio in XPS.

30 150 270 390 510 630

dN(E

)/dE

(c)

(b)


(a)

Si C

O

Figure 4.1. AES of (0001)Si 6H-SiC surfaces after (a) 200£ SiH4 at 750°C, (b) 200£ SiH4 at 820°C, and (c) 200£ SiH4 at 880°C.

80

If a larger a SiH4 exposure (> ˜ 500£) was used at Tsub > 900°C, a second surface

reconstruction, the (3x3), is observed with LEED (see Figure 4.2(c)). In this case, the Si

LVV/C KLL pph ratio was observed to increase from ≈ 3 to 5 suggesting that excess silicon

was deposited on the SiC surface. The excess Si was confirmed by XPS analysis of the Si 2p

core level from the (3x3) surface where two Si 2p peaks were observed at 99.5 and 101.3 eV.

The two different peaks are indicative of Si-Si and Si-C bonding respectively (see Figure

4.3). More detailed analysis of the intensity of these two peaks based on the attenuation of

the Si-C Si 2p peak indicated that the (3x3) reconstruction corresponded to a silicon coverage

of ≈ 1.5 monolayers (i.e. an incomplete bilayer). Higher SiH4 exposures (>2000£) in

temperature range of 900-1050°C were not typically observed to result in additional

reconstructions past the. However, annealing of the (3x3) surface at 1000°C in UHV did

lead to the loss of the excess Si resulting in the surface reverting back to the (1x1) LEED

pattern. Still further annealing, resulted in the observation in LEED of a third surface

reconstruction, the (√3x√3)R30° (see Figure 4.2(d)). Figure 4.4 displays XPS spectra of the

Si 2p core level from (0001)Si surfaces exhibiting the three different SiC reconstructions

observed in this study. As can be seen, the intensity of the Si-Si peak/shoulder decreases in

going from the (3x3) to the (1x1) to the (√3x√3)R30° surfaces. Further, the Si LVV/C KLL

pph ratio in AES was observed to decrease from ≈ 4-5 for the (3x3) surface to 3 and 2

respectively for the (1x1) and (√3x√3)R30° surfaces. Accounting for the 2:1 difference in

sensitivity to Si and C in AES [55], the (√3x√3)R30° surface appears to be the closest to bulk

stoichiometry. Similarly results were observed in XPS were the Si2p/C1s ratio was observed

to decrease from ≈ 1.3 for the (3x3) surface to 1.1 and 0.9 for the (1x1) and (√3x√3)R30°

surfaces. Other reconstructions past the (√3x√3)R30° such as a (6√3x6√3)R30° or (6x6)

[45,60] were not observed with extended annealing in the temperature range of 1000-1100°C.

These results are in excellent agreement with the results of Kaplan [55] who found that the

81

surface oxide could be removed at 850°C and (3x3) and (√3x√3)R30° reconstructions could

be obtained via annealing in evaporated silicon and/or UHV

For comparison purposes, (√3x√3)R30° 6H-SiC (0001)Si surfaces were also prepared

by thermal desorption at 1000°C. A comparison of the AES survey spectra acquired from

(√3x√3)R30° surfaces prepared by thermal desorption and SiH4 CVC is presented in Figure

4.5. It is clearly evident in Figure 4.5 that the Si LVV/C KLL pph ratio is < 1 for the thermal

desorption surface and > 1 for the CVC prepared surface. This suggests that the

(√3x√3)R30° surface prepared by thermal desorption is more carbon rich (or Si deficient)

than the (√3x√3)R30° surface prepared by CVC. This observation was additionally

supported by XPS analysis of the C 1s core level from these two surfaces (see Figure 4.6).

The spectra display a second C-C C 1s peak in the thermal desorption spectrum which is not

observed from the CVC surface. It should also be pointed out that in this, study removal of

oxygen below the detection limits of AES and XPS was rarely achieved using thermal

desorption.

82

(a) (b)

(c) (d)

83

(e) (f) (g)

Figure 4.2. LEED patterns from (a) HF dipped (0001)Si 6H-SiC, (b) (1x1) (0001)Si 6H-SiC, (c) (3x3) (0001)Si 6H-SiC, (d) (v3xv3)R30° (0001)Si 6H-SiC, (e) (1x1) (000-1)C 6H-SiC, (f) (11-20) 6H-SiC, (g) (10-10) 6H-SiC.

96 98 100 102 104 106

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

Si-Si

Si-C

Figure 4.3. XPS spectra of the Si 2p core level from a (3x3) reconstructed (0001)Si 6H-SiC surface.

84

95 97 99 101 103 105 107

(1x1)(¦3x¦3)(3x3)

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

Figure 4.4. XPS spectra of the Si 2p core level from (3x3), (1x1) and (v3xv3)R30° reconstructed (0001)Si 6H-SiC surfaces.

100 200 300 400 500 600 700

dN(E

)/dE


(a)

(b)

Si

C

N O

Figure 4.5. AES survey spectra from (v3xv3)R30° reconstructed (0001)Si 6H-SiC surfaces prepared by (a) SiH4 CVC, and (b) thermal desorption.

85

280 282 284 286 288

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

Figure 4.6. XPS spectra of the C 1s core level from (v3xv3)R30° reconstructed (0001)Si 6H-SiC surfaces prepared by (a) SiH4 CVC, and (b) thermal desorption. 4.4.2. (000-1)C 6H-SiC

Oxide removal from (000-1)C 6H-SiC surfaces via annealing in SiH4 in the SiC ALE

system was observed to exhibit a similar temperature dependence to (0001)Si surfaces with

temperatures > 850°C being generally required for complete oxide removal. However in

contrast to (0001)Si, (3x3) reconstructions were never observed from CVC cleaned (000-1)C

surfaces despite the observation of Si LVV/C KLL pph ratios as large as 3-5 in AES and Si-

Si bonded Si 2p peaks in XPS (see Figure 4.2(e)). Further, we were not able to obtain

(√3x√3)R30° reconstructions from CVC cleaned (000-1)C surfaces despite extended

annealing (≈ 1 hr.) at 1100°C. These observations may be related to the trace amounts of

nitrogen consistently observed from the (000-1)C wafers after CVC cleaning or thermal

86

desorption (see Figure 4.7(a)). Similar persistent traces of nitrogen have also been observed

from (000-1)C 6H-SiC surfaces by Bermudez [61] and have been attributed to preventing

reconstruction of the SiC surface.

Comparisons between SiH4 CVC and thermal desorption cleaned (000-1)C surfaces

were also made. Firstly, traces of oxygen were observed from (000-1)C surfaces prepared by

annealing in UHV at 1050°C for 15 min. which was similar to (0001)Si surfaces. However

for the (000-1)C orientation, significant amounts of graphitic C-C bonding were observed to

form on the (000-1)C surface after annealing in UHV at 1050°C to thermally desorb the

surface oxide. The graphite formation was clearly observed in EELS spectra via a loss peak

at 6 eV (see Figure 4.8(a)) and a second C 1s peak at 284.5 eV in XPS (see Figure 4.9(a)).

This observation is to be contrasted to (0001)Si surfaces where a second C 1s peak was

observed in XPS from (√3x√3)R30° surfaces prepared via thermal desorption but a 6 eV loss

peak in EELS was not observed (see Figure 4.10). However for (000-1)C surfaces which had

undergone a 1050°C SiH4 CVC treatment, no graphite bonded C was observed in EELS (see

Figure 4.8(b)) and no C-C bonding C 1s was observed in XPS. In fact, the SiH4 CVC

treatment was also found useful for removing (i.e. converting to SiC) graphite formed on

(000-1)C surfaces via thermal desorption.

87

30 130 230 330 430 530 630 730

(a)

(b)

(c)

dN(E

)/dE


SiC

Figure 4.7. AES survey spectra from (0001)C 6H-SiC surfaces prepared by (a) thermal desorption, (b) SiH4 CVC, and (c) SiH4/C2H4 CVC.

One interesting characteristic of the SiH4 CVC treated (000-1)C surfaces was an

inability to regraphitize this surface via high temperature annealing in UHV. Despite

annealing at 1100°C for extended periods of time (≈ 1-2 hr.), we were unsuccessful in

obtaining a 6 eV loss peak in EELS and/or a C-C bonding C 1s peak in XPS. Additionally, a

Si LVV/C KLL pph ratio of > 1 was maintained for all SiH4 CVC treated (000-1)C surfaces

annealed at 1100°C. The Si LVV/C KLL ratio of > 1 and similarities between the EELS

spectra of SiH4 CVC treated (0001)Si (see Figure 4.11) and (000-1)C (see Figure 4.8(b))

surfaces suggests that despite high temperature annealing, SiH4 CVC treated (000-1)C

surfaces are still silicon terminated. In order to reduce the Si LVV/C KLL ratio to < 1 (i.e.

carbon termination), it was found necessary to expose the SiH4 CVC treated (000-1)C

surfaces to C2H4.(see Figure 4.7(c)). However it was observed that C2H4 exposure at

88

temperatures > 950°C resulted in regraphitization of the (000-1)C surfaces (see Figure

4.8(c)). In order to reduce the Si LVV/C KLL ratio to < 1 and avoid graphite formation, it

was found necessary to decrease the sample temperature to ≈ 850°C (see Figure 4.12).

Unfortunately in this case, sharp (1x1) LEED patterns were not obtained.

-35 -30 -25 -20 -15 -10 -5 0 5

(a)

(b)

(c)

Cou

nts (

arb.

uni

ts)

Loss Energy (eV)E

last

ic P

eak

Figure 4.8. EELS spectra from (000-1)C 6H-SiC surfaces after (a) annealing in UHV at 1050°C, (b) annealing in 2000£ SiH4 at 1050°C followed by (c) annealing in 2000£ C2H4 at 950°C.

89

280 282 284 286 288

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

Figure 4.9. XPS spectra of the C 1s core level from (000-1)C 6H-SiC surfaces after (a) 2000£ SiH4 at 1050°C, (b) thermal desorption in UHV at 1050°C.

-40 -35 -30 -25 -20 -15 -10 -5 0

Cou

nts (

arb.

uni

ts)

Energy Loss (eV)

Ela

stic

Pea

k

Figure 4.10. EELS spectrum from (v3xv3)R30° reconstructed (0001)Si 6H-SiC surface prepared by annealing in UHV at 1050°C.

90

-40 -35 -30 -25 -20 -15 -10 -5 0

Cou

nts (

arb.

uni

ts)

Energy Loss (eV)

(a)

(b)

(c)

Ela

stic

Pea

k

Figure 4.11. EELS spectra of (3x3), (1x1), and (v3xv3)R30° reconstructed (0001)Si 6H-SiC surfaces prepared via SiH4 CVC.

-35 -30 -25 -20 -15 -10 -5 0 5

(a)

(b)

(c)

Cou

nts (

arb.

uni

ts)

Electron Energy Loss (eV)

Ela

stic

Pea

k

Figure 4.12. EELS spectra from (000-1)C 6H-SiC surfaces after (a) 2000£ SiH4 at 1000°C, (b) 400£ C2H4 at 850°C, and (c) 800£ C2H4 at 850°C.

91

4.4.3. (10-10) & (11-20) 6H-SiC

Oxide removal from (10-10) and (11-20) 6H-SiC surfaces via a SiH4 CVC clean was

observed to exhibit a similar temperature dependence to the (0001)Si and (000-1)C surfaces.

The LEED patterns displayed by these surfaces were observed to intensify and sharpen with

the SiH4 clean, but no reconstructions were observed to occur as a result of the clean or

deposition of excess silicon (see Figure 2(f),(g)). As with the (0001)Si and (000-1)C, the

SiH4 clean resulted in Si LVV/C KLL ratios > 1 which could not be reduced below one by

annealing at 1100°C. However, the Si LVV/C KLL ratio could be reduced to < 1 by

annealing in C2H4 at 900-1000°C. In this case, a 6 eV loss peak characteristic of graphite

formation was not observed from the (10-10) and (11-20) 6H-SiC surfaces (see Figure 13).

However, a C-C bonding C 1s bonding peak could be seen in XPS for large C2H4 exposures.

Similar to (0001)Si and (000-1)C, thermal desorption of the oxide from (11-20) and (10-10)

surfaces at 1050°C did result in incomplete thermal desorption and the observation of a C-C

bonded C 1s peak in XPS.

92

-35 -30 -25 -20 -15 -10 -5 0 5

(b)

(a)

Cou

nts (

arb.

uni

ts)

Electron Energy Loss (eV)

Ela

stic

Pea

k

Figure 4.13. EELS spectra of (10-10) 6H-SiC (a) after annealing in SiH4 at 1000°C, and (b) then annealing in C2H4 at 850°C.

4.4.4. Low vacuum CVC/LPCVD clean

In order to simulate conditions in lower vacuum processes (i.e. LPCVD and

OMVPE), (0001)Si 6H-SiC wafers were exposed to various fluxes of a 1% SiH4/H2 mixture

in the plasma CVD system. In this case, it was found that oxygen free, (3x3) reconstructed

(0001)Si 6H-SiC surfaces could be attained via annealing the SiC wafer in 70 sccm H2 and

0.01 sccm SiH4 at 900°C for 1 min. No other differences were observed between (0001)Si

6H-SiC surfaces prepared in low fluxes of SiH4 in the GSMBE and higher the fluxes of

H2/SiH4 in the plasma CVD.

93

4.5. Discussion

4.5.1. (0001)Si 6H-SiC

The results presented here demonstrate that the SiH4 CVC clean for (0001)Si 6H-SiC

yields surfaces similar to those prepared by the Si evaporation technique of Kaplan [55]. Our

Si 2p XPS spectra indicate an ≈ 1.5 monolayer coverage of Si (i.e. incomplete bilayer) for the

(3x3) reconstructed surface generated via the SiH4 CVC clean. This result is consistent with

the previously reported AES results and incomplete bilayer model of Kaplan [55] for the

(3x3) reconstruction. The (3x3) reconstruction has also been recently observed by Kulakov

et al [59] and Li and Tsong [60] using STM. Based on their observation of only one maxima

in the (3x3) unit cell, Kulakov et al [59] proposed a slight modification to the model of

Kaplan for the (3x3) reconstruction. In the model of Kulakov et al [59], the (3x3)

reconstruction still consists of an incomplete Si bilayer but with 1 adatom, 3 rest atoms, and 7

atoms directly over Si atoms in SiC as opposed to the 2 adatoms, 6 rest atoms and 8 second

layer Si atoms in the model of Kaplan [55]. In contrast, Li and Tsong [60] have proposed

that the (3x3) reconstruction consists of a 4/9 monolayer coverage of Si-C tetrahedra

organized in a (3x3) pattern. Our observation of a Si-Si bonding Si 2p peak in XPS with an ≈

1.5 monolayer coverage is clearly more consistent with the (3x3) models proposed by Kaplan

and Kulakov et al [59]. In a separate x-ray photoelectron diffraction study [75], we also

provide data which additionally supports the (3x3) model of both Kaplan and Kulakov et al

[59]. Finally, we note that similar Si-Si bonding Si 2p peaks have also been observed in XPS

94

from (3x2) [76] and c(4x2) [77] reconstructed (100) 3C-SiC surfaces prepared via annealing

in an evaporated Si flux.

For the (√3x√3)R30° reconstruction, it is interesting to note the differences in surface

stoichiometry observed between the two (√3x√3)R30° reconstructed surfaces prepared by

thermal desorption and SiH4 CVC. As previously noted, the Si LVV/C KLL pph ratio in

AES was < 1 for the (√3x√3)R30° surface prepared by thermal desorption whereas the SiH4

CVC (√3x√3)R30° surface was observed to have a Si LVV/C KLL ratio > 1. Additionally, a

C-C bonding C 1s peak was observed in XPS for the thermal desorption (√3x√3)R30°

surface. The observation of C-C bonding for (√3x√3)R30° (0001)Si surfaces prepared by

thermal desorption has been observed in several other studies [33,36,38-41]. For

(√3x√3)R30° (0001)Si surfaces prepared via annealing in an evaporated Si flux, Kaplan has

previously reported the observation of a 6 eV loss peak indicative of graphite in EELS.

However, in our case for (√3x√3)R30° surfaces prepared by the SiH4 CVC process, we did

not detect either a clear 6 eV loss peak in EELS or a C-C bonding C 1s peak in XPS

The above noted differences between the (√3x√3)R30° reconstructed (0001)Si

surfaces is important as various adatom models for this reconstruction have been proposed

based on recent STM investigations and theoretical calculations. Recent STM images by

Owman and Martensson [44,45] and Li and Tsong [60] have shown that the (√3x√3)R30°

(0001)Si surfaces prepared by thermal desorption are similar in nature to the group III

adatom (√3x√3)R30° Si (111) reconstructed surfaces [78,79]. These STM investigations

[44,45,60] have shown the existence of a 1/3 monolayer coverage of adatoms in threefold

symmetric sites. Unfortunately, neither investigation was able to resolve the chemical nature

95

of the adatom (i.e. silicon or carbon) or the adatom site (i.e. T4 or H3). However, the first-

principles total-energy calculations of Northrup and Neugebauer [80] indicate that Si

adatoms are preferred over C adatoms and that T4 sites are preferred over H3 sites. In

contrast, the quantum mechanical cluster calculations of Badziag [81,82] indicate that the

structure of the (√3x√3)R30° reconstruction on (0001) SiC surfaces should consist of

hydrogenated C3 triangles (resembling cyclopropane) saturating Si dangling bonds at the

surface. Based on the differences in surface stoichiometry, it is tempting to conclude that the

(√3x√3)R30° surface prepared by thermal desorption in this study and others [33,36,38-

41,43-45] is most appropriately described by the model of Badziag [81,82]. Similarly, for the

(√3x√3)R30° surface prepared by a SiH4 CVC, the model of Northrup and Neugebauer [80]

perhaps best describes the structure of this reconstructed surface. In the case of the (1x1)

reconstruction observed between the transformation from (3x3) to (√3x√3)R30°

reconstructions, the authors feel that this surface is disordered and is likely composed of a

mixture of (√3x√3)R30° and (3x3) reconstructions.

Finally additional proof of the effectiveness of the SiH4 CVC cleaning procedure has

been recently demonstrated by Kern [58] using secondary ion mass spectroscopy (SIMS)

analysis. In this case, SIMS analysis of the interface between a 3C-SiC epitaxial film grown

by GSMBE on a 6H-SiC wafer which had been cleaned by this procedure revealed absolutely

zero oxygen at the interface [58]. For comparison purposes the detection limit of AES and

XPS for O is typically 0.1 at%. SIMS on the other hand has a sensitivity limit in the ppm

range.

4.5.2. (000-1)C 6H-SiC

96

As noted above, we were unable to observe any surface reconstructions from the

(000-1)C surface for wafers undergoing either thermal desorption or SiH4 CVC treatments.

This is in contrast to the results of Nakanishi et al [33] where (3x3) reconstructions were

observed in LEED from the (000-1)C surface after thermal desorption at temperatures of

900-1100°C. Using STM, Li and Tsong [60] have additionally observed (3x3) and

(√3x√3)R30° reconstructions from (000-1)C surfaces prepared via annealing in an

evaporated Si flux. As previously noted by Bermudez [61], our inability to observe any

reconstructions from the (000-1)C surface with LEED may be related to the presence of

nitrogen on these surfaces. The presence of the nitrogen may be related to segregation of

nitrogen dopants to surface due to volatilization of silicon from the SiC surface leading to the

observed graphite formation. However, the authors have observed larger concentrations of

nitrogen on (0001)Si surfaces which have displayed (3x3) and (√3x√3)R30° reconstructions.

In this case, the source of nitrogen was attributed to reaction of the (0001)Si surface with

residual ammonia (NH3) in the system. Alternatively, the inability to observe surface

reconstructions from the (000-1)C surface with LEED may be related to an increased surface

roughness for the (000-1)C surface relative to the (0001)Si surface. This could occur due to

the lack of an optimized surface polishing procedure for the (000-1)C face. In the case of

Nakanishi et al [33], the (000-1)C surfaces used were the natural faces of Acheson crystals

which are speculated to be naturally atomically smooth. As STM samples a smaller region of

the surface compared to LEED, this may also explain the ability of Li and Tsong using STM

97

[60] to observe (3x3) and (√3x√3)R30° reconstructions from (000-1)C surfaces which have

undergone the same polishing/surface treatment as in our study.

The observation of graphite in EELS from (000-1)C surfaces prepared via thermal

desorption at 1050°C and not from (0001)Si surfaces similarly prepared is consistent with the

observations of Muehlhoff et al [32] where carbon segregation for the (000-1)C surface was

observed to occur ≈ 300°C lower than that for the (000-1)Si face. However, the inability to

regraphitize (000-1)C surfaces via annealing after a SiH4 CVC treatment was particularly

surprising, given the relative ease with which this occurred when SiH4 CVC treated (000-1)C

surfaces were exposed to C2H4. These results clearly indicate that the addition of a

monolayer of Si somehow stabilizes the (000-1)C surface against graphitization. Regardless

though, this surface termination is highly unstable in the presence of excess or free carbon

and immediately forms graphite. These results indicate that for (000-1)C surfaces a silicon

rich surface should be maintained during high temperature cleaning and growth in order to

avoid graphite formation. The ability to produce such a surface is an advantage for SiH4

CVC over thermal desorption treatments. Finally, the authors note that C2H4 has been

previously used for the preparation of carbon terminated c(2x2) reconstructed (001) 3C-SiC

surfaces [83,84]. In our, case it was found necessary to use low temperature C2H4 exposures

(≈ 850°C) in order to suppress graphite formation. However, surfaces prepared in such a

manner exhibited poor (1x1) LEED patterns indicating a disordered surface.

98

4.5.3. (11-20) and (10-10) 6H-SiC

The (11-20) and (10-10) orientations of 6H-SiC are non-polar surfaces with an equal

number of carbon and silicon atoms at the outermost surface. Accordingly, one may expect

these surfaces to exhibit properties intermediate to those of the (0001)Si and (000-1)C

orientations. In our limited studies here, though, we have observed the (11-20) and (10-10)

orientations to behave almost exactly like that of the (0001)Si face of 6H-SiC. However as

these are the first reported examinations of these orientations, the results reported here only

emphasize the need for a more detailed examination of the these surfaces. In particularly, a

more detailed examination of the electronic structure of the (11-20) and (10-10) surfaces

should be most beneficial as it should lead to a better understanding of the polar (0001)Si and

(000-1)C orientations.

99

4.6 Conclusions

A chemical vapor cleaning (CVC) procedure based on annealing in fluxes of SiH4

and C2H4 has been demonstrated for (0001)Si, (000-1)C, (11-20), and (10-10) 6H-SiC

surfaces. In comparison to SiC surfaces prepared by thermal desorption techniques, SiH4

CVC prepared surfaces were found to be of higher purity, free of both oxides and C-C

bonded carbon/graphite. For the (0001)Si orientation, the SiH4 CVC procedure was found to

produce (3x3) reconstructed surfaces which consisted of an incomplete bilayer of silicon on

top of the SiC surface. Reconstructed (√3x√3)R30° (0001)Si 6H-SiC surfaces could be

prepared by annealing the (3x3) SiH4 CVC surface in UHV at 1050°C. In contrast, no

reconstructions were observed for SiH4 CVC prepared (000-1)C, (11-20), and (10-10) 6H-

SiC surfaces. The SiH4 CVC procedure was found to be particularly effective in preventing

and removing graphite formation from (000-1)C surfaces. The stoichiometry of both (000-

1)C and (11-20) and (10-10) surfaces were easily controlled via exposure to C2H4.


The authors would like to thank Cree Research, Inc. for supplying the wafers used in

these experiments. This research was supported by the Office of Naval Research and by the

Department of Education through an Electronic Materials/GAANN fellowship.

100

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71. S.W. King, M.C. Benjamin, J.P. Barnak, R.J. Nemanich, and R.F. Davis, submitted to Journal of Applied Physics. 72. V.S. Smentkowski and J.T. Yates Jr., J. Vac. Sci. Technol. A, 7, 3325 (1989). 73. M.J. Bozack, L. Muehlhoff, J.N Russel Jr., W.J. Choyke, and J.T. Yates, Jr., J. Vac. Sci. Technol. A, 5, 1 (1987) 74. XPS Handbook, Perkin Elmer. XPD of (3x3) 6H-SiC (0001)Si 75. S.W. King, R.S. Busby, R.J. Nemanich, and R.F. Davis, submitted to Surface Science. (3x2) and c(4x2) 3C-SiC (100) 76. V.M Bermudez and J.P. Long, Appl. Phys. Lett., 66, 475 (1995). 77. M.L. Shek, Surface Science, 349, 317 (1996). Group III adatom (√3x√3)R30° Si (111) 78. X. Chen, T. Abukawa, S. Kono, Surface Science, 356, 28 (1996). 79. Y. Taguchi, M. Date, N. Takagi, T. Aruga, M. Nishijima, Applied Surface Science, 82/83, 434 (1994). (√3x√3)R30° Theory 80. J.E. Northrup and J. Neugebauer, Phys. Rev. B., 52, R17001 (1995). 81. P. Badziag, Surface Science, 337, 1 (1995). 82. P. Badziag, Surface Science, 352, 396 (1996). C2H4 - (001) 3C-SiC 83. J.M. Powers, A. Wander, P.J. Rous, M.A. Van Hove, and G.A. Somorjai, Phys. Rev. B, 44, 11159 (1991). 84. V.M. Bermudez and R. Kaplan, Phys. Rev. B, 44, 11149 (1991).

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5. X-ray Photoelectron Diffraction from (3x3) and (√3x√3)R30° (0001)Si

6H-SiC Surfaces.


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Surface Science

by

Sean W. King, *Richard S. Busby, *Robert J. Nemanich, and Robert F. Davis




Raleigh, NC 27695

106

5.1. Abstract

High resolution (±1°) x-ray photoelectron diffraction (XPD) patterns were obtained

along high symmetry azimuths of (3x3) and (√3x√3)R30° reconstructed (0001)Si 6H-SiC

surfaces. The data obtained were compared to previously reported XPD patterns from (7x7)

Si (111) as well as models proposed for the (3x3) and (√3x√3)R30° 6H-SiC reconstructions.

Forward scattering features similar to those observed from (7x7) Si (111) were also observed

from (√3x√3)R30° 6H-SiC (0001)Si surfaces. However, additional features not observed in

(7x7) Si (111) were observed in the (√3x√3)R30° 6H-SiC XPD patterns which were

attributed to the substitution of carbon atoms for silicon atoms on the diamond FCC lattice.

Unlike (1x1) and (7x7) Si (111) surfaces, differences were observed between the XPD

patterns of (3x3) and (√3x√3)R30° SiC (0001)Si surfaces. The most significant difference

observed between the (3x3) and (√3x√3)R30° reconstructions was the equivalence of the [10-

10] and [01-10] azimuths in the (3x3) structure. The differences between the (3x3) and

(√3x√3)R30° XPD patterns were attributed to the presence of an incomplete bilayer of Si the

(3x3) surface. The (3x3) SiC XPD patterns observed in this study are consistent with a

faulted Si bilayer stacking sequence recently proposed based on STM observations.

107

5.2. Introduction

X-ray photoelectron diffraction (XPD) is an exciting new technique for probing the

local atomic structure of metal and semiconductor surfaces with atomic specificity [1-3].

XPD experiments essentially consist of performing angle dependent x-ray photoelectron

spectroscopy (XPS) measurements. Anisotropies in the angular dependence of the intensity

of emitted photoelectrons in XPS are created by interference between other emitted

photoelectron waves and/or scattering by nearest neighbor atoms. For the high kinetic

energies (≈ 1 keV) typically employed, XPD spectra are dominated by forward scattering (or

focusing) of the emitted photoelectron by the potential of the atomic nucleus of nearest

neighbor atoms. This effect which can be viewed as a zeroth order approximation to XPD

creates intensity enhancements along crystallographic and surface-adsorbate bond directions

(see Figure 5.1). Accordingly, this technique has been successfully employed in the

determination of surface adsorption sites for various atoms and molecules on metals and

semiconductors as well as for studying a number of different epitaxial growth systems [1-6].

Most recently, further developments and enhancements of this technique have actually

resulted in the demonstration of holographic images of various metal [7-10] and

semiconductor surfaces [11]. In this paper, we apply XPD to study the atomic structure of

(3x3) and (√3x√3)R30° reconstructed (0001)Si 6H-SiC surfaces.

108

Emitter Scatterer

Intensit

y

Lost

Intensity Gained

Figure 5.1. Schematic illustrating forward focusing/scattering effects in x-ray photoelectron diffraction experiments.

SiC is a wide band gap compound semiconductor (Eg (6H-SiC) = 3.0 eV) which is of

considerable importance to the development of high temperature, high frequency, and high

power electronic devices [12]. The ability to develop SiC into the material of choice for

these applications, however, has been currently limited by the inability to control the types

and densities of a variety of line, planar, and macroscopic defects in SiC wafers and films

[12]. By analogy to a more thoroughly investigated material such as silicon [13-16], it is

conceivable that many of these defects originate and/or nucleate at defects on the SiC

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surface. Therefore in order to understand how these defects originate, a detailed

understanding of the atomic structure of the SiC surface is needed. For the (0001) surface of

6H-SiC, this has in part been provided by many recent scanning tunneling microscopy (STM)

studies [17-24] which have identified a variety of different surface reconstructions ranging

from (3x3), (√3x√3)R30°, (9x9), (6x6), and (6√3x6√3)R30°.

By analogy to the group III adatom (√3x√3)R30° Si (111) reconstructed surfaces [25-

28], many have proposed that (√3x√3)R30° 6H-SiC (0001) surface reconstructions are due to

bulk terminated (0001) 6H-SiC surfaces with a 1/3 ML (monolayer) coverage of silicon or

carbon adatoms in the T4 position (see Figure 5.2) [29,30]. Recent STM investigations by

Owman and Martensson [19] and Li and Tsong [21] have been able to confirm the three fold

symmetric unit cell, but unfortunately were unable to resolve the chemical identity of the

adatom or determine the exact position of the adatom (i.e. T4 or H3). However, Owman and

Martensson [19] were able to determine that the reconstruction was not composed of a

mixture of Si and C adatoms or a mixture of T4 and H3 sites (i.e. single adatom on a single

site). These findings by Owman and Martensson [19] are complementary to the theoretical

results of Northrup and Neugebauer [31]. Their recent supercell calculations using the

density functional method have shown that for (v3xv3)R30° (111) 3C-SiC surfaces, Si

adatoms are preferred over C adatoms and that the T4 site is favored over the H3 site by both

Si and C adatoms. These results are consistent with previous calculations by Northrup which

showed the T4 site to be preferred in Si (111) (√3x√3)R30°:Si adatom geometries [32

In contrast, semi-empirical, self consistent quantum mechanical cluster calculations

by Badziag [33,34] show that for the (0001)Si (√3x√3)R30° reconstructed surface a

hydrogenated triangle of C atoms (i.e. similar to cyclopropane) centered on the T4 position is

110

energetically more favorable than C or Si adatoms hydrogenated or unhydrogenated. This

hydrogenated C3 model is similar in nature to the milk stool model deduced for (√3x√3)R30°

Si (111):Sb surfaces where the Sb atoms form trimers centered on the T4 site [35,36]. The

validity of this model is somewhat questionable since it would require hydrogen to not

desorb from the SiC surface until temperatures of > 1150°C at which point the (√3x√3)R30°

reconstruction disappears. In defense of his model, Badziag points out that for diamond,

hydrogen desorbs at temperatures of 1000 and 1150°C for the (111) [37] and (100) [38]

surfaces which is consistent with the observed stability of this reconstruction. However,

Allendorf and Outka [39] observed two hydrogen desorption peaks from polycrystalline SiC

surfaces at ≈ 700 and 850°C which is well below the temperature at which this reconstruction

is observed to occur.

111

View Along [11-20] or [1-10][0001], [111]

[10-10], [1-21]

T4 H3b.)

H3

T4

<10-10>a.)

<11-20>

C atomSi atom

Figure 5.2. Schematics illustrating various adatom adsorption sites for (√3x√3)R30° reconstructions on (111)/(0001) surfaces. (a) Top down view along [000-1], (b) Side view along [11-20].

112

For the (3x3) (0001)Si 6H-SiC surface, Kaplan [29] originally proposed a model

based on AES data for a SiC surface terminated by a bilayer of silicon. Based on analogy to

the (7x7) Si (111) DAS model, Kaplan proposed a (3x3) unit cell which consisted of two

adatoms, six rest atoms (three dimers), and eight silicon atoms in the second layer positioned

approximately directly over the silicon atoms of the SiC substrate. However, the recent STM

results of Kulakov et al [22] detected only one maxima (i.e. one adatom) in the (3x3) unit cell

which is in contrast to the model proposed by Kaplan which would predict two maxima.

Based on this discrepancy, Kulakov et al [22] proposed a modified structure which was

consistent with the AES results of Kaplan and their STM data. The model for the (3x3)

surface proposed by Kulakov et al [22] consists of a unit cell with 1 adatom, 3 rest atoms, and

7 silicon atoms located approximately on top of the silicon atoms of the SiC surface (see

Figure 5.3). This model includes 3 dimers and three dangling bonds (two unsatisfied Si

bonds from the SiC substrate, and 1 dangling bond from the adatom) compared to the 4

dangling bonds in the model by Kaplan [29]. However, Kulakov et al [22] did observe

stacking faults in their (3x3) reconstructed surface which had a structure essentially like that

of the (3x3) model proposed by Kaplan [29] (see Figure 5.4). Using STM, Li and Tsong [21]

also confirmed the presence of one maxima in the (3x3) unit cell but, in contrast, concluded

that the (3x3) reconstruction consisted of only 4/9 ML coverage of silicon for the (0001)Si

6H-SiC surface. Accordingly they attributed the (3x3) surface to extra Si-C tetrahedra on the

surface distributed in a 3x3 pattern (see Figure 5.5) rather than a bilayer of silicon.

113

[11-20]

[10-10]

a.)

View Along [10-10] or [1-21] [001], [111]

[11-20], [1-10]

b.)

View Along [11-20] or [1-10] [0001], [111]

[10-10], [1-21]

C atomSi atomSi Dangling

Bond

c.)

114

Figure 5.3. Model proposed by Kulalov et al [22] for the (3x3) reconstructed (0001)Si 6H-SiC surface. (a) Top down view along [000-1], (b) side view along [11-20], and (c) [10-10].

[11-20]

[10-10]

View Along [11-20] or [1-10] [0001], [111]

[10-10], [1-21]

C atomSi atomSi Dangling

Bond

a.)

b.)

Figure 5.4 Model proposed by Kaplan [29] for the (3x3) reconstructed (0001)Si 6H-SiC surface. (a) Top down view along [000-1], (b) side view along [11-20].

115

[11-20]

[10-10]

Figure 5.5. Top down view of model proposed by Li and Tsong [21] for the (3x3) reconstructed (0001)Si 6H-SiC surface.

In this paper, we report the first XPD patterns obtained from (0001)Si 6H-SiC

surfaces. The XPD patterns obtained from the (3x3) and (√3x√3)R30° (0001)Si 6H-SiC

surfaces are compared with those obtained from (7x7) (111) Si surfaces and the proposed

models for these SiC reconstructions. Based on this data, we are able to support only some

of the proposed models for these reconstructions. Estimates of the experimental inaccuracies

116

possible in XPS experiments due to anisotropies in the angular dependence of the intensity of

Si 2p and C 1s photoemission are also provided by this data.

5.3. Experimental

The experiments described in this paper were conducted in an integrated surface

analysis and growth system which has been previously described [40,41]. In this study, only

the XPS and SiC ALE systems were used. The n-type (Nd=1018/cm3), off axis (4° toward

{11-20}) (0001)Si 6H-SiC wafers used in this research were supplied by Cree Research, Inc.

with an ≈ 1 µm n-type 6H epilayer (Nd=1017/cm3) and a 1000Å thermally grown oxide.

The back side of the SiC wafer was sputter coated with tungsten after removal of the thermal

oxide with a 10 min. dip in 10:1 HF. The back side tungsten coating was necessary in order

to improve the heating efficiency of the SiC wafer by our tungsten filament heater as SiC is

transparent in the infra-red. Prior to insertion into the SiC ALE system, the SiC wafers were

given an ex situ clean consisting of ultrasonification in trichloroethylene, acetone, and

methanol for 10 min. each, followed by a 10 min. 10:1 buffered HF vapor clean to remove

any native oxides. The SiC wafer was then loaded into the SiC ALE system and annealed in

10-6 Torr SiH4 for 15 min. at 1050°C. This produced an oxygen free (3x3) reconstructed

surface. The (√3x√3)R30° reconstruction was generated by annealing the (3x3) surface in

UHV in the ALE system at 1050°C for ≈ 10 min. Additional details regarding our sample

preparation and procedure have been previously published [40-42,43].

After either the (3x3) or (√3x√3)R30° surface had been prepared, the SiC wafer was

transferred in situ to the XPS system. XPD patterns were acquired in this system by rotating

the SiC wafer about various polar and azimuthal angles using a computer driven goniometer 117

with five degrees of freedom (x,y,z,θ, and φ) while the positions of the x-ray source and

electron energy analyzer were fixed. Though the angular acceptance of the lens of the

electron energy analyzer (VG CLAMII) was ± 7°, an angular resolution of ≈ ±1° was

achieved by geometric constraints via grounding the lens and using smaller channeltron

acceptance slits. The SiC XPD patterns were acquired by monitoring the Si 2p and C 1s core

levels photoexcited by Al Kα radiation (hν = 1486.6 eV). The kinetic energy of these

photoelectrons (≈ 1380 and 1203 eV respectively) was sufficiently high that forward

scattering effects should be dominant and probe ≈ 20Å of the SiC surface. Polar scans along

high symmetry azimuths were acquired in increments of 0.9° from -35° to 70°. The wafer

flat provided by Cree was used to locate the various azimuths and is of the {10-10} family of

planes. The data presented is the raw angular distribution of the measured intensity of the Si

2p and C 1s core levels. No attempts were made to correct for background, variation in

sampling depth, or surface area seen by the electron energy analyzer. To ensure that the

system was operating properly, XPD spectra were first acquired from Si (100) and Si (111)

surfaces and compared with previously published high resolution XPS spectra for these

surfaces [44-48]. Figure 5.6 displays an XPD spectrum from a (2x1) Si (001) surface along

the [110] azimuth. As illustrated, sharp features with FWHM ≅ 3° were easily resolved and

were found to be in excellent agreement with previously reported results for this surface [48].

118

-40 -30 -20 -10 0 10 20 30 40 50 60 70

45Þ15Þ

20Þ23Þ

Cou

nts (

arb.

uni

ts)

Polar Angle (ÞTheta)

30Þ

60Þ

Figure 5.6. XPD pattern from (2x1) Si (100) along the [110] azimuth.

