Synthesis and Properties of Van der Waals-bonded

Semiconductor Heterojunctions with Gallium Nitride

DISSERTATION

Presented in Partial Fulfillment of the Requirements for the Degree Doctor of

Philosophy in the Graduate School of The Ohio State University

By

Choong Hee Lee

Graduate Program in Electrical and Computer Engineering

The Ohio State University

2018

Dissertation Committee

Professor Siddharth Rajan, Advisor

Professor Wu Lu

Professor Tyler Grassman

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i

Abstract

This thesis explores the heterojunctions between van de Waals (vdW) bonded

materials and conventional semiconductors. Two-dimensional (2D) metal chalcogenide

have recently attracted considerable attention due to their unique optical and electronic

properties and great potential in a wide range of applications in flexible, low-power

electronics, and optical applications. The absence of out-of-plane dangling bonds in layered

materials can enable vdW epitaxy on highly lattice-mismatched substrate. Synthesis of

high-quality 2D materials in centimeter scale remains a key challenge. Among several

synthesis methods, chemical vapor deposition (CVD) method can provide a well ordered

lattice orientation of 2D films grown on a crystalline dielectric substrate, but it requires

extra transfer process to demonstrate various heterostructures. In contrast to CVD methods,

molecular beam epitaxy (MBE) can enable high-quality growth of 2D films with great

control of growth rate and doping profile for heterojunctions with no intermixing at the

interface. This can be exploited to combine dissimilar 2D or 3D semiconductors in large

scale, enabling new flexibility in epitaxy. In this dissertation, the synthesis of 2D metal

chalcogenides with a wide bandgap semiconductor, gallium nitride (GaN), and the

electrical transport of 2D/GaN heterojunction devices are investigated. The growth of GaN

on 2D substrate and related processing techniques are also developed.

ii

The growth study of 2D materials was initiated by growing gallium selenide (GaSe)

to gain a better understanding of vdW epitaxy of 2D films. A two-step growth method

involving high temperature nucleation of single crystalline domains and low temperature

growth to enhance coalescence is adopted to obtain continuous GaSe with an epitaxial

relationship with the substrate. The transport of 2D/GaN is then studied via vdW epitaxy

of 2D tin diselenide (SnSe2) on GaN substrate. Current transport behavior of SnSe2/GaN

diodes is quantitatively characterize by electrical and optical measurement and explained

using conventional 3D/3D heterojunction transport theory.

Epitaxial growth of GaN on 2D substrate presents significant epitaxial challenges,

especially on nearly lattice matched MoS2 substrate. MoS2 degradation during GaN growth

is a major problem. In the second part of the thesis, the growth conditions for forming GaN

on MoS2, identify the growth mechanism, and introduce GaN/MoS2 device fabrication

process are investigated. A comprehensive growth diagram for GaN/MoS2 is developed,

and the degradation of MoS2 during GaN growth is explained.

Along with the material growth, two essential processing techniques, layer transfer

and layer-by-layer etching of 2D materials are developed. Transferring the 2D films to

arbitrary semiconductor substrates offers large flexibility in fabricating various types of

heterostructures. Chemical free benign transfer process is introduced. In addition, the

process of precise thickness control of the MoS2 film using digital etching technique (i.e.,

oxidation or removal oxide layer) is then demonstrated. These studies provide a

comprehensive understanding of the MBE synthesis of 2D with 3D semiconductors and their

heterojunction devices, which can be used for high performance device applications.

iii

Dedication

Dedicated to my wife, Eun Jung

iv

Acknowledgments

Starting a doctoral program was not easy, and Prof. Siddharth Rajan was the one

who gave me the opportunity in a difficult situation. With his great guidance and support,

I was able to complete my Ph.D. He is one of the most wonderful people I have ever met,

and I am very grateful to him for sharing his knowledge and experience with me.

I also thank Prof. Wu Lu, Prof. Tyler Grassman, and Prof. ChunNing Lau for

serving on my defense committee and for having valuable discussions with me. I am also

thankful to a number of professors—including Prof. Roberto Myers and Prof. Steven

Ringel—who gave me great advice and experience. I thank Prof. Aaron Arehart and Pran

Paul for their help with measurements, and I also thank Jared Johnson and Prof. Jinwoo

Hwang for providing amazing images on 2D and III-Nitride samples.

I would like to thank Mark Brenner for helping me learn about MBE equipment

and helping me make the system run smoothly. I also like to thank the NTW staff members

(John Carlin, Derek Ditmer, Aaron Payne, Pete Janney, Aimee Price, Paul Steffen and

others) who helped with the clean room process, the NSL staff members (Camelia Selcu,

Denis Pelekhov and others), and CBC staff members (Lisa Alexander and Gordon Renkes)

who helped with the analysis.

Thanks to our group members who worked together during the long period, I had

great company and felt I was not alone. I would like to thank the seniors—including Pil

v

Sung Park, Digbijoy Nath and Fatih Akyol—who took care of many things. I really enjoyed

the coffee breaks and other past-times with my former group members, Sriram

Krishnamoorthy, Ting-Hsiang Hung, Sanyam Bajaj, Edwin Lee II, Omor Shoron, Sadia

Khandaker, and Yuewei Zhang. It has been a pleasure working with new recent members

and others, including Zhanbo Xia, Shahadat Hassan, Towhid Razzak, Gurudatt Rao, Caiyu

Wang, Zane Jamal-Eddine, Chandan Joishi, Wahidur Rahman, Nidhin Kurian and Dante

O’Hara. I am also glad that I can stay with them a little longer after my graduation.

Finally, I want to say that I love my wife and I am very thankful to her, as well as

my son who stayed with me, my parents, and my in-laws who have greatly supported me

when I was in a distant country.

vi

Vita

January 06, 1984 ............................................Born – Seoul, S. Korea

2002 to 2006 ..................................................B. S. in Electrical and computer engineering,

Yonsei University, Seoul, S. Korea

2006 to 2008 ..................................................M. S. in Electrical and computer engineering,

Yonsei University, Seoul, S. Korea

2008 to 2012 ..................................................Engineer, Samsung Electro-Mechanics,

Suwon, S. Korea

2013 to present ..............................................Graduate Research Associate,

Electrical and computer engineering,

The Ohio State University, Ohio, US

Publications

Journal publications:

Choong Hee Lee, Sriram Krishnamoorthy, Pran K. Paul, Dante J. O'Hara, Mark R.

Brenner, Roland K. Kawakami, Aaron R. Arehart, and Siddharth Rajan, “Large-area

SnSe2/GaN heterojunction diodes grown by molecular beam epitaxy” Applied Physics

Letters 111, 202101 (2017).

Choong Hee Lee, Sriram Krishnamoorthy, Dante J. O'Hara, Jared M. Johnson, John

Jamison, Roberto C. Myers, Roland K. Kawakami, Jinwoo Hwang, and Siddharth Rajan.

"Molecular Beam Epitaxy of 2D-layered Gallium Selenide on GaN substrates." Journal

of Applied Physics 121, 094302 (2017).

Choong Hee Lee, Edwin W. Lee II, William McCulloch, Zane Jamal-Eddine, Sriram

Krishnamoorthy, Michael J. Newburger, Roland K. Kawakami, Yiying Wu, and

Siddharth Rajan. "A self-limiting layer-by-layer etching technique for 2H-MoS2."

Applied Physics Express 10, 035201 (2017).

vii

Sriram Krishnamoorthy, Edwin Lee II, Choong Hee Lee, Yuewei Zhang, William D.

McCulloch, Jared M. Johnson, Jinwoo Hwang, Yiying Wu, and Siddharth Rajan. "High

Current Density 2D/3D Esaki Tunnel Diodes." Applied Physics Letters 109, 183505

(2016).

Choong Hee Lee, William McCulloch, Edwin W. Lee II, Lu Ma, Sriram

Krishnamoorthy, Jinwoo Hwang, Yiying Wu, and Siddharth Rajan. "Transferred large

area single crystal MoS2 field effect transistors." Applied Physics Letters 107, no. 19

193503 (2015).

Lee II, Edwin W., Choong Hee Lee, Pran K. Paul, Lu Ma, William D. McCulloch,

Sriram Krishnamoorthy, Yiying Wu, Aaron R. Arehart, and Siddharth Rajan. "Layer-

transferred MoS2/GaN PN diodes." Applied Physics Letters 107, no. 10 103505 (2015).

Lee II, Edwin W., Lu Ma, Digbijoy N. Nath, Choong Hee Lee, Aaron Arehart, Yiying

Wu, and Siddharth Rajan. "Growth and electrical characterization of two-dimensional

layered MoS2/SiC heterojunctions." Applied Physics Letters 105, no. 20 203504 (2014).

Ma, Lu, Digbijoy N. Nath, Edwin W. Lee II, Choong Hee Lee, Mingzhe Yu, Aaron

Arehart, Siddharth Rajan, and Yiying Wu. "Epitaxial growth of large area single-

crystalline few-layer MoS2 with high space charge mobility of 192 cm2 V− 1 s− 1."

Applied Physics Letters 105, no. 7 072105 (2014).

Laskar, Masihhur R., Digbijoy N. Nath, Lu Ma, Edwin W. Lee II, Choong Hee Lee,

Thomas Kent, Zihao Yang et al. "P-type doping of MoS2 thin films using Nb." Applied

Physics Letters 104, no. 9 092104. (2014).

Conference presentations:

Choong Hee Lee, Sriram Krishnamoorthy, Pran K. Paul, Dante J. O'Hara, Mark R.

Brenner, Roland K. Kawakami, Aaron. R. Arehart, and Siddharth Rajan, "Large-area

SnSe2/GaN heterojunction diodes grown by molecular beam epitaxy", 59th Electronic

Materials Conference (2017).

viii

Sriram Krishnamoorthy, Yuewei Zhang, Edwin Lee, Choong Hee Lee, William

McCulloch, Jared Johnson, Lu Ma, Jinwoo Hwang, Yiying Wu and Siddharth Rajan,

"Modeling and Demonstration of High Current MoS2/GaN Interband Tunnel Junctions",

International Workshop on Nitride Semiconductors (IWN 2016).

Sriram Krishnamoorthy, Edwin Lee, Choong Hee Lee, William McCulloch, Yuewei

Zhang, Jared Johnson, Lu Ma, Jinwoo Hwang, Yiying Wu and Siddharth Rajan, " MoS2/

GaN Inter-band Tunnel Junctions (2D/3D Tunnel Junctions)", WOCSEMMAD

(Workshop on Compound Semiconductor Materials and Devices) (2016).