5.4. Results

Figures 5.7-5.11 show various Si 2p and C 1s XPD patterns obtained from (7x7) Si

(111), and (3x3) and (√3x√3)R30° 6H-SiC (0001)Si surfaces. These patterns were acquired

along high symmetry azimuths such as [10-10], [11-20], and [01-10] azimuths (note <1-

21>cub = <10-10>hex and <1-10>cub = <11-20>hex). As shown in these figures, peaks of

varying FWHM were observed. Generally speaking, diffraction structures of FWHM ≈ 10-

20° are consistent with forward scattering features associated with crystallographic directions

or nearest neighbor atoms [1-3]. Narrow peaks or wide peaks with fine structure of ≈ 3° are

usually due to scattering from more distant atoms or complex/higher order interference

phenomena [1-3].

119

-10 0 10 20 30 40 50 60 70

Cou

nts (

arb.

uni

ts)


(a)

(b)

(c)

Figure 5.7. Si 2p x-ray photoelectron diffraction pattern along [1-21]/[10-10] azimuths from (a) (7x7) Si (111), (b) (3x3) 6H-SiC (0001)Si, and (c) (v3xv3)R30° (0001)Si 6H-SiC.

120

-10 0 10 20 30 40 50 60 70

Cou

nts (

arb.

uni

ts)


(a)

(b)

(c)

Figure 5.8. Si 2p x-ray photoelectron diffraction patterns from (a) (3x3) 6H-SiC (0001)Si along [01-10], (b) (3x3) 6H-SiC (0001)Si along [10-10], and (c) (v3xv3)R30° 6H-SiC (0001)Si along [01-10].

121

-10 0 10 20 30 40 50 60 70

Cou

nts (

arb.

uni

ts)


(a)

(b)

(c)

Figure 5.9. Si 2p x-ray photoelectron diffraction patterns along [-110]/[11-20] from (a) (7x7) Si (111), (b) (3x3) 6H-SiC (0001)Si, and (c) (v3xv3)R30° 6H-SiC (0001)Si.

122

-15 -10 -5 0 5 10 15 20 25 30 35 40 45 50 55 60 65 70

Cou

nts (

arb.

uni

ts)


(a)

(b)

Figure 5.10. C 1s XPD patterns from (v3xv3)R30° 6H-SiC (0001)Si along (a) [01-10], and (b) [11-20].

123

-10 0 10 20 30 40 50 60 70

Cou

nts (

arb.

uni

ts)


(a)

(b)

(c)

Figure 5.11. C 1s XPD patterns from (3x3) 6H-SiC (0001)Si along (a) [10-10], (b) [11-20], and (c) [10-10] azimuths.

124

In order to determine which peaks in the XPD spectra were due to forward scattering

effects, we next consider the atomic structure of the (0001)Si 6H-SiC surface. SiC is a

unique material that exhibits several different polymorphs which differ only in the stacking

sequence along the c axis. This particular phenomena is referred to as polytypism, and as

many as 256 different polytypes of SiC have been reported. However, only a few polytypes,

3C, 4H, 6H, and 15R are commonly observed. In the Ramsdell notation used to describe

these polytypes, the preceding number represents the number of Si-C bilayers needed to

repeat the stacking sequence along [111]/[0001] directions and the following letter describes

the crystal structure (i.e. C = cubic, H = hexagonal, and R = rhombohedral). In the case of

6H-SiC, the H is deceiving as 6H-SiC is actually 66.6% cubic and exhibits an ABCB'C'A'

stacking sequence which is similar to that of 3C-SiC differing only in the periodic stacking

fault in the 6H structure (see Figure 5.12). Accordingly, in a surface sensitive technique such

as XPD which effectively only samples the first 10Å of the surface, (0001) 6H and (111) 3C-

SiC should be essentially indistinguishable. Therefore for simplicity sake, we will treat the

(0001) 6H XPD spectra as if it were from (111) 3C-SiC. This is fortuitous as 3C-SiC and Si

and have similar crystal structures and therefore comparisons can be made between XPD

spectra from (111) Si and (111)/(0001) 3C/6H-SiC. Based on these considerations, the

expected peaks for forward scattering/focusing along certain crystallographic directions for

bulk terminated (111) 3C-SiC, and (0001)Si 6H-SiC surfaces are listed in Table 1 (see

Figures 5.13 and 5.14). Table 5.1 presents the forward scattering peaks expected from both

C 1s and Si 2p photoelectrons. In the case of Si (111), the C 1s forward scattering peaks

125

would be expected to appear in Si 2p XPD patterns as the carbon atoms have been replaced

by Si atoms.

3C-SiC

[111]

[1-21][1-10]

6H-SiC

[0001]

[10-10][11-20]

Stacking = ABCABC

Stacking = ABCB'C'A'

126

Figure 5.12. Schematic illustrating the differences in stacking along the [111]/[0001] direction for 3C and 6H-SiC.

View Along [11-20] or [1-10] [0001], [111]

[10-10], [1-21]

35.2Þ

54.7Þ70.5Þ

29.5Þ

0Þ

35.2Þ

Figure 5.13. Schematic illustrating expected forward scattering/focusing peaks in XPD along the [11-20] azimuth of 3C/6H-SiC.

127

58.5Þ

72.9Þ44.4Þ31.4Þ

View Along [10-10] or [1-21] [0001], [111]

[11-20], [1-10]

Figure 5.14. Schematic illustrating expected forward scattering/focusing peaks in XPD along the [10-10] azimuth for 3C/6H-SiC. Table 5.1. Expected forward scattering/focusing peaks from bulk terminated (111) Si, (111) 3C-SiC, and (0001)Si 6H-SiC surfaces along [10-10], [11-20], and [01-10] azimuths.

Si 2p [10-10] [11-20] [01-10] Scatterer 35.3° Si 54.7° Si 58.5° C 70.5° C-Si 72.9° C C 1s [10-10] [11-20] [01-10] Scatterer 29.5° Si 31.4° Si 35.3° C 44.4° Si 54.7° C 70.5° Si

5.4.1. (√3x√3)R30° (0001)Si 6H-SiC

128

For Si 2p XPD patterns obtained from (√3x√3)R30° 6H-SiC (0001)Si surfaces, most

of the expected forward scattering peaks were identified. In the [10-10] azimuth (see Figure

5.7(c)), a broad/intense peak at 36-37° was identified as expected. This peak was similar to

and consistent with forward scattering along the [011]/[10-11] crystallographic axes (i.e. the

Si-C atomic row). However unlike the (7x7) Si (111) surface, broad/intense peaks at 19 and

42° were symmetrically observed on both sides of the [011]/[10-11] forward scattering peak.

As mirror symmetry is expected about the [011] atomic row due to (100) glide planes [44],

these additional peaks are probably due to higher order interference phenomena or forward

scattering from larger emitter-scatterer distances. Similar to (7x7) Si (111), peaks of this

nature were also observed at 15 and 59° in Si 2p XPD patterns along the [10-10] azimuth.

However in contrast to (7x7) Si (111), a forward scattering peak at 70.5° was not observed.

In the case of Si (111), this peak is due to forward scattering along the (11-1) crystallographic

direction. The absence of this peak from (√3x√3)R30° 6H-SiC (0001)Si surfaces may be

related to the fact that (0001) and (000-1) are not equivalent directions in SiC but are

equivalent directions in Si.

C 1s XPD patterns obtained along the [10-10] azimuth showed a single sharp peak at

30° which is 5° off from the expected value of 35° for forward scattering along the [011]/[10-

11] atomic row (see Figure 5.10(a)). This discrepancy, however, may be related to the fact

that the 35° peak would be expected based on forward scattering by carbon atoms. Carbon

has a smaller atomic nucleus and should be expected to be a weaker scatterer. Therefore the

position of this forward scattering peak maybe determined more by silicon atoms along the

[011]/[10-11] row.

129

Like Si (111), asymmetry's were observed between Si 2p XPD spectra acquired along

<10-10> and <01-10> azimuths of (√3x√3)R30° 6H-SiC (0001)Si (i.e. <10-10> ≠ <01-10>).

In the [01-10] azimuth, the expected peak for forward scattering in the [100] direction was

observed at 55°C (see Figure 5.8(c)). Similar to Si 2p XPD patterns along [-12-1] azimuths

[44], peaks due to complex/higher order interference phenomena were also observed from

10-40° in the [01-10] azimuth of (√3x√3)R30° 6H-SiC (0001)Si surfaces. However, unlike

Si, symmetry was not observed about this peak.

In the [11-20] azimuth, a mosaic of broad diffraction peaks of equal intensity were

observed from Si 2p XPD patterns from the (√3x√3)R30° 6H-SiC (0001)Si surface (see

Figure 5.9(c)). Most of the peaks observed in the Si 2p XPD pattern from (7x7) Si (111)

observed along the [1-10] azimuth were also observed in the Si 2p XPD patterns along the

[11-20] azimuth from the (√3x√3)R30° 6H-SiC (0001)Si surface. However, the peak at

58.5° expected for forward scattering along the [-131] direction was observed to have a

volcano shape for (√3x√3)R30° 6H-SiC (0001)Si surface and a rounded shape for the (7x7)

Si (111) surface. In fact, rounded peaks from the Si (111) surface along [-110] were

observed to have a volcano shape in the [11-20] azimuth of the SiC surface (and vice versa).

The C 1s XPD patterns obtained along the [11-20] azimuth showed strong peaks at 25, 35,

and 52.5° (see Figure 5.10(b)). The sharpest/most intense peak at 35° is slightly off from the

expected value of 31.4° for scattering by a top layer carbon atom (see Figure 5.14).

However, this can be expected as the scattering carbon atom does not actually lie in the (10-

10) plane.

The maximum anisotropy in intensity observed in both the Si 2p and C 1s XPD

patterns was observed about the 0°/[0001] forward scattering peak. For Si 2p and C 1s the 130

maximum anisotropy (Imax-Imin)/Imax was ˜ 65 and 40 % respectively. Similar to (7x7) Si

(111), higher order diffraction effects were also observed at 10-15° on both sides of the Si 2p

and C 1s 0°/[0001] forward scattering peaks. However in contrast to Si (111), the Si 2p

0°/[0001] forward scattering peak from (0001)Si 6H-SiC surfaces does not exhibit a volcano

type shape but rather a flat sawtooth type shape (see Figure 5.8(a)). The C 1s 0°/[0001]

forward scattering peak from (0001)Si 6H-SiC does exhibit a volcano shape (see Figure

5.10(a)). A similar effect has been observed between Si 2p and C 1s spectra from (001) Si

and 3C-SiC [48-50]. The shape of this peak is strongly affected by the presence of scattering

atoms surrounding the [111]/[0001] direction. As silicon is the nearest neighbor atom to

carbon along the [0001] direction, scattering by these silicon atoms probably induces the

volcano shape observed in the C 1s XPD. For silicon atoms in SiC, carbon is the nearest

neighbor atom, but the scattering factor of carbon is much weaker, hence the sawtooth

structure. However in pure silicon and diamond, all the atoms are either silicon or carbon

and the volcano shape reappears [44-48,51]. The also explains many of the differences

between Si 2p XPD spectra from Si and SiC along the [11-20] azimuth. Finally, it should be

mentioned that the centroid of the Si 2p and C 1s 0°/[0001] forward scattering peaks were

observed to vary by ± 2°. As this variation was not observed from on axis (001) and (111) Si

substrates, the variation in the position of [0001] in the polar scans is probably related to the

fact that the SiC wafers were 4° off axis in the <11-20> direction and not related to drift in

our sample stage.

5.4.2. (3x3) (0001)Si 6H-SiC

131

In contrast to reconstructed and unreconstructed diamond and silicon surfaces [44-

48,51], significant differences were observed between the Si 2p XPD patterns from (3x3) and

(√3x√3)R30° 6H-SiC (0001)Si surfaces. For instance, the [011]/[101-1] forward scattering

peak at ≈ 33° was observed from [10-10] Si 2p XPD patterns from both (3x3) and

(√3x√3)R30° 6H-SiC surfaces (see Figure 5.7(b) and (c)). However, peaks centered

symmetrically at 29 and 46° about the [011]/[101-1] forward scattering feature were not

observed from the (3x3) surface (see Figure 5.7(b)) which is more similar to Si 2p XPD

patterns from Si (111) in this azimuth (see Figure 5.7(a)). Additionally, the sharp higher

order diffraction peaks at 9-18 and 58° observed from the (√3x√3)R30° surface were

observed to be more broad and less intense for the (3x3) surface. Similar differences were

also seen in the (3x3) C 1s XPD patterns in the [10-10] azimuth. In this case, a volcano

shaped peak centered at 35° was observed in [10-10] C 1s XPD patterns from the (3x3)

surface instead of the one sharp peak centered at 30° observed from the (√3x√3)R30° 6H-SiC

surface (see Figure 5.11(c)). Sharper peaks centered symmetrically about the 35° volcano

peak at 20 and 59° were also observed in the (3x3) [10-10] C 1s XPD patterns.

In the [11-20] azimuth, a mosaic of sharp higher order diffraction features were

observed in C 1s XPD patterns from the (3x3) surface instead of the single sharp peak

centered at 35° observed in the (v3xv3)R30° XPD patterns (see Figure 5.11(b)). However

for [11-20] Si 2p XPD patterns, there were not any significant differences between the (3x3)

and (√3x√3)R30° 6H-SiC surfaces (see Figure 9(b),(c)). Finally in Si 2p XPD patterns, the

0°/[0001] forward scattering peak was observed to be volcano shaped for the (3x3) surface

whereas this peak exhibited a sawtooth/rounded shape for the (√3x√3)R30° 6H-SiC surface

(see Figures 5.7(b),(c) and 5.8(b),(c)). The opposite, however, was observed in C 1s XPD

132

patterns where a volcano shaped 0°/[0001] forward scattering peak was observed from the

(√3x√3)R30° surface whereas a sawtooth/rounded peak was observed from the (3x3) surface

(see Figures 5.10 and 5.11).

Perhaps the largest differences between XPD of (3x3) and (√3x√3)R30° 6H-SiC

surfaces are found in the [01-10] azimuth. As can be seen in Figures 5.8(a),(b) and

5.11(a),(c), Si 2p and C 1s XPD patterns from the (3x3) 6H-SiC surface along the [01-10]

and [10-10] azimuths are identical. This is in complete contrast, to the (√3x√3)R30° 6H-SiC

surface in which the [10-10] and [01-10] azimuths were observed to be completely different

(i.e. asymmetric about [0001] along <10-10>). The equivalence of the [10-10] and [01-10]

directions for the (3x3) 6H-SiC surface suggests drastic changes in the surface structure of

the SiC which will be discussed further in the next section.

5.5. Discussion

5.5.1. (√3x√3)R30° (0001)Si 6H-SiC

As some peaks were observed in XPD patterns from (√3x√3)R30° (0001)Si 6H-SiC

which had not been previously observed from Si (111) surfaces, attempts were initially made

to see if these extra peaks could be assigned to forward scattering/diffraction due to silicon or

carbon adatoms on the surface. As previously mentioned, adatoms in T4 or H3 sites are

commonly believed to be the origin of the (√3x√3)R30° reconstruction. Table 5.2 lists the

additional expected forward scattering peaks due to Si or C adatoms based on various

different models proposed for the (√3x√3)R30° reconstruction [31-34]. Unfortunately, we

133

were not able to determine with any certainty whether any of these peaks truly existed.

Therefore, single scattering cluster simulations are probably necessary in order to determine

the exact structure of the (√3x√3)R30° reconstruction based on XPD data. However, the

authors note that Pirri et al [44] and Kuttel et al [51] have experienced similar difficulties in

distinguishing between XPD patterns from (7x7) and (1x1) Si (111) and (2x1) and (1x1)

Diamond (111) respectively.

Table 5.2. Expected Si 2p and C 1s photoelectron diffraction peaks for adatom scattering in T4 and H3 positions based on data of Northrop and Neugebauer [31] and Badziag [34] for (√3x√3)R30° (0001)Si 6H-SiC.

Si adatom, T4 Northrop & Neugebauer [31] <10-10> <11-20> <-1010> Si 2p C 1s Si 2p C 1s Si 2p C1s 66.9° 56.0 45.2° 65.2° 53.6° 52.1° 43.2° 20.1° 36.5° C adatom, T4 Northrop & Neugebauer [31] <10-10> <11-20> <-1010> Si 2p C 1s Si 2p C 1s Si 2p C1s 71.8° 59.1° 53.8° 55.3° 56.1° 22.3° 40.6° C3 Model Badziag [34] <10-10> <11-20> Si 2p C 1s Si 2p C 1s 68.8° 48° 33.3° 37.9° 27.3° 20.2° 17.0° 14.5° 9.5°

134

5.5.2. (3x3) (0001)Si 6H-SiC

As mentioned above, significant differences were observed between Si 2p and C 1s

XPD patterns from the (3x3) and (√3x√3)R30° 6H-SiC (0001)Si surfaces. Perhaps, the most

striking difference was the observed equality of the [10-10] and [01-10] azimuths for the

(3x3) reconstruction and the inequality of these azimuths for the (√3x√3)R30° reconstruction.

In order to gain further insight into the nature of these differences, comparisons were made to

previously proposed models for the (3x3) reconstruction based on recent STM images

[21,22]. Based on these models, a new set of forward scattering peaks were

calculated/estimated for the (3x3) reconstruction and which are presented in Table 5.3. As

discussed in the introduction, Li and Tsong [21] have proposed that the (3x3) reconstruction

is a result of 4/9 ML absorption of Si-C tetrahedra arranged in a (3x3) pattern (see Figure

5.5). As can be seen in Figure 5.5 and Table 5.3, this model does not predict equivalence of

the [10-10] and [01-10] azimuths. This model is also in disagreement with our previous

observation of a Si-Si Si 2p bonding peak in XPS which indicated a partial bilayer of silicon

on the (3x3) SiC surface [40].

Based on the previous AES data of Kaplan [29] and their STM data, Kulakov et al

[22] proposed a different model for the (3x3) reconstruction which consisted of an

incomplete bilayer of Si. We find this model for the (3x3) reconstruction to be in better

agreement with our observed XPS and XPD patterns. First, this is clearly consistent with

our observation of a Si-Si Si 2p bonding peak in XPS. As can be seen in Figure 5.3(c), this

model also specifically adds an additional Si-Si bilayer to the [011]/[10-11] atomic row. As

silicon has a larger nucleus it should be a more effective scatterer than carbon. Therefore the

135

addition of the Si-Si bilayer should enhance the intensity along the [011]/[10-11] chain due to

increased forward scattering. This is exactly what we observe in both our C 1s and Si 2p [10-

10] XPD patterns. Unfortunately, the model proposed by Kulakov et al fails to explain the

observed equality of our [10-10] and [01-10] Si 2p and C 1s patterns. However, we do note

that the model originally proposed by Kaplan [29] for the (3x3) reconstruction would explain

the equivalence of the [10-10] and [01-10] XPD patterns (see Figure 5.4). This is primarily a

result of the stacking fault in this structure which produces Si-Si bilayers oriented in both the

[10-10] and [01-10] directions. The presence of this structure on SiC surfaces has actually

been confirmed by Kulakov et al [22] were they observed faults or domains of different

orientation in STM images of the (3x3) surface. The stacking structure in these domains are

consistent with the model originally proposed by Kaplan for the (3x3) reconstruction [29].

Unfortunately however, at this stage it is difficult to determine with certainty whether our

observations of the equivalence between [10-10] and [01-10] in XPD are due to these

stacking faults.

Table 5.3. Estimated forward scattering peaks for (3x3) reconstructed (111)/(0001) 3C/6H-SiC surfaces based on models proposed by Kulakov et al [22] and Li and Tsong [21].

Si 2p Kulakov Li & Tsong [11-20] [10-10] [01-10] [11-20] [10-10] [01-10] 59.3° 66.2° 66.2° 59.3° 66.2° 48.6° 51.5° 53.7° 53.7° 52.6° 29.5° 42.9° 29.5° 48.6° 42.9° 38.2° 33.2° C 1s Kulakov Li & Tsong [11-20] [10-10] [01-10] [11-20] [10-10] [01-10] 59.4° 43.3° 38.0° 58.5° 43.3° 58.5°

136

39.3° 30.8° 30.8° 39.3° 30.8°

27.3° 21.3° 25.3° 25.3°

137

5.6. Conclusion

High resolution (±1°) x-ray photoelectron diffraction (XPD) patterns were obtained

along high symmetry azimuths of (3x3) and (√3x√3)R30° reconstructed (0001)Si 6H-SiC

surfaces. The data obtained were compared to previously reported XPD patterns from (7x7)

Si (111) as well as models proposed models for the (3x3) and (√3x√3)R30° 6H-SiC

reconstructions. Forward scattering features similar to those observed from (7x7) Si (111)

were also observed from (√3x√3)R30° 6H-SiC (0001)Si surfaces. However, additional

features not observed in (7x7) Si (111) were observed in the (√3x√3)R30° 6H-SiC XPD

patterns which were attributed to the substitution of carbon atoms for silicon atoms on the

diamond FCC lattice. Unlike (1x1) and (7x7) Si (111) surfaces, differences were observed

between the XPD patterns of (3x3) and (√3x√3)R30° SiC (0001)Si surfaces. The most

significant difference observed between the (3x3) and (√3x√3)R30° reconstructions was the

equivalence of the [10-10] and [01-10] azimuths in the (3x3) structure. The faulted (3x3)

structure proposed by Kulakov et al [22] was found to be consistent with the (3x3) XPD

patterns presented here.


The authors would like to thank Cree Research, Inc. for supplying the 6H-SiC wafers.

.8 References

The authors would also like to thank the REU program for partially sponsoring this effort.

Meaningful discussions regarding experimental setup with Dr. Egelhoff are also noted. The

research was supported by ONR under contract and the Department of Education via an

Electronic Materials/GAANN fellowship..

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141

6. Interaction of Atomic Hydrogen with (3x3) 6H-SiC (0001)Si Surfaces.


to

Surface Science

by

Sean W. King, *Mark C. Benjamin, John P. Barnak,

*Robert J. Nemanich, and Robert F. Davis




Raleigh, NC 27695

142

6.1. Abstract

X-ray photoelectron spectroscopy (XPS), Auger electron spectroscopy (AES), low

energy electron diffraction (LEED), and temperature programmed desorption (TPD) were

used to examine the interaction of atomic hydrogen with (3x3) 6H-SiC (0001)Si surfaces. It

was found that atomic hydrogen exposure selectively removes silicon from the SiC surface

converting the (3x3) surface to a (1x1) surface. Selective removal of silicon was witnessed

by the reduction and removal of the Si-Si bonding Si 2p XPS peak from (3x3) (0001)Si 6H-

SiC surfaces exposed to atomic hydrogen. Additional etching of the SiC surface was

indicated by the reduction in the Si LVV/C KLL ratio in AES from 1.3 to 0.4 following

exposure of (3x3) surfaces to a remote rf H plasma. TPD of atomic H treated (3x3) SiC

surfaces showed weak hydrogen desorption in the range of 400-600°C where desorption from

silicon atoms would be expected by analogy to (111) Si. However, the hydrogen desorption

signal increased at higher temperatures where hydrogen desorption from carbon sites would

be expected based on analogy to (111) diamond surfaces. C-H termination of the SiC surface

was supported by the observation of some C-C bonding after thermal desorption of rf plasma

treated SiC surfaces at T > 1000°C.

143

6.2. Introduction

Hydrogen is a common constituent in many semiconductor processes including

chemical vapor deposition (CVD), reactive ion etching (RIE), wet chemical cleaning, gas

source molecular beam epitaxy (GSMBE), and rapid thermal annealing (RTA) [1-12].

Knowledge of the interaction and chemistry of hydrogen at semiconductor surfaces is

therefore of great importance in order to understand the fundamentals of the above processes.

For these reasons, many surface analytical studies concerned with the interaction of hydrogen

with semiconductor surfaces such as silicon, diamond, and gallium arsenide [13-46] have

been conducted. However, there have been relatively few surface analytical studies

concerned with the interaction of hydrogen and silicon carbide (SiC) [47-51]. SiC is a wide

bandgap semiconductor (Eg (6H-SiC) = 3.0 eV) which is of interest for high power, high

frequency, and high temperature electronic devices due to its excellent oxidation resistance,

high saturation electron drift velocity (vsat = 2x107 cm/s), high breakdown voltage (EB = 2.5

MV/cm) , high thermal conductivity (κ = 4.9 W/cm K), and high melting point (Tmelt ≈

3000°C) [52,53]. Due to moderately close lattice matching (?a/ao AlN/SiC = 0.8%, GaN/SiC

= 3.5%), SiC is also of interest as a heteroepitaxial substrate for growth of III-V nitride

compounds which in turn are of interest for blue/UV optoelectronic applications as well as

high power and high frequency devices [54]. Unfortunately though, many of the same

properties which make SiC of interest in these demanding conditions also makes it a

challenging material to work with from a processing point view. Therefore in order for SiC

to succeed in many of these applications, advances must be made in SiC processing such as

144

growth, etching, contact formation, and surface cleaning [52]. As many of these processes

are currently based on using hydrogen or hydrogenated species, an increased understanding

of the interaction of hydrogen with SiC surfaces should assist in the further development of

these processes.

Some of the first investigations of the interaction of hydrogen with SiC revealed that

etching of SiC by hydrogen occurred at high temperatures (>1500°C) [55-58]. The work of

Chu and Campbell [55] in particular found hydrogen to be a non-preferential etchant for

single crystal hexagonal SiC yielding useful etch rates of 2-4 microns/min. at temperatures of

1600-1700°C. Further, Bartlett and Mueller [57], found an H2 etch prior to SiC CVD to be

instrumental in obtaining good homoepitaxy. Subsequent hydrogen/SiC studies focused on

the interaction of high energy H+ and D+ ions (i.e. sputtering and implantation) with

polycrystalline 3C-SiC which in these studies was being considered as a first wall material

for thermonuclear fusion reactors [59-65]. However, more recent studies have investigated

the interaction of low energy (i.e. thermally generated) atomic hydrogen with polycrystalline

3C-SiC surfaces [47-48]. The first such study by Allendorf et al [47] used temperature

programmed desorption (TPD) and Auger electron spectroscopy (AES) to investigate the

adsorption and desorption of thermally generated hydrogen on sputter cleaned polycrystalline

3C-SiC surfaces. They observed > 1 ML adsorption of hydrogen which was observed via

TPD to desorb in a broad temperature range from 400-1000°C. Analysis of the broad

desorption feature indicated two desorption peaks at 700 and 850°C characterized by first

order desorption with activation energies of 63 and 72 kcal/mole respectively. Due to the

polycrystalline nature of these surfaces, Allendorf et al [47] unfortunately were not able to

assign these two desorption features to desorption from specific sites. However, they did

145

show a reduction in the AES Si/C ratio from 1.31 to 0.46 after atomic H exposure suggesting

etching or selective etching of the SiC surface. A similar effect was observed by Lannon et

al [49] for polycrystalline 3C-SiC films exposed to (H/H2)+ ions of various energies (10-

2000 eV). In this case, the 3C-SiC films were grown in situ via carbonization of (001)

silicon wafers with C2H4 and were not exposed to any sputtering. Finally, the reaction of

thermally generated atomic hydrogen with polycrystalline 3C-SiC films was also studied by

Kim and Olander [48] using modulated molecular beam mass spectrometry. In this case,

they were able to observe etching of SiC by atomic hydrogen at temperatures ranging from

300-1100K via the detection of SiH4, CH4, and C2H2 reaction products. Based on their

study and those of Allendorf et al [47], Kim and Olander proposed a precursor model for

SiH4 and CH4 formation based on a surface composed of adsorbed atomic hydrogen

overlaying mono and dihydrides of silicon and carbon on the SiC surface. In their model, the

dihydrides act as the precursors for SiH4 and CH4 generation and the production of these

species follows a first order reaction between dihydrides and the overlayer of adsorbed

atomic hydrogen. However, the results of Kim and Olander and those of Kim and Choi [58]

showed pronounced and enhanced etching of polycrystalline 3C-SiC surfaces at grain

boundaries suggesting that their results may be more indicative of processes occurring at

grain boundaries rather than at crystalline surfaces. Therefore in this study, we have chosen

to examine the reaction of atomic hydrogen with single crystal (3x3) reconstructed (0001)Si

6H-SiC surfaces.

The (3x3) reconstructed (0001)Si 6H-SiC surface is rapidly becoming a well

characterized semiconductor surface [66-73] and common starting point for most MBE

growth of SiC, AlN, and GaN on (0001)Si 6H-SiC substrates [73-75]. We and others have

146

shown this surface to consist of an incomplete bilayer of silicon overlaying the SiC surface

[66-68,76]. STM investigations have additionally shown this bilayer to be arranged in a

structure similar to the (7x7) Si (111) DAS model [70-72]. In our case, this surface was

prepared by chemical vapor cleaning processes and was not exposed to any sputtering

processes which are know to create a number of surface defects which can control surface

chemistry. Surface analytical techniques such as UV and x-ray photoelectron spectroscopy

(UPS and XPS), Auger electron spectroscopy (AES), low energy electron diffraction

(LEED), and temperature programmed desorption (TPD) were used to study not only the

effects of atomic hydrogen on the chemistry of SiC surfaces but on the electronic structure of

(3x3) 6H-SiC (0001)Si surfaces as well. The details of the effect on the electronic structure,

however, well be presented in a separate paper [77].

In this paper, we show that exposure of (3x3) 6H-SiC (0001)Si surfaces to atomic

hydrogen from a remote rf plasma source results in complete removal of all Si-Si bonded

silicon from the surface leaving a (1x1) surface. In turn, TPD showed weak H2 desorption in

the range of 400-600°C where desorption from silicon atoms would be expected based on

analogy to (111) Si surfaces [13-19]. A corresponding rise in H2 desorption was observed at

higher temperatures where H2 desorption from carbon sites would be expected based on

analogy to (111) diamond surfaces [32,34] and which in turn suggests C-H termination of the

SiC surface. This was supported by the observation of some C-C bonding after thermal

desorption of the rf plasma treated SiC surface at T > 1000°C. Similar to results previously

reported by other researchers, we also observed removal of silicon from single crystal (3x3)

6H-SiC (0001)Si SiC surfaces using smaller fluxes of atomic hydrogen generated from a hot

rhenium filament. In this case, complete removal of all Si-Si bonded silicon was not

147

observed, but TPD was able to detect H2 desorption in the range of 300-500°C indicative of

Si-H desorption. However, the hydrogen desorption in this temperature range was broad and

irregular resembling desorption from (2x1) B/Si (001) surfaces [78].

6.3. Experimental




having a base pressure of 9x10-10 Torr (see Refs. 76 and 79 for details of the transfer line,

and many of the associated systems). The experiments described in this paper employed the

SiC atomic layer epitaxy (ALE)/temperature programmed desorption (TPD), Auger electron

spectroscopy (AES), low energy electron diffraction (LEED), x-ray photoelectron

spectroscopy (XPS), and remote H plasma systems. A brief description of these systems is

provided below.

The SiC ALE system consisted of a UHV chamber with a base pressure of 3x10-10

Torr and was equipped with a residual gas analyzer (RGA) and a variety of gas dosers. For

TPD experiments, the RGA (a 0-200 amu quadrapole gas analyzer from Hiden Analytical

Ltd.) was housed in a separate differentially pumped cylindrical chamber (similar in design

to that of Smentkowski and Yates [80] ). The RGA chamber had a 0.5 cm diameter orifice at

the head of the RGA for TPD experiments and an approximately 50 cm2 "sunroof" which

could be opened to allow monitoring of residual gases in the system. The sample heating

stage for the TPD experiments consisted of a wound tungsten heating filament positioned

148

close to the back of the sample and mounted on a boron nitride disk [79]. A W/6%Re-

W/26%Re thermocouple was employed to measure the temperature of the backside of the

wafer. Heating profiles to 1100°C were easily obtained using a programmable

microprocessor and 20 amp SCR power supply. Actual surface/sample temperatures (i.e.

those reported herein) were measured using an infra-red pyrometer with a spectral response

of 0.8 to 1.1 µm and a emissivity setting of 0.5. The estimated experimental accuracy for the

latter temperatures was estimated to be ± 25°C..

Low fluxes of atomic hydrogen were generated in the ALE system using a hot

filament fabricated from 0.25 mm diameter rhenium wire. Temperatures >1700°C as

measured by the previously mentioned optical pyrometer were used to generate the atomic

hydrogen. The rhenium filament was positioned approximately 3 1/2" away from the SiC

wafer. No attempts were made to try and accurately measure the atomic H flux at the SiC

surface, and exposures were quoted in units of Langmuirs (10-6 Torr sec.). All hot filament

atomic H exposures were conducted without heating of the SiC by the sample heater (i.e.

room temperature). Any heating of the SiC surface by the hot filament is felt to be minimal

and at most could have raised the surface temperature by 100°C.

The XPS experiments were performed in a stainless steel UHV chamber (base

pressure = 2x10-10 Torr) equipped with a dual anode (Mg/Al) x-ray source and a 100 mm


were obtained using Al Kα radiation (hν = 1486.6 eV) at 12 kV and 20 mA emission



was achieved by periodically taking scans of the Au 4f7/2 and Cu 2p3/2 peaks from

149


peaks (83.98 and 932.67 eV, respectively) [81]. Curve fitting of most data was performed


with a linear background was found to best represent the data. The Auger electron

spectrometer and the low energy electron diffraction optics were mounted on a six way cross

off the transfer line and pumped through the transfer line. In the AES analysis, a 3 keV, 1mA

beam was used. Each Auger electron spectrum was collected in the undifferentiated mode

and numerically differentiated. In LEED an 80 eV, 1mA beam was used.

The plasma system consisted of an all metal seal stainless steel vacuum chamber

pumped by a 330 l/s turbomolecular pump. The base pressure of this system was 4x10-9

Torr and was limited by the double o-ring sealed quartz tube attached to the top of the system

where the rf discharge was produced. The process gases flowed through this quartz tube and

an inductively coupled plasma was generated using an rf power supply (13.56 MHz) and rf

matching network attached to a copper coil wrapped around the quartz tube. The sample was

located 40 cm down from the center of the rf coil. An inline Nanochem purifier and filter

was used for point of use purification of hydrogen. Sample heating in the plasma system

was conducted using a sample heater similar in design to the one previously described in the

ALE system. Depending on the chamber pressure and rf power, the plasma could be

maintained in the quartz tube or extended down toward the sample region. For the

experiments described in this study, an rf power of 20 W and a chamber pressure of 10-15

mTorr was used which confines the plasma to the quartz tube (i.e. a remote plasma). The

plasma system was also equipped with a differentially pumped 0-100 amu RGA which

allowed direct analysis of the purity of the process gases. RGA analysis of the hydrogen

150

(99.999% purity) used in these experiments after in situ purification revealed that the

impurity level of these gases were below the baseline of the system (<1ppm).

The n-type (ND=1018/cm3), 1" diameter, off axis (4° toward {11-20}) (0001)Si 6H-

SiC wafers used in this research were supplied by Cree Research with an ≈ 1 µm n-type

(ND=1017/cm3) 6H epilayer and a 1000Å thermally grown oxide. The thermal oxide was

removed by a 10 min. dip in 10:1 HF. The unpolished back side of each wafer was

subsequently coated via RF sputtering with tungsten to increase the heating efficiency of the

SiC, as the latter is partially transparent to the infrared radiation emitted from the tungsten

filament heater. After coating the backside of the SiC wafer with tungsten, the SiC wafers

were ultrasonically rinsed in trichloroethylene, acetone and methanol each for 5-10 min. and

then exposed to the vapor from a 10:1 buffered HF solution for 10 min. The wafers were

then mounted to a 1" diameter ring shaped Mo sample holder using Ta wire and inserted into

the transfer line load lock for experimentation. The SiC wafer was then loaded into the SiC

ALE system and annealed in 10-6 Torr SiH4 for 15 min. at 1050°C. This produced an

oxygen free (3x3) reconstructed surface [76]. For comparison purposes, Si (111) wafers

were also examined in this study. In this case, n-type (0.8-1.2 ? cm) Si (111) wafers were

cleaned by dipping in 10:1 HF to remove the thermal oxide and then annealed in the ALE

system at 950°C. This produced an oxygen free (7x7) reconstructed surface.

6.4. Results

6.4.1. Interaction of rf Atomic H with (3x3) vicinal (0001)Si 6H-SiC surfaces

151

Figure 6.1(a) shows an AES spectrum obtained from a (0001)Si 6H-SiC wafer after

annealing in SiH4 to produce the (3x3) reconstructed SiC surface. As illustrated, the

treatment removes oxygen from the SiC surface to levels below the detection limit of AES

and produces a silicon rich surface. Accounting for the 2:1 difference in sensitivity between

Si and C in AES [47], the Si/C peak to peak height (pph) ratio for the 3x3 surface shown here

was 1.35. Figure 6.2(a) shows an XPS spectrum of the Si 2p core level obtained from the

same surface. As displayed, two peaks at 99.5 and 101.5 eV were detected by XPS and

which were indicative of Si-Si and Si-C bonding respectively. This figure indicates the > 1

Si/C pph ratio in AES for the (3x3) 6H-SiC surface is primarily due to excess silicon

deposited on the surface. In previous studies, it has been shown that the (3x3) reconstruction

corresponds to an incomplete bilayer coverage of silicon on top of the SiC surface [67,68,76].

Figure 6.1(b) shows an AES spectrum obtained from the (3x3) 6H-SiC (0001)Si

wafer after a 1 min., 450°C remote rf plasma treatment (20W, 15 mTorr). As can be seen,

the intensity of the Si LVV transition is greatly reduced by the rf plasma exposure and the

corresponding Si/C pph ratio is reduced to 0.4. Correspondingly, a reduction/elimination of

the Si-Si bonding peak in the Si 2p XPS spectrum was also observed (see Figure 6.2(b)).

(Note: the shift in energy of the Si 2p anc C 1s core levels with H plasma processing is

related to band bending effects discussed in a seperate paper [77]). Additionally, the LEED

pattern from this surface was observed to switch from (3x3) to (1x1). The (1x1) pattern was

composed of broad dots suggestive of increased surface disorder. This etching phenomena

was observed to occur throughout the temperature range investigated (25-800°C). At this

point, we also note that the pph Si/C ratio obtained in AES here was very similar to the value

152

of 0.47 obtained by Allendorf et al [47] from polycrystalline 3C-SiC films exposed to

thermally generated atomic H.