Sriram Krishnamoorthy, Edwin Lee, Choong Hee Lee, William McCulloch, Yuewei

Zhang, Jared Johnson, Lu Ma, Jinwoo Hwang, Yiying Wu and Siddharth Rajan, "High

Current Density p-MoS2/n-GaN Inter-Band 2D/3D Tunnel Junctions", 58th Electronic

Materials Conference, June 22-24, Newark, Delaware (2016).

Choong Hee Lee, Sriram Krishnamoorthy, Dante J. O'Hara, Roberto C. Myers, Roland

K. Kawakami and Siddharth Rajan, "Molecular Beam Epitaxy of GaSe on c-

Sapphire(0001) Using Valved Se Cracking Source", 58th Electronic Materials

Conference, June 22-24, Newark, Delaware (2016).

Edwin W. Lee II, Choong Hee Lee, P. K. Paul, L. Ma, W. D. McCulloch, S.

Krishnamoorthy, Y. Wu, A. R. Arehart, S. Rajan, "Electrical and Optical Characterization

of MoS2/GaN Heterojunctions Formed by Film Transfer", 11th Topical Workshop on

Heterostructure Microelectronics, Takayama, Japan (2015).

Edwin W. Lee II, Choong Hee Lee, P. K. Paul, L. Ma, W. D. McCulloch, S.

Krishnamoorthy, Y. Wu, A. R. Arehart, S. Rajan, "Electrical and Optical Characterization

of MoS2/GaN Heterojunctions Formed by Film Transfer", International Symposium on

Compound Semiconductors, University of California-Santa Barbara (2015).

Edwin W. Lee II, Choong Hee Lee, P. K. Paul, L. Ma, W. D. McCulloch, S.

Krishnamoorthy, Y. Wu, A. R. Arehart, S. Rajan, "Electrical and Optical Characterization

of MoS2/GaN Heterojunctions Formed by Film Transfer", 57th Electronic Materials

Conference (EMC), The Ohio State University (2015).

Choong Hee Lee, William McCulloch, Edwin W. Lee II, Lu Ma, Sriram Krishnamoorthy,

Jinwoo Hwang, Yiying Wu, and Siddharth Rajan, "Transferred large area single crystal

ix

MoS2 field effect transistors", 57th Electronic Materials Conference (EMC), Columbus,

OH (2015).

Edwin W. Lee II, M. R. Laskar, D. N. Nath, L. Ma, Choong Hee Lee, T. Kent, Z. Yang,

R. Mishra, M. A. Roldan, J.-C. Idrobo, S. T. Pantelides, S. J. Pennycook, R. Myers, Y.

Wu, S. Rajan, "p-type conductivity in MoS2 by Nb doping", 56th Electronic Materials

Conference (EMC), Univeristy of California-Santa Barbara, June (2014).

Edwin W. Lee II, M. R. Laskar, D. N. Nath, L. Ma, Choong Hee Lee, T. Kent, Z. Yang,

O. F. Shoron, R. C. Myers, Y. Wu, S. Rajan, "P-Type Conductivity in MoS2 and WS2 by

Nb Doping", 2014 MRS Spring Meeting, San Francisco, CA, April (2014).

Fields of Study

Major Field: Electrical and Computer Engineering

x

Table of Contents

Abstract ............................................................................................................................... i

Dedication ......................................................................................................................... iii

Acknowledgments ............................................................................................................ iv

Vita .................................................................................................................................... vi

List of Figures .................................................................................................................. xii

Chapter 1 Introduction..................................................................................................... 1

1.1. Van der Waals heterostructures............................................................................ 1

1.2. Van der Waals Epitaxy ......................................................................................... 5

1.2.1. Molecular beam epitaxy of 2D materials ...................................................... 5

1.2.2. Molecular beam epitaxy of GaN on 2D ...................................................... 10

1.3. Application – 2D/3D heterojunction for HBTs .................................................. 11

1.4. Overview of the thesis ........................................................................................ 14

Chapter 2 Properties of 2D semiconductors on GaN heterojunctions ................... 16

2.1. GaSe grown on GaN using MBE ....................................................................... 17

2.1.1. Experimental methods ................................................................................ 19

2.1.2. Growth of mechanisms of GaSe on GaN .................................................... 21

2.1.3. Two-step growth of GaSe on GaN substrate .............................................. 27

2.1.4. Microstructure of GaSe on GaN substrate .................................................. 33

2.2. SnSe2 grown on GaN using MBE ...................................................................... 37

2.2.1. Experimental methods ................................................................................ 38

2.2.2. Structural properties of SnSe2/GaN junction .............................................. 39

2.2.3. Electrical and optical properties of SnSe2/GaN junction ............................ 41

2.3. MoSe2 on GaN using MBE ................................................................................ 48

2.3.1. Experimental methods ................................................................................ 49

xi

2.3.2. MoSe2 growth on GaN substrate................................................................. 50

2.3.3. Nb-doping of MoSe2 film ........................................................................... 53

2.4. Conclusions ........................................................................................................ 57

Chapter 3 Properties of 3D semiconductor/MoS2 Heterojunctions ....................... 59

3.1. Properties of MBE-Grown GaN on MoS2 Heterojunctions ............................... 60

3.1.1. GaN growth on MoS2 using MBE .............................................................. 61

3.1.2. Fabrication of GaN/MoS2 diode and experimental results ......................... 72

3.2. Properties of ZnO/MoS2 heterojunction ............................................................. 84

3.2.1. Growth optimization of ZnO using ALD.................................................... 85

3.2.2. Properties of ZnO/MoS2 diode.................................................................... 88

3.3. Conclusions ........................................................................................................ 94

Chapter 4 Device fabrication technique for 2D/3D devices .................................... 95

4.1. Layer transfer of MoS2 films .............................................................................. 96

4.1.1. Transfer process using DI water ................................................................. 99

4.2. Digital etching of MoS2.................................................................................... 110

4.2.1. Layer-by-layer etching mechanism........................................................... 110

4.3. Conclusions ...................................................................................................... 123

Chapter 5 Conclusions and future work ................................................................. 124

5.1. Conclusions ...................................................................................................... 124

5.2. Future work ...................................................................................................... 126

5.2.1. Demonstration of 3D/2D/3D HBTs .......................................................... 126

5.2.1. Direct growth of p-type 2D materials on 3D semiconductors .................. 127

Bibliography .................................................................................................................. 129

xii

List of Figures

Figure 1. (a) The spectrum of electromagnetic radiation of various 2D materials and their

band gaps [16]. (b) Band structure of MoS2 depicting indirect-to-direct bandgap transition

from bulk to monolayer [13] and (c) a crystal structure of MoS2. ...................................... 2

Figure 2. Bandgap energies of various 2D materials as a function of lattice parameters.

Blue bold line indicates the lattice constants of III-Nitrides [17]. ...................................... 3

Figure 3. (a) Energy band line up of various 2D materials as a function of lattice mismatch

with respect to GaN, and (b) a comparison of John’s figure of merit for InP (blue line) and

GaN (red line) based HBTs............................................................................................... 12

Figure 4. Schematic illustration of quasi van der Waals epitaxy [36]. ............................... 8

Figure 5. Veeco Gen930 system for 2D and III-Nitride MBE system ............................. 11

Figure 6. Side and top view of GaSe unit cell with different polytypes. .......................... 18

Figure 7. RHEED patterns observed along [1010] azimuth for (a) sapphire substrate and

(b) GaSe film. (c) XRD spectra of GaSe grown at 400 oC with Ga:Se of 1:200 (black) and

1:100 (red). Se flux was at maintained at 1×10-6 Torr. ..................................................... 22

Figure 8. (a) AFM images of GaSe as a function of growth temperature and Ga:Se BEP

flux ratio. RMS surface roughness is marked in the image. Red boxes indicate growth

conditions to obtain relatively smooth surface morphology. (b) Surface morphology of

GaSe film grown at the optimized condition (400 oC, Ga:Se=1:100) showing atomic steps.

(c) Step height of GaSe film taken from the red line in (b). ............................................. 23

Figure 9. Se sticking coefficient and growth rate at different growth temperatures......... 25

Figure 10. AFM scan image showing particles removed on GaSe film after HCl etching.

........................................................................................................................................... 26

xiii

Figure 11. (a) RHEED pattern of GaSe showing coexistence of a- and m-planes. (b) XRD

pattern, and (c) Surface morphology of GaSe film grown on GaN substrate. .................. 27

Figure 12. (a) The XRD spectra for GaSe film grown at different condition with Se flux at

1x10-5 Torr. The asterisks indicate the substrate peaks of GaN (002) and Sapphire (006) at

34.5 and 42 degree, respectively. (b-e) RHEED patterns of GaN and GaSe along the [1120]

and [1010] azimuth showing basal plane alignment. (f) AFM image of the GaSe film

showing aligned triangular domains. ................................................................................ 28

Figure 13. (a) Schematic of the two-step growth of GaSe on GaN substrates. (b) RHEED

patterns of GaSe after the two-step growth. (c) XRD scan of GaSe after first nucleation

step (black) and second (red) low temperature growth step. (d) XRD phi scan at GaSe (103)

and GaN (102) planes confirming basal plane alignment. ................................................ 30

Figure 14. (a) AFM image of the GaSe after two-step growth. (b) Raman spectra of the

GaSe film grown after first (red) and second (blue) step. Substrate is also shown for

comparison. (c) Raman intensity mapping of the A11g peak over 20 μm by 20 μm. ........ 31

Figure 15. Photoluminescence spectra of two-step grown GaSe film. PL spectra for

Substrate (GaN/sapphire) is also included for reference. ................................................. 33

Figure 16. (a) Cross-sectional TEM image of GaSe film growth after first step at 575 oC.

(b) Magnified image from the boxed area in (a). (c) GaSe TEM image taken from the same

sample but different area. (d) and (e) Magnified images from (c). (f) Ball-and-stick model

of 𝜀- and 𝛽-GaSe types. 60o rotation of every other layer in GaSe structure in 𝜀-type turns

out to be 𝛽-type. ................................................................................................................ 34

Figure 17. (a) and (b) Cross sectional STEM images of GaSe after two-step growth taken

from different region. (c) Magnified image from boxed in (b). Ball-and-stick models of

GaSe and GaN are also presented. (c) Defects formed in the middle of GaSe film marked

with an arrow. ................................................................................................................... 35

Figure 18. (a)-(d) RHEED patterns of SnSe2 and GaN along the [1120] and [1010]

azimuth. (e) XRD spectra for SnSe2 on GaN/sapphire substrate exhibiting (001) family of

diffraction peaks. (f) XRD φ scan of GaN (103) and SnSe2 (101). .................................. 40

xiv

Figure 19. (a) 2 𝜇𝑚 × 2 𝜇𝑚 atomic force microscopy image of SnSe2 after growth with the