As hydrogen can not be detected by either AES or XPS, TPD was used to confirm

that hydrogen adsorbed or desorbed from the SiC surface. Figure 6.3 shows a TPD spectrum

obtained from the 450°C rf plasma treated SiC surface. In our case, we were only able to

observe a weak desorption feature centered around 600°C. Beyond this feature, the intensity

of the H2 signal was observed to gradually increase to the endpoint of our temperature data

acquisition hardware which is 950°C. This gradual hydrogen desorption at higher

temperatures could be related to the hydrogen desorption observed from 450-950°C by

Allendorf et al [47]. However, it could also be due to desorption of subsurface hydrogen or

outgassing from our heater. Allendorf et al [47] have previously noted desorption/outgassing

of subsurface hydrogen at 1000°C from their polycrystalline 3C-SiC films due to residual

hydrogen trapped in the SiC during CVD growth of the SiC film. However in their case, the

outgassing feature was observed to be quite abrupt and intense whereas we observed a

gradual rise in the H2 signal. Additionally, we typically observe higher levels of outgassing

of subsurface hydrogen from rf plasma treated silicon wafers at much lower temperatures

(data not shown). This could however be related to the observation of Keroak and Terreault

[65] that implanted deuterium desorbs from silicon at relatively lower temperatures than SiC

(1200 vs. 2700°C).

To test outgassing of our heater as a source of hydrogen in TPD experiments on

plasma treated SiC, we acquired a TPD spectrum of the heater just prior to the SiC TPD

experiments (see Figure 6.4(a)). As can be seen, an essentially featureless spectra was

obtained with the detected H2 signal being two orders of magnitude lower than that from the

153

plasma treated SiC. To test outgassing of the moly sample holder, we additionally performed

TPD on a 1" diameter molybdenum plate exposed to the same plasma conditions as the SiC

wafer (see Figure 6.4(b)). In this case, we observed a large and broad desorption spectrum

centered at ≈ 550°C which is typical for surface and subsurface desorption.

100 200 300 400 500 600 700

dN(E

)/dE


(a)

(b)

SiC

Figure 6.1. AES of (3x3) reconstructed (0001)Si 6H-SiC (a) before remote H plasma, and (b) after remote H plasma (1 min., 20 W, 15 mTorr, and 450°C)

154

96 98 100 102 104 106

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

Si-Si

Si-C

(a)

(b)

Figure 6.2. XPS of the Si 2p core level from (3x3) reconstructed (0001)Si 6H-SiC (a) before remote H plasma, and (b) after remote H plasma (1 min., 20 W, 15 mTorr, and 450°C)

300 400 500 600 700 800 900

a.m

.u #

2 (a

rb. u

nits

)

Temperature (ÞC)

Figure 6.3. TPD of (1x1) 6H-SiC (0001)Si after remote H plasma exposure (1 min., 20 W, 15 mTorr, and 450°C), (ß = 1°C/sec.)

155

200 300 400 500 600 700 800 900

a.m

.u. #

2 (a

rb. u

nits

)

Temperature (ÞC)

(b)

(a)

Figure 6.4. TPD of (a) sample heating stage after outgassing, and (b) molybdenum plate after remote H plasma exposure (1 min., 20 W, 15 mTorr, and 450°C), (ß=1°C/sec.)

Finally, XPS of the C 1s core level before and after TPD revealed the presence of a

second C 1s peak after TPD (see Figure 6.5). However after TPD, the SiC surface still

displayed a (1x1) LEED pattern. Before TPD, the XPS showed only one C 1s peak. After

TPD, a second peak at higher binding energy appeared and which is indicative of some C-C

bonding at the surface. In the TPD experiments, the SiC wafer dwells at the maximum

temperature (≈1000°C) for less than a minute. In our experience, the time at this temperature

is not sufficiently long enough to result in the volatilization of a enough silicon to produce

this much C-C bonding at the surface alone. Therefore, the appearance of some C-C bonding

must be related to the loss of hydrogen from the SiC surface.

156

280 282 284 286 288

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

Figure 6.5. XPS of the C1s core level from (0001)Si 6H-SiC (a) before remote H plasma, (b) after remote H plasma (1 min., 20 W, 15 mTorr, and 450°C), and (c) after annealing at 1000°C.

6.4.2. Interaction of Thermal Atomic H with (3x3) vicinal (0001)Si 6H-SiC surfaces

In order to separate out subsurface hydrogen outgassing and other plasma related

induced effects, atomic H generated via cracking H2 over a hot rhenium filament was also

used as a source of atomic hydrogen. For comparison purposes, TPD spectra were first

acquired from (7x7) Si (111) surfaces exposed to atomic hydrogen generated by the hot

filament. In these experiments, it was observed that the surface switched to (1x1) after the

atomic H exposure. Figure 6.6 shows a TPD spectrum acquired from the hydrogen

terminated (1x1) Si (111) surface. As can be seen, a sharp desorption peak centered at ≈

475°C typical of monohydride (ß1) desorption from silicon was observed [13,15-18]. The

157

shoulder at lower energies was assigned to dihydride and trihydride (ß2,3) desorption from

silicon [13,15-18]. The ability to detect these features demonstrates the ability of the hot

filament to produce atomic hydrogen as well as validates our TPD apparatus.

200 300 400 500 600

a.m

.u #

2 (a

rb. u

nits

)

Temperature (ÞC)

Figure 6.6. TPD of Si (111) after room temperature exposure to 2000£ H2 with rhenium filament at > 1700°C (ß=1°C/sec.)

Figure 6.7 shows an AES spectrum of a (3x3) 6H-SiC (0001)Si surface before and

after exposure to atomic H (2000 Langmiur H2) from the hot rhenium filament. In this case,

the hot filament H exposure was observed to still maintain a Si/C ratio of > 1 but the ratio did

decrease from 1.35 to 0.9 and the LEED pattern was observed to switch from (3x3) to (1x1)

(however, in this case the LEED pattern displayed sharp dots). This observation was

supported by a similar reduction in the intensity of the Si-Si Si 2p peak in XPS. Unlike the 158

(111) Si surface however, a sharp desorption feature was not observed from the SiC surface

exposed to atomic H (see Figure 6.8) with only broad H2 desorption in the range of 350-

650°C being detected. In this case, the SiC TPD spectra more closely resembled the TPD

spectra of Kim at al [78] obtained from (2x1) B/Si (001) surfaces. Additionally, XPS did not

detect any C-C bonding at the SiC surface after TPD of the SiC surfaces treated with atomic

H from the hot filament.

30 130 230 330 430 530 630 730

(a)

(b)

dN(E

)/dE


Si

C

Figure 6.7. AES of (3x3) 6H-SiC (0001)Si (a) before atomic H exposure and (b) after room temperature exposure to 2000£ H2 with hot filament at > 1700°C.

159

100 200 300 400 500 600 700 800

a.m

.u. #

2 (a

rb. u

nits

)

Temperature (ÞC)

Figure 6.8. TPD of (0001)Si 6H-SiC after room temperature exposure to 2000£ H2 with rhenium filament at > 1700°C (ß=1°C/sec.)

As a final check, TPD spectra were also acquired from chemically vapor/silane

cleaned (3x3) SiC surfaces which had been cooled in the silane to ≈ 300°C (see Figure 6.9).

The samples were cooled in SiH4 to 300°C to hopefully maintain some hydrogen termination

of the SiC surface without preferentially losing silicon from the surface. In this case, more

pronounced hydrogen desorption at 475-525°C was observed and which was more

comparable to hydrogen desorption from the Si (111) TPD spectra. However, the intensity of

this hydrogen desorption feature was not nearly intense or sharp as that observed from the Si

(111) wafer. This could be related, though, to the inability of the silane flux to maintain a

hydrogen terminated surface at low temperatures.

160

300 400 500 600 700

a.m

.u. #

2 (a

rb. u

nits

)

Temperature (ÞC)

Figure 6.9. TPD of (3x3) 6H-SiC (0001)Si after cooling to 300°C in 10-6 Torr SiH4 (ß=1°C/sec.)

6.5. Discussion

In order to gain a better understanding of the observed etching of SiC surfaces

presented above, the authors feel that it is necessary to first separately consider the etching of

the two elemental components of SiC (i.e. silicon and diamond). Previous studies have

reported etching of silicon [14,24,28], diamond [32,34,37,39], and amorphous silicon carbon

alloys [82] by atomic hydrogen produced either via plasma excitation or thermal

decomposition. In the case of silicon, atomic hydrogen etching has been associated with the

formation of di and trihydrides (SiH2(a) and SiH3(a)) on the silicon surface. These species

have been shown to be the precursors for the final etch product, silane (SiH4(g)) [14,24].

However, SiH2(a) and SiH3(a) species were only observed to form at low temperatures (0-

161

200°C) where hydrogen desorption is negligible. At higher temperatures (>350°C) were

hydrogen desorption from di and trihydride sites is more appreciable, the silicon surface was

observed to be terminated by mostly monohydrides (SiH(a)) and the yield of etch products

such as SiH4(g) and Si2H2(g) decreased. Correspondingly, studies on H plasma cleaning of

silicon have observed etching/roughening of silicon surfaces at temperatures < 450°C and no

surface roughening at temperatures > 450°C [28].

In the case of diamond, etch products such as CHx and C2Hx have not been observed

from (001) or (111) single crystal diamond surfaces exposed to atomic H. TPD experiments

on single crystal (001) [32,34] and (111) [38] diamond surfaces have only shown

recombinative desorption of hydrogen at ≈ 1000°C. Etch precursors such as CHx or C2Hx at

≈ 425°C have only been observed via TPD from (001) and (111) oriented polycrystalline

CVD diamond films exposed to atomic H [39]. However, Kuttel et al [37] clearly have

observed etching of (001) and (111) single crystal diamond surfaces at temperatures up to

870°C in a microwave H plasma via a measured decrease in RMS surface roughness from 7

nm to 1 nm after plasma processing. The long etch time (17 hr.) required by Kuttel to

produce these results indicates that the formation of CHx species on diamond surfaces by

atomic H is perhaps slow or inefficient.

The simple observation that H plasma etching of diamond surfaces occurs at higher

temperatures than that observed in the case of silicon suggests that perhaps the atomic H etch

rates for silicon and carbon in SiC will be different. In fact, this has already been observed

for amorphous SiC films. Using XPS and glancing incidence XRD, Kalomiros et al [82]

have shown preferential etching of silicon from an a-SiC surface and the formation of a

polycrystalline hydrogenated carbon layer by a rf plasma process at 230°C. This behavior is

162

exactly what should be predicted based on analogy to the observed etching characteristics of

diamond and silicon. This behavior is also consistent with our observation of the selective

removal of excess silicon from (3x3) (0001)Si 6H-SiC surfaces by atomic H generated both

by plasma excitation and thermal decomposition.

A second observation to be made based on analogy to silicon and diamond is that

atomic H etching of surfaces generally occurs at low temperatures where hydrogen

desorption is low and a fully hydrogenated surface can be maintained. Therefore in order to

better understand the etching characteristics of SiC surfaces in atomic H, we have attempted

to estimate the hydrogen surface coverage of silicon carbide surfaces in an atomic H flux. To

do this, we have used a simple model in which we consider hydrogen adsorption and

desorption from silicon and carbon sites separately. The adsorption/desorption processes at

silicon and carbon sites are modeled using published data for the kinetics of these processes

on (001) silicon [23,31] and (001) diamond surfaces respectively [32,34]. Although the

experimental data presented in this study is for (111)/(0001) oriented surfaces, the choice to

use kinetic parameters from (001) orientations is primarily for sake of consistency as kinetic

parameters for hydrogen desorption from (111) diamond surfaces are not available. The

values for Edes and ν used in these calculations are presented in Table 6.1.

163

Table 6.1. Kinetic Parameters used to model hydrogen adsorption/desorption from Si and C sites on SiC surfaces. (Note: all data the data shown below are for first order processes).

Si C ß1 ß2,3 Hamza [32] Thomas [34] Edes (kcal/mol) 47 [23] 25 [31] 37 72.7

ν (#/sec) 8x1011 107 3x105 1013

For the carbon sites, hydrogen desorption was modeled using the kinetic parameters

determined for hydrogen desorption from (001) diamond surfaces by both Hamza et al [32]

and Thomas et al [34]. This is primarily a result of the extremely low value of 37 kcal/mol

for Edes reported by Hamza et al [32] which is lower than that reported for silicon. This

value for Edes is extremely surprising given the simple fact that hydrogen desorption from

diamond surfaces occurs at temperatures 500°C higher than on silicon. Therefore, the data of

Thomas et al [34] was used as well since in this case Edes = 72.7 kcal/mol which is higher

than most values reported for H2 Edes from silicon surfaces. Additionally, Allendorf et al

[47] have also reported Edes = 72 kcal/mol and ν = 1013/sec for hydrogen desorption from

polycrystalline SiC surfaces.

To estimate the SiC hydrogen surface coverage, we employ the method of Schulberg

et al [94] in which we assume a steady state equilibrium between the incoming flux of atomic

H and surface desorption of H2. Desorption kinetics are typically described by the general

Polyani-Wigner rate expression [31]:

desorption rate = -dØ/dt = νn exp(-Edes/RT) (1)

164

where: n = the reaction order Ø = the adsorbate surface coverage ν = the pre-exponential factor, νo = 1028/cm2 sec,

ν1 = 1013/sec, ν2 = 10-2 cm2/sec Edes = the activation energy for desorption

In steady conditions, the flux of adsorbates leaving the surface via desorption will be

equal to the incoming flux of adsorbate times the adsorbate sticking coefficient. The sticking

coefficient is described by:

S = So(1 - Ø/Ømax)n (2) where S = Sticking Probability So = Initial Sticking Probability (i.e. S at Ø = 0) Ømax = Maximum Surface Coverage

The combination of equations 1 and 2 allows the determination of the steady state surface

coverage of the adsorbate [94]. Based on this model and the desorption kinetic data tabulated

above, we have therefore estimated the surface coverage of hydrogen on silicon and carbon

sites both as a function of temperature and flux. The results are presented in Figures 6.10-

6.13 and are based on the assumption of a unity initial sticking coefficient (i.e. So = 1)

Figures 6.10 and 6.11 show the estimated monohydride and di/trihydride surface

coverages for silicon sites as a function of temperature and flux (ML/sec). Figure 6.10

indicates that in fluxes typical of low pressure hydrogen plasma processes, all silicon sites

should be saturated with hydrogen up to temperatures of ≈ 800°C. However, Figure 6.11

indicates that the concentration of di/trihydrides should start to decrease around 500-600°C.

165

Although the temperature range is slightly higher, this is clearly consistent with the

previously reported etching behavior of silicon in a remote H plasma. Accordingly, etching

of silicon in SiC should stop or decrease at temperatures above ≈ 600°C. However, as shown

in Figures 6.12 and 6.13, carbon sites on SiC surfaces should remain saturated with hydrogen

up to temperatures of 1000-1100°C indicating that etching of carbon in SiC could continue

up to temperatures these very same temperatures. We also note the significant difference in

predicted concentration of occupied carbon sites based on the data of Hamza et al [32] and

Thomas et al [34]. Based on the data of Hamza et al [32], carbon sites would be expected to

be saturated with hydrogen up to 1200°C in a 100 ML/sec H flux. However, the data

Thomas et al would predict that at least of half of these sites would be empty under the same

conditions.

0

0.2

0.4

0.6

0.8

1

200 400 600 800 1000 1200 1400 1600

Ø/Ø

max

Temperature (ÞC)

1 100 10 4

10 6

Figure 6.10. Mono-hydride surface coverage on silicon sites of SiC.

166

0

0.2

0.4

0.6

0.8

1

200 400 600 800 1000 1200 1400 1600

Ø/Ø

max

Temperature (ÞC)

1100

10 4

Figure 6.11. Di-hydride surface coverage on silicon sites of SiC.

0

0.2

0.4

0.6

0.8

1

200 400 600 800 1000 1200 1400 1600

Ø/Ø

max

Temperature (ÞC)

1

10

100

Figure 6.12. Hydrogen surface coverage on carbon sites of SiC based on kinetic data of Hamza [32].

167

0

0.2

0.4

0.6

0.8

1

200 400 600 800 1000 1200 1400 1600

Ø/Ø

max

Temperature (ÞC)

1

100

10 4

10 6

Figure 6.13. Hydrogen surface coverage on carbon sites of SiC based on kinetic data of Thomas et al [34].

Based on analogy to silicon and Figures 6.10-6.13, one would expect etching of SiC

in atomic H at temperatures < 1000°C and no etching at temperatures ≈ > 1000°C. In reality,

SiC is probably etched by atomic H at all temperatures. However, the etch rate probably

exhibits some temperature dependence. At low temperatures (i.e. RT-500°C), the etch rate is

probably low due to limited thermal activation and perhaps surface mobility. As the

temperature is increased into the range of 500-1000°C, the etch rate should decrease due to

significant desorption of hydrogen through silicon sites and the reduction of etch precursors

such as SiH3(a). However, the etch rate in this temperature range will probably still be

measurable due to limited desorption from carbon sites and hence the ability to form CHx

etch products. At temperatures > 1000°C, the etch rate should decrease due to increased

desorption of hydrogen from both carbon and silicon sites but may eventually increases due

to increased thermal activation and possible volatilization of silicon at temperatures >1500°C.

168

This type of etch rate dependence has been partially observed by Kim and Olander [48] in

their modulated molecular beam mass spectrometry studies. In these studies, they observed

SiH4, CH4, and C2H2 etch products from polycrystalline 3C-SiC films during atomic H

exposure over the temperature range of RT-800°C. However, the yield of SiH4 from SiC

was observed to initially increase with temperature and then start to decrease at ≈ 500°C.

The yield of CH4, however, was observed to gradually increase over the entire temperature

range investigated (0-800°C). Though Kim and Olander [48] explain the decrease in SiH4

production to depletion of silicon from the surface, their results are also clearly consistent

with our explanation based on enhanced hydrogen desorption from silicon sites. However,

our inability to avoid selective removal of silicon from the (3x3) (0001)Si 6H-SiC surface in

the rf H plasma over the temperature range of 0-800°C is also consistent with the results of

Kim and Olander [48].

In the higher temperature range (1000-1700°C), it has been previously noted that the

etch rate of SiC in molecular hydrogen (i.e. H2) increases with increasing temperature. In

this case, we feel that the actual etching of SiC is due to atomic H produced by thermal

decomposition of molecular H2. Figure 6.14 shows the predicted percent dissociation of

molecular hydrogen into atomic hydrogen based on thermodynamic calculations (note figure

produced by HSC program). As can be seen, significant production of atomic H is only

predicted to occur at temperatures of 1500-1700°C. So these higher temperatures for

molecular H2 etching are at least to some extent probably necessary to produce a significant

concentration of atomic hydrogen. The higher temperatures probably also assist in the

volatilization of silicon from the surface as well as increasing the surface mobility.

169

0.0

5.0

10

15

1000 1500 2000 2500 3000

% H

2 Dis

soci

atio

n

Temperature (ÞC)

Figure 6.14. Percent dissociation of H2 into H as function of temperature.

At this point, it is worth considering some of the potential errors in this simplified

model. First for this model to be valid, diffusion of hydrogen from carbon to silicon sites

(and vice versa) must be minimal. For Si-Ge alloys this subject has sparked much debate. In

this system, a lowering of Tmax for monohydride (ß1) desorption from silicon was observed

with the addition of germanium to the surface [83,84]. This lowering of Tmax has been

argued to be due to weakening of the Si-H bond due to electronic matrix effects [83]. Others,

however, have argued that the lowering of ß1 Tmax can be described by considering

hydrogen diffusion from silicon sites to germanium sites where the activation energy for

desorption is lower (Edes Si-H = 55 kcal/mol vs. Ge-H = 35 kcal/mol) [84,85]. This is

reasonable for silicon surfaces where the activation energies for hydrogen diffusion on (100)

and (111) surfaces have been determined to be ≈ 30-41 kcal/mol [86,87] and 35 kcal/mole

170

[88] respectively. Unfortunately for the case of SiC, activation energies for hydrogen surface

diffusion have not been determined. However, activation energies for surface diffusion of Si

and C atoms on SiC surfaces have been reported. Based on surface diffusion length

measurements in CVD, Kimoto et al [89,90] determined an activation energy of 82 kcal/mol

for surface diffusion of Si or C atoms on (0001)Si 6H-SiC surfaces. These measured values

are in agreement with the calculations of Takai et al [91], which found activation barriers of

106-126 kcal/mol for carbon diffusion on (111) and (-111) 3C-SiC surfaces. As much lower

activation energies for self diffusion of Si on Si (111) [92,93] and (001) [93] surfaces have

been determined (18-35 and 6-25 kcal/mole respectively) it seems reasonable to expect the

mobility of H atoms on SiC surfaces to be lower than on Si surfaces. Accordingly, H surface

diffusion should be negligible in the temperature range studied here and should not affect our

model.

Another possible source of error in our model is that Tmax for hydrogen desorption

from Si and C sites on SiC could be significantly different from that observed from silicon

and diamond respectively. As previously mentioned, it has been suggested that the addition

of Ge to Si lowers Tmax for ß1 H2 desorption from Si sites due to a weakening of the Si-H

bond by germanium [83]. It has additionally been observed that boron doping (1019/cm3) of

Si also lowers Tmax ß1 whereas dosing Si (001) surfaces with diborane (B2H6) produces

two ß1 desorption features [78]. These effects have been explained based on varying

electronic effects in which the Si-H bond is weakened by the more electronegative dopant or

due to changes in the work function by the dopant or substitutional atom. Clearly, the

addition of carbon to the silicon lattice (and vice versa) could have similar effects. The

addition of carbon changes both the work function (and band gap) as well as creating a more

171

polar bond with silicon to weaken any Si-H bonds at the surface. This line of reasoning

could explain our inability to observe sharp desorption features from atomic H treated (3x3)

(0001)Si 6H-SiC surfaces.

In the case of (3x3) SiC surfaces exposed and cooled in SiH4 to 300°C, H2

desorption at 475°C was observed and which was consistent with the observed H2 desorption

from Si (111) surfaces. This can be explained by the fact that for this surface all of the

silicon atoms terminating the SiC surface have Si-Si backbonds as the (3x3) surface consists

of a bilayer of silicon atoms. However, in the case of the hot filament atomic H treated (3x3)

SiC surfaces, a broad range of H2 desorption in the temperature range of 200-600°C was

observed and which was similar in appearance to that observed from B2H6 treated Si (001)

surfaces [78]. In this case, some of the Si bilayer has been removed via etching by atomic H

and hence desorption occurs from Si atoms with Si-Si and Si-C backbonds (i.e. C-Si-H or Si-

Si-H). For the silicon atoms with carbon backbonds, the Si-H bond is weakened as the

underlying carbon atom is more electronegative and withdraws charge which the silicon atom

would share with hydrogen. Accordingly, the weaker Si-H bonds translates into a lower

Edes. This range of Edes is what gives Figure 6.8 its broad nature. However, we note that in

the studies of Ascherl et al [95] dosing silicon surfaces with trimethylsilane (SiH(CH3)3) did

not result in any change in Tmax for ß1.

In the case of the H plasma treated (0001)Si 6H-SiC surfaces, TPD showed a weak

H2 desorption feature at ≈ 625°C with a gradual increase in H2 signal at higher temperatures.

This feature is intermediate to what would be expected from carbon and silicon sites based on

analogy to silicon or diamond surfaces. This feature could be related to H2 desorption from

172

CHx sites, but is higher in temperature than previous reports of CHx desorption from

polycrystalline CVD diamond films [39]. This TPD spectrum, however, could be greatly

effected surface roughening induced by the atomic H etching of the SiC surface. This could

produce greatly affect Edes.

Setting aside our simple desorption model, etching of SiC by atomic hydrogen can

also be predicted based on simple comparison of the Si-C, Si-H, and C-H bond energies

which are 76.5, 70.8, and 98.8 kcal/mol respectively [97]. As can be seen, the C-H bond is

actually stronger than the Si-C bond and energy can be gained by breaking a Si-C bond and

forming a C-H bond. Although, the reaction between molecular hydrogen (H2) and SiC to

form Si-H and C-H bonds is not energetically favorable, i.e.

Si-C + H-H <=> Si-H + C-H (i) 76.5+104 <===> 70.4 + 98.8 179.7 <====> 169.2 kcal/mol

The reaction between pre dissociated atomic H and SiC is, i.e.

Si-C + 2H <=====> Si-H + C-H (ii)

76.5 <====> 70.4 + 98.8

Thus for the (0001)Si 6H-SiC surface, it is more energetically favorable for atomic

hydrogen to insert itself into a Si-C bond forming C-H bonds rather than simply terminating a

Si dangling bond and forming a relatively weak Si-H bond. Thus a C-H terminated SiC

surface would be expected after atomic hydrogen exposure. This is in agreement with the < 1

Si/C pph ratio we and other have observed from SiC surfaces exposed to atomic H. This also

agrees with our observation of the formation of some C-C bonding in XPS after TPD at

1000°C. At temperatures of 1000°C it would be expected that any hydrogen adsorbed on

173

carbon would desorb leaving behind some C-C bonding. Additionally, we have

previously noted that in HF wet chemical processing H termination of silicon atoms at

(0001)Si 6H-SiC surfaces is highly unstable due to the polarity of the underlying Si-C bonds

[96]. OH- termination is instead favored due to the ability of the Si-OH bond to cancel the

dipole produced by the Si-C bond below. For these same reasons, Si-H termination of

(0001)Si 6H-SiC surfaces in vacuum are equally unstable and hence there is an additional

driving force for the H atom to insert itself into the Si-C bond. It is also worth mentioning

that in the (3x3) structure there are actually a significant number of Si-Si bonds. We note

that in this case, these bonds are compressed and distorted to values far from the their

equilibrium value due to the differences in the lattice constants between Si and SiC (≈ 20 %

[10,70-72]). As such, these bonds are probably more reactive with atomic hydrogen and it

may be more energetically favorable to form SiHx species rather than to remain in the

distorted (3x3) structure.

Finally, the above results and discussion indicate that in H plasma cleaning of SiC,

silicon should be added to the plasma chemistry in order to compensate for the selective

removal of silicon from the SiC surface due to etching. The addition of silicon to the plasma

(via SiH4 or Si2H6) should also assist in the reduction of silicon oxides which are

particularly difficult to remove in H plasmas. Lin et al [51] have previously investigated

cleaning of HF dipped SiC surfaces using a 1:1 H2:He mixture in an ECR plasma source

(650°C, 90 min., 5x10-4 Torr). In their case, they reported the removal of C-C, C-O, and C-

F species from the SiC surface but were not successful in completely removing silicon oxides

(Si-O) from the surface. The addition of silicon to the plasma chemistry therefore could

assist in the removal of silicon oxides through chemical reduction and formation of more

174

volatile sub oxides. Initial investigations in our lab using H2/1%SiH4 mixtures in plasma

cleaning of HF dipped SiC surfaces have shown an enhanced removal of silicon oxides.

However, deposition of silicon was a particular problem indicating that such a cleaning

process requires delicately balancing the silicon deposition rate with its etching rate. This

will require tight process control and perhaps only a very narrow processing window will be

available.

6.6. Conclusion

In conclusion, we have shown that atomic hydrogen exposure selectively removes

silicon from (3x3) 6H-SiC (0001)Si surfaces. Atomic hydrogen exposures reduces and

removes the Si-Si bonding Si 2p XPS peak and converts the (3x3) LEED pattern to (1x1).

Additional etching of the SiC surface was indicated by the reduction in the Si LVV/C KLL

ratio in AES from 1.3 to 0.4 following exposure of (3x3) surfaces to a remote rf H plasma.

TPD of atomic H treated (3x3) SiC surfaces showed weak hydrogen desorption in the range

of 400-600°C where desorption from silicon atoms would be expected by analogy to (111)

Si. However, the hydrogen desorption signal increased at higher temperatures where

hydrogen desorption from carbon sites would be expected based on analogy to (111)

diamond surfaces. C-H termination of the SiC surface was supported by the observation of

some C-C bonding after thermal desorption of rf plasma treated SiC surface at T > 1000°C.

Based on these observations, we conclude that atomic H processing of SiC surfaces

selectively removes silicon from the surface and favors C-H termination.

175

6.7 Acknowledgments

The authors would like to thank Cree Research, Inc. for supplying the 6H-SiC wafers.

The research was supported by the Office of Naval Research under contract and through the

Department of Education through an Electronic Materials/GAANN Fellowship.

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7. EX SITU AND IN SITU SURFACE CLEANING PROCESSES FOR

(0001) AlN, GaN, AND AlxGa1-xN SURFACES


to the


by

Sean W. King, Laura L. Smith, John P. Barnak, *James A. Christman,

Michael D. Bremser, *Robert J. Nemanich, and Robert F. Davis




Raleigh, NC 27695

182

7.1 Abstract

Exposure to numerous acids and bases and UV/O3 oxidation were used to determine

the best ex situ cleaning techniques for the (0001) surfaces of AlN and GaN. HF and HCl

were found to produce surfaces with the lowest coverage of oxygen after wet chemical

cleaning of AlN and GaN, respectively. However, AES and XPS revealed the surfaces to be

terminated with F and Cl which inhibited re-oxidation prior to loading into vacuum. TPD

showed that temperatures of 650 and 850°C were necessary to thermally desorb the Cl and F,

respectively. UV/O3 oxidation was observed to reduce the surface carbon coverage for both

AlN and GaN though incompletely. This was related to the relative inert chemical nature and

resistance to oxidation of GaN and AlN surfaces. In situ remote hydrogen plasma exposure

at 450°C removed halogens and hydrocarbons remaining after ex situ cleaning of both AlN

and GaN surfaces; however, oxide free surfaces were not be achieved. Complete thermal

desorption of the surface oxide from (0001) GaN in UHV was only achieved at temperatures

> 800°C where some GaN decomposition occurred. Annealing GaN in NH3 at 800°C

reduced the surface oxide without loss of surface stoichiometry.

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7.2. Introduction

GaN and AlN are completely miscible semiconductors with wide direct band gaps of

3.40 and 6.2 eV, respectively [1-3]. Combined with the 1.9 eV direct band gap of InN [1],

the III-V nitride materials system is accordingly of considerable interest for many visible-UV

optoelectronics device applications [4]. The recent demonstration of a blue laser diode based

on a InGaN quantum well (QW) structure [5] highlights many of the recent advances which

have been made in this field. However, GaN and AlN are also of interest for high power,

high frequency, and high temperature device applications due to the other extreme materials

properties that they exhibit including: high temperature stability (Tmelt AlN ≅ 2850°C), high

saturation electron drift velocity (vsat GaN = 2.7x107 cm/sec), high thermal conductivity (κ

AlN = 3.0 W/cm K), and high breakdown voltage (GaN = 2.0 MV/cm) [2-3]. The recent

observation of a negative electron affinity for AlN [6] and AlxGa1-xN [7] alloys also makes

these materials candidates for field emitters in cold cathode electron devices. However, the

advancement of GaN and AlN toward use in these applications will demand the continued

development of new and improved processes to fabricate these materials and devices. As

surface cleaning processes are the foundations on which most semiconductor fabrication

steps are built [1-3], it is likely that improvements in GaN and AlN processing can be further

assisted by optimization of surface cleaning processes for GaN and AlN.

Experience gained in silicon and gallium arsenide technology has shown that the

criteria for surface cleanliness must include removal of not only native oxides and organic 184

contaminants but also metallic impurities, particulate contaminants, adsorbed molecules, and

residual species left by previous processes [8-10]. The deleterious effects of incomplete

removal of all these various contaminants are numerous. In Si homoepitaxy, for example,

improper removal of surface oxides and organic contaminants have been shown to result in

an increased density of line and planar defects in epitaxial films from <104/cm2 to >

1010/cm2 [11-16]. In fact, recent studies on the heteroepitaxial growth of SixGe1-x alloys on

Si (100) have shown that surface defects produced in the Si substrate by residual

organic/carbon contamination act as the preferred sites for misfit dislocation generation [17].

Most importantly however, these increased epitaxial defect densities have been shown to

decrease device yield and performance [18-19]. Surface cleaning is also important in other

semiconductor processes including metal contact formation and gate dielectric formation. In

the case of metal contact formation, improper removal of surface oxides prior to metal

deposition has been demonstrated to result in higher contact resistances and lack of contact

parameter uniformity [20-26]. Control of the metal contact Schottky barrier height has also

been shown to be dependent on surface preparation [27-33]. In this case, control of the

Schottky barrier height is dependent on controlling the existence and density of

surface/interface states which determine the surface Fermi level position. Surface and

interface states can frequently be related to surface dangling bonds which can be passivated

by various surface cleaning processes. Control of interface states is also of importance in

MISFET devices where interface states can produce changes in threshold voltage and must

be controlled to ensure uniformity in device operation. Drift in MIS device parameters has

also been linked to the presence of alkali ions at the interface which can, in part, be

controlled by surface cleaning prior to insulator deposition [34].

185

Due to the above observations, numerous studies have been conducted on Si [8-10]

and GaAs [35-71] which have been concerned with obtaining clean, structurally well ordered

surfaces. For GaAs, these studies have covered a wide range of ex situ and in situ processes

including wet chemical [35-53], UV/O3 oxidation [54-62], thermal desorption 63,64],

chemical beam [65], and atomic H cleaning [66-71]. In the case of AlN and GaN, however,

there have been comparatively fewer investigations of processes to obtain clean surfaces of

these materials [72-88]. In fact to the authors knowledge, there have been no such studies for

AlN. Most studies of the surface properties of AlN in vacuum have relied either on ion

bombardment/sputtering [72-74] or in situ preparation of polycrystalline AlN films on Al via

nitriding with hydrazine [75] or N2+ ions [76,77].

For GaN, perhaps the first surface cleaning study was that of Hedmann and

Martensson [78] who used XPS to examine single crystal HVPE GaN/Al2O3 (0001) films

etched in H3PO4 at 100°C and then annealed in situ at 300°C. In this case, they reported that

the in situ anneal removed some oxygen and carbon contaminants but were unsuccessful in

completely removing these contaminants at this temperature. Since the early reports by

Hedmann and Martensson, there have only recently been more investigations of surface

cleaning processes for GaN. These stuides have been conducted on high quality MOCVD

and OMVPE GaN films grown on (0001) Al2O3 and 6H-SiC substrates [79-88].

The first such investigation by Khan et al [79], examined in situ cleaning of MOCVD

GaN (0001) films in MBE via annealing in an evaporated flux of Ga (2-5x1015/cm2 min) at

600 or 900°C. In situ Auger electron spectroscopy (AES) analysis of GaN films cleaned in

this fashion showed the CKLL peak intensity to be 2% of the NKLL signal while the O

contamination was close to the AES sensitivity limit. However, low energy electron

diffraction (LEED) of these surfaces displayed only unreconstructed (1x1) diffraction

186

patterns. This technique has since been used by Bermudez et al to prepare clean GaN

surfaces to study the interfaces and interaction of GaN with Ni [80], Al [81], and O2 [82]. In

the latter study, Bermudez concluded that surfaces prepared via annealing in a Ga flux show

upward bending of 0.9 eV due to surface Fermi level pinning by surface states of unidentified

character. These states were removed by re-exposure of the clean GaN surfaces to molecular

oxygen (O2) which decreased the band bending by ˜ 0.15 eV. In contrast, Bermudez also

found that surfaces prepared by only wet chemical cleaning in 1:10 NH4OH:H2O showed

only 0.4±0.2 eV upward band bending. In addition, Bermudez also investigated N2+

sputtering for cleaning GaN surfaces in situ and found the technique to be equivalent to

annealing in a beam of Ga [82]. Nitrogen ion sputtering has also been used by Sung et al

[83] to examine the polarity of GaN films grown on (0001) Al2O3. Their time-of-flight

scattering and recoiling spectrometry (TOF-SARS) and classical ion trajectory simulations

indicate that (0001) GaN surfaces prepared in this fashion are nitrogen terminated with Ga

atoms comprising the second layer. These findings, however, are contrary to those of Daudin

et al [84] who using ion channeling and convergent beam electron diffraction techniques

concluded that smooth MOCVD GaN films grown on (0001) Al2O3 are Ga terminated

whereas rough/pyramidal GaN films are N terminated.

In a previous study [85], we have investigated cleaning of GaN surfaces using HF

and HCl wet chemical processes followed by in situ thermal desorption. In this study, it was

found that HCl:DI wet chemical processes produced the lowest coverages of oxygen and

carbon contaminants but that HF wet chemistries resulted in GaN surfaces for which thermal

desorption of carbon contaminants in situ was more efficient [85]. More recent studies by

others have investigated cleaning of GaN using wet chemical treatments based on warm

187

NH4OH [86] and HF solutions buffered with (H2FNH4) [87]. In situ cleaning of GaN

surfaces using a H2:He plasma has also been demonstrated by Hughes et al [88].

In this study, we have chosen to investigate both ex situ and in situ cleaning of AlN

and GaN surfaces using numerous surface analytical techniques including Auger electron

spectroscopy (AES), x-ray photoelectron spectroscopy (XPS), ultra-violet photoelectron

spectroscopy (UPS), low energy electron diffraction (LEED), and temperature programmed

desorption (TPD). A broad range of chemistries and techniques for cleaning both GaN and

AlN surfaces was selected in order to better facilitate a direct comparison of the various

different cleaning processes available. In this study, we have chosen to investigate UV/O3

oxidation for ex situ carbon contamination removal and a variety of standard wet chemistries

and ex situ chemical vapor exposures for oxide removal. Wet chemistries based on H2SO4

and H3PO4 solutions common in GaAs technology for chemical oxide growth were also

investigated [41-44]. The in situ cleaning processes examined included thermal desorption,

exposure to hydrogen plasmas, and annealing in fluxes of Al, Ga, NH3 and SiH4.

Additionally, the use of GaN and In as a passivating/protective layer for AlN surfaces was

additionally investigated.

7.3. Experimental

7.3.1. Integrated Surface Preparation and Analysis System.




188

having a base pressure of 9x10-10 Torr (additional details of the transfer line, and many of

the associated systems are provided in Refs. 89,90). The experiments described in this paper

employed the III-V nitride gas source molecular beam epitaxy (GSMBE), Auger electron

spectroscopy (AES), low energy electron diffraction (LEED), x-ray/UV photoelectron

spectroscopy (XPS/UPS), and remote H2 plasma CVD systems. A brief description of these

systems is provided below.