RMS roughness of 0.99 nm. (b) Raman spectra of SnSe2 on GaN/sapphire substrate with

characteristic Eg and A1g peaks. ........................................................................................ 41

Figure 20. (a) The I-V characteristic of SnSe2/GaN junction shows 9 orders of magnitude

rectification and ideality factor of 1.1. Inset of (a) shows optical microscope image of the

devices with different dimensions. (b) C-V characteristic of SnSe2/GaN diode. The 1/C2 is

linear with respect to voltage and 0.99 V of built-in potential was extracted. ................. 42

Figure 21. I-V data for multiple GaN/SnSe2 diodes. ........................................................ 44

Figure 22. Extracted doping density profile of GaN layer from C-V curve. .................... 44

Figure 23. (a) Temperature dependent I-V characteristics of SnSe2/GaN diode. (c)

Measured ideality factor as a function of temperature. 𝜂 exhibits an 15.8 meV characteristic

energy. (d) Forward-bias I-V-T data plotted with fits of theoretical TFE I-V relationship.

A characteristic energy of 15.8 meV and barrier height of 0.84 V best fit the data over the

temperature ranging from 100K to 400K.......................................................................... 45

Figure 24. (a) Internal photoemission (IPE) measurement results of SnSe2/GaN diode at

300 K. The linear fit corresponds to a barrier height of 1.03 eV. (b) Band diagram of

SnSe2/GaN with 1 eV barrier. The Fermi level (orange dots) is shown at zero energy. .. 47

Figure 25. Mo flux as a function of e-beam power and flux current and corresponding

deposition rate ................................................................................................................... 50

Figure 26. RHEED patterns of GaN (0 min) and MoSe2 along the [1120] azimuth showing

basal plane alignment. ....................................................................................................... 51

Figure 27. XRD and Raman spectra for MoSe2 film grown at various temperatures. ..... 52

Figure 28. AFM images of MoSe2 film grown at different temperatures. ........................ 53

Figure 29. RHEED patterns of Nb-doped MoSe2 along the [1120] and [1010] azimuth.

........................................................................................................................................... 54

Figure 30. XRD and Raman spectra for Nb-doped MoSe2 films. .................................... 55

Figure 31. AFM images of MoSe2 with different Mo:Nb flux ratio. ................................ 56

Figure 32. Raman spectra for GaN on MoS2 after GaN growth under different growth

condition of (a) Ga-rich, (b) N-rich varying growth temperature, and (c) Ga/N ratio. .... 61

xv

Figure 33. Growth diagram of GaN growth on MoS2. ..................................................... 62

Figure 34. (a) RHEED patterns for GaN with different growth thickness. (b) XRD spectra

and (c) AFM image of 500-nm GaN grown on MoS2. ..................................................... 63

Figure 35. (a) X-sectional TEM images of GaN/MoS2. Both mixed hexagonal and

zincblende structure of GaN are clearly shown. Corresponding structures with Ball-stick

models are shown in (b), and (c) is zoomed from (a). ...................................................... 64

Figure 36. XRD spectra for two-step growth of GaN on MoS2........................................ 65

Figure 37. AFM images for two-step-grown GaN film under different conditions. ........ 67

Figure 38. (a) STEM image of Ga-rich GaN on N-rich GaN. (b) Microstructure of unknown

material underneath N-rich GaN. Inset shows undamaged 10-nm-MoS2 film. ................ 68

Figure 39. STEM EDS mapping of MoS2 area after Ga-rich GaN growth ..................... 69

Figure 40. MoS2 degradation process depicted schematically along with STEM image . 70

Figure 41. Process for GaN and MoS2 etch rate calibration test. ..................................... 73

Figure 42. Surface morphologies of SiO2 and SiNx on MoS2........................................... 74

Figure 43. Schematic illustration of the GaN lift-off process using SiNx mask. Optical

microscope images after GaN growth and after the lift-off process. ................................ 75

Figure 44. Patterning Mo film on MoS2. Properties of various metal films as a hard mask.

........................................................................................................................................... 76

Figure 45. Schematically illustrated GaN lift-off using Mo mask patterned by wet etch. 77

Figure 46. Optical microscope images of Mo film with different thickness. ................... 78

Figure 47. A schematic diagram showing the process sequence of the GaN lift-off and an

optical microscope images corresponding to each step. ................................................... 78

Figure 48. Optical microscope image of final device structure and the I-V characteristics.

........................................................................................................................................... 80

Figure 49. A schematic diagram of a process for lift-off of GaN film thicker than SiO2/Mo

mask layers........................................................................................................................ 81

Figure 50. A schematic and microscope image showing the problem that occurs when lift-

off is performed using BOE before GaN on SiO2 is sufficiently removed. ...................... 82

xvi

Figure 51. A schematic and microscopic image showing the damaged GaN side wall

damaged after the lift-off and Ti/Au contact shorted to MoS2 ......................................... 82

Figure 52. Optical microscope image and the I-V characteristic of the GaN/MoS2 diode.

........................................................................................................................................... 83

Figure 53. Thickness of ZnO as a function of number of pulsing cycles at different growth

temperatures and the extracted growth rate. ..................................................................... 85

Figure 54. Carrier concentration, resistivity, and Hall mobility of ZnO films grown at

different temperatures. ...................................................................................................... 86

Figure 55. AFM images and roughness as a function of temperature. ............................. 87

Figure 56. Schematic process flow for ZnO/MoS2 diode and optical microscope image. 88

Figure 57. Optical microscope images of device fabrication process .............................. 89

Figure 58. I-V characteristics of ZnO/MoS2 devices. ....................................................... 91

Figure 59. Ball-stick models for binding of H2O and TMA on MoS2 [114]..................... 92

Figure 60. Schematic process flow of a bubbling transfer process and the photographs of

the actual bubbling transfer of MoS2 film on Sapphire. ................................................... 97

Figure 61. Optical microscope images, AFN and XRD spectra for transferred MoS2 film

onto SiO2/Si substrate. Cracks and wrinkles were induced during the process. ............... 98

Figure 62. Schematic of MoS2 transistor fabrication process and corresponding optical

microscopy images.......................................................................................................... 100

Figure 63. (a) Schematic of the large-area MoS2 transfer using DI water. (b) Centimeter

scale MoS2 film floating in the DI water. (c) The film was transferred to SiO2/Si wafer and

(d) TLM structures were fabricated on the film.............................................................. 101

Figure 64. XRD spectra of the as-grown and transferred MoS2 films. The clear (0002)

diffraction peak can be seen for transferred MoS2 sample. ............................................ 103

Figure 65. (a) AFM image of transferred MoS2 layer on SiO2/Si wafer with RMS of 1.23

nm. (b) Cross-sectional plot of the transferred MoS2 layer showing thickness of 10 nm.

......................................................................................................................................... 104

xvii

Figure 66. (a) Raman spectra for as-grown and transferred MoS2 film. The transferred

MoS2 sample retains similar peak position and intensity ratio. (b) Contour map of Raman

peak intensity ratio of E12g and A1

g for transferred MoS2 film over 100 μm2 area. ........ 105

Figure 67. I-V measurement from MoS2 device transferred from sapphire to another

identical sapphire substrate showing 15% reduction in current. ..................................... 106

Figure 68. (a) Output and (b) transfer characteristics from the MoS2 transistor device. (c)

Transfer characteristics in linear scale. (d) Rsh and Rc extracted from TLM structure. .. 107

Figure 69. The transfer characteristics with a dual sweep at different drain bias. The

hysteresis in this case is measured in terms of the voltage difference at (Left) half the

maximum current and (Right) Vth shift. 3.4V of hysteresis was obtained from the transfer

curves at low drain bias by sweeping the gate voltage from -8V to 8V. ........................ 108

Figure 70. Schematic diagram of layer-by-layer etching process using the cycles of

oxidation/removal of oxide layer. ................................................................................... 113

Figure 71. AFM image of selectively etched MoS2. The step height measurement displays

around 10 nm of etch depth. Optical microscope image (left top) also shows distinct color

contrast where the MoS2 survived in the bright region................................................... 114

Figure 72. Etch depth as a function of oxygen plasma exposure time. .......................... 115

Figure 73. XPS spectra of core levels of (a) Mo 3d and (b) S 2p and (c) O 1s obtained from

oxygen plasma treated and HCl wet etched MoS2 films. ................................................ 116

Figure 74. 30 um by 30 um AFM scans of selectively etched MoS2 film with (a) 1 cycle

and (b) 2 cycles of digital etch. (c) Etch depth for different number of cycles of digital etch

as a function of oxidation time under O2 plasma. ........................................................... 118

Figure 75. (a) Raman spectra of MoS2 film with thickness ranging from bulk (13ML) to

2ML after multiple cycles of digital etch. (b) Frequency difference of E12g and A1g (red)

and corresponding thickness (blue) as a function of number of layers. .......................... 119

Figure 76. (a) Optical microscope image of selectively etched MoS2 film. The dark areas

denote selectively etched regions in MoS2. (b) Line scan Raman spectra of the MoS2 film

along the dashed line in (a). (c) Contour map of frequency difference and (d) peak intensity

ratio of the E12g and A1g peaks taken from dashed area in (a). ....................................... 120

xviii

Figure 77. (a) AFM images of surface morphologies of MoS2 film etched down from bulk

to 2ML. (b) RMS roughness as a function of number of MoS2 layers. Orange diagonal lines

indicate the range of the RMS values. ............................................................................ 122

1

Chapter 1

Introduction

Since the discovery of graphene, interest in the various physical properties of a new

class of two-dimensional (2D) materials has increased greatly in recent years [1-3]. Even

though graphene has been extensively studied in the past few years for high frequency

electronics, the lack of a band gap hinders its application in logic and switching devices [4-

6]. Therefore, the research has been extended to other 2D materials that have finite band

gaps, such as transition metal dichalcogenides (TMDs) (e.g., MoS2, MoSe2, WS2, WSe2)

[3, 7-9], black phosphors (BP) [10, 11], and boron nitride (BN) [12]. These materials,

including graphene, have band gaps that span from infrared to ultraviolet wavelengths, as

shown in Figure 1(a). These 2D materials also have exceptional optical properties,

including the indirect-to-direct bandgap transition shown in Figure 1(b) [13], giant spin-

valley coupling [14], and the control of valley polarization and coherence [15]. In addition,

due to the wide range of bandgaps ranging from 0 to 5 eV as shown in Figure 1(b), there

are 2D materials that are electrically metallic, semiconducting, or insulating.