The III-N GSMBE system consisted of a UHV chamber with a base pressure of

3x10-10 Torr and was equipped with a residual gas analyzer (RGA) and a variety of gas

dosers, and Knudsen cells. The RGA (a 0-200 amu quadrapole gas analyzer from Hiden

Analytical Ltd.) was housed in a separate differentially pumped cylindrical chamber (similar

in design to that of Smentkowski and Yates [91]) which had a 0.5 cm diameter orifice at the

head of the RGA for TPD experiments and an approximately 50 cm2 "sunroof" for

monitoring residual gases in the system. The sample heating stage in the ALE system

consisted of a wound tungsten heating filament positioned close to the back of the sample

and mounted on a boron nitride disk [89]. A W/6%Re-W/26%Re thermocouple was

employed to measure the temperature of the backside of the wafer. Surface temperatures and

heating profiles to 1100°C were easily achieved using a programmable microprocessor and

20 amp SCR power supply. Actual surface/sample temperatures (i.e. those reported herein)

were recorded using an infra-red pyrometer with a spectral response of 0.8 to 1.1 µm and a

emissivity setting of 0.5. The estimated experimental accuracy for the latter temperatures

was estimated to be ± 25°C.

Source materials in the GSMBE included Al, Ga, NH3, and SiH4. Al (99.9999%)

was evaporated from a 25 cc "cold lip" Knudsen cell and Ga (99.99999%) was evaporated

189

from a 25 cc dual filament Knudsen cell. The NH3 (99.9995%) was further purified via an in

line purifier connected directly to a leak valve mounted on the GSMBE chamber. The SiH4

(99.995%) was used as received with out any additional purification. Sample exposure to the

NH3 and SiH4 was obtained using "molecular beam" dosers similar to the design of Bozack

et al [92]. Collimation of the ammonia and silane into a molecular beam focused onto the

sample was achieved with this doser using a 13 mm diameter x 2 mm thick glass capillary

array with a ten micrometer pore size (Galileo Electro Optics Inc.). Though, the doser to

sample distance was fixed at 2", no attempts were made to accurately measure the flux of

NH3 or SiH4 and hence all exposures are quoted as Langmuirs (£ = 10-6 Torr sec).

The XPS and UPS experiments were performed in a stainless steel UHV chamber

(base pressure = 2x10-10 Torr) equipped with a dual anode (Mg/Al) x-ray source, He I UV

lamp, and a 100 mm hemispherical electron energy analyzer (VG CLAM II). All XPS

spectra of AlN reported herein were obtained using Al Kα radiation (hν = 1486.6 eV) at 12

kV and 20 mA emission current. For GaN and AlGaN all XPS spectra were acquired using

Mg Kα radiation (hν = 1253.6 eV) at 10 kV and 20 mA emission. XPS analysis typically

required less than 1 hour during which time the pressure never increased above 9x10-10 Torr.

Calibration of the binding energy scale for all scans was achieved by periodically taking

scans of the Au 4f7/2 and Cu 2p3/2 peaks from standards and correcting for the discrepancies

in the measured and known values of these two peaks (83.98 and 932.67 eV, respectively

[74]). Significant sample charging, however, was observed during XPS of bulk

polycrystalline AlN wafers. To correct for these charging effects, the C 1s peak from these

AlN surfaces was assigned a value of 285.7 eV and all the other core levels (O1s, Al 2p, N

1s, F 1s) shifted accordingly. The value of 285.7 eV for the C 1s core level was based on the

190

observation that adventitious carbon on thin AlN surfaces (˜30Å) was observed to occur at

this energy. Curve fitting of most data was performed using the software package GRAMS

386. A combination Gaussian-Lorentzian curve shape with a linear background was found to

best represent the data.. The Auger electron spectrometer and the low energy electron

diffraction optics were mounted on a six way cross off the transfer line and pumped through

the transfer line. In the AES analysis, a 3 keV, 1 mA beam was used. Each Auger electron

spectrum was collected in the undifferentiated mode and numerically differentiated. In

LEED an 80 eV, 1 mA beam was used.

The remote plasma CVD system consisted of a metal seal stainless steel vacuum

chamber pumped by a 330 l/s turbomolecular pump. The base pressure of this system was

4x10-9 Torr. The process gases flowed through a quartz tube mounted at the top of the

chamber, and the plasma is excited by rf (13.56 MHz) applied through a copper coil wrapped

around the quartz tube. The sample was located 40 cm below the center of the rf coil. An in

line Nanochem purifier and filter was used for point of use purification of hydrogen and

silane. Sample heating in the plasma system was conducted using a sample heater similar in

design to the one previously described in the ALE system. The plasma system was also

equipped with a differentially pumped 0-100 amu RGA which allowed direct analysis of the

purity of the process gases. RGA analysis of the hydrogen and silane (both 99.999% purity)

used in these experiments after in situ purification revealed that the impurity level of these

gases were below the baseline of the system (<1 ppm)

7.3.2. Samples and Ex Situ Preparation:

191

The surfaces of a variety of AlN and GaN samples were investigated. The AlN

samples were derived from films (1) epitaxially grown on 6H-SiC (0001)Si by

organometallic vapor phase epitaxy (OMVPE) [93] or gas source-molecular beam epitaxy

(GSMBE) [94], (2) deposited via reactive ion sputtering on Si (111) and (3) hot pressed

polycrystalline AlN wafers. The GaN surfaces investigated were those of GaN films

epitaxially deposited on epitaxial AlN buffer layers grown on 6H-SiC (0001) by OMVPE

[93] and GSMBE [94]. For GaN and AlN films on SiC substrates, the back sides of the

substrates were RF sputter coated with tungsten. This was done to improve the heating

efficiency of these samples as SiC, GaN, and AlN are all partially transparent to the infrared

radiation emitted by our tungsten filament heaters.

The experimental system employed for the ex situ UV/O3 exposures described in this

study employed a high intensity Hg lamp positioned in close proximity (˜ 1 cm) too the AlN

and GaN samples. In order to increase the concentration of generated O3 (i.e. to increase the

oxidation rate), the UV/O3 box was purged with 1 L/sec O2 during some UV exposures.

Further details of this process have been previously described [54-62]. CMOS grade acids

and bases and high resistivity (18.4 M?) de-ionized water were used in all the ex situ wet

chemical cleaning processes examined. The wet chemical cleans investigated included

various mixtures of the following acids and bases: HCl, HF, NH4F, HNO3, H2SO4, H3PO4,

H2O2, NH4OH, NaOH, KOH, acetic acid, and lactic acid. The selection of these chemicals

was based on their extensive usage in standard microelectronic processes. Unless otherwise

noted, all AlN and GaN samples were rinsed in DI water and blown dry with N2 after any wet

chemical processes. After each ex situ cleaning process, the AlN or GaN sample was

mounted to a molybdenum sample holder and loaded into the transfer line load lock for

subsequent surface analysis or in situ cleaning 192

7.4. Results

7.4.1. Ex situ cleaning of AlN

Figure 7.1(a) shows an AES spectrum of the surface of an as-received OMVPE AlN

sample on which a thin oxide layer has formed during ambient exposure. Additional

oxidation via a UV/O3 treatment was investigated initially as a possible ex situ method for C

removal from the nitride surfaces. Both AES and XPS were used to examine an OMVPE

AlN film which had been previously cleaned in trichloroethylene, acetone, and methanol for

5 min. in each solvent and then exposed to UV/O3 for 10 min. at room temperature. As

shown in Figure 7.1(b), the combination of solvent cleaning and UV/O3 exposure reduced

the intensity of the CKLL peak due to surface C by ˜ 50%. A similar decrease in the intensity

of the C 1s core level was also observed in XPS.

Longer UV/O3 exposures (30 min. - 1 hr.) with or without a solvent preclean did not

further appreciably decrease the surface C coverage. However, Figure 7.1(b) does show that

the AlN surface was further oxidized by the UV/O3 treatment. This was confirmed by XPS

analysis of the O 1s core level from bulk polycrystalline AlN wafers (see Figure 7.2). XPS

of the O 1s core level from the AlN wafer after solvent cleaning showed a broad peak which

was deconvoluted into two peaks at 531.3 eV and 533.0 eV. The intensity of both O 1s

peaks was observed to increase after the UV/O3 exposure (see Figure 7.2(a,b)). The reported

binding energies of the O 1s core level from various forms of aluminum oxide (Al2O3) have

ranged from 530.7-532.5 eV [95-100], while the reported binding energy positions for the O

193

1s core level from various nitrate compounds (i.e. MNO3) have ranged from 532.7-533.6 eV

[100]. Therefore, it is tempting to assign the O 1s peak at 531.3 eV to Al-O bonding and the

O 1s peak at 533.0 eV to N-O bonding. However, the second O 1s peak at 533.0 eV could

alternatively be due to aluminum hydroxides [95-99]. Previous XPS examinations of various

aluminum oxides and hydroxides (sapphire, gibbsite, bayerite, bauxite, boehmite, and

diaspore) by Tsuchida et al [97], have shown that the binding energy of the O 1s core level of

OH- species (hydroxides) is typically 532.0-532.3 eV while for O2- species it is typically

530.7-531.5 eV. Additionally, the binding energy of the O 1s core level for H2O has been

reported to be 533.3 eV [100]. Further, Tsuchida et al [97] were successful in deconvoluting

the broad O 1s spectrum from boehmite (AlO(OH)) and diaspore (AlO(OH)) into two

separate peaks located at 530.7 and 532.2. eV which they attributed to O2- and OH- species.

By analogy to these previous observations, it therefore seems likely that the native oxide and

UV/O3 generated oxides on AlN surfaces may alternatively be composed primarily of oxygen

bonded to aluminum in both Al-O and AlO-OH states.

The issue of N-O bonding versus AlO-OH bonding could be resolved simply by the

detection of chemically shifted Al 2p and N 1s core levels which would indicate the presence

of either Al-O or N-O bonding. Interestingly though, no chemically shifted peaks were

observed in XPS spectra of the Al 2p and N 1s core levels from these surfaces (see Figure 7.3

and 7.4). Closer inspection of the literature, however, reveals that relative to pure Al, the

chemical shift of the Al 2p core level for both Al-O and Al-N bonding is approximately the

same = 1.4-1.5 eV [100]. However, the reported binding energy for the N 1s core level for

N-O bonding is 401.4-402.8 eV [100] which is significantly different from the reported

binding energy of 397.3 eV for N-Al bonding [76]. Since in our case only a single N 1s peak

at 398.0 eV indicative of N-Al bonding was observed (Figure 7.4), it seems likely that very

194

little N-O bonding is present at the surface and that the native and UV/O3 oxides are

composed mostly of aluminum oxides and hydroxides. Comparing the relative intensities of

the two O 1s peaks, it was also observed that the higher binding energy peak increased in

intensity with UV/O3 exposure whereas the N 1s intensity decreased (data not shown). This

suggests also that the O 1s peak at 533.0 eV is not related to N-O bonding but to AlO-OH

bonding. Therefore it appears that these surface oxides are predominantly composed of

hydroxide (OH-) species (see Figure 7.2 (a,b)). However, it should be pointed out that in this

case the energy of the photoemitted Al 2p and N 1s core levels are ˜ 1410 and 1000 eV

respectively and accordingly are not very sensitive to the outermost surface layer of AlN (λ ˜

20Å). Hence, the presence of some N-O bonding at the surface can not be absolutely ruled

out. In fact, more surface sensitive photoemission is needed to completely resolve this issue.

30 130 230 330 430 530 630 730

(a)

(b)

(c)

Al C

N

O

dN(E

)/dE

F


Figure 7.1. AES survey spectra of OMVPE AlN: (a) as received, (b) solvent cleaned and 20 min. UV/O3 exposure, and (c) 3 min. dip in 10:1 buffered HF (BHF).

195

528 530 532 534 536 538

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

O2- OH -

Figure 7.2. XPS of O 1s core level from bulk AlN wafer (a) as received, (b) UV/O3 exposure, and (c) 10:1 BHF.

70 72 74 76 78

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

Al-N

Figure 7.3. XPS of Al 2p core level from bulk AlN wafer after a 10:1 BHF dip.

196

392 396 400 404 408 412 416 420 424

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

N-Al

AlNPlasmon

Figure 7.4. XPS of N 1s core level from bulk AlN wafer after a 10:1 BHF dip.

As AlN is reasonably inert, oxidation of an AlN surface in a typical laboratory

ambient was not observed to proceed rapidly. Therefore, UV/O3 exposure was used to

repeatedly grow a thin oxide on an AlN surface to assess the efficacy of wet chemical

removal of this oxide. Among the wet chemicals investigated, 1:1 HCl:DI, 1:1

NH4OH:H2O2, RCA SC1 and SC2 solutions were observed to significantly reduce the surface

oxide on UV/O3 treated AlN surfaces. However as shown in Figure 7.1(c), 10:1 buffered HF

(7:1 NH4F:HF) processes were observed to most dramatically reduce/remove the UV/O3

oxide on AlN surfaces. A comparison of these various wet chemical cleans is provided in

Table 7.1 and Figure 7.5. For comparison, the CKLL/NKLL and OKLL/NKLL AES peak to peak

height (pph) ratios for UV/O3 treated AlN surfaces etched in NH4OH:H2O2 was ˜ 0.32 and

0.58 respectively, while for UV/O3 treated AlN surfaces etched in 10:1 buffered HF (BHF),

CKLL/NKLL and OKLL/NKLL ratios of 0.22 and 0.12 respectively were obtained (see Table 7.1).

197

In terms of carbon coverage, the RCA SC1 clean was observed to be equivalent to the BHF

clean, however, the RCA SC2 clean was observed to leave a slightly larger carbon coverage.

As displayed in Figure 7.1, it is also interesting to note that a change in the AlLVV line shape

from that typical of aluminum oxide to that of aluminum nitride was also observed after the

10:1 BHF clean (see Figure 7.6) [85]. Further examination of Figure 7.2(c) also reveals that

the 10:1 BHF treatment reduces the intensity of the OH- O 1s core level to that of the O2- O

1s core level. This suggests that HF primarily attacks hydroxide (OH-) species on AlN

surfaces. Although, a detailed comparison between solutions was not made, similar results

were also obtained with both 10:1 HF, 10:1 BHF, and 40% NH4F.

Table 7.1. OKLL/NKLL and CKLL/NKLL AES pph ratios from OMVPE AlN surfaces given various wet chemical treatments following a UV/O3 oxidation (uncorrected for differences in sensitivity).

C/N O/N Al/N UV/O3 0.27 2.57 0.66 10:1 BHF 0.22 0.12 0.24 1:1 HCl:DI 0.29 0.36 0.27 1:1 NH4OH:H2O2 0.32 0.58 0.27 RCA SC1 0.20 0.21 0.30 RCA SC2 0.33 0.21 0.30

198

100 200 300 400 500 600 700

dN(E

)/dE


(a)

(b)

(c)

(d)

(e)

AlC

NO

F

Figure 7.5. AES of (0001) OMVPE AlN after UV/O3 oxidation and oxide removal with (a) 1:1 NH4OH:H2O2, (b) 1:1 HCl:DI, (c) 10:1 BHF, (d) RCA SC1, and (e) RCA SC2. (spectra normalized to NKLL).

30 50 70 90 110 130

dN(E

)/dE


(a)

(b)

Figure 7.6. Close up of AlLVV from AES of OMVPE AlN after (a) UV/O3 and (b) 10:1 BHF. 199

It is also important to note that the presence of a small concentration of fluorine was

detected in the AES spectra shown in Figure 7.1(c) and 7.5(c). Relative to XPS, AES is

particularly insensitive to detecting the presence of fluorine and electron stimulated

desorption effects are also a problem [101,102]. Therefore, XPS was employed to further

investigate AlN surfaces cleaned using HF based solutions. For this purpose, a thin (˜ 30Å)

AlN film was intentionally grown on 6H-SiC (0001) by GSMBE [94] in order to avoid any

possible charging effects. Figure 7.7(a) shows an XPS spectrum of the F 1s region from a

GSMBE AlN film after dipping in 10:1 buffered HF (BHF) for 10 min. A broad peak was

detected and which was deconvoluted using a Gaussian-Lorentzian distribution into two

peaks having peak positions of 686.8 and 688.5 eV (see Table 7.2). These peaks were

assigned to Al-F and N-F bonding based on analogy to previous reports of XPS from

AlF3.H2O and NF3 [100,103,104]. The carbon and oxygen contamination on this surface was

studied as well. Figure 7.8(a) shows that after HF dipping most of the carbon is located at

285.7 eV which is typical of adventitious carbon and is indicative of a mixture of C-O and C-

H bonding [100]. The O 1s core level from the HF treated surface was observed to be quite

broad (Figure 7.9(a)) and in this case fitted to a single Gaussian-Lorentzian line shape

centered at 533.1 eV. Thermal desorption of the F, C, and O contamination from this surface

was also studied and is described in the following section.

200

Table 7.2. XPS core level positions and full width half maxima (Γ, FWHM) from a GSMBE AlN surface after dipping in 10:1 BHF and annealing at various temperatures.

Al 2p (Γ) N 1s (Γ) O 1s (Γ) C 1s (Γ) F 1s (Γ)____ ) ) ) BHF Dip 75.4 (1.9) 398.6 (1.8 533.1 (3.2 285.8 (2.3 686.8 (2.2)

688.5 (2.2) 400°C 75.3 (1.9) 398.6 (1.8) 533.1 (2.9) 285.7 (2.1) 686.7 (2.1) 530.7 (1.5) 688.7 (2.3) 600°C 75.2 (1.9) 398.5 (1.8) 533.1 (2.4) 285.5 (1.8) 686.5 (2.0) 531.4 (1.9) 688.5 (2.6) 800°C 75.2 (1.9) 398.4 (1.8) 532.9 (2.6) 285.5 (1.8) 686.5 (1.9) 531.1 (1.4) 688.3 (3.0) 950°C 75.2 (1.9) 398.4 (1.8) 532.8 (2.9)

682 684 686 688 690 692 694

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

(d)

(e)

Al-F N-F

Figure 7.7. XPS of the F 1s core level from a 30Å AlN GSMBE film on (0001) 6H-SiC after (a) dipping in 10:1 BHF, and annealing for 15 min. at: (b) 400°C, (c) 600°C, (d) 800°C, and (e) 950°C.

201

280 282 284 286 288 290

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

(d)

(e)

C-Si

Adventitious Carbon

Figure 7.8. XPS of the C 1s core level from a 30Å AlN GSMBE film on (0001) 6H-SiC after (a) dipping in 10:1 BHF, and annealing for 15 min. at: (b) 400°C, (c) 600°C, (d) 800°C, and (e) 950°C.

526 528 530 532 534 536

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

(d)

(e)

Figure 7.9. XPS of the O 1s core level from a 30Å AlN GSMBE film on (0001) 6H-SiC after (a) dipping in 10:1 BHF, and annealing for 15 min. at: (b) 400°C, (c) 600°C, (d) 800°C, and (e) 950°C.

202

Other wet chemistries based on H2SO4, H3PO4, and NaOH etc. were also investigated

for both oxide and carbon removal from AlN surfaces as these chemicals are commonly used

in GaAs processing [35-53] as well as for wet chemical etching of III-V nitride compounds

[105-108]. Treatments in concentrated H2SO4 and H3PO4 were observed to leave residual

sulfate and phosphate on the surface which was probably related to difficulties in rinsing

these chemicals off the AlN surface due to their viscous nature. H2O2:H2SO4 (Piranha etch)

was observed to be good for removing gross carbon contamination from AlN surfaces

(photoresist, etc.). NaOH was observed to leave traces of Na on the surface, which however,

were successfully removed below the detection limits of XPS with an RCA clean (see Figure

7.10). It should be noted, though, that the detection limits of XPS is ˜ 0.1 atomic % and Na

contamination levels below this level have been historically noted to wreak havoc with Si

MOSFET devices [34]. More dilute levels of H3PO4 were moderately successful for oxide

removal, but it was observed that when etching AlN in H3PO4 at higher temperatures (100-

150°C) the, AFM RMS surface roughness increased from as low as 20Å to as high as 200Å.

203

1060 1065 1070 1075 1080 1085 1090

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

Figure 7.10. XPS of Na 2p from polycrystalline AlN surface after (a) etching in NaOH and (b) an RCA clean.

7.4.2. In Situ Processing of AlN

The chemistry and thermal of desorption of F, C, and O contaminants on AlN

surfaces after HF processing was further investigated using AES, XPS, and temperature

programmed desorption (TPD). Fig. 7.7(b) shows that after annealing an HF dipped AlN

surface at 400°C for 15 min., the two F 1s peaks become more distinguishable and positioned

at 686.7 and 688.7 eV. Annealing at 600°C resulted in a reduction in intensity of the higher

binding energy peak, as shown in Figure 7.7(c); this peak was further reduced after annealing

at 800°C. Complete elimination of the low binding energy peak at 686.5 eV was not

achieved until annealing at 950°C for 15 min. (Fig. 7.7(e)). During this series of thermal

desorptions, the C 1s and O 1s core levels were also monitored in XPS. The intensity of the

204

C 1s core level was observed to decrease appreciably after the 600°C anneal, but complete

removal of carbon from the AlN surface was not observed until after the 950°C anneal

similar to fluorine (see Figure 7.8(b-e)). The intensity of the O 1s peak initially slightly

decreased after the 400 and 600°C anneals presumably due to water and CO desorption.

However. the O1s intensity almost doubled in intensity after the 950°C anneal (see Figure

7.9(b-e)). This was attributed to reaction of the AlN surface with water desorbing from the

chamber wall due to heating of the walls during the anneal. It is also important to note that

after the 400°C anneal, two O 1s peaks can again be resolved. However after the 950°C

anneal, only one O 1s peak could be resolved suggesting that the hydroxides dehydrogenate

forming O2- species. Similar desorption effects were also seen for carbon and oxygen from

solvent cleaned and UV/O3 processed AlN surfaces.

In a separate study, TPD was performed on an amorphous AlN film sputtered on Si

(111) which had been subsequently dipped in 10:1 BHF (see Figures 7.11). For this surface,

a strong desorption peak for m/e- 16 and 18 (H2O) was observed to occur at temperatures of <

200°C (see Figure 7.4(a)) which is agreement with the observed decrease in O 1s intensity in

XPS for AlN surfaces annealed at 400°C. A large TPD peak was detected for desorption of

m/e- 19 and 20 (F and HF) at approximately 400°C while a small peak for m/e- 38 (F2) was

detected at 500°C (see Figure 7.11(b,c)). This is also in agreement with the observed decease

in intensity of the 688.7 eV F 1s peak after annealing fluorinated AlN surfaces at this

temperature. Desorption features at 400-500°C for m/e- 2, 12, 16, 18, and 28 were also

detected and are probably related to desorption of H2O, CO, and various organic and

hydrocarbon contaminants. A general increase in m/e- 2 (H2), m/e- 19 (F), and m/e- 28 (N2 or

CO) occurred after 600°C and continued until 1000°C after which annealing was stopped.

205

100 200 300 400 500 600

Cou

nts (

arb.

uni

ts)

Temperature (ÞC)

(a)

(b)

(c)

Figure 7.11. TPD of m/e- (a) 18, (b) 20 and (c) 38 from 10:1 BHF dipped AlN (ß = 20°C/min.).

As complete thermal desorption of fluorine from AlN occurred only at elevated

temperatures, the lower temperature process of remote H plasma cleaning was investigated

for this purpose and for the removal of other contaminants. The details regarding the RF

plasma system and its operation are described elsewhere [97]. Atomic H has been previously

shown to be extremely efficient for the removal of halogens from Si (001) surfaces

[109,110]. Figure 7.12 shows an XPS spectrum of the F 1s region taken from a

polycrystalline AlN wafer before and after exposure to a remote H plasma. Figure 7.12(a)

shows the presence of a large quantity of F on the AlN surface after dipping in 10:1 BHF.

Figure 7.12(b) reveals that almost complete removal of fluorine is achieved after exposure of

the AlN wafer to a 15 mTorr, 20W, remote H plasma at 450°C for 5 min. This technique was

206

also efficient for removal of C from AlN surfaces, as shown by the AES spectra in Figure.

7.13(a) and (b). However, it was not found efficient for removing oxygen from AlN

surfaces.

Annealing AlN in fluxes of Al (˜.1 ML/sec), Ga (˜.1 ML/sec), and NH3 (˜ 1 sccm)

was investigated. Though effective for removing fluorine and carbon, none of these

processes was found particularly affective in further removing oxygen from the AlN surface

(see Figure 7.14). No attempts were made to evaporate a thin layer of Al or Ga onto the AlN

surface at room temperature and then thermally desorb this layer at higher temperatures via

the method of Khan [79]. Exposure to silane at high temperatures (1000°C) was the only in

situ clean found capable of appreciably removing oxygen from an AlN surface (see Figure

7.15). Unfortunately, the loss of oxygen produced by the silane exposure was at the expense

of some deposition of silicon presumably due to the formation of Si-N bonding at the surface

(data not shown).

207

685 687 689 691 693 695 697

(a)

(b)

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

Figure 7.12. XPS of F 1s core level from polycrystalline AlN wafer cleaned in 10:1 BHF (a) before and (b) after remote H plasma exposure at 450°C (15 mTorr, 20W). Note F 1s shifted due to sample charging.

30 130 230 330 430 530 630 730

(a)

(b)

dN(E

)/dE


Al CN O

Figure 7.13. AES survey spectra of OMVPE AlN after: (a) UV/O3 and 10:1 BHF dip and (b) remote H plasma at 450°C. 208

526 528 530 532 534 536

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

Figure 7.14. XPS of O 1s from (0001) GSMBE AlN after (a) annealing at 1000°C and (b) annealing in a 0.1 ML/sec flux of Al at 1000°C.

526 528 530 532 534 536 538

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

Figure 7.15. The XPS O 1s core level from an AlN surface (a) before and (b) after annealing in a SiH4 flux.

209

As previous studies on GaAs have demonstrated that a thin (<1Å) In layer can be

used as a passivation layer to protect GaAs surfaces in air and then be later thermally

desorbed in vacuum [111], In was investigated as a passivation layer for AlN. These films

were deposited on OMVPE and GSMBE AlN films in situ immediately after growth (i.e. no

exposure to atmosphere). Unfortunately, In was not found to be a suitable passivation layer

for AlN as the In films were observed to ball up on the AlN surface instead of wetting the

AlN surface. This effect left part of the AlN surface exposed to ambient and therefore some

of the AlN surface was allowed to oxidize. As illustrated in Figure 7.16, after thermal

desorption of the In layer at 750°C, a significant concentration of oxygen was present on the

AlN surface. The level of this oxygen contamination was equivalent to that left after wet

chemical processing in HF. Therefore, 200Å of GaN was also investigated as an alternative

passivation/protecting layer for AlN. In this case, complete surface coverage was less of a

problem and thermal desorption at ˜ 950°C resulted in an essentially oxygen and carbon free

surface (see Figures 7.17 and 7.18). However, complete desorption of the GaN film was not

observed to occur at this temperature. In order to more completely desorb the GaN

passivation layer, extended annealing at temperatures of > 1000°C were required. However

even after extended annealing at >1000°C, AES and XPS still detected a persistent trace of

Ga on the surface (see Figure 7.19, and Tables 7.3 and 7.4). Additionally, the extended high

temperature annealing was observed to result in increased oxidation of the AlN surface due to

heating of the chamber walls causing water to desorb and react with the AlN surface (see

Figure 7.17(e)).

210

Table 7.3. XPS core level positions and FWHM (Γ) from a 200Å OMVPE GaN capping layer on OMVPE AlN after annealing at various temperatures. Al 2p (Γ) Ga 3d (Γ) Ga 2p3/2 (Γ) O 1s (Γ) C 1s (Γ)____ As Received 20.5 (1.6) 1118.4 (1.8) 532.8 (2.4) 285.7 (2.1) 531.1 (1.5) 500°C 20.2 (1.6) 1118.0 (1.9) 533.2 (2.0) 284.8 (2.0) 531.2 (2.0) 750 75.2 (1.8) 20.1 (1.5) 1117.9 (1.8) 531.8 (2.8) 284.5 (1.9) 950 75.3 (1.7) 20.3 (1.3) 1118.2 (1.9) 532.4 (3.8) 21.5 (1.9) 1119.6 (2.2) >1000°C 75.3 (1.6) 20.2 (1.5) 1118.8 (2.7) 532.6 (3.5)

Table 7.4. Ratio of integrated intensity of Al 2p to Ga 3d and Ga 2p3/2. Al2p/Ga3d Al2p/Ga2p3/2 As Received 0.0 0.0 500°C 0.0 0.0 750°C 0.5 0.06 950°C 3.33 0.11 >1000°C 7.2 0.31

526 528 530 532 534 536 538

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

211

Figure 7.16. XPS of O 1s from In capping layer on OMVPE AlN (a) as received, (b) after annealing at 600°C, and (c) after annealing at 750°C.

526 528 530 532 534 536 538

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

(d)

(e)

Figure 7.17. XPS of O 1s core level from 200Å GaN capping layer on (0001) AlN buffer layer, (a) as received, (b) after annealing at 500°C, (c) 750°C, (d) 950°C, and (e) > 1000°C.

280 282 284 286 288

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

(d)

(e)

Figure 7.18. XPS of C 1s core level from 200Å GaN capping layer on (0001) AlN buffer layer, (a) as received, (b) after annealing at 500°C, (c) 750°C, (d) 950°C, and (e) >1000°C. 212

1114 1116 1118 1120 1122 1124

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

(d)

(e)

Figure 7.19. XPS of Ga 2p3/2 core level from 200Å GaN capping layer on (0001) AlN buffer layer, (a) as received, (b) after annealing at 500°C, (c) 750°C, (d) 950°C, and (e) >1000°C.

7.4.3. Ex situ Cleaning of GaN

Figure 7.20 shows an AES spectrum from an as received GaN film grown via

GSMBE. As this sample was kept on a laminar flow bench for the period of time that it was

removed from vacuum (1 day), the level of C contamination was quite small. To investigate

the ability of UV/O3 to remove carbon from GaN surfaces, the GaN film was exposed to

UV/O3 for 10-20 min. and investigated with AES. Figure 7.20(b) shows that this exposure

resulted in additional oxidation of the GaN surface without significantly affecting the level of

C contamination. In fact, the C level actually increased, probably as a result of the increased

handling of the sample during unmounting and remounting for the UV/O3 treatment.

213

Examination of the UV/O3 treated GaN surfaces with LEED did not detect any diffraction

patterns, indicating that the O3 generated oxide is likely amorphous. To test that our UV/O3

apparatus was operating correctly, a GaAs substrate was examined simultaneously with an

OMVPE GaN film. As shown in Figure 7.21, AES showed complete removal of carbon

from the GaAs surface after a UV/O3 treatment in agreement with the results of others

[54,58]. Further examination of more contaminated (0001) OMVPE GaN surfaces using

XPS, however, did show a reduction in surface carbon coverage on these surfaces after

UV/O3 exposure (see Figure 7.22). As in the case of AlN, complete removal of carbon was

not detected but a reduction in carbon was observed in XPS and AES with the C 1s peak in

XPS shifting to a higher binding energy indicative of oxidation of the surface carbon (see

Figure 7.22 and Table 7.5).

30 230 430 630 830 1030 1230

(a)

(b)

(c)

dN(E

)/dE


GaON

C

Cl

Ga

Figure 7.20. AES survey spectra from GSMBE GaN after (a) 1 day in air on laminar flow bench, (b) UV/O3 oxidation, and (c) 5 min. etch in 1:1 HCl:DI.

214

Table 7.5. XPS core level positions for OMVPE GaN surfaces after various ex situ treatments (Γ= FHWM). Treatment Ga 3d (Γ) N 1s (Γ) Ga2p3/2 (Γ) O 1s (Γ) C 1s (Γ)____ As received 20.1 (1.4) 397.7 (1.3) 1118.2 (1.8) 532.4 (3.1) 285.3 (2.0) Solvents 20.1 (1.4) 397.7 (1.3) 1118.2 (1.8) 532.4 (3.1) 285.3 (2.0) UV/O3 20.3 (1.4) 397.9 (1.3) 1118.5 (1.9) 532.7 (3.1) 285.8 (2.3) UV/O3 24hr. 20.8 (1.7) 398.2 (1.4) 1119.0 (2.1) 533.0 (3.2) 285.8 (3.9) 531.5 (1.9) HCl:DI 20.2 (1.4) 397.9 (1.4) 1118.3 (1.7) 532.7 (2.9) 285.5 (2.1) BHF Vapor 20.8 (1.4) 398.2 (1.2) 1118.8 (1.8) 533.2 (3.6) 286.1 (2.3) 403.8 (2.0) 11120.5 (1.9)

30 230 430 630 830 1030 1230

(a)

(b)

dN(E

)/dE


GaC O

GaAs

Figure 7.21. AES survey spectra from (001) GaAs (a) before and (b) after UV/O3.

215

280 282 284 286 288 290

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

Figure 7.22. XPS of C 1s core level from (0001) OMVPE GaN after (a) ultrasonification in trichloroethylene, acetone, and methanol, and (b) UV/O3 exposure.

To further assist in the removal of carbon from the GaN surface in later experiments,

the UV/O3 box was purged with 1 L/sec. of oxygen to increase the concentration of ozone

generated within the box. The O2 purge did result in a further decrease in the surface carbon

coverage by the UV/O3 treatment, but complete carbon removal was still not achieved (see

Figure 7.23). However, the oxygen purge also enhanced the oxidation rate of the GaN

surfaces during the UV/O3 exposures. This was observed by an almost complete

disappearance of the NKLL and N 1s peak in AES and XPS respectively (see Figure 7.23). In

addition, the Ga and N core levels were observed to broaden and shift to higher binding

energies by ˜ 0.7-0.8 eV and 0.5 eV respectively (see Table 7.5). However despite the high

degree of oxidation of this GaN surface, chemically shifted N or Ga core levels could not be

216

detected/resolved to indicate the existence of either Ga-O or N-O bonding. As with AlN, the

observed chemical shifts for Ga core levels in XPS for Ga-N and Ga-O bonding are both

approximately the same. For the Ga 3d core level, binding energies of 19.6-21.0 eV [100]

have been reported for Ga2O3 and for GaN the reported Ga 3d core levels are 19.2-20.3

[81,100]. However as mentioned previously, the reported chemical shifts for the N 1s core

level for N-Ga (BE=397.2 [100]) and N-Ox (BE=400-405 [100]) bonding are much larger. In

fact, for oxidized InN a large chemically shifted N 1s core level at ˜ 405 eV has been

observed and attributed to NO and NO2 species [112]. As a large chemically shifted N 1s

core level was not observed in XPS of the UV/O3 treated GaN surfaces and the N 1s core

level was observed to broaden and decrease in intensity relative to the Ga core levels, it

seems likely that the oxide formed on GaN by the UV/O3 treatment is composed mostly of

Ga-O bonded oxygen.

30 230 430 630 830 1030 1230

dN(E

)/dE


O

GaNCGa

217

Figure 7.23. AES survey scan of OMVPE GaN after a 24 hr UV/O3 exposure with 1L/sec. flowing O2. XPS of the O 1s core level from an OMVPE GaN surface undergoing a 24 hr. UV/O3

exposure with 1 L/sec flowing O2 showed the development of a second O 1s peak at 531.5

eV possibly due to the formation of stoichiometric Ga2O3 (O1s = 530.8 eV [100]). Prior to

UV/O3 exposure or after a short UV/O3 exposure, only a single broad O 1s core level

(FWHM = 3.1 eV) centered at 532.4-532.7 eV could be detected from these GaN surfaces

(see Figure 7.24 and Table 7.5). Therefore by analogy to AlN surfaces, it seems likely that

the oxides formed on air exposed and UV/O3 treated GaN surfaces are also composed mostly

of O2- and OH- species (see Figure 7.24).

526 528 530 532 534 536 538

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

(c)

O2-

OH -

Figure 7.24. XPS of O 1s core level from (0001) OMVPE GaN (a) after solvent cleaning, (b) after UV/O3 oxidation for 25 min., and (c) after UV/O3 oxidation for 24 hr. with 1 L/sec flowing O2.

218

The combination of UV/O3 oxidation and exposure to the same acids and bases

investigated above for oxide removal from AlN was also used to determine the best wet

chemical method for removing oxides from GaN surfaces. In this case, solutions of HCl,

NH4OH, and HF were found to be very effective for oxide removal. Table 7.6 summarizes

the AES OKLL/NKLL and CKLL/NKLL pph ratios obtained from GaN surfaces undergoing

UV/O3 oxidation followed by numerous wet chemical treatments. As illustrated in Table 7.6,

1:1 HCl:DI was found to produce the lowest OKLL/NKLL ratio of the wet chemistries

investigated and as illustrated in Fig. 7.20(c) and Fig. 7.25(a) significant coverages of Cl

were observed on HCl treated GaN surfaces. Slightly higher OKLL/NKLL ratios were obtained

with 10:1 HF and 10:1 BHF solutions, but F was not detected by either AES or XPS.

However, the lowest CKLL/NKLL ratios were obtained with these HF based solutions. It

should be noted, however, that the values presented in Table 7.6 represent the very best

values obtained with each clean and as illustrated in Figure 7.25, on average similar results

were obtained for both 1:1 HCl:DI, 10:1 BHF, and 1:1 NH4OH:H2O2 wet chemical

treatments. In fact we have previously reported slightly higher values for OKLL/NKLL and

CKLL/NKLL ratios for HCl and HF treated GaN surfaces. These higher values could be related

to longer times required to mount samples, laboratory ambient and cleanliness (i.e. pollen

count, smoke, air circulation) and the fact that Teflon beakers were used for HF processes

and Pyrex/glass beakers were used for all other wet chemical processes. It should be noted,

however, that for HCl treatments, the oxygen surface coverage was observed to be inversely

related to the amount of Cl detected on the surface (i.e. higher Cl coverage equals lower

oxygen coverage). A similar relation was also observed between carbon and oxygen (i.e.

higher carbon coverage equals lower oxygen coverage). All GaN surfaces treated in either 219

HCl:DI, HF, BHF, or NH4OH:H2O2 displayed (1x1) LEED diffraction patterns indicating

removal of the UV/O3 oxide. Other wet chemistries including , RCA SC1 and RCA SC2,

1:1:7 H2SO4:H2O2:DI, H3PO4, and Acetic were also investigated. RCA SC1 and SC2 were

found to reduce the UV/O3 oxide on GaN surfaces but (see Figure 7.26) the SC2 clean was

found to leave more carbon on the surface relative to the SC1 clean which is similar to what

was observed for the AlN surfaces. The H2SO4 and H3PO4 cleans were found to leave

residual sulfates and phosphates on the GaN surfaces with accordingly higher oxygen levels.