2

Figure 1. (a) The spectrum of electromagnetic radiation of various 2D materials and their

band gaps [16]. (b) Band structure of MoS2 depicting indirect-to-direct bandgap transition

from bulk to monolayer [13] and (c) a crystal structure of MoS2.

3

Furthermore, 2D materials have a unique structural feature—that is, each atomic

plane is covalently-bonded and sandwiched between two planes held together by weak van

der Waals (vdW) forces, as seen in Figure 1(c). Due to this feature in 2D materials, these

layered materials can be isolated and stacked with various semiconductors to produce a

variety of artificial heterostructures, which are completely different from conventional

three-dimensional (3D) heterostructures.

Figure 2. Bandgap energies of various 2D materials as a function of lattice parameters.

Blue bold line indicates the lattice constants of III-Nitrides.[17]

This can be exploited to combine dissimilar 2D or 3D semiconductors, enabling

new flexibility in fabricating various types of heterostructures. Figure 2 depicts the lattice

constants and bandgaps of various 2D materials. The blue bold lines indicate the lattice

4

constants of III-Nitrides and overlaps with some 2D materials with the bandgap of 1-2 eV,

as a result the epitaxial registry between the 2D materials and III-Nitrides is expected. For

instance, the lattice mismatch between GaN and MoS2 is below 1%, whereas lattice

mismatch between GaN and sapphire is 16%.[18] Before looking into 2D/GaN

heterostructure, there was an initial work on bulk GaN growth on lattice matched substrate

of MoS2 mineral rock to reduce dislocation density and residual stress by utilizing vdW

epitaxy [19]. Then the scalability of MoS2 has been improved by the development of

synthesis process which further improved the feasibly of using GaN epitaxy on MoS2

platform. Recently, Tangi et al. implemented the MBE growth III-Nitride on CVD-grown

monolayer MoS2 substrate as shown in Figure 3 (a). In addition to the bulk GaN epitaxy,

MoS2/GaN heterojunction was also focused on different applications such as highly

efficient photocatalytic water splitting (Figure 3(b)) and vertical diode for heterojunction

bipolar transistor (HBT) device (Figure 3(c)).

Figure 3. (a) Raman spectra for GaN/MoS2 heterojunction.[20] (b) MoS2/GaN for

photocatalytic water splitting.[21] (c) Band lineup of MoS2/GaN heterojunction.[22]

5

1.2.1. Molecular beam epitaxy of 2D materials

A majority of the research on synthesizing 2D materials has begun with top-down

methods of micromechanical and solution-based exfoliation from bulk minerals. Though

single crystal flakes can be obtained using this method, the lateral dimension of flake is

only few tens of micrometer. The other approach of bottom-up method is synthesizing 2D

films using various growth technique such as sputtering [23], vapor-solid growth [24, 25],

CVD [26, 27], metal-organic CVD (MOCVD) and atomic layer deposition (ALD) as

shown in Figure 4.

Figure 4. Various growth methods for TMDs [28]

6

In particular, the CVD process can produce large-area single-crystalline triangular

islands, which then can merge together and coalesce. Most studies on CVD synthesis use

SiO2 substrate and usually the grown 2D films suffer from random orientation of domains

and high-density grain boundaries due to the amorphous nature of the SiO2 substrate. On

the other hand, well ordered (60o rotation) lattice orientation of MoS2 film (Figure 5 (b),

(c)) was grown on sapphire substrate with terrace morphology. It was pointed out that the

van der Waals interaction can control the lattice orientation of epi-layer even if the

interaction is relatively weak. This indicates that even though vdW epitaxy circumvents

the strict requirement of lattice matching, the crystallographic orientation of the overlayer

is still defined by that of the substrate.

Figure 5. (a) Randomly oriented MoS2 triangular domains grown on SiO2 [29]. Schematic

drawing (b) and SEM image (c) showing the lattice orientations [30].

7

Although the CVD-grown films provide uniform coverage over the large scale

substrate, high growth temperature (> 700 oC) and intermixing at the heterojunction

interface still remains a challenge [31]. In contrast to CVD methods, MBE offers great

control of growth rate, which results in abrupt and sharp interface and precise doping

control for high-quality heterojunction. The illustration of MBE system and the photograph

of 2D MBE system that used for the growth study here is represented in Figure 6.

Figure 6. Illustration and photograph of MBE system

The initial work of vdW epitaxy was first done by Atsushi Koma in 1984 by

demonstrating the MBE growth of NbSe2 on MoS2 with ~10% lattice mismatch [32].

Unlike conventional MBE growth of 3D materials whose material quality is greatly

influenced by the lattice mismatch, such constraint is no longer applied to 2D materials

growth. Due to low density of out-of-plane bonds on the surface, these layered materials

have a relaxed lattice matching condition. The concept of vdW epitaxy extended from the

growth of 2D materials on 2D substrates to the heteroepitaxial growth of 2D materials onto

8

3D substrates or vice versa (i.e., quasi-vdW epitaxy) lately as shown in Figure 7. Recently,

epitaxial growth of GaSe on mica substrate was demonstrated even with a significant lattice

mismatch of 35%. [33] In addition to 2D material growth on 3D substrate, the use of vdW

epitaxy was also applied to the growth of 3D semiconductor on an identical substrate that

is covered with monolayer graphene. It was shown that the strong potential field from

substrate penetrates the weak vdW potential of graphene and enables epitaxial growth, so

called remote epitaxy.[34] In addition, the 2D GaAs growth on graphene/ Si substrate was

also demonstrated in a similar manner.[35]

Figure 7. Schematic illustration of quasi van der Waals epitaxy [36].