Table 7.6. AES pph ratios of UV/O3 and wet chemical processed OMVPE GaN surfaces Treatment OKLL/NKLL CKLL/NKLL GaLMM/NKLL 3 hr. in air 0.22 0.07 As Received 0.15 0.21 0.59 Solvents 0.20 0.24 0.61 UV/O3 0.85 0.19 0.81 HCl:DI 0.12 0.23 0.62 NH4OH:H2O2 0.29 0.22 0.73 10:1 HF 0.15 0.10 0.69 40% NH4F 0.23 0.14 0.77 10:1 BHF 0.23 0.20 1:1:7 H2SO4:H2O2:DI 0.33 0.22 0.63 85% H3PO4 @125°C 0.27 0.15 0.76 RCA SC1 0.25 0.24 0.60 RCA SC2 0.16 0.35 0.62 BHF Vapor 0.43 0.35 0.77

220

30 130 230 330 430 530 630 730

dN(E

)/dE


(a)

(b)

(c)

Cl CONGa

Figure 7.25. AES of (0001) OMVPE GaN after UV/O3 oxidation and oxide removal with (a) 1:1 HCl:DI, (b) 10:1 BHF, and (c) 1:1 NH4OH:H2O2 (spectra normalized to NKLL).

30 230 430 630 830 1030 1230

dN(E

)/dE


Ga C

NO

Ga

(a)

(b)

(c)

Figure 7.26. AES of (0001) OMVPE GaN after UV/O3 oxidation followed by (a) RCA SC1 and (b) RCA SC2 (spectra normalized to NKLL). 221

As we have previously found an ex situ exposure to the equilibrium vapor from a

BHF solution to be useful for UV/O3 oxide removal from (0001)Si 6H-SiC surfaces, this

approach was also investigated for oxide removal from (0001) GaN surfaces. As displayed

in Table 7.6 and Figure 7.27, the oxygen surface coverages observed in AES from UV/O3

and BHF vapor treated GaN surfaces were not found to be lower than those for typical wet

chemical approaches. However, XPS analysis seemed to indicate a lower oxygen coverage

for the vapor cleaned GaN surfaces compared to the wet chemical cleaned surfaces. The

reasons for this are currently unclear but may be related to enhanced electron beam oxidation

of the GaN surface due to physisorbed water left on the surface by the vapor treatment. One

other significant difference between the BHF vapor treatment and the wet chemical cleans

was the detection of large amounts of fluorine on the BHF vapor treated surfaces (see Figure

7.27) compared to the BHF wet chemical cleans for which no fluorine was detected. XPS of

the F 1s core level from the BHF vapor treated surfaces showed a broad F 1s peak which

could be deconvoluted into two separate peaks at 686.6 and 688.2 eV similar to BHF wet

chemical processed AlN surfaces (see Figure 7.28). In addition, chemically shifted Ga 2p

and N 1s core level peaks at 1120.5 and 403.8 eV presumably due to Ga-F and N-F bonding

were also detected indicating the possible formation of both GaF3 and NF3 on the surface (see

Figure 7.29, 7.30 and Table 7.5). It is interesting to note that the intensity of the Ga 2p3/2 Ga-

F peak was observed to decrease in intensity along with the 686.6 eV F1s peak after rinsing

in DI water whereas the N 1s peak at 403.8 eV was not effected. Both Ga-F and N-F Ga2p3/2

and N1s peaks were completely removed by cleaning in 1:1 NH4OH:H2O2. As a final note,

preliminary AFM examinations indicated that all of the above wet chemical treatments (with

the exception of H3PO4) did not increase the GaN RMS surface roughness above that of the

as grown surface ˜ 10Å.

222

30 230 430 630 830 1030 1230

C

N

OGa

Cl

F

Ga

dN(E

)/dE

(a)


(b)

(c)

Figure 7.27. AES of (0001) OMVPE GaN (a) HCl:DI dip, (b) UV/O3 oxidation, and (c) HF vapor cleaning.

684 686 688 690 692 694

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

Ga-F

N-F

Figure 7.28. XPS of F1s core level from (0001) GaN after UV/O3 oxidation followed by (a) 10:1 BHF vapor clean, and (b) a DI rinse. 223

1114 1116 1118 1120 1122

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(b)

(a)

Ga-F

Ga-N

Figure 7.29. XPS of Ga2p core level from (0001) OMVPE GaN after UV/O3 oxidation and (a) BHF vapor clean and (b) DI rinse and NH4OH:H2O2 clean.

393 396 399 402 405

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

N-Ga

N-F

Figure 7.30. XPS of N 1s core level from (0001) OMVPE GaN after UV/O3 oxidation and (a) BHF vapor clean and (b) DI rinse and NH4OH:H2O2 clean. 224

7.4.4. In Situ Processing of GaN

In a previous study, we used AES to examine the thermal desorption of O, C, and Cl

contaminants left on GaN surfaces after various DI:HCl:HF:MeOH wet chemical processes

[85]. In this study, it was observed that annealing these GaN surfaces at 800°C resulted in

only incomplete desorption of carbon and oxygen contaminants [85]. In the GaN/AlN

capping layer study, however, complete desorption of carbon from the GaN surface was

achieved after annealing at 950°C (see Figure 7.18). In the case of oxygen, almost complete

thermal desorption of oxygen was also achieved in this temperature range (see Figure 7.17).

However in separate TPD experiments, it was observed that at Tsub >˜ 800°C an increase in

m/e- 69 (Ga) could be detected and which exponentially increased with temperature (see

Figure 7.31). This suggests that the oxygen and carbon contaminants were removed from the

GaN surface through sublimation of the GaN film. Unfortunately due to experimental

difficulties we were not successful in observing Cl desorption from GaN surfaces using TPD.

However, the concentration of the former was observed to be decreased below the detection

limit of AES by annealing at 600°C or greater (see Figure 7.32). This is contrast to GaAs

surfaces where GaCl was observed to desorb at ˜ 220°C [113].

Further investigation of the removal of native and UV/O3 oxides from GaN via

annealing in fluxes of Ga and NH3 was also conducted. Using AES, both procedures were

observed to result in a reduction in the amount of oxygen on GaN surfaces (see Fig. 7.33(a-

d)). However, oxygen could not be removed below the detection limit (˜ 0.1%) of our AES

system. This however, was later found to be related to electron beam oxidation of the GaN

surface by the Auger e-beam during AES analysis. Figure 7.34 shows a series of AES

spectra taken sequentially for differing periods of time from an as grown GSMBE GaN 225

surface. As this figure illustrates, the intensity of the OKLL peak was observed to increase

with the number of scans. When scanning very quickly or performing only one scan,

essentially no oxygen could be detected from the as grown GaN surface in AES. A detailed

examination of Ga vapor flux cleaning was not investigated here, as it has been previously

thoroughly examined by Kahn [79] and Bermudez [80-82]. However, as shown in Figure

7.33(d) essentially atomically clean GaN surfaces were obtained by annealing in NH3 at

800°C. Further, (2x2) reconstructions were observed in LEED from OMVPE and GSMBE

GaN films annealed in ammonia which has not been observed from GaN surfaces cleaned via

annealing in a Ga flux.

Table 7.7. XPS core levels from GSMBE GaN surfaces after various treatments. Treatment Ga 3d (Γ) Ga 2p3/2 (Γ) N 1s (Γ) O 1s (Γ) C1s (Γ) UV/O3 20.3 (1.6) 1118.1 (1.9) 387.7 (1.3) 532.0 (2.7) 285.0 (1.8) 650°C UHV 20.1 (1.6) 1117.8 (1.9) 397.5 (1.2) 531.0 (2.2) 283.9 (2.3) NH3-650°C 20.1 (1.6) 1117.8 (1.9) 397.5 (1.3) 531.2 (2.7) 283.9 (2.3) Ga-650°C 20.0 (1.8) 1117.4 (2.2) 397.6 (1.6) 532.0 (2.5) NH3-800°C 20.0 (1.6) 1117.5 (1.8) 397.5 (1.4) 531.0 (2.6)

226

0 10 0

1 10 -11

2 10 -11

3 10 -11

4 10 -11

5 10 -11

200 300 400 500 600 700 800 900 1000

m/e

- 69

(arb

. uni

ts)

Temperature (ÞC)

Figure 7.31. m/e- 69 (Ga) signal from GSMBE GaN as a function of surface temperature.

30 130 230 330 430 530 630 730

dN(E

)/dE


Ga C

N

O

Cl(a)

(b)

(c)

(d)

Figure 7.32. AES of (0001) OMVPE GaN after (a) HCl vapor clean, and annealing at (b) 300°C, (c) 450°C, and (d) 600°C. (spectra normalized to NKLL).

227

30 230 430 630 830 1030 1230

dN(E

)/dE


(a)

(b)

(c)

(d)

Ga CN

O Ga

Figure 7.33. AES survey spectra from GSMBE GaN (a) exposed to UV/O3, and after annealing in: (b) UHV at 650°C, 20 min., (c) NH3 (5x10-6 Torr) at 650°C, 20 min, and (d) NH3 (5x10-6 Torr) at 800°C, 25 min. (spectra normalized to NKLL).

30 230 430 630 830 1030 1230

dN(E

)/dE


O

Ga

N

Ga

(a)

(b)

(c)

(d)

Figure 7.34. AES survey spectra from GSMBE GaN after various sequential scans (a) 1, (b) 3, (c) 8, and (d) 9 scans. (spectra normalized to NKLL). 228

The surface electronic structure of (0001) OMVPE GaN surfaces after wet chemical

processing and in situ processing was also investigated using UPS. UPS spectra of a (0001)

OMVPE GaN surface after cleaning in methanol, 1:1 HCl:DI, and annealing in NH3 are

displayed in Figure 7.35 along with a UPS spectrum of an as grown (2x2) GSMBE GaN film.

It was observed that after wet chemical processing the VBM of the GaN surface was located

at -3.8 eV below the system Fermi level perhaps suggesting that the GaN surface Fermi level

is inverted as the bandgap of GaN is only 3.4 eV. After annealing in NH3, however, the

VBM of the GaN surface was observed to move closer to the Fermi level to ˜ 3.0 eV below

EF. In comparison, the VBM of an as grown (2x2) GSMBE GaN film was observed to be

located ˜ 2.7 eV below the Fermi level. However the position of the GaN core levels relative

to the GaN VBM were observed to be the same for the as grown GSMBE GaN surface and

the OMVPE GaN film annealed in NH3 at 800°C (see Table 7.8).

-10 -8 -6 -4 -2 0

Cou

nts (

arb.

uni

ts)


-2.7

-3.0

-3.8

(d)

(c)

(b)

(a)

Figure 7.35. UPS of OMVPE GaN after (a) rinsing in methanol, (b) 1:1 HCl:DI, and (c) annealing in NH3 at 800°C, (d) as grown (2x2) GSMBE GaN.

229

Table 7.8. Ga and N core levels from GaN relative to the GaN VBM after various processes. Treatment Ga 3d-VBM Ga2p3/2-VBM Ga3p3/2-VBM N1s-VBM As Received 17.5 eV 1115.5 eV 102.8 eV 395.0 eV 1:1 HCl:DI 17.7 eV 1115.8 eV 103.1 eV 395.2 eV 800°C-NH3 18.4 eV 1116.3 eV 103.6 eV 395.8 eV (2x2) 18.4 eV 1116.3 eV 103.6 eV 395.8 eV

The use of a remote hydrogen plasma for cleaning GaN was also investigated. Figure

7.36(a) shows an AES spectrum from a HCl:DI cleaned OMVPE GaN film before plasma

exposure. Exposure to a 15 mTorr, 20W remote H plasma for 5 min. at 100°C resulted in the

complete removal of both Cl and C from the surface within the detection limits of AES.

However, the level of oxygen was found to increase dramatically, due to the low temperature

plasma exposure (see Figure 7.34(b)). A second exposure of the same sample to a H plasma

at 450°C for 10 min. resulted in a reduction of the oxygen level close to that prior to the first

H plasma exposure. Higher sample temperatures during plasma processing were investigated

(600-800°C), but further reduction of the surface oxygen levels was not realized.

230

200 400 600 800 1000 1200

dN(E

)/dE


Cl CN

O Ga

(a)

(b)

(c)

Figure 7.36. AES survey spectra from OMVPE GaN (a) after oxide removal with 1:1 HCl:DI, (b) 100°C H plasma exposure, and (c) 450°C H plasma exposure.

7.4.5. Ex Situ Cleaning of AlGaN Surfaces

Many of the wet chemical processes found effective for oxide and carbon removal for

GaN and AlN surfaces were also investigated for surfaces of OMVPE AlxGa1-xN (x˜0.5)

films. As shown in Figure 7.37, AlGaN surfaces were observed to retain both Cl and F on

the surface for both HCl and HF wet chemical processes. HF vapor cleaning was also

observed to result in the formation of a large amount of F on the surface leading to AlF3

formation (see Figure 7.36 and 7.37).

231

30 230 430 630 830 1030 1230 1430 1630

(a)

(b)

(c)

(d)

(e)

dN(E

)/dE


CN

O

Cl

F

Ga Al

Figure 7.37. AES of AlGaN (a) as received, (b) after solvent cleaning in trichloroethylene, acetone, methanol, and isopropanol, (c) 1:1 NH4OH:H2O2, (d) 1:1 HCl:DI, and (e) 10:1 BHF.

30 230 430 630 830 1030 1230 1430 1630

(a)

(b)

Al C

N

O

dN(E

)/dE


F Ga Al

Figure 7.38. AES of AlGaN after (a) UV/O3 exposure and (b) HF vapor oxide removal.

232

70 72 74 76 78 80 82

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

(a)

(b)

Al-N

Al-F

Figure 7.39. XPS of Al 2p core level after UV/O3 oxidation and (a) HF vapor oxide removal, and (b) DI rinse and NH4OH:H2O2.

7.5. Discussion

7.5.1. As Received and UV/O3 surfaces

Prabhakaran et al [86] have previously used XPS to examine the native oxide formed

on GaN and observed two O 1s core levels located at 531.3 and 532.7 eV which is in

agreement with our results from native and UV/O3 oxides on both AlN and GaN surfaces.

Prabhakaran et al [86] assigned the lower binding energy peak to Ga2O3 based on the

observation that this peak was removed via etching in NH4OH which is known to dissolve

Ga2O3. However, Prabhakaran et al [86] did not offer an explanation as to the origin of the

233

higher binding energy O 1s core level. Although, Prabhakaran et al [86] did observe a

decrease in intensity on the lower binding energy side of the Ga 2p and N 1s core levels after

etching in NH4OH which they speculated may be due removal of an oxynitride phase at the

surface. Using Zr Mξ radiation (hν = 151.6 eV), Bermudez [82] has recently studied the

interaction of O2 with UHV prepared atomically clean (0001) GaN surfaces. In this case, he

was not able to identify whether oxygen bonded with either N or Ga on the surface but was

able to identify the appearance of a satellite on the Ga 3d core level due to oxygen exposure

which may indicate the formation of Ga-O bonds.

In our XPS analysis of the O 1s core level from as received and UV/O3 treated AlN

and GaN surfaces described here, the presence of oxygen in two different chemical states was

attributed to oxygen bonded to Al (AlN) or Ga (GaN) in both O2- and OH- chemical states.

This was primarily based on the lack of observation of a chemically shifted N 1s core level at

˜ 401-402 eV. The formation of aluminum hydroxides on both aluminum and aluminum

nitride has been previously observed using XPS and FTIR [95-99]. However in both cases,

the hydroxides were observed to form only after suspension in water (H2O) for extended

periods of time (days-weeks). Additionally, Nylund and Olefjord [95,96] observed the

spontaneous decomposition of hydroxides on aluminum into aluminum oxides when inserted

into UHV. Obviously, in our case the "hydroxide" species seem to be much more stable and

were easily detected in UHV by XPS. For these reasons, the higher binding energy O 1s core

level may alternatively be due to oxygen bonded to Al or Ga in a chemical state somewhere

between O2- and OH-. However, we do not feel that the higher binding energy O 1s peak is

due to N-O bonding. In our case the NKLL and N 1s signal in AES and XPS respectively

were observed to decrease in intensity with increased UV/O3 exposure and accordingly the

intensity of both O 1s core levels was observed to increase and the Ga core levels 234

correspondingly were observed to broaden. This argues against the formation or inclusion of

oxynitride phases in the UV/O3 oxide and perhaps the native oxide. In fact one would expect

the more volatile N-O species to be removed from the GaN surface during the UV/O3

treatment which may be the limiting factor in growing an oxide on GaN by this technique.

Finally in our case, we did not observe a decrease in intensity on the low binding energy side

of the N 1s core levels after etching UV/O3 or native oxides in HCl, HF, or NH4OH. Perhaps

during the initial oxidation of GaN surfaces, NOx species form and then are gradually

removed from the surface allowing the formation of a more stoichiometric gallium oxide on

the surface. More detailed analysis of oxidized GaN and AlN surfaces using FTIR could

help resolve this issue.

As for the nature of the carbon contaminants accumulated on both AlN and GaN

surfaces in ambient and wet chemical processing, the XPS C 1s core level was generally

observed to be located at 285.2-285.8 eV with a broad FWHM of 2.2-3.0 eV. Prabhakaran et

al [86] has made similar observations for GaN surfaces and concluded that the adventitious

carbon is composed mostly of C-H bonding. However, in our case we feel that the surface

carbon is really a mixture of C-H and C-O bonding as Miyauchi et al [115] has previously

shown that the C 1s core level for CH2, C-O, and O-C=O bonded carbon contaminants on

silicon surfaces are located at 284.6, 286.3 and 288.4 eV respectively. The exact energy

position of the C 1s core level for these contaminants could be affected by the differences in

band bending at the surfaces of Si, GaN, and AlN. However, we feel that the observation of

a decrease in intensity of the OKLL transition after annealing at > 500°C (see Figure 7.32)

simultaneous with the decrease in intensity and shift to lower binding energy for the C 1s

core level (see Figure 7.18) supports the contention that significant concentrations of C-O

bonded carbon are present on the GaN and AlN surfaces.

235

The ability to use UV/O3 oxidation to completely remove carbon below the detection

limits of AES from GaAs surfaces and not from GaN and AlN surfaces was initially

troublesome. Except for freshly grown surfaces, the UV/O3 exposure was always observed

to reduce the level of carbon detected by both AES and XPS from both GaN and AlN

surfaces. Additionally, the surface carbon C 1s peak was always observed to be shifted to

higher binding energies after the UV/O3 exposure which is indicative of oxidation of the

surface carbon and the formation of more C-O bonds (see Figure 7.22 and Table 7.5).

However as mentioned, complete carbon removal was never achieved and the authors have

accordingly postulated several ideas as to why this is so. It should be first pointed out,

though, that Baunack et al [62] and Fominski et al [57] have previously reported incomplete

removal of carbon from silicon surfaces using UV/O3 oxidation, and we have also observed

this for (0001) 6H-SiC surfaces (see Chapter 3). In the case of GaN, AlN, and SiC, the

inability to completely remove carbon by UV/O3 oxidation may be related to several factors.

The first factor is the extreme chemical inertness of these materials and their resistance to

oxidation which prevents the formation of a complete passivating oxide to prevent carbon

accumulation during sample mounting and transfer. A similarly related factor is the wide

band gap of these materials (especially compared to GaAs). It may be that in addition to the

generation O3, the UV light from the Hg lamp may also create electron hole pairs at the

surface of GaAs which assists in the oxidation of surface carbon via transfer of electrons/hole

pairs from the GaAs to the surface carbon. As the nitrides have much larger bandgaps, fewer

electron/hole pairs may be generated at the surface by the Hg lamp or transfer of electrons

from the Nitride and SiC surface may be not as efficient. Finally in the case of GaN and

AlN, the bonding in these materials is much more ionic than in the case of GaAs and

therefore adventitious carbon may be more strongly bonded or attracted to the nitride

236

surfaces. In summary, the UV/O3 oxidation treatment was found most useful when trying to

remove gross carbon contamination left by photoresist, silver paste, extensive storage in

Fluoroware containers, or (in the case of SiC) polishing treatments.

One further point regarding the oxidation of GaN surfaces that the authors would like

to address, is that as grown GSMBE GaN surfaces were not observed to oxidize rapidly in air

as detailed in Table 7.6. It was generally observed that freshly grown GSMBE GaN surfaces

acquired very few contaminants during overnight storage on a laminar flow bench, and a

similar resistance to contamination was also observed for freshly grown OMVPE GaN films

stored in Flouroware containers (note: Ma et al. [114] have also observed this for fresh

HVPE films). It was also generally observed that the surface carbon initially accumulated on

the nitride surfaces was more closely related to CHx species which was indicated by the

lower binding energy of 284.8-285.0 eV for the XPS C 1s peak from these surfaces.

However, with further aging and processing more C-O bonded carbon was observed to

accumulate on the nitride surfaces shifting the C 1s peak to higher binding energies of 285.6-

286 eV. As previously mentioned, it was sometimes observed that the UV/O3 and wet

chemical cleaning processes described above deposited/accumulated more oxygen and

carbon contaminants on the surface than were previously detected from these surfaces prior

to cleaning. As can be imagined, this greatly complicated our cleaning experiments.

Therefore, most of the experiments described above were conducted on GaN and AlN

surfaces which had been allowed to "age" for several weeks in their Fluoroware containers.

One exception, is the data presented in Table 7.7 for which the GSMBE film was "aged" on a

laminar flow bench for only one day.

237

7.5.2. Wet Chemical and HF Vapor Processing

Ex situ wet chemical cleaning and oxide removal from (0001) GaN and AlN surfaces

with HCl and HF based solutions were observed to produce surfaces of these materials with

the lowest levels of oxygen but with significant concentrations of Cl and F, respectively. It is

interesting to note that in HF and HCl wet chemical processing, F was exclusively detected

on AlN surfaces whereas Cl was exclusively detected on GaN surfaces. This suggests that F

and Cl exclusively bond with Al and Ga atoms at AlN and GaN surfaces respectively. In

fact, this may be expected simply based on comparison of the bond strengths of halogens

with Ga, Al, and N [116]. As Table 7.9 illustrates, the bond strengths of Al-F and Ga-Cl are

much larger than those for N-Cl and N-F. Thus preferential bonding of F and Cl with Al and

Ga over N should be expected based simply on bond strengths. However, XPS of HF

processed AlN surfaces showed two F 1s core levels at 686.5-686.8 eV and 688.3-688.7 eV

which were attributed to both Al-F bonding and N-F bonding. The lower binding energy F

1s peak at ˜ 686.6 eV is in excellent agreement with the value of 686.3 eV previously

reported for fluorine in AlF3.H2O [100]. As the authors are currently not aware of any

reported values for the F 1s core level from NFx species, the assignment of the F 1s peak at

688.5 eV to N-F bonding is based primarily on the observation of a chemically shifted N 1s

core level at ˜ 404 eV from BHF vapor treated OMVPE GaN and AlGaN surfaces (see Figure

7.30).

238

As displayed in Figures 7.28-30 and 7.39, chemically shifted N 1s, Ga 2p3/2, and Al

2p core levels were observed from GaN and AlGaN surfaces treated with a combination

UV/O3-BHF vapor treatment. These chemically shifted core levels were attributed to the

formation of NF3, GaF3 and AlF3 species respectively [117-119]. As only two F 1s peaks

positioned at ˜ 686.5 and 688.5 eV were detected from these surfaces it seems natural to

assign these two peaks to N-F bonding and either Al-F or Ga-F bonding. To further support

our assignment of the higher binding energy F 1s peak at ˜ 688.5 eV to N-F bonding, we note

that for the BHF vapor treated GaN surface the F 1s peak at 686.5 eV and the chemically

shifted Ga 2p3/2 core level were reduced in intensity by rinsing in DI water while the F ls

peak at 688.5 eV and the chemically shifted N 1s peak were not (see Figure 7.28 and 29).

Table 7.9. Bond energies of Cl, F, and H with Al, Ga, and N [116]. Bond Bond Energy (kJ/mol) Al-F 664 Ga-F 577 N-F 238 H-F 566 Al-Cl 511 Ga-Cl 481 N-Cl 381 H-Cl 428 H-H 432 Al-H 285 Ga-H 275 N-H 339

It is important to note, however, that the F 1s peak at ˜ 688.5 eV may also be due to

physisorbed HF. TPD analysis of BHF treated AlN surfaces showed a desorption peak for

both m/e- 20 and 38 at 400°C and 500°C respectively which corresponds will with the

observed decrease in intensity of the 688.5 eV F 1s peak in XPS of fluorinated AlN surfaces.

239

However, the 688.5 eV F 1s was not observed to completely disappear until after

temperatures of 950°C were reached which is were the 686.5 eV was also observed to

disappear. As the authors are not currently aware of any reported values of the binding

energy of the F 1s core level from HF, the authors therefore must leave open the possibility

that N-F and H-F bonded fluorine is present on BHF processed AlN surfaces.

The observation of N-F bonding on HF wet chemically processed AlN surfaces raises

the question of whether some N-Cl bonding is also present on HCl wet chemically treated

GaN surfaces. The observation of both Ga-F bonding and N-F bonding on BHF vapor

treated GaN surfaces also raises the question as to why F is exclusively detected from HF wet

chemically processed AlN surfaces and not from GaN surfaces as well. Unfortunately, we

were not able to detect any chemically shifted Ga or N peaks for HCl wet chemical or vapor

treated GaN surfaces. Further, Cl was extremely difficult to detect with XPS due to week

sensitivity and hence we can't say a lot about what Cl is bonded to on GaN except based on

our previous bond strength arguments. However, we do note that the bond strength of N-Cl

is ˜ 150 kJ/mol larger than the N-F bond strength so N-Cl bonding on GaN should perhaps be

expected based on the observation of N-F bonding on AlN. As for the specificity for F

adsorption on AlN surfaces in HF and Cl adsorption on GaN in HCl, the authors speculate

that this behavior may be actually related to the differences in bandgap between these two

materials. Ohmi [120] has previously explained the ability of HF to hydrogen terminate

silicon surfaces based on the excellent alignment of the H+ ion with the valence maximum of

silicon in HF solutions. This close alignment allows for efficient transfer of electrons and the

formation of a covalent bond. Ohmi also states that the energy position of an ion in an

aqueous solution is a direct function of the electronegativity of the ion with more

electronegative ions having a lower (more negative) energy. As fluorine is more

240

electronegative than chlorine (Pauling electronegativity for F=4 and Cl=3.2), the F- ion

should lie at a lower energy than the Cl- ion. As the bandgap of AlN is much larger than

GaN, it should follow that the VBM of AlN lies below that of GaN. Accordingly, interaction

of F- with the VBM of AlN would be expected whereas the Cl- ion would be expected to

better interact with the VBM of GaN (see Figure 7.40).

GaN AlN1:1 HCl 10:1 HF

H+

OH-

F-SO42-

OH-

F-

K+

O3-

Cl- Cl-

SO42-

EF

Figure 7.40. Schematic illustrating alignment of Cl- and F- ions with VBM of GaN and AlN in 1:1 HCl:DI and 10:1 HF respectively.

One important aspect about the scenario illustrated in Figure 7.40 is that the

termination of GaN and AlN in aqueous solutions may be dependent on the doping level and

type (i.e. the position of the Fermi level). Accordingly, termination of GaN surfaces by H or

OH ions may be more favored for p-type material leading to less chlorine termination. In

+

-

241

contrast however, increased Cl and F termination of GaN and AlN surfaces could perhaps be

alternatively achieved by adjustment and optimization of the pH of HCl and HF solutions

respectively which would control the energy position of Cl and F ions in these solutions. - -

Complete Cl and F termination of GaN and AlN surfaces is desirable as these

contaminants have been observed to desorb from these surfaces at much lower temperatures

than both carbon and oxygen contaminants. In fact for GaN surfaces, complete desorption of

oxides and carbon contaminants was only observed at temperatures where some GaN surface

decomposition was observed to occur (see Figure 7.17-19 and 7.31). In the case of AlN, a

completely fluorine terminated AlN surface is even more desirable as thermal desorption of

oxides from this surface was not even observed to be possible at temperatures of 1000-

1100°C (see Figure 7.9). Additionally, as the bond strengths of N, Ga, and Al with Cl and F

are so strong, tying up dangling bonds at these nitride surface with Cl and F should stabilize

and inhibit re-oxidation of the surface in air. It is also important to note that we have

empirically observed a direct correlation between the halogen surface coverage and the

carbon and oxygen surface coverage on these surfaces. Typically, the larger the halogen

coverage the lower the oxygen and carbon coverage. We have also empirically observed that

larger oxygen coverages correspond to lower carbon coverages and vice versa. This is in

agreement with the previous observations of Ingrey [35] for III-V arsenide and phosphide

surfaces. Ingrey [35], has previously noted that there are a finite number of adsorption sites

on semiconductor surfaces for which oxygen, carbon, and halogen contaminants compete.

Saturation of these sites with one particular specie hinders the adsorption or contamination of

the surface by other species.

Putting the scenario in Figure 7.40 aside, it is interesting to note that HF processes

were generally observed to leave surfaces with fewer carbon contaminants relative to HCl

242

processes. This may be related to the fact that HCl and other processes were conducted in

glass (Pyrex) beakers whereas HF processes were conducted in Teflon beakers (note: HF

etches glass). However, the RCA SC1 clean which was also conducted in a Pyrex beaker

was observed to leave a similar level of carbon contaminants to HF processed GaN and AlN

surfaces. In contrast, the RCA SC2 clean which is an HCl based solution was observed to

always leave a higher carbon coverage. This perhaps suggests that the chemical state of

carbon species in HF are different from those in HCl. In fact, it was observed that the C 1s

peak for surface carbon on HF treated GaN surfaces was typically 0.2-0.4 eV higher in

binding energy than that for HCl treated GaN surfaces. This in turn implies that the HF

processes may leave more C-O bonded carbon whereas HCl processes leave more C-H

bonded carbon. The importance of C-H bonded carbon vs. C-O bonded carbon will be more

fully discussed in the subsequent section on thermal desorption. However, it is worth

mentioning here that combinations of HF and HCl may produce cleaner GaN surfaces due to

enhanced replacement of carbon for chlorine on the GaN surface.

Removal of the UV/O oxide from AlN surfaces using HF was initially surprising

given the known chemical inertness of sapphire (Al O ) [121]. However, there have been

previous reports of etching amorphous aluminum oxide films by HF [121,122] H PO [123],

and CF RIE [124]. Additionally surfaces of Al [95,96], Al O [97], and AlN [98,99] are also

all known to be composed to some extent of aluminum hydroxides (AlO-OH) which may be

more chemically reactive. In fact, our XPS results show that HF attacks primarily the oxide

associated with the higher binding energy O 1s core level which we have previously

attributed to oxygen in a OH chemical state (see Figure 7.2). On the other hand, oxide

removal from GaN surfaces using HCl, HF, and NH OH chemistries was not surprising given

that they already have been previously reported and are commonly used to remove oxides

3

2 3

3 4

x 2 3

-

4

243

from GaAs surfaces prior to metal contact deposition [35-53]. Solutions of 1:1:7

H SO :H O :DI are also used to clean GaAs surfaces, but do so by oxidizing the surface and

contaminants, forming a thin passivating oxide layer on the GaAs surface which is then

thermally desorbed in situ typically prior to epitaxy [40-44]. In our case, the H SO :H O

cleans were observed to remove carbon from the surface but were not observed to form an

oxide layer which again is not suprising given the relative chemical inertness of GaN

compared to GaAs.

2 4 2 2

2 4 2 2

As the AES spectra displayed in Figures 7.27 and 7.38 illustrate, the BHF vapor

treatment does not appear to be very effective for removal of oxides from AlN and GaN

surfaces relative to other wet chemical processes. However as previously mentioned, the

oxide coverage for the BHF vapor treated surfaces appeared to be much lower in XPS.

Although we have no clear explanation for this anomalous behavior, we do note that both

AlF and GaCl are extremely hygroscopic and like to form AlF H O and GaCl H O [117-

119]. Accordingly, these BHF vapor treated surfaces may have retained several monolayers

of physisorbed water on the surface which when excited by the electron beam during AES

analysis may have oxidized the GaN and AlGaN surfaces. In the case of XPS analysis, these

ebeam induced oxidation effects would not be expected to occur. Accordingly, better results

with BHF vapor cleaning may be achieved under parameters where GaF or AlF formation

is prohibited (i.e. by heating substrate or not using a UV/O exposure prior to the vapor

treatment).

. .3 3 3 2 3 2

3 3

3

7.5.3. Thermal Desorption and Capping Layers

244

In a previous paper, we examined the thermal desorption of oxygen and carbon

contaminants on GaN surfaces after wet chemical cleaning in HCl:MeOH, HCl:DI, HF:DI,

and HF:MeOH solutions [85]. It was observed that for all surfaces incomplete removal of

these contaminants was not achieved even after annealing at 800°C [85]. However, thermal

desorption of carbon contaminants at 800°C was observed to be greater for HF treated

surfaces relative to HCl treated surfaces [85]. In particular, the carbon desorption from the

HF:MeOH (MeOH = methanol) treated GaN surfaces was observed to be the most complete

of the wet chemical treatments investigated [85]. These observations suggest perhaps that the

chemical state of the carbon contaminants left on GaN surfaces after HCl and HF wet

chemical processing are somehow different. In fact as mentioned above, we have previously

observed that the binding energy of the C 1s peak for surface carbon on GaN surfaces is

typically located at ˜ 0.4 eV higher binding energy for HF processed GaN surfaces compared

to HCl processed GaN surfaces. In turn, this difference in binding energy for the surface

carbon C 1s suggests that HF wet chemical processes leave more C-O bonded carbon

contaminants on surfaces whereas HCl processes leave more C-H bonded carbon

contaminants. This is important as in XPS examinations of the thermal desorption of carbon

contaminants from both GaN and AlN surfaces, we have observed that the intensity of the C

1s core level decreases and shifts to lower binding energy with higher annealing

temperatures. Generally, after wet chemical processing, the C 1s core level is located at ˜

285-286 eV and shifts to 284-285 eV after annealing at 500-600°C (see Table 7.2 and 7.7 and

Figure 7.18). Ignoring the possibility of band bending effects, this indicates that most of the

C-O bonded carbon desorbs at temperatures = 500-600°C leaving behind only C-H bonded

carbon which apparently desorbs at much higher temperatures. As HF processes tend to

x

245

leave more C-O bonded carbon contaminants, it should therefore not be surprising that

thermal desorption of carbon contaminants at 800°C from these surfaces should be more

complete. Additionally as carbon contaminants left by MeOH should also be expected to be

composed predominantly of C-O bonding, it should not be surprising that the thermal

desorption of carbon contaminants from MeOH:HF treated GaN surfaces is also more

complete at 800°C. These observations and conclusions further emphasize the need for the

investigation of HF:HCl and even HF:HCl:MeOH wet chemical processes for cleaning GaN

surfaces.

Despite all of the above observations, no wet chemical treatments were observed to

have a significant effect on the desorption of oxides from GaN (or AlN) surfaces. Unlike

GaAs, complete thermal desorption of the surface oxide for GaN surfaces was not observed

to occur at 650°C [63,64]. In our case, complete thermal desorption of the surface oxide

from GaN in UHV was not observed to occur until temperatures of 900-950°C where

achieved at which point significant decomposition of the GaN film was observed to occur

with our RGA (see Figure 7.17, and 7.31). Additionally, the XPS spectra of the Ga 2p

core level from a GaN surface annealed in UHV at 950°C (see Figure 7.19) displayed two

core levels perhaps indicating the preferential loss of nitrogen from the surface and hence the

stoichiometry of the GaN film/surface (note: this in contrast to the results of Munir and

Searcy [125] which concluded that GaN decomposes/sublimes congruently).

3/2

The discrepancy between oxide desorption from GaAs and GaN surfaces can be

explained, however, by considering the fact that the oxides from both GaAs and GaN

probably leave the surface as either Ga-O, As-O, or N-O species instead of O . In order for

this to happen, Ga-As or Ga-N bonds must be broken. As the Ga-N bond is much stronger

than the Ga-As bond, higher temperatures will be required to break the Ga-N bond and hence

2

246

desorb the oxide. Therefore it should not be surprising that thermal desorption of oxygen

from GaN surfaces occurs at slightly higher temperatures than for GaAs and that the

temperature at which this occurs is also the temperature at which some decomposition of the

GaN film can be observed. In fact in the case of GaAs, As incongruently sublimes from the

GaAs in UHV at 650°C and GaAs surfaces accordingly have to be annealed in a flux of As in

order to maintain a stoichiometric surface during oxide desorption. Accordingly, annealing

GaN in fluxes of N or Ga may also be necessary to maintain a stoichometric surface during

thermal desorption of the surface oxide (annealing GaN in fluxes of Ga and NH will be

discussed further in the following section). The above line of reasoning also explains the

inability to thermally desorb the oxide from AlN surfaces. The Al-N bond is significantly

stronger than the Ga-N bond and accordingly higher temperatures than were investigated

here would be necessary in order to break the Al-N bond and allow desorption of the oxide as

either Al-O or N-O.

3

As displayed in Figures 7.9 and 7.17, it is interesting to note that the two O 1s core

levels attributed to oxygen in O and OH chemical states merge into one peak of smaller

FWHM at ˜ 800°C. Previous analysis of the thermal decomposition of aluminum hydroxides

(Al(OH) ) using thermogravimetric analysis (TGA), has shown that these materials

decompose into aluminum oxides at temperatures of ˜ 500°C [126]. This is 300°C lower than

the temperature at which we observe the two O 1s chemical states to merge into one and

suggests that perhaps the higher binding energy O 1s peak is due oxygen bonded to

aluminum in some other chemical state than OH . In fact, this higher binding energy O 1s

state could be related to the non-bridging vs. bridging oxygen state seen in silicon oxide films

[127].

2- -

3

-

247

Based on the large bond strengths and the observation that both AlF and GaCl

sublime at 1300°C and 800°C [116-119], the rather high temperature stability of F and Cl on

AlN and GaN surfaces should not be surprising. For F on AlN, complete thermal desorption

of fluorine was not observed until temperatures of 950°C were achieved. However, TPD and

XPS experiments revealed the loss of fluorine from AlN surfaces at much lower temperatures

of T <500°C and which was correlated with desorption of either physisorbed HF or

desorption from nitrogen sites. Desorption of fluorine from AlN at T > 500°C was

correlated mainly with desorption from Al sites, however, desorption of fluorine from N sites

may also occur at T > 500°C. Complete desorption of Cl from GaN was observed to occur

at a slightly lower temperature of ˜ 700-800°C. Due to the inability to detect Cl with XPS

and experimental difficulties with the TPD system, we were not able to correlate Cl

desorption to specific sites (i.e. N or Ga). However as illustrated in Figure 7.32, a significant

decrease in intensity of Cl detected by AES was observed after annealing an HCl vapor

treated GaN film at 450°C. Similar to AlN, Cl desorption from GaN at T < 450°C may be

related to desorption of physisorbed HCl or desorption of Cl from N sites. Based on the Ga-

Cl bond strength, however, it should be expected that Cl desorption from GaN surfaces at

T > 500°C should occur predominantly from Ga sites.

3 3

sub

sub

sub

sub

sub

As for the capping layers on AlN, our results indicate that the success of this idea is

dependent on finding a capping layer which (1) uniformly covers the AlN surface (i.e. no

exposed AlN), (2) can be desorbed at T < 800°C, and (3) can be deposited in situ via either

MBE or OMVPE. Clearly, GaN is too thermally stable to be used for this purpose and

complete coverage with In and other group III (Ga, Th), IV (Sn, Pb), and V (As, Sb, Bi)

elements is likely to always be a problem. Group III-As, P compounds, however, are likely

to exhibit both the lower temperature stability and coverage needed. Unfortunately, most

248

OMVPE and MBE systems do not currently have the capability of depositing both nitride,

phosphide, and arsenide compounds. Therefore, InN may be the best alternative. InN has

been reported to decompose in UHV at ˜ 600°C [128] and should wet/cover the AlN surface.

Unfortunately, growth of pure InN is extremely difficult and has not been frequently

reported. However, In Ga N alloys may work just as well as pure InN capping layers. x 1-x

As for the difficulty in completely desorbing OMPVE GaN capping layers from AlN,

the authors speculate that Ga may be perhaps trapped at the steps of the AlN surface.

Alternatively, residual TMA or TEA may have been left in the OMVPE reactor at the start of

the GaN capping layer growth and which lead to the formation of a thin Al Ga N layer at

the AlN/GaN interface and which is stable at > 1000°C. In fact, the authors note that the

problem of complete GaN desorption was not observed for GaN films grown on AlN buffer

layers by GSMBE. In this case, there is no chance of intermixing of Al and Ga during

growth. Interdiffusion of the AlN and GaN during growth is another alternative but unlikely

for the growth temperatures used in this study (˜1000°C).

x 1-x

7.5.4. Chemical Vapor Cleaning and H Plasma Processes

As mentioned above, annealing GaN surfaces in a flux of Ga or N in order to

maintain a stoichiometric surface may be a more appropriate method for removal of surface

oxides as opposed to simple thermal desorption where some decomposition and or

sublimation of the GaN film is required in order to completely remove the oxide. Annealing

GaN surfaces in fluxes of reactive gas/vapor species such as Ga, N , and NH , may also be +2 3

249

beneficial in that these species may assist in the removal of carbon and oxygen via reacting

with the contaminants and forming more volatile species which desorb at lower temperatures.

In fact, it is now common to anneal silicon and silicon carbide surfaces in reactive fluxes of

Ga, Ge, Si, SiH , Si H , and GeH to assist both in maintaining a stoichiometric surface as

well in removing the surface oxide [129-135]. Further, Kahn et al [79] and Bermudez [80-

82] have demonstrated cleaning of GaN via annealing of GaN surfaces in a flux of Ga.

However, our results have shown that annealing in a NH flux can also be used for obtaining

an atomically clean GaN surface. This approach is more closely analogous to flux cleaning

of other III-V compounds where it has been found necessary to anneal in a flux of the group

V component during oxide desorption in order to counteract incongruent sublimation of the

group V component [65].

4 2 6 4

3

Figure 7.33 illustrates that NH was found to be effective for removal of surface

carbon at temperatures < 600°C which is especially impressive in comparison to thermal

desorption in which carbon removal was incomplete even at temperatures of 800°C.

Ammonia has been previously reported to be an excellent scavenger of hydrocarbons [137].

In Figure 7.34, it is also shown that annealing GaN in 5x10 Torr NH3 at 800°C completely

removes carbon from the GaN surface and leaves only a submonlayer coverage of oxygen.

Although the submononlayer coverage of O in AES can be attributed to electron beam

induced oxidation of the GaN surface during the AES analysis as indicated in Figure 7.34

[136], small traces of oxygen were almost always detected by XPS from GaN surfaces

annealed in NH at 800-900°C. However, NH cleaned GaN surfaces were observed to

display (2x2) reconstructions in LEED. The (2x2) reconstruction has only been previously

observed from as grown RF and ECR MBE samples [138,139].

3

-6

3 3

250

UPS analysis of the electronic structure of NH cleaned GaN surfaces also indicates

no band bending at the GaN surface. For NH cleaned, undoped, OMVPE GaN samples the

VBM was observed to be located at ˜ 3.0 eV below the system Fermi level (see Figure 7.35).

Given that one would expect the Fermi level of undoped GaN to lie 0.1-0.2 eV below the

GaN CBM (E = 3.4 eV), this result indicates an essentially flat band condition for NH

prepared GaN surfaces. This is in contrast to the UPS results obtained by Bermudez [82] for

GaN surfaces prepared via annealing in a Ga flux. In this case, Bermudez observed the GaN

VBM to be located at ˜ 2.4 eV below E indicating significant band bending at the surface

[82]. This result is in much better agreement with our observation of E -VBM = 2.7 eV for

as grown (2x2) GSMBE GaN films examined in situ (see Figure 7.35) It is important to note,

however, that in the case of GaN surfaces prepared by Ga flux cleaning, Bermudez observed

the presence of surface states in his UPS spectra [82] which had previously lead him to

conclude that the GaN VBM was 0.5 eV closer to E [81]. In our case, surface states were

not clearly visible in our UPS spectra and assuming even that we are overestimating the

VBM by 0.5 eV as Bermudez previously did indicates that the (2x2) as grown surface would

be flat band and the NH cleaned OMVPE surfaces are inverted. At this point it is important

to note photovoltage effects could be strongly influencing our results and leading to our

discrepancies with the results of Bermudez. In fact, Bermudez has previously reported the

observation of a photovoltaic shift of ˜ 0.2 eV for the GaN VBM and Ga 3d core level for

GaN surfaces irradiated with a high pressure Hg arc lamp [81]. Finally, we note that

Bermudez observed that the GaN VBM was ˜ 3.0 eV below E for GaN surfaces prepared by

only wet chemical cleaning in 1:10 NH OH:DI.

3

3

g 3

F

F

F

3

F

4

Our plasma cleaning results indicate that H plasmas are very efficient for in situ

removal of halogen and carbon species at temperatures of 450°C which is 300-400°C lower

251

than the temperature required to desorb these contaminants in UHV or annealing in a flux of

NH or Ga. Halogen removal from silicon surfaces using atomic H has been shown to be via

an Ely-Rideal mechanism [110]. In this mechanism, atomic H is able to extract halogens

from silicon without being thermally accommodated at the surface. The H atom can

accomplish this due to the 220 kJ/mol (relative to 1/2H2(g)) of excess potential energy

residing in the H atom. In the case of GaN and AlN, a detailed study has not yet been made

to ascertain whether Ely-Rideal is operable in halogen extraction from these surfaces.

However in Figure 7.12, it was shown that a small amount of F was still present on the AlN

surface after H plasma exposure at 450°C. Given the large flux of atomic H (1016/cm2s

[109]) produced in the plasma, complete removal of F from AlN would be expected if Ely-

Rideal were operable.

3

In one of the ground breaking papers by Nakamura et al [140] it was shown that

annealing p-type GaN in ammonia led to high resistivity material due to compensation of the

p-type dopants with hydrogen. Therefore, one natural concern with regard to cleaning GaN

surfaces via annealing in NH or a H plasma is whether these processes will lead to

compensation of p-type dopants in GaN films. H plasma processes have been previously

shown to lead to compensation of other III-V compounds (InGaAlP and InGaP [141]).

Additionally, Pearton et al [142-143] have intentionally used an ECR H plasma at 250-

400°C to implant H into GaN, AlN, and InN to study the effect of compensation of these

materials. For ECR plasma exposure at 250°C they have observed significant compensation

of both n and p type material. As such H plasma cleaning processes should perhaps operate

at temperatures > 500°C in order to avoid incorporation of H in the subsurface of GaN films.

In the case of NH , Nakamura et al [141] has previously shown that annealing p-type GaN in

3

3

252

1 atm. NH will lead to compensation of the p-type dopants. However for the NH cleaning

processes we have described here, a much lower flux of ammonia was used and hence

hydrogen incorporation and p-type dopant compensation should be much less. In fact, we

note that Kim et al [144] have recently achieved the growth of p-type GaN via NH -GSMBE

with out post growth annealing. In this case the NH flux and temperature used to grow the

GaN is equivalent to the NH flux and temperature we have used to clean the GaN surface.

3 3

3

3

3

Finally, it is also interesting to examine a few cases of homoepitaxial growth on GaN.

In the first such report by Gassmann et al [145], GaN was grown on bulk GaN single crystals

by MBE. In this case the GaN crystal was annealed/cleaned at only 675°C where MBE GaN

growth occurred. Correspondingly, cross sectional TEM showed the presence of a 50 nm

highly defective area at the interface between the MBE GaN film and the GaN single crystal.

As we have shown, thermal desorption at 675°C still leaves a significant amount of carbon

and oxygen on the GaN surface. In contrast, Ponce et al [146] did not observe such a

defective layer in cross section TEM of the interface between a GaN epilayer grown by

MOCVD on a GaN single crystal. In this case, Ponce et al [146] did not specify their surface

cleaning/preparation except that growth occurred at 1050°C. In our thermal desorption

studies, we have shown that thermal desorption at 950°C in UHV is sufficient for complete

desorption of carbon and oxygen from the GaN surface. However, Ponce et al [146] still

note the presence of dislocations originating at the interface between the GaN film/wafer

interface which they attribute to non-optimized surface prep conditions.

7.6. Conclusions

253

Of the numerous acids and bases examined, HF and HCl solutions produced AlN and

GaN surfaces with the lowest coverages of oxygen respectively. However using AES and

XPS, significant amounts of F and Cl were detected on these surfaces after dipping in HF

and HCl, respectively. It is hypothesized that these halogens tie up dangling bonds at these

nitride surfaces hindering re-oxidation of the surface. Fluorine was very thermally stable

requiring temperatures of > 850°C for desorbtion. Remote H plasma exposure was effective

for removing halogens and hydrocarbons from both AlN and GaN surfaces at temperatures of

450°C, but was not particularly efficient for oxide removal. Annealing GaN in NH3 at 700-

800°C produced clean as well as stoichiometric GaN surfaces.

7.7 Acknowledgments

The work described herein was supported by the ONR under contract N00014-91-J-

1416. Appreciation is expressed to Cree Research, Inc. for the 6H-SiC wafers.

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263

8. X-ray Photoelectron Spectroscopy analysis of GaN/(0001)AlN and

AlN/(0001)GaN

Growth Mechanisms


to the


by

Sean W. King, William G. Perry, Eric P. Carlson, Robert J. Therrien, Robert J. Nemanich,

and Robert F. Davis



Raleigh, NC 27695

264

8.1. Abstract

The growth mechanisms of GaN on (0001) AlN and AlN on (0001) GaN have been

investigated using x-ray photoelectron spectroscopy (XPS), low energy electron diffraction

(LEED), and Auger electron spectroscopy (AES). It has been found that GaN growth on

(0001) AlN at low temperatures (650-780°C) occurs via a Stranski-Krastanov 3D type

growth mechanism. The GaN on (0001) AlN growth mechanism, however, switches to a

Frank van der Merwe/layer by layer type growth mechanism at higher temperatures

(>800°C). AlN growth on (0001) GaN was observed to occur via a FM/layer by layer

growth mechanism in the temperature range investigated (750-900°C). We propose a model

based on the interaction of atomic hydrogen with GaN/AlN surfaces which shows that the

surface kinetics of hydrogen desorption/ammonia decomposition is the determining factor for

the GaN growth mechanism.

265

8.2. Introduction

GaN and AlN are completely miscible semiconductors with wide band gaps of 3.40

and 6.2 eV, respectively. Many potential applications for these materials including UV-

visible/optoelectronics, high-power, high-frequency, and high-temperature electronic devices

have been recently realized [1-3]. To date, the two most popular techniques for growth of

these materials has been organometallic vapor phase epitaxy (OMVPE) and electron

cyclotron resonance-molecular beam epitaxy (ECR-MBE). The success and improved

understanding of the OMVPE and ECR-MBE techniques has been greatly aided by several

studies of the growth mechanisms of GaN and AlN on various substrates by these techniques

[4-10]. An alternative to OMVPE and ECR-MBE growth of GaN, is reactive molecular

beam epitaxy (RMBE) which essentially replaces an ECR N2 source with NH3 in an MBE

system (i.e. NH3-GSMBE). NH3-GSMBE is currently gaining increased attention due to its

inherent simplicity relative to ECR-MBE and the improved electrical and optical properties

of GaN films grown by this technique [11-14].

The first reports of growth of AlN, GaN, and AlxGa1-xN alloys by NH3-GSMBE

were those by Yoshida et al [15-19]. They reported successful growth of highly resistive

single crystals of AlN (as determined by reflection high energy electron diffraction

(RHEED)) on (0001) and (11-20) Al2O3 substrates at temperatures of 1000 and 1100°C

respectively [15,16]. The GaN films grown at 700°C on (0001) Al2O3 were also single

crystalline though conductive with high carrier concentrations of 1019-1020/cm3 and

mobilities of 30 cm2/Vsec [17]. Use of an AlN buffer layer for growth of GaN on Al2O3,

266

did not significantly improve the electrical properties of their films. However, Yoshida et al

did make the key observation that the band edge cathodoluminescence (CL) intensity from

their GaN films grown on AlN/Al2O3 were 25 times more intense than those films grown

directly on Al2O3 [18,19].

Since the preliminary work by Yoshida et al, a steady improvement in the quality of

GaN films grown by NH3-MBE has been achieved. GaN films grown directly on (0001)

Al2O3 by Powell et al [20] at 760-780°C were found to exhibit carrier concentrations as low

as 1-4x1018/cm3 and mobilities as high as 100-110cm2/Vsec. A further reduction in carrier

concentrations to 2x1017/cm3 was additionally demonstrated by Yang et al [11] and Kamp

[14] et al by the use on an AlN buffer layer on (0001) Al2O3 [11,14]. Finally, through

optimization of the NH3 flux, Kim et al [13] have been able to grow at 850°C highly resistive

GaN films with carrier concentrations < 1014/cm3 and mobilities as high as 200 cm2/Vsec

[13]. The reduction of the background carrier concentrations to these levels, has additionally

allowed Kim et al and Yang et al to recently achieve Mg p-type doping of GaN without post

growth annealing [11,13].

As improvements and increased understanding of OMVPE and ECR-MBE

techniques were aided by several studies on the growth mechanisms of GaN and AlN on

various substrates [4-10], it should be expected that further improvement in NH3-GSMBE

growth of GaN and AlN should also be aided by such studies. In this paper, we have used

surface analytical techniques such as x-ray photoelectron spectroscopy (XPS), Auger electron

spectroscopy (AES), and low energy electron diffraction (LEED) to study the initial growth

mechanisms of GaN on AlN and AlN on GaN by the NH3-GSMBE technique. We show

267

that at low growth temperatures (Tsub<800°C) GaN growth on AlN proceeds via a Stranski-

Krastanov (SK) growth mechanism (2D-3D) and at higher temperatures (Tsub>800°C) GaN

growth on AlN proceeds via a Frank van der Merwe (FM)/layer by layer mechanism. The

change from a SK to a FM growth mechanism is attributed to a transition from a fully

hydrogen terminated GaN surface to a partially hydrogen terminated surface with increasing

temperature. AlN was observed to grow on GaN in a FM/layer by layer mechanism

throughout the temperature range investigated.

8.3. Experimental

8.3.1. Thin Film Growth and Analysis System


(UHV) configuration which integrates several completely independent UHV thin film growth

and analysis systems via a 36 ft. long transfer line having a base pressure of 9x10-10 Torr

(see Ref. 21 for details of the transfer line, and many of the associated systems). The

experiments described in this paper employed the III-V nitride gas source molecular beam

epitaxy (GSMBE), Auger electron spectroscopy (AES), low energy electron diffraction

(LEED), and x-ray photoelectron spectroscopy (XPS) systems.

The GSMBE system with a base pressure of 3x10-10 Torr was designed and

constructed specifically for the growth of III-V nitride thin films. The sample heating stage

consisted of a wound tungsten heating filament positioned close to the back of the sample

268

and mounted on a boron nitride disk [21]. A W/6%Re-W/26%Re thermocouple was

employed to measure the temperature of the backside of the wafer. Heating profiles to

1100°C were easily achieved using a programmable microprocessor and 20 amp SCR power

supply. Actual surface/sample temperatures (i.e. those reported herein) were measured using

an infra-red thermometer with a spectral response of 0.8 to 1.1 µm and a emissivity setting of

0.5. The experimental accuracy for the substrate temperatures was estimated to be ± 25°C.

Source materials in the GSMBE included Al, Ga, and NH3. Al (99.9999%) was

evaporated from a 25 cc "cold lip" Knudsen cell and Ga (99.99999%) was evaporated from a

25 cc dual filament Knudsen cell. The NH3 (99.9995%) was further purified via an inline

purifier connected directly to a leak valve mounted on the GSMBE chamber. Sample

exposure to the NH3 was obtained using "molecular beam" dosers similar to design of Yates

et al [22]. Collimation of the ammonia into a molecular beam focused onto the sample was

achieved with this doser using a 13 mm diameter x 2 mm thick glass capillary array with a

ten micrometer pore size (Galileo Electro Optics Inc.). The doser to sample distance was

fixed at 2". This doser arrangement enhanced the ammonia flux to the sample by a factor of

10-100 relative to the background ammonia flux. Analysis of the ammonia using a

quadrapole residual gas analyzer (RGA) revealed extraneous peaks at 28 and 44 indicating

that CO, N2, and CO2 were the principal contaminants in the gas.

The XPS experiments were performed in a stainless steel UHV chamber (base

pressure = 2x10-10 Torr) equipped with a dual anode (Mg/Al) x-ray and a 100 mm


were obtained using Mg Kα radiation (hν = 1253.6 eV) using 12 kV and 20 mA emission


269


was achieved by periodically recording scans of the Au 4f7/2 and Cu 2p3/2 peaks from


peaks (83.98 and 932.67 eV, respectively) [23]. Curve fitting of most data was performed


with a linear background was found to best represent the.

The Auger electron spectrometer and the low energy electron diffraction optics were

mounted on a six way cross off the transfer line and pumped through the transfer line. In the

AES analysis, a 3 keV, 1mA beam was used. Each Auger electron spectrum was collected

and numerically differentiated. In LEED an 80 eV, 1mA beam was used.

8.3.2. Substrate and Thin Film Preparation and Analysis

The substrates used in this research were ≈ 1.5x1.5 cm2 pieces cut from 1 3/16"

diameter on and off-axis (4° toward (11-20)), n-type (Nd=1018/cm3) 6H-SiC (0001)Si

wafers obtained from Cree Research, Inc. All wafers were received with an ≈ 1 µm n-type

epitaxial layer (Nd=5x1017) on which was grown ≈ 500-1000Å of thermal oxide. After

removal of the thermal oxide with 10:1 HF, the unpolished back side of each wafer was

subsequently coated via RF sputtering with tungsten to increase the heating efficiency of the

SiC, as the latter is partially transparent to the infrared radiation emitted from the tungsten

filament heater. All wafers were then ultrasonically rinsed in acetone and methanol, exposed

to the vapor from a 10:1 buffered HF solution for 10 min, and mounted using Ta wire to a 1"

270

diameter Mo disk with an approximately 1.5 cm2 square hole cut in the center. Each

wafer/Mo assembly was then fastened to a ring shaped Mo sample holder using Ta wire and

inserted into the transfer line load lock. The in situ procedure used for the final cleaning step

of the 6H-SiC substrates was similar to that described by Kaplan and Kern et al [24,25] and

is described in detail elsewhere [26]. Briefly, each SiC wafer was annealed in the GSMBE

system in a flux of 10-6-10-5 Torr SiH4 for ≈ 15-20 min at 950-1050°C. Analysis via AES

and XPS revealed oxygen-free, silicon terminated SiC surfaces which displayed either (1x1)

or (3x3) LEED patterns. If a (3x3) LEED pattern was obtained (indicative of the formation

of a ˜ bilayer of free silicon on the surface [24,26]), the sample was annealed in UHV at

1050°C for 5-10 min. to desorb the excess silicon. This procedure resulted in a (1x1) LEED

diffraction pattern. Growth of the AlN buffer layer was always initiated on an oxygen free

(1x1) 6H-SiC (0001) surface.

To initiate the deposition of AlN, the 6H-SiC wafer was raised to a temperature of

1050°C at which point the shutter to the Al Kcell (at 1150°C) was opened. A few seconds

later, ammonia was admitted into the system which created a total pressure of ≈ 10-5 Torr.

Growth proceeded at a rate of 1000 Å/hr. for approximately 15 min. after which the Al cell

was shuttered. The sample was allowed to cool in ammonia until approximately 800-900°C

at which point the ammonia valve was closed. The AlN films displayed (2x2) reconstructed

surfaces in LEED immediately after growth. This reconstruction was sensitive to either

contamination or temperature, as a (1x1) LEED pattern was observed several hours after

growth. SEM analysis showed the films to be free of surface topography at 10 kX. The

films were too resistive for electrical measurements. Additional details regarding this

research have been published elsewhere [27,28].

271

To achieve the growth of the GaN films on the AlN buffer layer, the latter was heated

to 650-800°C in 10-4 Torr (˜ 50 sccm) ammonia for 10 min., after which the Ga cell (at

1020°C) was opened and growth allowed to proceed. The growth rate for these conditions

was determined via cross sectional SEM to be ≈ 2000 Å/hr. After the desired GaN thickness

had been achieved, the Ga cell was closed and the GaN film allowed to cool in ammonia to

approximately 600°C after which the ammonia valve was closed. Films of AlN were also

deposited at 800°C and 10-5 Torr NH3 on undoped (0001) GaN films previously grown by

both GSMBE and organometallic vapor phase epitaxy (OMVPE) [29]. Other parameters

were similar to those used in the growth of the AlN buffer layer. The OMVPE GaN films

were cleaned by annealing in 10-4 Torr NH3 at 800°C for 15 min. Auger and x-ray

photoelectron spectroscopies did not detect the presence of any oxygen or carbon [30].

8.3.3. Growth Mode Analysis

The experimental procedure and analysis used to study the growth modes of GaN

(AlN) on AlN (GaN) was similar to that used by Sitar et al. [31] to study the growth modes

of AlN and GaN on (0001)Si 6H-SiC and (0001) Al2O3. Therefore, only a brief description

of the procedure will be given here.

A series of sequential depositions each having a thickness of ≈ 0.5-5Å of GaN (AlN)

on AlN (GaN) were conducted until a GaN (AlN) film thickness of ≈ 35Å was achieved.

Following each deposition, XPS, LEED, and AES analysis were performed. Each series of

depositions and analysis was completed within 12-14 hours to ensure the cleanliness of the 272

surfaces and interfaces. To determine the growth mode of the GaN (AlN) film, the ratio of

the initial integrated intensity of the Al 2p (Ga 3d) core level from the AlN buffer layer

(OMVPE GaN film) (Io) was measured against the integrated intensity of the Al 2p (Ga 3d)

core level from the GaN/AlN interface (Is). Is/Io was plotted against the calculated GaN

thickness (= growth rate x growth time). Theoretical curves for the expected Al 2p (Ga 3d)

attenuation from layer-by-layer (Frank van der Merwe (FM)), layer-by-layer plus island

(Stranski-Krastanov (SK)), and island (Volmer-Weber (VW)) growth modes were

simultaneously plotted and compared with the experimentally determined attenuation to

elucidate the growth mode(s). The following relations were used for FM, SK, and VW

growth modes:

FM: Is/Io = exp (-t/λ) (1)

where:

t = thickness of the growing film

λ = mean free path of the photoelectron being measured

Io = Initial intensity of the substrate core level

Is = Intensity of the substrate core level with a overlying film of thickness t

SK: Is/Io = (1-θ) + exp(-t/λ) (2)

where θ is the surface coverage of the film/islands.

VW: Is/Io = (1-θ)exp(-q/λ) + exp(-t/λ) (3)

where q = thickness of the film before onset of 3D growth 273

The following relation from Briggs and Seah [32] was used to calculate the mean free

paths of the core levels of interest:

λ = 0.41(aE)1/2 + 538E-2 monolayers (4)

E = Kinetic energy of the photoelectron

a = (ρMw/NA)1/3 where ρ = density of overlying film

Mw = Molecular weight of film

NA = Avogadro's Number

From the above relation, the mean free path of Al 2p photoelectrons in GaN was determined

to be 19Å, and the mean free paths of Ga 3d, 3p, and 2p photoelectrons in AlN were

determined to be 19, 18 and 6Å respectively.

8.4. Results

8.4.1. Growth of GaN films.

Figure 8.1 shows a plot of the attenuation of the Al 2p core level as a function of the

overlying GaN film thickness for growth at 650°C. Curves for the expected attenuation for

Frank van der Merwe (FM), Stranski-Krastanov (SK), and Volmer-Weber (VW) growth

modes are also included. The Al 2p core level intensity actually increases after the first few

274

GaN depositions. This effect is believed to be due to forward scattering effects where the

trajectory of a photoelectron emitted by an underlying atom is scattered toward the direction

of the overlying atom due to the positive charge of the nucleus of the latter [33,34]. This can

cause an increase in the intensity of the core levels along certain crystallographic directions

[33]. However, the forward scattering effect diminishes with further depositions, and the

experimental data starts to follow the curve expected for FM growth. However, at a GaN

thickness of 10-12 Å the Al 2p attenuation starts to exhibit a positive deviation again from

the curve for FM growth and approaches the curve expected for SK growth.

Unreconstructed (1x1) LEED patterns were displayed throughout the entire sequence

of GaN depositions at 650°C (see Fig. 8.2(a)). The oxygen levels measured by AES and

XPS throughout the series of 650°C GaN depositions were found to be either undetectable or

less than 1% of a monolayer. The microstructure of the surface of the GaN films grown at

650°C is shown in Figure 8.3(a). These films were n-type and extremely conductive,

exhibiting four point probe sheet resistances of 10-2 ?/sq which is indicative of free carrier

concentrations of 1019/cm3 or greater [35]. Photoluminescence spectra of these films

showed very broad donor bound exciton emission of weak intensity (see Figure 8.4(a)).

275

0.20

0.40

0.60

0.80

1.00

1.20

0 5 10 15 20 25 30

I/Io

I/Io FM

I/Io VW

I/Io SK

I / I

o

GaN Thickness (Å)

Figure 8.1. Attenuation of Al 2p core level from AlN buffer layer as a function of overlying GaN film thickness for T =650°C. Filled circle = experimental I/I , filled diamond = theoretical I/I for Frank van der Merwe layer by layer growth, filled triangles = theoretical I/I for Volmer Weber and Stranski-Krastanov growth.

sub o

o

o

(a) (b) Figure 8.2. LEED diffraction patterns from (a) (1x1) (0001) GaN, and (b) (2x2) (0001) GaN.

276

(a) (b)

(c)

Figure 8.3. SEM micrographs at 10 kX from GaN films grown in GSMBE at (a) 650°C, (b) 750°C, and (c) 800°C.

277

(a)

(b)

Figure 8.4. Photoluminesence (PL) at 4K of NH -GSMBE GaN grown at (a) 650°C and (b) 800°C.

3

278

Figure 8.5 shows a plot of the Al 2p attenuation as a function of GaN thickness for

films grown on AlN at 800°C. The SEM micrograph in Figure 8.3(b) shows a very smooth

surface containing occasional "pits" similar in appearance to those observed in thin OMVPE

GaN surfaces which have been speculated to be due to the incomplete coalescence of flat

GaN islands [29]. The RMS surface roughness of this film as determined by AFM was ˜ 40Å

which is comparable to the < 20Å typically measured from OMVPE GaN surfaces.

0

0.2

0.4

0.6

0.8

1

1.2

0 5 10 15 20 25 30 35 40

I/Io (Al 2p)

I/Io (FM Theory)

o

GaN Thickness (Å)

Forward ScatteringEffects

Figure 8.5. Attenuation of Al 2p core level from AlN buffer layer as a function of overlying GaN film thickness for T =800°C. Filled circle = experimental I/I , filled triangle = theoretical I/I for Frank van der Merwe layer by layer growth.

sub o

o

279

Reconstructed (2x2) LEED patterns were displayed throughout the series of GaN

depositions on AlN at 800°C (see Fig. 8.2(b)). These surfaces were similarly sensitive to

surface contamination and temperature as with the (2x2) reconstructed AlN surface.

Typically, (1x1) LEED patterns developed after 3-4 hours in a vacuum of 10-9 to 10-8 Torr.

The (2x2) reconstructed surfaces could be restored by annealing in 10-5 Torr NH3 for 5 min.

No attempts were made to determine the actual structure of the (2x2) reconstruction.

However, the work of Bernholc et al [36] indicate that a (2x2) N adatom reconstruction is

energetically most favorable for Ga terminated (0001) GaN and a (2x2) N vacancy

reconstruction is energetically favorable for the nitrogen terminated (0001) GaN surface. For

the experiments described heree, gallium termination. For AlN growth on the silicon face of

6H-SiC (0001), one expects the formation of Si-N bonds at the AlN/SiC interface which

inturn implies Al termination of the AlN film. Accordingly, Ga termination should be

expected for GaN growth Al terminated AlN. The oxygen levels detected by both AES and

XPS from GaN films grown at 800°C were similarly found to be undetectable or less than

1%.

The GaN films grown at 800°C were found to be more resistive (0.2-1M?) in contrast

to those grown at 650°C. Sheet resistances for the former were typically too high for four

point probe measurement. The carrier concentrations (ND-NA) of the 800°C films

determined by CV measurements were in the range of 1-5x1017/cm3. GaN films grown at

825°C were more resistive with a carrier concentration of 2-5x1016/cm3. However, Hall

measurements found n = 2x1017/cm3 and µ = 60cm2/Vsec. PL for these films displayed

sharp (4 meV) donor-bound exciton emission and very little D-A emission. (see Figure

8.4(b)). 280

8.4.2. Growth of AlN films on GaN.

A plot of the attenuation of the Ga 3p, 2p, and 3d core levels from an OMVPE GaN

as a function of overlying AlN thickness is shown in Figure 8.6. As can be seen, FM growth

of AlN on OMVPE GaN was observed to occur (similar results were obtained for AlN

growth on GSMBE GaN). No surface topography was observed on the AlN surface via

SEM, but the AlN/GaN was observed to have cracked on occassion. The cause of the

cracking is currently not known. The resistivity of these AlN films were beyond the range of

our experimental capabilities and CV did not indicate any charge.

0.0

0.2

0.4

0.6

0.8

1.0

1.2

0 5 10 15 20 25 30 35 40

I/Io Ga 3d

3d Theory

I/Io Ga3p

3p Theory

I/I o

AlN Thickness (Å)

Figure 8.6. Attenuation of Ga 3d and 3p core levels from OMVPE GaN as a function of overlying GaN film thickness for Tsub = 800°C. Filled circle = experimental I/I for Ga 3d, filled square = experimental I/I for Ga 3p, empty circle and square = experimental I/I for Frank van der Merwe layer by layer growth.

o

o o

281

8.5. Discussion

8.5.1. GaN on AlN Growth Mechanisms

8.5.1.1 Strain Effects:

The above XPS CL attenuation analysis shown in Figures 1 & 5 clearly illustrates

that the growth of GaN on monocrystalline AlN by NH3-GSMBE initially proceeds in a

FM/layer by layer mechanism for the temperature range investigated. However using XPS

and SEM, we have observed that growth can either continue in a FM/layer by layer fashion

or switch into a three dimensional SK type growth regime. Further inspection of Figure 1

reveals that the deviation from the expected Al 2p attenuation for FM growth and the

experimentally observed Al 2p attenuation occurs at ≈ 10-12Å which is also reported

Matthews-Blakeslee critical thickness for GaN on AlN [37]. This observation is clearly in

line with the classical interpretation of SK type growth where interfacial strain forces 3D

growth.

The differences between the different types of growth modes (FM, SK, and VW) are

typically explained in terms of macroscopic surface and interfacial energies. Classically,

Frank van der Merwe (FM) growth is predicted when the sum of the film surface energy (σf)

and the substrate-film interfacial energy (σi) is less than the substrate surface energy (σs) (i.e.

σf +σi < σs). Volmer-Weber (VW) growth is predicted for the opposite case, i.e. when σf +

σi > σs. Prediction of Stranski-Krastanov (SK) growth requires a more detailed analysis of

the substrate-film interfacial energy, σi. Several different factors are lumped together in σi.

282

The two most important factors contributing to the interfacial energy are: σb, the energy due

to the chemical bonds formed between the two materials at the interface, and σst, the strain

formed at the interface due to imperfect lattice matching. The interfacial energy, σi, can then

be described as σi = σst + σb. Thus, Stranski-Krastanov growth is described by the situation

where the expression σf + σb + σst < σs is satisfied initially at the start of film growth.

However, as growth continues the interfacial strain energy term increases dramatically and

switches the equality of the expression to σf + σb + σst > σs. At this point, the interfacial

stain energy can be lowered by either generation of misfit dislocations (as envisioned by

Matthews and Blakeslee [38-40]) or switching into three dimensional growth (i.e. surface

roughening [41]). Clearly, stress relaxation via switching to 3D growth is clearly in line with

our observations of SK type growth of GaN at 650°C.

In contrast to GaN growth at 650°C, no deviations between the Al 2p attenuation and

theory for FM growth was observed for GaN growth at 800°C. Correspondingly, SEM and

AFM analysis of thicker films (2000-5000Å) revealed smoother surfaces with RMS

roughness ˜ 40 Å consistent with FM/layer by layer growth. Clearly for strain relaxation via

surface roughening to be the cause of SK type GaN growth at 650°C, a different strain

relaxation mechanism must be in operation during GaN growth at 800°C. A different strain

relaxation mechanism at 800°C could occur as a result of a lowering of strain due to better

lattice matching between GaN and AlN at 800°C. However, the thermal expansion

coefficients (c and a) for GaN are both larger than those of AlN, and hence the lattice

matching and strain is actually poorer at 800°C [42].

283

Clearly, an alternative explanation is that the interfacial strain at the GaN/AlN surface

is relieved via misfit dislocation generation. The switch from strain relaxation via surface

roughening to strain relaxation via misfit dislocation would imply that some sort of thermal

activation is required to nucleate a misfit dislocation. Owing to the large moduli of the III-V

nitrides, this is reasonable as the energy to create a dislocation is quite high (i.e. dislocation

energy ≅ Gb2, where G=E/2(1+ν) an EGaN=300 GPa [43]). Therefore, one could

alternatively argue that at 800°C the GaN film grows psuedomorphically on the AlN and that

at the MB critical thickness strain relaxation occurs via misfit dislocation generation rather

than via surface roughening. However, the authors note that the underlying AlN buffer layer

although thin (≈250Å) and monocrystalline is thicker than the critical thickness for AlN on

6H-SiC (≈40Å) and is highly defective with a large density of misfit dislocations which

certainly propagate into the GaN film. Accordingly, there are already plenty of defects at the

GaN/AlN interface to act as low energy nucleation sites for misfits in the GaN film. (In fact

in low strain SixGe1-x alloy growth on Si, misfits are found to preferentially nucleate at

dislocations and other defects at the Si surface [44]).

It is worth noting that a switch from 3D to 2D GaN growth with substrate

temperature for NH3-GSMBE growth of GaN directly on (0001) Al2O3 has additionally

been observed using RHEED by Powell et al [20]. Their RHEED studies showed that films

grown at Tsub < 760°C exhibited well defined transmission spots indicative of three

dimensional island growth and at Tsub > 780°C, RHEED exhibited sharp Kikuchi lines and a

streaky (1x1) patterns indicative of two dimensional growth [20]. As the lattice matching

and interfacial bonding between GaN and Al2O3 are completely different to the case of GaN

on AlN, one would not expect to see a switch from 3D to 2D growth in the same temperature 284

range if strain relaxation mechanisms were responsible for the changes in growth mechanism.

This in turn suggests that surface processes such as adsorption, desorption, and diffusion are

playing an important role in determining the growth mechanism of GaN on AlN in NH3-

GSMBE. In fact, in Si and Ge GSMBE growth using SiH4, Si2H6, and GeH4 a transition

from 3D to 2D growth with increasing substrate temperature has also been observed [45-49].

In this homoepitaxial case, the transition from 3D to 2D growth was explained in terms of

sluggish hydrogen desorption kinetics/sight blocking at low growth temperatures [50]. At

this point, the authors provide a simple model for the interaction of NH3 with GaN surfaces

in MBE and which shows that hydrogen desorption/ammonia decomposition may also be the

determining factor for the NH3-GSMBE GaN growth mechanism on all substrates including

Al2O3, AlN, Si, and GaN.

8.5.1.2. Hydrogen desorption

Based on RGA studies of the surface cracking of ammonia, Kamp et al [14] proposed

the following sequences of reactions for describing the reaction of ammonia with GaN

surfaces:

NH3g ---- NH3ad (i) NH3ad ---- NH2ad + Had (ii) NH2ad +Had ----- NHad + 2Had (iii) NHad + 2Had ------ Nad + 3Had (iv) 2Had --- H2g (v) 2Nad --- N2ad (vi) - No growth

285

Gaad + Nad --- GaNs (vii) - GaN growth

For simplicity, we first start by modeling reaction (v), i.e. adsorption and desorption

of hydrogen from GaN surfaces. Our approach is to estimate the steady state surface

coverage of hydrogen on GaN and AlN surfaces in the presence of an atomic hydrogen flux.

This will allow us to estimate the hydrogen surface coverage during NH3-GSMBE growth in

the case where ammonia decomposition on GaN and AlN surfaces is rapid and not the

limiting reaction. The authors also note that this model is a reasonably accurate description

for the interaction of GaN/AlN surfaces with hydrogen plasmas in which there is known to be

a large concentration of atomic hydrogen [51].

Desorption kinetics are typically described by the general Polyani-Wigner rate

expression [52]:

rate = -dθ/dt =νnnexp(-Edes/RT) (5) where: n = the reaction order θ = the adsorbate surface coverage ν = the pre-exponential factor,νo=1028/cm2 sec,ν1=1013/sec,

ν2=10-2 cm2/sec Edes = the activation energy for desorption

In steady state conditions, the flux of adsorbates leaving the surface via desorption will be

equal to the incoming flux of adsorbate times the adsorbate sticking coefficient. The sticking

coefficient is described by:

S = So(1 - θ/θmax)n (6)

286

The combination of equations 5 and 6 allows the determination of the steady state surface

coverage of the adsorbate [53].

For hydrogen on GaN, desorption of hydrogen can be expected from both Ga and N

sites. Chiang et al [54] using isothermal analysis and direct recoil time of flight (TOF) mass

spectroscopy have recently determined that deuterium desorption from Ga sites on Ar+

sputter cleaned polycrystalline GaN surfaces is second order with an activation energy of 9

kcal/mol. The findings of Chiang et al for H desorption from Ga are similar to those found

for hydrogen desorption from Ga in GaAs where a second order desorption activation energy

of 13 kcal/mol was found [55]. Second order desorption is consistent with the following

mechanism:

2Dads ---- GaD2 -----D2(g) + Ga. (viii)

Chiang et al speculate that the rate limiting reaction in this expression is actually diffusion of

the deuterium to Ga sites and that D2 desorption is actually rapid. Unfortunately, Chiang et

al did not provide any information regarding hydrogen desorption from nitrogen sites except

that hydrogen desorption from nitrogen sites is complete by 600°C. Despite this we have

estimated that hydrogen desorption from nitrogen in GaN is second order with an activation

energy of 55 kcal/mol. This estimation is based on the similar bond energies of hydrogen to

nitrogen and silicon in ammonia and silane (431.8 and 398.3 kJ/mol respectively [56]) and

the determination of second order desorption kinetics of hydrogen from Si (111) with Edes =

57.5 kcal/mol by Schulze and Henzler [57]. (Note: desorption of hydrogen from nitrogen

sites using 1st order desorption kinetics with Edes=45 kcal/mol were also calculated). To the

287

authors knowledge, there have been no reports on Edes for hydrogen desorption specifically

from a nitrogen site to support this assumption. However, hydrogen desorption sites from

polycrystalline MoN with Edes ≅ 50 kcal/mol have been reported by Choi et al [58-60]. For

hydrogen desorption from AlN we considered hydrogen desorption from nitrogen sites in

AlN to be similar to hydrogen desorption from nitrogen sites on GaN surfaces and hence

assumed Edes ≅ 50 kcal/mol. For hydrogen desorption from Al sites on AlN we used the

experimentally determined activation energies of hydrogen desorption from metallic (111) Al

surfaces which were 17 kcal/mol for 0th order desorption kinetics[61,62]. In our

calculations, we assumed an activation energy of 17 kcal/mol and 2nd order desorption

kinetics.

Using Edes(Ga) = 9 kcal/mol, Edes(Al) = 17 kcal/mol, Edes(N) = 55 kcal/mol,ν2=

10-2cm2/sec, So=1 and Ømax = 2.3x1015/cm2, we were able to generate plots of hydrogen

surface coverage on Ga, Al, and N sites of GaN and AlN surfaces as a function of

temperature for various atomic H fluxes (see Figures 8.7, 8.8, 8.9). As can be seen in Figure

8.7, very little hydrogen remains adsorbed to Ga at any temperature except for extremely

large fluxes of atomic H (1024/cm2sec.). However, for H fluxes of 1017-1018/cm2sec (i.e.

typical NH3 MBE conditions) hydrogen is stable on Al sites on AlN to temperatures of

100°C (see Figure 8.9). In contrast, saturation of nitrogen sites with hydrogen is observed to

occur for all fluxes up to temperatures of 400-550°C (see Figure 8.8). These figures also

show that desorption of hydrogen from nitrogen sites is nearly complete at 600°C for low H

fluxes. This is consistent with the observations of Chiang et al. [55] and lends support to our

assumptions for the activation energy for desorption of hydrogen from nitrogen. Further

288

inspection of Figure 8.8, reveals that for atomic H fluxes of 1017-1018/cm2sec. hydrogen

desorption does not become significant until temperatures of 700-800°C. This is of

importance in that H fluxes of this magnitude are comparable to the NH3 fluxes used in MBE

growth of GaN. Therefore, it is not unreasonable to expect similar hydrogen surface

coverages for GaN surfaces exposed to an ammonia flux. In fact, the authors argue that the

simple model we have presented here for the interaction of atomic hydrogen with a nitrogen

site is exactly analogous to the case of H desorption/decomposition of chemisorbed NH3 (i.e.

reactions iii-v).

0

0.2

0.4

0.6

0.8

1

0 200 400 600 800 1000

10e1510e1710e1810e2010e2210e24

Ø/Ø

max

Temperature (ÞC)

Edes

= 7 kcal/mol, 2 nd Order

v = 0.02 cm 2/sec

Figure 8.7. Hydrogen surface coverage of Ga sites as a function of temperature and flux.

289

0

0.2

0.4

0.6

0.8

1

300 400 500 600 700 800 900 1000

10e1510e1710e1810e2010e22

Ø/Ø

max

Temperature (ÞC)

Edes

= 50 kcal/mol, 1 st Order

v = 10 13 /sec

Figure 8.8. Hydrogen surface coverage of N sites as a function of temperature and flux.

0

0.2

0.4

0.6

0.8

1

0 200 400 600 800 1000

10e1510e1710e1810e2010e2210e24

Ø/Ø

max

Temperature (ÞC)

Edes

= 17 kcal/mol, 2 nd Order

v = 0.02 cm 2/sec

Figure 8.9. Hydrogen surface coverage of Al sites as a function of temperature and flux.

290

Based on the above findings, it now seems reasonable to ascribe the observed change

in GaN growth mechanism from SK to FM with increasing growth temperature as being due

to reactive site blocking by incomplete desorption (decomposition) of hydrogen (ammonia)

from GaN surfaces. At temperatures below 700°C, the GaN surface is so saturated with

hydrogen that the hydrogen effectively blocks reactive sites at steps, kinks, etc. preventing

incorporation of Ga into the GaN lattice. As unincorporated Ga may not be able to desorb

from the GaN surface at these temperatures Ga clusters may form leading to 3D growth

(alternatively, the unincorporated Ga clusters could grow to form Ga droplets which have

also been observed [13,20]). At higher temperatures (Tsub>780°C), hydrogen

desorption/ammonia decomposition is more complete leaving more open reactive sites with

which to incorporate the incoming Ga and thus preventing Ga cluster formation.

Additionally, it is intriguing to note that figure 8.8 shows that for higher atomic H/NH3

fluxes of 1020-1022/cm2sec characteristic of CVD growth, hydrogen saturated surfaces are

maintained up to temperatures of 950-1050°C. Correspondingly, this is the temperature

range in which OMVPE growth of high quality GaN is observed to occur. This suggests that

the need for high temperatures in CVD growth of GaN is not to decompose GaN per se but to

speed the kinetics up to the point that the surface can handle the large fluxes of species

impinging on it.

Figures 8.7 and 8.8 may additionally explain the extremely low activation energies

for Ga desorption from GaN observed by Jones et al [63] during NH3-MBE growth. For

growth in the temperature range of 625-740°C, Jones et al using desorption mass

spectroscopy (DMS) observed the activation energy for Ga desorption was 1.4 eV (32.2

291

kcal/mol). At higher temperatures (740-875°C), a lower activation energy for Ga desorption

of 0.4 eV (9.2 kcal/mol) was observed. Both of these values are extremely low in

comparison to those observed E for Ga desorption from GaAs which are observed to range

from 2.8 eV (64.5 kcal/mol) to 4.9 eV (112 kcal/mol) [64-67] These values correspond to the

activation energies for evaporation of Ga from liquid Ga and sublimation of Ga from GaAs

respectively [64]. Jones et al explained their low values as being to due to the presence of a

"hydrogen terminated liquid Ga pool which is microscopic in size" at low temperatures and a

hydrogen terminated nitrogen site at high temperatures. From Figure 8.7, it obvious that in

typical NH3-MBE growth conditions, hydrogen termination of Ga sites is not likely. Further

at higher temperatures where they observe a change in the activation energy, Figure 8.8

shows that there is very little hydrogen termination of either Ga or N sites. Instead, we

suggest that at high temperatures Ga desorption occurs as GaHx species which would be

expected to have the very low activation energies observed (i.e. Edes for AlHx on Al (111) =

27 kcal/mol). At lower temperatures, Ga desorption is an average of desorption as GaHx and

desorption from Ga droplets (i.e. (2.8 + 0.4)/2 = 1.6 ≈ 1.4 eV).

act

Finally, the authors recognize that several other detailed studies of the growth

modes/mechanisms of GaN and AlN have been previously performed [4-10] Unfortunately,

most of these studies were concerned with growth of AlN and GaN on other substrates such

as Si, Al2O3, and MgO, or low temperature/amorphous OMVPE/CVD GaN/AlN buffer

layers deposited on Al2O3. Further, most of these studies were conducted using other

growth methods such as OMVPE/CVD and ECR-MBE both of which are sufficiently

different from our growth method. With this is consideration, the authors do not feel that

these studies merit further discussion.

292

In summary, it has been shown that the growth mechanism of GaN on AlN is

dependent on growth temperature switching from SK type growth to FM type growth with

increasing temperature. Based on previous observations in Si/Ge GSMBE [45-49], we

provide a model which indicates that the surface kinetics of hydrogen desorption/ammonia

decomposition may be the determining factor in the change of growth kinetics. These

findings indicate that in order for 2D growth of high quality GaN to be achieved at

temperatures lower than Tsub ≈ 780°C, time must be allowed for hydrogen to desorb from

the surface. Hence, migration enhanced or ALE type growth schemes may be necessary for

growth of GaN in this regime.

8.5.2. AlN growth mechanisms on (0001) GaN

As shown in Figure 8.6 AlN growth on GaN is observed to occur in a layer by layer

fashion. This complements the observations by Sitar et al [30] in which it was observed that

ECR-MBE AlN grows in a layer by layer fashion on both Al2O3 and SiC. By the same

method, we have additionally observed 2D FM/layer by layer growth of AlN on (0001) 6H-

SiC by NH3-GSMBE .

8.5.3. Surface Reconstructions

Typically reports of most RHEED analysis of GaN films has consisted of the

observation of either "spotty" or "streaky" unreconstructed (1x1) GaN surfaces. Only, a few

reports of the observation of (2x2) reconstructed GaN surfaces have been made [68-71], and

293

until recently, all such reports of (2x2) reconstructions came from GaN films grown on

(0001) 6H-SiC substrates by either ECR or RF MBE (perhaps indicating the better quality of

epitaxy achieved on these substrates). However, Iwata et al have recently reported the

observation of (2x2) and (4x4) reconstructed GaN surfaces grown on (0001) sapphire using

an improved ECR source [71]. Our observation of (2x2) reconstructed GaN surfaces is the

first such case for GaN films grown via NH3-MBE. In our case, the observation of (2x2)

reconstructions coincided with a corresponding decrease in surface roughness for our GaN

films. Based on this, the authors are skeptical of the claims of layer by layer growth and

"smooth" GaN surfaces obtained by other RF, ECR, and NH3-GSMBE in which (1x1)

RHEED patterns were observed.

8.6 Conclusions

The following conclusions can be drawn:

1.) Growth of GaN at 650°C on AlN occurs via a Stranski-Krastanov (SK) growth

mechanism resulting in a relaxed GaN film.

2.) Growth of GaN at 800°C on AlN occurs via a Frank van der Merwe growth

mechanism resulting in a strained GaN film.

3.) Growth of AlN at 800°C on GaN occurs via a Frank van der Merwe growth

mechanism (but apparently does not result in a strained film).

294

4.) The surface kinetics of hydrogen desorption/ammonia decomposition is the

determining factor in which type of GaN growth mechanism occurs.

8.7. References

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33. W.F. Egelhoff, Crit. Rev. in Solid State and Material Sciences 16 213 (1990). 34. S.A. Chambers, Surface Science Reports 16 261 (1992). 35. A. Skfo, R. Thegra, P. Epeap, E. At, J. Irrep. Results, 57 530 (1996). 36. J. Bernholc, Spring MRS Symp. 1996. 37. S. Krishnankutty, R.M. Kolbas, M.A. Khan, J.N. Kuznia, J.M. Van Hove, and D.T. Olson, J. Electron. Mater. 21 437 (1992). 38. J.W. Matthews and A.E. Blakeslee, J. Cryst. Growth 27 118 (1975). 39. J.W. Matthews and A.E. Blakeslee, J. Cryst. Growth 29 273 (1975). 40. J.W. Matthews and A.E. Blakeslee, J. Cryst. Growth 32 265 (1975). 41. A.G. Cullis, MRS Bulletin 21 21 (1996). 42. H.P. Maruska and J.J. Tietjen, Appl. Phys. Lett., 15 327 (1969). 43. I. Akasaki and H. Amano, in Properties of Group III Nitrides, J.H. Edgar, Ed., Inspec., London (1994). 44. F.K. LeGoues, MRS Bulletin 21 38 (1996). 45. S.M. Mokler, W.K. Liu, N. Ohtani, and B.A. Joyce, Appl. Phys. Lett. 59 3419 (1991). 46. S.M. Mokler, W.K. Liu, N. Ohtani, and B.A. Joyce, Appl. Surf. Sci. 60/61 92 (1992). 47. S.M. Mokler, W.K. Liu, N. Ohtani, and B.A. Joyce, J. Cryst. Growth 120 290 (1992). 48. K. Sakamoto, H. Matsuhata, K. Miki, T. Sakamoto, J. Cryst. Growth 157 295 (1995). 49. L. Chen, T. Chou, W. Tsai, G. Huang, H. Tseng, H. Lin, and C. Chang, Jpn. J. Appl. Phys. 34 L869 (1995). 50. S. Gates, and S. Kulkarni, Appl. Phys. Lett. 60 53 (1992). 51. J. Barnack, Foo Foo, and G. Rabbit, Ph.D. Dissertation, NCSU (1997). 52. D.H. Parker, M.E. Jones, and B.E. Koel, Surface Science 233 65 (1990). 53. M. Schulberg, M. Allendorf, and D. Outka, Surface Science 341 262 (1995).

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298

9. XPS MEASUREMENT OF THE SiC/AlN BAND-OFFSET AT THE

(0001) INTERFACE

Presented at the 1995 Fall MRS Conference

III-V Nitride Symposium

by

Sean W. King , Mark C. Benjamin , Robert J. Nemanich , Robert F. Davis , * ** ** *

and Walter R.L. Lambrecht†

*Department of Materials Science and Engineering

**Department of Physics


Raleigh NC 27695

†Department of Physics

Case Western Reserve University

Cleveland, OH 44106

299

9.1 Abstract

X-ray photoelectron spectroscopy has been used to determine the band-offset at the

SiC/AlN (0001) interface. First, the valence band spectra were determined for the bulk

materials and analyzed with the help of calculated densities of states. Core levels were then

measured across the interface for a thin film of 2H-AlN on 6H-SiC which allowed the

determination of a band offset of 1.4 ± 0.3 eV. The analysis of the discrepancies between

measured peak positions and densities of states obtained within the local density

approximation provides information on self energy corrections in good agreement with

independent agreement with independent calculations of the latter.

9.2. Introduction

Silicon carbide wafers are being used increasingly as substrates for the growth of III-

V nitride thin films. In particular, SiC is rather closely lattice matched to AlN (?a/a = 0.9%)

which is often used as a buffer layer for GaN growth. The availability of bulk 6H-SiC

substrate wafers of high quality is instrumental for this purpose. Since SiC can also be grown

on AlN layers on SiC [1], one may also consider the use of SiC as an active quantum well

layer in a AlN/SiC/AlN heterostructure device. From both points of view, the band-offset at

the SiC/AlN interface is of obvious interest. To date, only two previous values are available:

a theoretical value by Lambrecht and Segall [2] which was for the (110) interface between

zincblende SiC and AlN; and an experimental value obtained indirectly from measurements

o

300

of the Fermi level of 2H-AlN grown on 6H-SiC (0001) by Benjamin et al. [3]. The

investigations described here provide a more direct experimental determination of the band

offset at the basal plane interface between 6H-SiC and wurtzite AlN. The procedure consists

of measuring the core levels at the interface between a thin film of AlN (0001) grown on a

6H-SiC (0001) substrate and separately determining the energy of the valence band edges

with respect to the core levels for thick films. Calculated densities of states were used to aid

in the determination of the valence band edge and allowed us to obtain additional information

on the electronic structure of the materials. In particular, we obtained results for the

difference in the quasi-particle self energy shifts of the N 2s and C 2s bands with respect to

those of the upper N 2p and C 2p like valence bands.

9.3. Experiment

A unique and integrated ultra high vacuum (UHV) system consisting of a 36 ft. long

UHV transfer line to which several thin film deposition and surface analysis units were

connected was employed in this research. The details of this integrated system have been

previously described [4]. The as-received, n-type (N ˜ 10 /cm ), vicinal 6H-SiC (0001)

substrate wafers containing a one micron thick, n-type (N ˜ 10 /cm ) 6H-SiC (0001)

epitaxial layer were sequentially dipped in 10% HF for 5-10 min. to remove the thermally

grown 750Å silicon oxide surface layer, rinsed in 18.4 M? de-ionized water, blown dry with

N , mounted to a molybdenum sample holder, loaded into ultra high vacuum (UHV) and

degassed at 250, 500, 700, and 900°C for 30 min. each and annealed in a 10 -10 Torr flux

of silane at 950°C for ˜ 20 min. X-ray photoelectron spectroscopy (XPS) and Auger electron

spectroscopy (AES) analyses of the SiC surface revealed that oxygen and non-carbidic

17-18 3d Si

16-17 3

-7

d Si

2

-6

301

carbon had been removed below the detection limits of these instruments. LEED displayed a

sharp (1x1) pattern. The SiC bulk core levels and valence band of the SiC were measured via

XPS.

Each AlN film was subsequently grown via gas-source molecular beam epitaxy

(GSMBE), at 700°C and 10-6 Torr total pressure using ULSI (99.9995%) NH and a flux of

high purity Al (99.999%) evaporated from a Knudsen cell at 1050°C as sources. The

temperature of 700°C was chosen to minimize any reaction between the SiC substrate and the

AlN. In order to prevent the formation of Si N at the SiC/AlN interface, the SiC wafer was

exposed to the Al flux for 5 min. at 700°C prior to the introduction of NH into the system.

Very thin films (10-20Å) were deposited to investigate the AlN/SiC heterojunction/interface.

Thicker films (200Å) were then deposited to measure the bulk AlN core levels and valence

band. The films were then subsequently transferred within a UHV environment to the

chambers containing the XPS, AES, and LEED units for analyses of the surface chemistry

and structure. Analysis via AES and XPS indicated that the films were stoichiometric and

contained < 5% ML of surface oxygen. LEED displayed a sharp (1x1) pattern. Further

details of the growth and cleaning procedures are described elsewhere [5]. All XPS analysis

was performed using the Al anode (hν = 1486.6 eV) at 20 mA, and 12 kV (240W). Due to

the inherently poor signal/noise ratio in XPS valence band spectra, 50 or more scans of this

region were acquired and summed together. All AES spectra were taken using a beam

voltage of 3 keV and an emission current of 1 mA. LEED was performed using rear view

optics, a beam voltage of ˜ 100 eV, and emission current of 1 mA. Calibration of the XPS

binding energy scale was performed by measuring the position of the Au 4f and shifting the

spectrums such that the peak position occurred at 83.98 eV.

3

3 4

3

7/2

302

9.4. Theory

The densities of states (DOS) used for the analysis of the valence band spectra were

calculated using the linear muffin-tin orbital [6] and density functional methods in the local

density approximation (LDA) [7]. It is important to realize that strictly speaking the band

structures obtained in this theory are not corresponding to the energies for extracting an

electron from the material as measured in photoemission. The latter are quasiparticle

energies and differ from the LDA Kohn-Sham eigenvalues by a self energy correction [8].

This is, among other things, responsible for the well known underestimate of the bandgaps by

the LDA. While these corrections have been found to be rather insensitive to the structure

[9], they are expected to depend on the amount of localization of the states involved. We will

show below that these corrections shift the C 2s and N 2s bands from the LDA calculated

positions with respect to the valence band edge. The available calculations of these

corrections using the GW method (i.e. using the leading term in Hedin's many body theory

[10] with G the one-electron Green's function and W the screened Coulomb interaction) for

SiC [11-13] and AlN [14] show that they are about constant (but not quite, see below) over

the upper valence band but are discontinuous across the ionicity gap. The present

comparison between LDA calculated DOS and measured valence band spectra confirms this

picture, and can be used to obtain an experimental value for these self-energy corrections.

9.5. Results

A comparison of the valence band spectra for 6H-SiC (0001) measured from the

substrate and the calculated DOS is presented in Figure 9.1. The latter is shown unbroadened

Si

303

as obtained from the highly accurate tetrahedron method and with a Gaussian broadening by

0.5 eV. The calculated and measured DOS are aligned to each other in the upper valence

band region. The major atomic orbital character of these peaks is indicated. The reference

level in these spectra is based on the Au 4f standard and is thus not directly related to any

intrinsic materials property of SiC. Thus only relative energy differences are meaningful.

The major reason for using the comparison to calculated DOS is that this allows for a more

precise determination of the actual valence band edge. First, we note that good agreement is

obtained between theoretical and calculated peak positions to about 10 eV binding energy.

The peak intensities of the spectra are influenced by matrix elements and details of the

experimental set-up such as collection solid angle and emission angle with respect to the

surface normal, not accounted for by the DOS. These intensities also depend slightly on the

background subtraction procedure. Here a linear background subtraction was used. The

recent discovery of the presence of surface states in the SiC band gap may also confuse the

direct location of the valence band maximum [15].

7/2

304

Figure 9.1. XPS spectra (arbitrary units) and theoretical densities of states (in states per unit cell per eV) of 6H-SiC.

The major requirement for an accurate band-offset determination is E - E where E

is the energy of the valence band maximum (VBM) and E is any core level of SiC.

However, the broadening near the valence band edge hinders an unambiguous determination

of this edge. In this case, it is preferable to determine the energy separation of the core levels

from the well defined C 2p-Si 3s peak and take the position of the VBM with respect to that

peak from the calculation. If we make use of the calculated GW corrections, the alignment

can be done even more accurately as explained below.

v c v

c

On closer inspection, it has been observed that the experimental C 2s peak is shifted

by about 1.0 eV to lower energy from its theoretical position and the C 2p-Si 3s peak is

shifted by about 0.4 eV. The reason for aligning the spectra in this manner is precisely the

self energy effects mentioned above. Indeed, the GW calculations by Rohlfing et al. [11] and

305

Backes et al. [12] predict a 0.4 eV shift from the valence band for the X3v eigenvalue in 3C-

SiC, which in that case, is closely associated with the C 2p-Si 3s peak. Wenzien et al. [13]

obtain only a 0.2 eV for this shift, so that this must be considered the approximate

uncertainty for this alignment procedure. The former calculations reveal that the C 2s band

self-energy correction is approximately 1.0 eV larger than that of the valence band maximum

while the latter shows this correction to be 1.4-1.6 eV. The present measurement indicates

that the former two are in better agreement with experiment. In summary then, we find that

the valence band maximum on the energy scale of Figure 9.1 lies at 2.2 ± 0.2 eV. We also

see that this agrees well with a simple straight line extrapolation from the half height point of

the experimental valence band edge.

Figure 9.2 shows a similar analysis for the valence-band spectrum of AlN. In this

case, we note that the upper valence band width is only 6 eV wide (in the theoretical DOS).

Rubio et al's [14] GW calculations predict a 2.0 eV shift for the N 2s peak and a 0.6 eV shift

for the N 2p-Al 3s like peak. We see that if we align one of these including the above

correction, the shift for the other is well reproduced. Due to the insulating nature of AlN,

some shifting of the AlN core levels and valence band spectra may expected due to charging.

However, the AlN was deposited thin enough (200Å) such that electrons tunneling from the

conducting SiC substrate into the AlN should minimize and make negligible any charging

affects. For these reasons, our indirect approach of comparing experimental valence band

data with the theoretical DOS is justified. With respect to the same Au 4f based reference

level, we then find the valence-maximum of AlN lies at 4.1 eV and is again in good

agreement with a direct straight line extrapolation of the edge. The core levels for the bulk

and interface system shown in Table 9.1 have also been measured on the same energy scale

as employed for the studies described above.

7/2

306

Figure 9.2. XPS spectra (arbitrary units) and densities of states (in states per unit cell per eV) of 2H-AlN

These data are then substituted into the expression

-?E =(E - E ) - (E - E ) + (E - E ) (1) SiC v v c

SiC AlN AlN SiC b v c b c c

AlNi

where the subscripts b and i indicate bulk and interface respectively. While in the above, E

and E are all positive electron binding energies, it is customary to give ?E in terms of the

energy levels which are the negative of the binding energies. Hence the minus sign on the

left of Eq.1. Using different core levels and remembering the uncertainties in the alignment

of each valence band spectrum we finally arrive at a value of ?E = 1.4 ± 0.3 eV.

v

c v

v

307

Table 9.1. Valence-band maxima and core levels measured on the same Au 4f7/2 based reference scale.

bulk SiC Si 2p 101.3 C 1s 283.5 VBM 2.2 bulk AlN Al 2p 75.3 N 1s 398.5 VBM 4.1 AlN/SiC heterojunction Si 2p 101.5 C 1s 283.6 Al 2p 74.9 N 1s 398.2

9.6. Discussion

The value obtained for the band-offset is in quite good agreement with the previously

calculated offset of 1.5 eV for the (110) zincblende interface [2]. This is perhaps somewhat

surprising since the latter is a non-polar interface while here we deal with a polar

heterovalent interface. In fact, from simple electron counting rules, one expects that a purely

N terminated surface would have an excess of 1/4 electron and thus must reconstruct its

surface for example by having one N vacancy every 4 N atoms in order to maintain charge

neutrality. In reality one may have a missing dimer every 4 instead of a simple vacancy or

any other arrangement which is equivalent in net charge. At present, it is not known on an

308

atomic scale what the interface structure is like, but we may note that if 1/4 of the N are

missing at the interface, this is for electron counting purposes equivalent to mixing the

nitrogen layer with C (group IV) anions. One expects that this would lower the dipole from

that of a non-polar interface [16] by a few 0.1 eV. However, a slightly larger degree of

intermixing may completely wipe out this interface dependence. This is probably indirectly

indicative that there are an equal number of Al-C and N-Si bonds at the interface.

The value of 1.4 eV obtained for the SiC/AlN band offset is larger than the

previously reported experimental value of 0.8 eV [3]. The discrepancies between these two

values may be related to the experimental techniques or to the preparation of the SiC/AlN

interface. The experiments described in this report employed recently developed surface

preparation processes that result in atomically clean and ordered SiC prior to AlN deposition.

In contrast, the SiC surface preparation of the prior study would typically exhibit a small

amount of oxygen at the interface ˜ 25-50% ML. In addition, the gas source MBE employed

in the present study results in a higher quality interface as opposed to the ECR technique

employed previously. The ECR technique has been shown to result in more damage and an

excess of Si-N bonding at the interface [17]. In the study presented in this paper, no oxygen

was detected at the SiC/AlN interface. Furthermore, the use of an initial Al flux prior to

ammonia exposure avoids the formation of a large amount of Si-N bonding. These factors

can strongly influence the band alignment between two semiconductors. In addition, the

initial study also assumed flat bands in the SiC near the interface while upward band bending

was noted as a distinct possibility. Since that band-offset value was based on the assumption

of alignment of the measured Fermi level of AlN and the bulk n-type doped SiC, it indeed

probes the macroscopic band alignment (affected by band bending) rather than the offset in

309

the immediate vicinity of the interface. This differs from the present XPS investigation

because of the limited escape depth of the photoelectrons.

9.7 Conclusion

By combining XPS studies of the valence band spectra with calculated DOS and a

careful analysis of the alignment between the two, taking into account known self-energy

corrections to the LDA band structures, the position of the valence band maxima of the 6H-

SiC (0001) and 2H-AlN (0001) with respect to their core levels has been determined. A

subsequent measurement of core levels at the heterojunction between a thin film of 2H-AlN

grown on top of 6H-SiC then allowed us to extract a band offset of 1.4 ± 0.3 eV. The latter is

in good agreement with the calculated value of the (110) zincblende SiC/AlN which indicates

that polar interface specific effects were averaged out by atomic level reconstructions of the

interface leading to an equal amount of Si-N and C-Al interface bonds.

9.8 Acknowledgments

The work at CWRU was supported by NSF (DMR-92-22387): the research at NCSU

was supported by ONR (NOOO14-92-J-1477). Appreciation is expressed to Cree Research

Inc. for the 6H-SiC wafers used in this study.

9.9 Addendum

310

Since the presentation of this paper, several additional experiments have been

conducted to examine interface effects on the band alignment between 2H-AlN and 6H-SiC.

As we have previously determined the position of the SiC core levels relative to the 6H-SiC

VBM and the AlN core levels relative to the AlN VBM, measurement of the valence band

discontinuity for various different interface conditions requires only measuring the difference

between the AlN core levels and the SiC core levels at the SiC/AlN interface. Table 9.2

summarizes the Si2p-Al2p results obtained from thin AlN films grown on 6H-SiC substrates

under various conditions. In these experiments, both on and off axis 6H-SiC substrates were

used, and both Al and NH3 pre-exposures were used to initiate growth of AlN on the SiC

substrates. For consistency, all the measurements reported here were taken from 20Å

AlN/6H-SiC interfaces. As this table displays, values for Si 2p - Al 2p were observed to

range from ˜ 26.5 - 27.0 eV. This in turn implies a variation in the 2H-AlN/6H-SiC valence

band discontinuity of ˜ 0.5 eV. This is somewhat consistent with the recent first principle

calculations by Majewski et al [18] which showed a variation in the valence band

discontinuity for the (001) SiC/AlN interface of 0.9 eV (?Ev = 1.5 - 2.4 eV). This variation

was found to correspond to differences in the bonding at the SiC/AlN interface with the

purely mixed C/N interfaces resulting in a discontinuity of 1.5 eV and a mixed Al/Si interface

resulting in a discontinuity of 2.4 eV. Unfortunately, Table 8.2 does not display a clear trend

between the substrate orientation, surface reconstruction, or pre-growth orientation.

Additionally, our calculations indicate the valence band discontinuity to range from 0.9-1.4

eV.

311

Table 9.2. Various measurements of Si2p-Al2p from a AlN/6H-SiC interface for different substrates, surface reconstructions, and pre-growth treatments. Orientation Reconstr. Al/NH3 Tsub(TC) Si2p-Al2p (0001), off (1x1) Al 950°C 26.6 (0001), off (1x1) Al 1250°C 26.6 (0001), on, with epi (1x1) Al 950°C 26.6 (0001), off (1x1) Al 950°C 26.8 (0001), off (1x1) Al 950°C 26.95 (0001), off (1x1) Al 1350°C 26.96 (0001), on, wit epi (1x1) Al 1400°C 26.95 (0001), off (1x1) NH3 1250°C 26.45 (0001), on (3x3) NH3 1250°C 27.3 (0001), on (3x3) NH3 1250°C 26.5 (0001), off (vx3v3)R30° NH3 1250°C 26.7 (0001), off (1x1) NH3 1250°C 26.7 (0001), off (3x3) NH3 1250°C 26.9 (0001), on, with epi (v3xv3)R30° NH3 1250°C 26.7 (000-1), on (1x1), Si NH3 1250°C 26.5 (000-1), on (1x1), graphiteNH3 1250°C 26.8 (1-100) (1x1) NH3 1250°C 27.0 (1-100) (1x1) Al 1250°C 27.0

9.10 References

1. L.B. Rowland, R.S. Kern, S. Tanaka, and R.F. Davis, Appl. Phys. Lett. 62, 3333 (1993). 2. W.R.L. Lambrecht and B. Segall, Phys. Rev. B 43, 7070 (1991). 3. M.C. Benjamin, C. Wang, R.F. Davis, and R.J. Nemanich, Appl. Phys. Lett., 64, 3288 (1994). 4. Jacob van der Weide, Ph.D. dissertation, North Carolina State University, 1993. 5. S.W. King, R.J. Nemanich, and R.F. Davis, unpublished. 6. O.K. Andersen, O. Jepsen, M. Sob, in Electronic Band Structure and its

312

Applications, ed. M. Yussouf (Springer, Heidelberg 1987), p. 1. 7. W. Kohn and L. J. Sham, Phys. Rev. 140, A1133 (1965). 8. L.J. Sham and W. Kohn, Phys. Rev. 145, A561 (1966). 9. W.R.L. Lambrecht, B. Segall, M. Yoganathan, W. Suttrop, R.P. Devaty, W.J. Choyke, J.A. Edmond, J.A. Powell, and M. Alouani, Phys. Rev. B 50 10722 (1994). 10. L. Hedin, Phys. Rev. 139, A796 (1965). 11. M. Rohlfing, P. Kruger, J. Pollmann Phys. Rev. B 48, 1791 (1993). 12. W.H. Backes, P.A. Bobbert, W. van Haeringen Phys. Rev. B 51, 4950 (1995). 13. B. Wenzien, P. Kackell, F. Bechstedt Phys. Rev. B 52, (1995). 14. A. Rubio, J.L. Corkill, M.L. Cohen, E.L. Shirley, and S.G. Louie, Phys. Rev. B 48, 11810 (1993). 15. M.C. Benjamin, S.W. King, R.F. Davis, and R.J. Nemanich, unpublished. 16. W.R.L. Lambrecht and B. Segall, Phys. Rev. B 41, 2832 (1990). 17. Z. Sitar, L.L. Smith, and R.F. Davis, J. Cryst. Growth 141, 11 (1994). 18. J.A. Majewski, M. Stadele, and P. Vogl, to be published in the proceedings of the Fall 1996 Materials Research Society.

313

10. Dependence of (0001) GaN/AlN Valence Band Discontinuity on

Surface Reconstruction

and Growth Temperature

Presented at the Fall 1996 MRS Conference

in the III-V Nitride Symposium

by

Sean W. King, Mark C. Benjamin*, Robert J. Nemanich*, and Robert F. Davis




Raleigh NC 27695

314

10.1 Abstract

X-ray and ultraviolet photoelectron spectroscopies have been used to determine the

heterojunction valence band discontinuity at the (0001) GaN/AlN interface. Discontinuity

values of 0.6±0.2 eV were determined for GaN grown on AlN at 650°C, and 0.9±0.2 eV for

GaN grown on AlN at 800°C.

10.2. Introduction

The semiconductor compounds of GaN and AlN are completely miscible with band

gaps of 3.40 and 6.2 eV, respectively. Many potential applications based on heterostructures

and bandgap engineering of these two materials and their alloys have been recently realized

including UV-visible/ optoelectronics and high-frequency devices [1-3]. Charge transport

and quantum confinement are a few of the many important parameters which can affect these

devices [4]. Therefore, reliable knowledge of the valence band discontinuity, ?Ev for the

GaN/AlN interface is extremely important to the advancement of III-V nitride technology.

Accordingly, several authors using a variety of characterization techniques have reported ?Ev

values for GaN/AlN heterojunctions fabricated by different growth techniques including

ECR-MBE [5-7], NH3-Gas Source-MBE [8], and OMVPE [9,10] on Al2O3 [9,10] and 6H-

SiC [5-8] substrates. Unfortunately, the electrical, optical, and microstructural characteristics

315

of these materials and the associated interfaces prepared/deposited by the various techniques

are quite different, and the corresponding values reported for ?Ev range from 0.5 eV [9] to

1.4 eV [7,8] (see Table 10.1). In a separate paper [11], the authors have shown that for GaN

grown via NH3-gas source MBE on a high temperature monocrystalline AlN buffer layer, a

3D to 2D growth transition can be observed with a corresponding change in the electrical,

optical, and microstructural properties of the resulting GaN films. In this paper, we show that

the valence band discontinuity for the (0001) GaN/AlN interface also changes with this

transition.

Table 10.1. Published data for GaN/AlN ?Ev.

?Ev GaN/AlN Orientation Authors Technique 0.8±0.3 eV (0001) - G. Martin et al, [5,6] - XPS 0.5 eV (0001) - J. Baur et al, [9,10] - Metal Impurities 1.36±0.07 eV (0001) - Waldrop & Grant, [8] - XPS 1.4 eV (0001) - Z. Sitar et al, [7] - CL/PL 0.65 eV (0001) - M. Wang, et a; [13] - Au SBH/Theory 0.85 eV (110) - Alabensi et al, [14] - Theory 0.44-0.75 eV (110) - Bernholc, et al [15] - Theory 0.48-1.12 eV (0001) - Ke, et al [16] - Theory 0.75 eV (001) - " " - Theory 0.81 eV (110) - " " - Theory 0.77 eV (111) - " " - Theory

10.3. Experimental

10.3.1. Thin Film Growth and Analysis

316

All experiments described below utilized a UHV system which integrates several

completely independent UHV thin film growth and analysis units via a 36 ft. long transfer

line having a base pressure of 9x10-10 Torr [12]. The experiments described in this paper

employed only the III-N GSMBE, XPS/UPS, and LEED units [11]. The details of our

sample preparation and growth have been described elsewhere [11] and, therefore, only a

brief description of these components of the research will be given here.

The GSMBE system (base pressure = 3x10-10) Torr was designed and constructed

specifically for the growth of III-V nitride thin films. Source materials in the GSMBE

included Al (99.9999%), Ga (99.99999%), and NH3 (99.9995%). Al and Ga were

evaporated from 25 cc "cold lip" and dual filament Knudsen cells respectively. The XPS and

UPS experiments were performed in a UHV chamber (base pressure = 2x10-10 Torr) and

equipped with a dual anode (Mg/Al) x-ray source, a differentially pumped helium resonance

UV lamp, and a 100 mm hemispherical electron energy analyzer (VG CLAM II). All XPS

spectra were obtained using Mg Kα radiation (hν = 1253.6 eV) at 12 kV and 20 mA

emission current. A combination Gaussian-Lorentzian curve shape with a linear background

best represented the XPS data. All UPS spectra were acquired using the unmonochromated

He I line (hν = 21.2 eV) from the UV lamp. The LEED patterns were obtained using an 80

eV, 1 mA beam.

317

The substrates used in this research were ≈ 1.5x1.5 cm2 pieces cut from 1 3/16 in.

diameter off-axis (4° toward (11-20)) n-type (Nd=1018/cm3) 6H-SiC (0001)Si wafers

obtained from Cree Research, Inc. All wafers were received with an ≈ 1 µm n-type epitaxial

layer (Nd=5x1017cm3) on which was grown ≈ 500-1000Å of thermal oxide. They were

ultrasonically and sequentially rinsed in trichloroethylene, acetone, and methanol, dipped in

10:1 buffered HF for 10 min., and mounted to a Mo sample holder using Ta wire. An

opaque W coating was deposited via RF sputtering on the unpolished side of each wafer

which allowed substrate temperatures of 1100°C to be easily achieved with a tungsten

filament heater. The in situ procedure used for the final cleaning step of the 6H-SiC

substrates was similar to that described by Kaplan and Kern [17,18]. Briefly, each SiC wafer

was annealed in the GSMBE system in a flux of 10-6-10-5 Torr SiH4 for ≈ 15-20 min at 950-

1050°C. Analysis via AES and XPS revealed oxygen-free, silicon terminated SiC surfaces

which displayed (1x1) LEED patterns.

An approximately 250Å monocrystalline AlN film grown at 1050°C in ≈ 10-5 Torr

NH3 on each (0001)Si 6H-SiC wafer was used as the buffer layer for growth of GaN. The

AlN films/buffer layers displayed (2x2) reconstructed surfaces in LEED immediately after

growth. This reconstruction was sensitive to either contamination or temperature, as a (1x1)

LEED pattern was observed several hours after growth. To achieve the growth of the GaN

films on the AlN buffer layer, the latter was heated to 650-800°C in 10-4 Torr (˜ 50 sccm)

ammonia, after which the Ga cell was opened and growth allowed to proceed. After the

desired GaN thickness had been achieved, the Ga cell was closed and the GaN film allowed

to cool in ammonia to approximately 600°C after which the ammonia valve was closed.

10.3.2. GaN/AlN ?Ev analysis

The method used in this research for calculating the GaN/AlN valence band discontinuity

was similar to that of Grant and Waldrop [8,20]. The basic scheme of this approach is to

318

reference the valence band maximum energy to a core level energy from each semiconductor

and then use the measured difference between the two core level energies from a junction

between the two semiconductors to indirectly determine the discontinuity. Specifically, the

position of one core level (CL) from the substrate (AlN) is measured with respect to the

substrate valence band maximum (VBM) i.e. (VBM-CL)AlNbulk. Subsequently, a thin layer

(≈15-20Å) of the second semiconductor (GaN) is deposited on the substrate and the

difference between the substrate and film core levels are measured, i.e. (CLAlN -

CLGaN)interface. Finally, the thickness of the overlying film is increased beyond the

sampling depth of XPS (≈ 250Å) and the CL-to-VBM energy is measured for the film, i.e.

(VBM-CL)GaNbulk. The valence band discontinuity between the two semiconductors is

then given as:

-?Ev(GaN/AlN)= (VBM-CL)AlNbulk - (VBM-CL)GaNbulk + (CLAlN-CLGaN)int. (1)

In the measurements by Grant & Waldrop [8] and Martin et al [5,6], XPS was used to

determine both the core level and valence band maxima energies. In the measurements

described herein, UPS was used to determine the VBM of AlN and GaN due to the increased

signal to noise ratio

(S/N) in the UPS VB spectra. Core level energies were measured via XPS.

319

0 2 4 6 8 10 12Binding Energy (eV)

He I VBM = 3.5 eVHe Iß

UPSx 10

Figure 10.1. UPS Spectra of (2x2) (0001) AlN surface.

10.4. Results

Determination of the energy of the AlN VBM was the most difficult aspect in the

measurement of the GaN/AlN valence band discontinuity. The AlN VBM position in the

UPS data was complicated by the presence of artifacts created by significant amounts of

emission from He Iß radiation (see Figure 10.1). Location of the VBM in the UPS VB

spectra was determined by extrapolating a straight line through the leading edge of the

spectra to the energy axis. A value of 71.5±0.1 eV was determined for Al2p-VBMAlN after

320

analysis of several AlN films. Due to complications with He Iß emission, Al2p-VBMAlN

was also determined using XPS VB spectra. In this case, a lower value of 71.3±0.3 eV was

found for Al2p-VBMAlN (data not shown). Therefore, in these studies, a value for Al2p-

VBMAlN = 71.4 ±0.2 eV was used (see Table 10.2). This number is slightly higher than our

previously reported value of 71.2±0.3 eV determined by aligning the theoretical valence band

density of states (VBDOS) of Lambrecht et al. [21] to the AlN VB XPS spectra.

Determination of the VBM for GaN was relatively straightforward. As shown in

Figure 10.2, extrapolation of the leading edge of the UPS VB spectra for GaN grown at

800°C to the energy axis yields a value of 2.4 ± 0.1 eV which is in excellent agreement with

the value of 2.4 ± 0.2 eV additionally obtained from XPS VB spectra (data not shown).

However, a slight change in the Ga and N core levels relative to GaN VBM occurred as a

function of GaN growth temperature. Table 10.2 lists the Ga and N core levels with respect

to the GaN VBM. Comparison of UPS VB spectra from 650 and 800°C GaN indicates that

the position of the GaN VBM moves only 0.2 eV closer to the Fermi level when increasing

the growth temperature from 650°C to 800°C. This indicates that the 0.6-0.7 eV change in

the CL-VBMGaN values with growth temperature resulted primarily from changes in the

positions of the GaN core levels.

321

Table 10.2. Al 2p and N 1s Core levels referenced to AlN VBM.

1050°C AlN Al2p-VBM 71.4±0.2 eV 650°C GaN Ga 3d-VBM 17.8±0.1 eV Al2p-Ga3d 54.1 800°C GaN Ga 3d-VBM 18.4±0.1 eV Al2p-Ga3d 53.9

0 2 4 6 8 10

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

He IHe I VBM = 2.4 eV

β

UPS x 10

Figure 10.2. UPS spectra of (2x2) (0001) GaN surface.

The positions of the Al 2p, N 1s, and Ga 3p, 3d, and 2p core levels were recorded as a

function of thickness of GaN. Examination of the data revealed that the CLGaN-CLAlN

were similar to within ±0.2 eV regardless of film thickness. However, for purposes of

322

comparison, values of ?Ev were calculated using AlNCL-GaNCL data taken at film

thicknesses in the range of 18-22Å. As noted by Grant and Waldrop, this thickness is beyond

the reported critical thickness for GaN on AlN [22] and should minimize strain effects.

Values for the (0001) GaN/AlN valence band discontinuity for GaN on AlN were

calculated using the data in Table 10.2,. For GaN grown at 800°C on AlN,

?Ev (GaN/AlN) = (VBM-CL)AlNbulk - (VBM-CL)GaNbulk + (CLAlN-CLGaN)int

= -71.4 + 18.4 + 53.9

= 0.9 ± 0.2 eV

Similarly, for GaN grown at 650°C, a ?Ev of 0.6±0.2 eV was calculated.

10.5. Discussion

Two reports of the (0001) GaN/AlN band alignment based on XPS measurements

have been published by Martin et al. [5,6] and Grant and Waldrop [8]. Both groups of

investigators arrived at a type I band alignment between GaN and AlN but the with

significantly different valence band discontinuities of 0.8±0.3 eV [5,6] and 1.36±0.07 eV [8].

As similar values for CLAlN-CLGaN

323

Table 10.3. CL-VBM data for AlN and GaN reported by various investigators. Al 2p-VBMAlN Ga 3d-VBMGaN King et al [21] 71.2±0.3 eV Bermudez et al [23] 71.9±0.2 eV 18.4±0.2 eV This Paper 71.4±0.2 eV 17.8±0.1 eV (650°C) 18.4±0.1 eV (800°C) Waldrop & Grant [8] 70.6±0.07 eV 17.76±0.07 eV Martin et al [5,6] 70.6±0.3 eV 17.1±0.3 eV

were measured in this paper and by Martin et al. and Grant and Waldrop, most of the

discrepancy in ?Ev arises from the values determined for CL-VBMAlN and CL-VBMGaN.

Table 10.3 summarizes the Al2p-VBMAlN and GaNCL-VBMGaN data reported in the ?Ev

measurements of this paper and those by Martin et al and Grant & Waldrop. Values for

Al2p-VBMAlN and Ga3d-VBMGaN from studies by Bermudez et al. [23] concerned with

the UHV reaction between Al and GaN are also included.

A comparison of the data in Table 10.3 shows a large distribution in the reported

Al2p-VBMAlN and Ga3d-VBMGaN values. As these values are extremely sensitive to the

location of the AlN/GaN VBM, it is possible that the observed discrepancy is a result of the

different methods used to locate the VBM. In the study reported here and in the research of

Bermudez et al. [23], the energy position of the VBM for AlN and GaN was determined by a

straight-line extrapolation of the leading edge of the XPS/UPS valence band spectra to the

energy axis. If a significant number of occupied states in the bandgap of AlN or GaN exist at

the surface, photoemission could occur from these states and would cause the valence band to

falsely appear closer to the Fermi level. This would result in a larger Al2p-VBMAlN or

324

Ga3d-VBMGaN value which would agree with the observation that the Al2p-VBM and

Ga3d-VBM values reported here and by Bermudez et al. are much larger than those by

Martin et al. and Grant & Waldrop. Theoretical calculations have yet to indicate the presence

of any types of states in the gap (i.e. surface states), and experimental work has yet to provide

direct evidence of surface states. We also note that for GaN films grown at 650°C (i.e.,

conditions identical to Grant and Waldrop) we obtained a value of 17.8±0.1 eV for Ga3d-

VBMGaN which is identical to the value reported by Grant and Waldrop. This indicates that

photoemission from surface states is probably not the source of discrepancy among the

various measurements.

However, the presence of surface states in sufficiently small densities as to be

undetected in photoemission can still pin the Fermi level in the gap and create band bending

at the surface. The presence of band bending at the surfaces of wide band gap

semiconductors could result in photo-voltage effects which could also account for the

observed discrepancies in the Cl-VBM data. This may explain why our value of 17.8±0.1 eV

for Ga3d-VBMGaN from (1x1) unreconstructed GaN films grown at 650°C and our value of

18.4±0.1 eV from (2x2) reconstructed GaN films grown at 800°C. However, we observed no

evidence of photovoltage effects. Without knowledge or direct evidence of surface states, the

observed discrepancies in ?Ev cannot be ascribed to photovoltage effects.

One final and possible source of the discrepancies in the reported ?Ev measurements

could result from differences in the types and densities of defect, strain and stoichiometry in

the GaN and AlN films used in the various studies. The recent theoretical calculations by Ke

et al. [16] have shown a strong variation in ?Ev (0001) GaN/AlN from 0.48 to 1.12 eV based

on the internal relaxation parameter, u (defined as d/c where d is the bond length along the

325

stacking direction and c is the lattice constant for the c axis). Their results show a very small

?Ev=0.48 eV when the GaN bond length along the c axis is forced into that of the ideal

hexagonal structure and a large ?Ev=1.12 eV when GaN is allowed to adopt the c axis bond

length of its bulk structure. For the more realistic intermediate case of the two extremes, they

calculated a ?Ev of 0.83 eV. The results of Ke et al. clearly illustrate the magnitude of the

effect that strain and defects at the GaN/AlN interface can play on the band alignment

between these materials. In a prior report [11], we have noted that the electrical, optical, and

microstructural properties of GaN films grown by the NH3-GSMBE technique at

temperatures of 650 and 800°C are different. In particular we have observed that GaN

growth at 650°C occurs via a 2D/Stranski-Krastanov type growth mechanism with the

corresponding films exhibiting rough surfaces, extremely high carrier concentrations

(N=1019-20/cm3), and broad, weak PL spectra. In contrast, GaN growth at 800°C was

observed to occur via a two dimensional/Frank van der Merwe growth mechanism with the

corresponding films exhibiting smooth surfaces, lower carrier concentrations (n=1016/cm3)

and sharp PL. Clearly, the defect levels in these films and at the interfaces are significantly

different and could be the cause for the increase in ?Ev from 0.6 to 0.9 eV with the increase

in growth temperature. Unfortunately, without detailed knowledge of the microstructure of

the GaN/AlN interfaces examined in this and other studies, it is not currently possible to link

these variations.

326

10.6. Conclusions

In conclusion, XPS/UPS has been used to determine the heterojunction valence band

discontinuity for the (0001) GaN/AlN interface. For GaN grown on AlN at 650°C, a valence

band discontinuity of 0.6±0.2 eV was determined, while a discontinuity of 0.9±0.2 eV was

determined for GaN grown on AlN at 800°C.

10.7. Addendum

Since the completion of this work, there have been several new experimental and

theoretical findings published which point to the existence of surface states on GaN and AlN

surfaces. First, Bermudez [24] has recently corrected his reported value for Ga 3d - VBMGaN

of 18.4 ± 0.2 eV to 18.0 ± 0.2 eV. He attributed the higher value of Ga 3d- VBMGaN to Fermi

level pinning by the existence of undetected surface states at the edge of the GaN VBM in the

prior study [23]. These surface states are apparently extinguished by O2 exposure which

unpins the GaN surface Fermi level and allowed for the more accurate measurement of Ga 3d

- VBMGaN.

327

The above findings by Bermudez [24] are additionally complemented by recent local

density approximation (LDA) supercell calculations by Northrup and Neugebauer [25] for

nonpolar (10-10) and (11-20) 2H-GaN surfaces. Their calculations show the presence of N

derived states (SN) just above the GaN VBM and Ga dangling bond states (SGa) just below

the GaN conduction band minimum (CBM) for the unrelaxed "ideal" (11-20) and (10-10)

surfaces. However when these surface are allowed to relax, both states move out of the gap

with the SN state dropping below the VBM and the SGa state moving above the CBM.

Similar results have also been obtained for the (11-20) surface of 2H-AlN by Kadas et al [26]

in their ab initio Hartree-Fock total energy calculations. Their results showed a sharp

localized state in the AlN band gap well below the AlN CBM associated with Al surface

atoms and states just above the AlN VBM associated with N surface atoms. Both of these

theoretical calculations therefore suggest that the surface state observed by Bermudez on

(0001) GaN could be associated with a nitrogen derived state. This suggestion is supported

by the recent time of flight scattering and recoiling spectrometry (TOF-SARS) studies of

Sung et al [27] which have found that GaN surfaces prepared similarly to Bermudez are N

terminated with Ga comprising the second layer of atoms. Unfortunately, these results are in

contradiction with the ion channeling and convergent beam electron diffraction studies of

Daudin et al [28] which have found smooth OMVPE GaN surfaces on (0001) Al2O3 to be Ga

terminated and rough GaN surfaces to be N terminated.

Although the films examined in this study are expected to be either Al or Ga

terminated (based on the fact that growth occurred on the Si face of (0001) 6H-SiC), the

theoretical calculations of Bernholc [15] have shown that for (2x2) Ga terminated GaN a N

adatom reconstruction is the most energetically favorable. Therefore, the above findings

reemphasize that both our values for Al2p - VBM for (2x2) (0001) AlN and Ga 3d - VBM

for (2x2) (0001) GaN may be both over estimated due to the unknowing inclusions of

nitrogen derived surface states in the location of the AlN and GaN VBM. This would bring

the calculated ?Ev GaN/AlN into closer agreement with the results of Grant and Waldrop [8]

328

and further away from theory. However if the overestimation of the values for both the

Al2p-VBMAlN and Ga3d-VBMGaN are the approximately the same, then the values for ?Ev

GaN/AlN reported above will not be affected. Clearly, all of the above indicates the need for

theoretical investigations of the electronic structure of both (0001) AlN and GaN surfaces.

More experimental investigations are also needed to determine the polarity of (0001) III-V

nitride films grown on both 6H-SiC and Al2O3 as well as to determine what effect various

processes may have the surface termination of these films.


The authors would like to thank Cree Research, Inc. for supplying the SiC substrates in this

research. This work was supported by the ONR under Contract N00014-92-J-1477.

10.9. References

1. S. Strite and H. Morkoc, J. Vac. Sci. Technol. B 10 p. 1237 (1992). 2. J.H. Edgar, J. Mater. Res. 7 p. 235 (1992). 3. R.F. Davis, Proc. of IEEE 79 p. 702 (1991). 4. A. Morgan and J. Williams, Physics and Technology of Heterojunction Devices, Wiley & Sons, New York 1985, pp. 241-244. 5. G. Martin, S. Strite, A. Botchkarev, A. Agarwal, A. Rockett, H. Morkoc, W.R.L. Lambrecht, and B. Segall, Appl. Phys. Lett. 65 p. 610 (1994). 6. G. Martin, S. Strite, A. Botchkarev, A. Agarwal, A. Rockett, W.R.L. Lambrecht, B. Segall, and H. Morkoc, J. Electronic Materials 24 p. 225 (1995). 7. Z. Sitar, M.J. Paisley, B. Yan, R.F.Davis, J. Ruan, and J.W. Choyke, Thin Solid Films, 200 p. 311 (1991). 329

8. J.R. Waldrop and R.W. Grant, Appl. Phys. Lett. 68 p. 2879 (1996). 9. J. Baur, K. Maier, M. Kunzer, U. Kaufmann, and J. Schneider, Appl. Phys. Lett. 65 p. 2211 (1994). 10. J. Baur, K. Maier, M. Kunzer, U. Kaufmann, Materials Science and Engr. B 29 p. 61 (1995). 11. S.W. King, M.C. Benjamin, R.J. Nemanich, and R.F. Davis, to be published. 12. J. van der Weide, Ph.D. Dissertation, NCSU (1994). 13. M.W. Wang, J.O. McCaldin, J.F. Swenberg, T.C. McGill, R.J. Hauenstein, Appl. Phys. Lett. 66 p. 1974 (1995). 14. E.A. Albanesi, W.R.L. Lambrecht, and B. Segall, J. Vac. Sci. Technol. B 12 2470 (1994). 15. Bernholc, Spring MRS Symposium 1996. 16. S. Ke, K. Zhang, and X. Xie, J. Appl. Phys. 80 p. 2918 (1996). 17. R. Kaplan, Surface Science, 215 p. 111 (1989). 18. R.S. Kern, Ph.D. dissertation, NCSU (1996). 19. S. Tanaka, Ph.D. dissertation, NCSU (1995). 20. E.A. Kraut, R.W. Grant, J.R. Waldrop, and S.P. Kowalczyk, Phys. Rev. Lett. 44 p. 1620 (1980). 21. S.W. King, M.C. Benjamin, R.J. Nemanich, R.F. Davis, MRS Proceedings 22. S. Krishnankutty, R.M. Kolbas, M.A. Khan, J.N. Kuznia, J.M. Van Hove, and D.T. Olson, J. Electron. Mater. 21 437 (1992). 23. V.M. Bermudez, T.M. Jung, K. Doverspike, and A.E. Wickenden, J. Appl. Phys. 79 p. 110 (1996). 24. V.M. Bermudez, J. Appl. Phys., 80 p. 1190 (1996). 25. J.E. Northrup and J. Neugebauer, Phys. Rev. B., 53 p. R10477 (1996). 26. K. Kadas, S. Alvarez, E. Ruiz, P. Alemany, Phys. Rev. B., 53 4933 (1996). 27. M.M. Sung, J. Ahn, V. Bykow, D.D. Koleske, A.E. Wickenden, and J.W. 330

Rabalais, Phys. Rev. B, 54 p. 14652 (1996). 28. B. Daudin, J.L. Rouviere, and M. Arlery, Appl. Phys. Lett., 69 p. 2480 (1996).

331

11. Gas Source Molecular Beam Epitaxy Growth of

Scandium Nitride

on (111) 3C and (0001) 6H-SiC


to the

Journal of Crystal Growth

by

Sean W. King, Kieran M. Tracy, David W. Bray, Eric P. Carlson, Robert J. Therrien,

William G. Perry, Robert J. Nemanich, and Robert F. Davis

Department of Materials Science & Engineering


Raleigh NC 27695.

332

11.1 Abstract

Films of ScN were grown via gas source molecular beam epitaxy (10-5 Torr NH3,

800-1050°C) on (111) 3C-SiC epilayers previously deposited on vicinal and on axis (0001)Si

6H-SiC substrates. Analysis via SEM on the former revealed featureless surfaces at a

magnification of 10 kX, and a slight step morphology on the latter which presumably

mimicked the steps on the vicinal surface. All films exhibited hexagonal (1x1) LEED

patterns indicating growth of (111) oriented ScN. However, TEM showed the films to be

polycrystalline with columnar grains oriented at ˜ 15° to the (0001) direction of the SiC

substrate. Omega-2θ XRD scans additionally showed a FWHM of 1047 arc sec. The

conductivity of the films decreased with decreasing growth temperature. The sheet resistance

of films grown at 1050°C were = 5 ? cm. Hot probe measurements showed all ScN films to

be n-type. Ultra-violet photoelectron spectroscopy of the ScN films grown at 1050°C and

800°C showed the valence band maximum of ScN to be positioned 1.6 and 1.2 eV

respectively below the system Fermi level, indicating a minimum band gap of 1.6 eV.

11.2. Introduction

Scandium nitride (ScN) possesses the NaCl crystal structure and a lattice constant,

ao, of ≈ 4.503Å [1,2]. Early reports [3-5], of the growth of single crystals of ScN reported

333

large deviations (1-5%) from stoichiometry (i.e. ScN1-x), presumably due to incomplete

nitridation of the scandium charge. Possibly for this reason, it is still the subject of much

debate as to whether ScN is a semiconductor or a semi-metal [1,2,4,6-11]. Theoretical

calculations indicate that the fundamental band gap of ScN is 0.0 eV or very small (< 0.1 eV)

[8-10]. However, optical transmission experiments of ScN fabricated by a variety of

techniques have measured bandgaps ranging from 1.5-2.1 eV [1,2,4,7]. Many other

properties of ScN are not known or in dispute. Table 11.1 summarizes some of the reported

values for the physical properties of ScN.

Despite the lack of detailed knowledge, ScN is of interest in III-V nitride technology

due to its reasonably close lattice matching with cubic and hexagonal GaN (ao (0001) GaN =

3.189Å, a (111) ScN = 3.139Å [17]) and its possible band gap of 2.1 eV [2]. The reported

low resistivity [16] of ScN makes it a viable alternative to TiN for use as a low resistivity

ohmic contact to n-type GaN. Further, the achievement of p-type ScN could lead to the

development of a low resistivity contact to p-type GaN. Finally, the moderately close lattice

matching of ScN to SiC could also allow ScN to be used as a conducting buffer layer to

replace the insulating AlN or AlGaN buffer layers currently employed in the growth of GaN

on SiC substrates. Successful growth of GaN films on a Sc/GaN/Al2O3 structure has already

been recently demonstrated [18]. Finally, the high temperature/thermodynamic stability and

possible bandgap of 1.5-2.1 eV makes ScN an attractive replacement for InN in the

fabrication of blue LEDS and Lasers. However, a serious impediment to the use of ScN in

both of the last two applications is the difference in crystal structure between ScN (NaCl) and

GaN (ZB or wurtzite). In order to determine whether ScN can be used in any of these

applications, the authors have begun an investigation regarding the growth of ScN and

334

GaScN alloys on (0001) 6H-SiC and (0001) GaN. The results of initial research regarding

the growth of ScN on (111) 3C-SiC and (0001) 6H-SiC via gas source-molecular beam

epitaxy (GSMBE) are described in the following sections.

Table 11.1. Properties of ScN

Crystal Structure NaCl [1] Lattice Constant 4.503Å [1] Bandgap, Eg 0.0-2.6 eV [1,4,5,10,11] Conductivity type. n or p? n-type [1,2,11] Carrier Concentration 9x1019-8.3x1020/cm3 [1,2,11] Electron Mobility 28-150 cm2/Vsec [1,2] Doping Si, C, Zn, and Mg, no p-type ScN [2] Resistivity 25.4 µ? cm [16] Electron Effective Mass 0.1-0.2mo [14] Dielectric Constant ε = 5.2 [9], ε = 10.8 [14] Hardness 1170±150 kg/mm2 for 50 gm load [16] Thermal Expansion Coefficient 8.68x10-6/°C [9], 8.1x10-6/°C [2] Tmelt 2550±50°C [16] Oxidation Resistance Inert in air to temperatures of 600°C [16] Etchants Completely dissolves in boiling HCl and HNO3 [16]

11.3. Experimental

A gas source-molecular beam epitaxy (GSMBE) system, designed and constructed

specifically for the growth of ScN, AlN, and GaN thin films, was used in these studies. The

GSMBE was a component of an ultrahigh vacuum transfer line which allowed for in situ

analysis of the ScN films by low energy electron diffraction (LEED), Auger electron

spectroscopy (AES), x-ray photoelectron spectroscopy (XPS), and ultra-violet photoelectron

spectroscopy (UPS). Source materials in this system consisted of NH3 (99.9995%), Sc

335

(99.99%), Ga (99.99999%), and Al (99.99999%). The metals were evaporated from standard

Knudsen cells. Impurities in the scandium charge including, Al, Ca, Cr, Cu, Mn, Si & Y,

were each = 30 ppm. However, significant amounts of fluorine were also detected by XPS

from Sc films evaporated onto 6H-SiC (0001). The source of the fluorine is currently

undetermined but is believed to be the scandium charge. A base pressure of 10-10 Torr was

achieved in the GSMBE system via a 400 l/s turbo pump and a 500 l/s ion pump. During

nitride growth, the GSMBE system was pumped by the turbo pump only in order to prevent

irreversible damage to the ion pump. Substrate temperatures of 1100°C were easily achieved

via a hot tungsten filament heater.

Vicinal and on axis, 6H-SiC wafers (0001) (Nd ≈ 1018/cm3) with a 1 micron 3C-SiC

epitaxial layer (Nd≈1017/cm3) were provided by Cree Research Inc. The unpolished side of

these wafers were coated with an opaque tungsten film via RF sputtering to increase the

thermal heating efficiency of the SiC, as the latter is transparent to the infra-red radiation

emitted by the tungsten filament heater. After sputter coating, the wafers were ultrasonicated

in trichloroethylene, acetone, and methanol for 10 min. each and then dipped in 10:1 buffered

HF for 10 min. to remove the surface oxide. The wafers were subsequently

degassed/annealed at 1050°C in the GSMBE system for 10-15 min. to further desorb the

surface oxide. This surface displayed a (√3x√3)R30° reconstructed surface in LEED.

Growth of ScN was initiated on this surface in 10-5-10-4 Torr NH3 at temperatures ranging

from 800-1050°C.

336

11.4. Results

11.4.1. Thermodynamics

In an initial attempt to understand the growth of ScN and deduce a set of optimized

growth parameters, an MBE/CVD equilibrium diagram for ScN was computed using the

software package HSC. The computed equilibrium diagram shown in Figure 11.1, shows

that under GSMBE conditions (i.e. 10-5-10-4 Torr NH3), ScN is the equilibrium phase to

900°C. This phase is stable at higher temperatures, but is in equilibrium with a significant

amount of Sc(v). At still higher temperatures ScN decomposes completely into Sc(v) and

N2(g). For comparison purposes, equilibrium diagrams for AlN, GaN, and InN were also

calculated using HSC (see Figures 11.2,11.3, and 11.5). The equilibrium diagram of ScN

closely resembles the equilibrium diagram determined for AlN (see Figure 11.2) indicating

the possibility that the optimized growth conditions for ScN and AlN should be

approximately the same. In contrast, the equilibrium diagrams computed for GaN and InN

(see Figures 11.3 and 11.5) showed these materials to become thermodynamically unstable at

much lower temperatures than ScN and AlN. In fact in order to stabilize these materials

under typical experimental conditions, it was found necessary to exclude N2 formation from

the HSC equilibrium calculations. In a sense, this is a crude way of imposing/including

kinetic limitations in thermodynamic calculations. However, this approach was shown to be

337

valid as the computed equilibrium for GaN and InN with N2 removed from the equilibrium

products was found to be in close agreement with the experimentally determined equilibrium

of GaN and InN in vacuum (see Figure 11.4 and 11.5).

10 -9

10 -7

10 -5

0.001

0.1

10

1000

10 12

10 14

10 16

10 18

10 20

10 22

10 24

600 800 1000 1200 1400 1600

NH

3 Pre

ssur

e (T

orr)

NH

3 Flux (#/cm2sec)

Temperature (ÞC)

ScN

Sc(v) + N2

ScN + Sc(v)

OMVPE/CVD

GSMBE

Flux = P/¦2šmkT

Figure 11.1. ScN equilibrium diagram computed using HSC.

338

10 -7

10 -5

0.001

0.1

10

1000

10 14

10 16

10 18

10 20

10 22

10 24

600 800 1000 1200 1400 1600 1800

NH

3 Pre

ssur

e (T

orr)

NH

3 Flux (#/cm2sec)

Temperature (ÞC)

AlN

Al + N2

Al + AlNOMVPE

GSMBE

Figure 11.2. AlN equilibrium diagram computed using HSC.

10 -3210 -2910 -2610 -2310 -2010 -1710 -1410 -1110 -810 -5

0.0110 110 410 7

10 1010 13

10 -1110 -810 -50.011010 410 710 1010 1310 1610 1910 2210 2510 2810 3110 34

200 400 600 800 1000 1200 1400 1600

NH

3 Pre

ssur

e (T

orr)

NH

3 Flux (#/cm2sec)

Temperature (ÞC)

Ga + N2

GaN HSC: Ga-NH3 Equilibrium with N

2

HSC: Ga-NH3 Equilibrium without N

2

GSMBE

OMVPE

Figure 11.3. GaN equilibrium diagram computed using HSC.

339

10 -10

10 -8

10 -6

10 -4

10 -2

10 0

10 2

10 4

10 6

10 8

10 10

10 12

10 14

10 16

10 18

10 20

10 22

10 24

10 26

10 28

10 30

600 700 800 900 1000 1100 1200

NH

3 Pre

ssur

e (T

orr)

NH

3 Flux (#/cm2sec)

Temperature (ÞC)

GaN Decomp. Rate(Munir & Searcy)

Equil. NH3 Flux

(Thurmond & Logan) OMVPE

GSMBE

GaN(s)

Ga + N2

Figure 11.4. GaN equilibrium diagram computed using experimental data of Thurmond and Logan [19] and Munir and Searcy [20].

10 -18

10 -16

10 -14

10 -12

10 -10

10 -8

10 -6

10 -4

10 -2

10 0

10 2

10 4

10 4

10 6

10 8

10 10

10 12

10 14

10 16

10 18

10 20

10 22

10 24

200 400 600 800 1000

3 Pre

ssur

e (T

orr)

NH

3 Flux (#/cm2sec)

Temperature (ÞC)

InN

In(l) +N2

GSMBE

OMVPE

HSC: In-NH3 Equilibrium with N

2

HSC: In-NH3 Equilibrium without N

2

InN Decomposition rate/N2

Vapor, Jones and Rose

Figure 11.5. InN equilibrium diagram computed using HSC and experimental data of Jones and Rose [21].

340

11.4.2. Growth

In situ AES and XPS analysis of ScN films grown in 10-5 Torr NH3 at 800-1050°C

detected the presence of only Sc and N (see Figure 11.6 and 11.7). Small traces of oxygen

were detected in AES (see Figure 11.6) but were attributed to electron beam induced

oxidation during AES. Other contaminants (most notably fluorine) were not detected. A

more detailed XPS analysis of the Sc2p1/2,3/2 and N 1s core levels showed the ScN films to

be stoichiometric within the detection limit of this analytical technique ( ≈ 0.1 at.%) (see

Figure 11.7). No unreacted Sc was detected on the surface. The XPS spectrum shown in

Figure 11.7 is extremely similar to that obtained by L. Porte [13] from ScN films prepared by

argon sputtering Sc onto Ta in 3x10-7 Torr N2. In situ UPS analysis showed a shift in the

ScN valence band maximum (VBM) from 1.6 eV below the Fermi level for the 1050°C films

to 1.2 eV below EF for the 800°C films (see Figure 11.8). Correspondingly, an increase in

the resistance of the ScN films with decreasing growth temperature was observed. For ScN

films grown at 1050°C, typical sheet resistances measured by four point probe were ˜ 5 ? cm.

Films grown at lower temperatures < 950°C were too resistive for four point probe

measurements. Hot probe measurements revealed that all ScN films were n-type. Doping of

ScN with Si and C from SiH4 and C2H4 respectively was investigated. The introduction of

Si did not appreciably decrease the resistivity of the ScN films. Capacitance-Voltage

measurements indicated ND-NA ≈ 5x1017/cm3. In contrast, the resistivity of C doped films

showed a dramatic decrease to 1.2 ? cm and ND-NA = 1019-1020/cm3 (CV measurements).

341

Unfortunately, hot probe measurements indicated these films were still n-type. Subsequent

SIMS analysis showed no incorporation of carbon into the ScN films but an increase in the

oxygen contamination by 2-3 order of magnitudes. The source of the oxygen is presumably

H2O or O2 impurities in the C2H4. However, these results indicate that oxygen may be an

excellent n-type dopant for ScN.

Hexagonal (1x1) LEED patterns were obtained at Ep ≈ 50 eV from all ScN films

indicating growth of (111) oriented ScN. Examination of these films via SEM revealed

featureless surfaces to magnifications of 10,000X. However, some films grown on off-axis

(0001) 6H-SiC exhibited a step like structure apparently mimicking the steps on the SiC

substrate. Transmission electron microscopy of 800°C ScN showed the films to be

polycrystalline with the grains oriented at an angle of ˜ 15° to the (0001) direction of the SiC

substrate (see Figure 11.9). This is in agreement with the wide FWHM of 1047 arc sec

observed in XRD Omega-2θ scans of similar films. A similar microstructure was observed

for ScN films grown on a 5000Å GaN/AlN/6H-SiC structure (see Figure 11.10).

342

100 200 300 400 500 600 700

dN(E

)/dE


O KLL

Sc LMM

Sc MNN

N KLL

Figure 11.6. AES survey scan of ScN film grown at 800°C on (111) 3C-SiC.

392 396 400 404 408

Cou

nts (

arb.

uni

ts)

Binding Energy (eV)

N 1s

Sc 2p3/2

Sc 2p1/2

Figure 11.7. XPS spectrum of N 1s and Sc 2p3/2,1/2 2p3/2,1/2 core levels from a ScN film grown at 800°C.

343

-14 -12 -10 -8 -6 -4 -2 0 2

Cou

nts (

arb.

uni

ts)

Energy Below Fermi Level (EF = 0) (eV)

He Iß

He I VBM = -1.6 eV

Figure 11.8. UPS spectrum of ScN film grown at 800°C on (111) 3C-SiC.

344

Figure 11.9. TEM of ScN film grown at 800°C on (111) 3C-SiC.

Figure 11.10. TEM of 800°C ScN on 5000Å (0001) GaN/AlN/6H-SiC.

11.5. Discussion

The above thermodynamic calculations and growth studies show that ScN films can

be easily grown/deposited via NH3-GSMBE in the temperature range of 800-1050°C.

Dismukes et al [1,2] have previously demonstrated growth of single crystal films of ScN on

(11-20) Al2O3 via hydride vapor phase epitaxy in the temperature range of 850-930°C. As

800 and 1050°C have been determined to be our optimized growth temperatures for GaN and

AlN respectively, ScN should be much easier to incorporate into these compounds than InN

which is highly unstable at these temperatures in MBE. Previous attempts in this research to

grow InN via NH3-GSMBE at these temperatures have been completely unsuccessful.

Unfortunately, the difference in crystal structure between ScN and AlN/GaN will limit the

345

range over which these compounds can be successfully combined to form equilibrium alloys.

The results of S. Lee et al. [22], however, have shown that single phase, NaCl structure

AlxTi1-xN with x as large as 0.8 can be fabricated via plasma enhanced chemical vapor

deposition. The observation by these investigators of metastable AlN particles with the NaCl

structure is also extremely encouraging.

Our UPS measurements on the 1050°C ScN films which show the ScN VBM to be

1.6 eV below the system Fermi level are in agreement with the UPS results of Porte [13]

which also found the VBM of ScN to be 1.5-2.0 eV below EF. These measurements indicate

that ScN has a minimum bandgap of 1.5 eV. The observation that the ScN VBM shifted 0.4

eV closer to EF with decreasing growth temperature indicates a decrease in donors

presumably due to nitrogen vacancies.

11.6. Conclusions

Scandium nitride films have been successfully deposited on (111) 3C/(0001)6H-SiC

substrates by NH3-GSMBE in the temperature range of 800-1050°C. All films were n-type

with a maximum conductivity of 5 ? cm for films grown at 1050°C. The conductivity of

these films decreased with decreasing temperature. Analysis via SEM showed featureless

surfaces at a magnification of 10 kX. Analysis via TEM indicated the films were

polycrystalline with grains oriented at ˜ 15° to the (0001) direction. A minimum bandgap of

1.6 eV for ScN was deduced via UPS measurements.

11.7 Acknowledgments

346

The authors would like to thank Cree Research, Inc. for supplying the SiC substrates

used in this research. This work was supported by the ONR under contracts N00014-91-J-

1410 and N00014-92-J-1477.

11.8 References

1. J.P. Dismukes, W.M. Yim, J.J. Tietjen, and R.E. Novak, RCA Review, 31 p. 681 (1970). 2. J.P. Dismukes, W.M. Yim, and V.S. Ban, J. Cryst. Growth, 13/14 p. 365 (1972). 3. R. Didchenko and F.P. Gortsema, J. Phys. Chem. Solids, 24 p. 863 (1963). 4. G. Busch, E. Kaldis, E. Schaufelberger-Teker, and P. Wachter, in Les Elements des Terres Rares, Edition du CNRS, Colloque Internationale No. 180, Tome I (1970), p. 359. 5. M.D. Lyutaya, A.B. Goncharuk, and I.I. Timofeeva, Z. Prik. Khimii, 48 p. 721 (1975). 6. N. Sclar, J. Appl. Phys., 33 p. 2999 (1962). 7. N. Sclar, J. Appl. Phys., 35 p. 1534 (1964). 8. K. Schwarz, P. Weinberger, and A. Neckel, Theor. Chim. Acta, 15 p. 159 (1969). 9. A. Neckel, P. Rastl, R. Eibler, P. Weinberger, and K. Schwarz, J. Phys. C, 9 p. 579 (1976). 10. R. Monnier, J. Rhyner, T.M. Rice, and D.D. Koelling, Phys. Rev. B, 31 p. 5554 (1985). 11. G. Travaglini, F. Marabelli, R. Monnier, E. Kaldis, and P. Wachter, Phys. Rev. B, 34 p. 2876 (1986). 12. W. Lengauer, J. Solid State Chemistry, 76 p. 412 (1988). 13. L. Porte, J. Phys. C: Solid State Phys., 18 p. 6701 (1985). 14. G. Harbeke, E. Meier, and J.P. Dismukes, Optics Communications, 4 p. 335 (1972). 347

15. M.D. Lyutaya and V.F. Bukhanevich, Russian J. Inorganic Chem., 7 p. 1290 (1962). 16. G.V. Samsonov, M.D. Lyutaya, and V.S Neshpor, Z. Prik. Khimii, 36 p. 2108 (1963). 17. S. Strite and H. Morkoc, J. Vac. Sci. Technol. B, 10 p. 1237 (1992). 18. D. Koleske, A. Wickenden, J. Freitas, R. Kaplan, S. Prokes, this symposium. 19. C.D. Thurmond and R.A. Logan, J. Electrochem. Soc., 119 p. 622 (1972). 20. Z.A. Munir and A.W. Searcy, J. Chem. Phys., 42 p. 4223 (1965). 21. R.D. Jones and K. Rose, J. Phys. Chem. Solids, 48 p. 587 (1987). 22. S. Lee, B. Kim, H. Kim, J. Lee, J. Appl. Phys. 80 p. 1469 (1996).

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