However, the MBE growth technique for 2D material growth suffered from the

ability to obtain high-quality materials. With the exception for some cases, many groups

reported polycrystalline of MBE-grown 2D materials [37, 38]. Mostly, it seems that the

selection of substrates was not suitable or that it did not meet the proper process conditions.

Although the growth technique for large-area high-quality 2D materials using MBE is not

fully matured, compatibility of the fabrication process with existing silicon technology

9

encourages exploring new aspects of 2D/3D heterostructures. In this context, the

integration of 2D materials with wide bandgap semiconductors can create a new class of

transistors.

Figure 8. Binary phase diagrams of various 2D materials.

In order to implement the 2D/3D heterostructure, basic guidelines for growth of 2D

materials can be provided by the equilibrium phase diagram which can determine possible

reactions in reactants. Few phase diagrams of previously investigated 2D materials are

shown in Figure 8. For instance, Mo-S phase diagram shows broad region of

10

semiconducting 2H-MoS2 phase around 65% of Mo atomic percent and stoichiometric

phase as the temperature reaches 1600 oC. This indicates MoS2 can be grown in high-

quality low-defect material at the elevated growth temperature. In addition, other 2D

semiconductors such as WS2, GaSe, SnSe2 show lines of phase boundaries indicating the

feasibility of high quality film growth.

1.2.2. Molecular beam epitaxy of GaN on 2D

The difficulty of depositing 3D materials on 2D substrates has arisen in the process

of depositing high-k dielectrics for transistor fabrication due to the lack of chemical bonds

on 2D surface [39, 40]. Several studies have demonstrated uniform and conformal

deposition of Al2O3 on MoS2 using atomic layer deposition (ALD) by creating chemical

states using plasma or ozone treatment [41] or a deposition seed layer such as Al. However,

this technique is only limited to the deposition of oxide materials, and it is not a universal

approach to grow 3D materials. In the case of GaN—which is of interest in this work—it

is very challenging to grow III-Nitrides itself on MoS2 material. For instance, Gupta et al.

[42] demonstrated GaN growth on MoS2, and they showed complete decomposition of

MoS2 after growth. Although the phenomenon of etching of MoS2 after GaN growth is still

unclear, successful growth of GaN under N-rich growth condition has been reported [20].

Nevertheless, a comprehensive growth study of GaN on MoS2 has not yet been conducted,

and growth mechanism is still unclear. Understanding the growth mechanism is critical for

studying the intrinsic properties of GaN and GaN/MoS2 heterojunction as an emitter-base

11

in HBTs. Photographs of the Veeco Gen930 system used to grow 2D material and MBE

material are shown in Figure 9.

Figure 9. Veeco Gen930 system for 2D and III-Nitride MBE system

Compared to field effect transistors (FETs) utilizing a lateral geometry with in-

plane transport, the vertical structure of HBTs are preferred for linear power amplifiers due

to higher current density, higher power density, a linear current gain, and higher device

integration. In addition, unlike FET, HBT does not experience trap-induced frequency

dispersion, and has higher device transconductance as well as uniform threshold voltage.

So far, III-V based HBT devices are among the fastest devices ever fabricated and are

utilized in commercial wireless mobile communication [43, 44]. However, a low

12

breakdown voltage is an obstacle to further improving device performance due to the small

bandgap of the devices. This limitation can be removed by adapting wide bandgap

materials such as III-Nitrides, which provides higher power handling capability and

bandgap engineering. However, GaN HBTs are limited by the low free hole concentration

in p-GaN [45, 46] and surface damaging during dry etching [47]. Some attempts have been

made to overcome the problems by employing re-growth schemes [48, 49] or combining

III-Nitrides with III-Vs via wafer fusion technique [50, 51], which can circumvent the

lattice mismatch problem. However, due to the complex re-growth process and the

interface issues that appeared during wafer fusion, the development of a new type of

process/technology is needed.

Figure 10. (a) Energy band line up of various 2D materials as a function of lattice

mismatch with respect to GaN, and (b) a comparison of John’s figure of merit for InP

(blue line) and GaN (red line) based HBTs.

13

A successful demonstration of 2D/3D heterojunctions exploiting p-MoS2 with n-

GaN [22] points out the feasibility terahertz (THz) applications using p-MoS2/n-GaN as

the base-collector junction of HBTs.

Figure 10(a) shows the bandgap energy of 2D materials and their conductivity types

with their band lineup with respect to GaN. This suggests the possibility of the fabrication

of HBT devices with an ideal band lineup through proper material selection and GaN

bandgap engineering. In particular, both the high saturation velocities and the large break

down voltage due to large bandgap of GaN are suitable for use as collectors in HBTs, while

the narrow gap p-type 2D material can be used as a base due to low sheet resistance.

The following figure of merit (FOM) calculations estimate the quantitative

advantage of 2D/GaN HBTs over conventional InP based devices shown in Figure 10(b).

This FOM is called Johnson’s figure of merit (JFOM) and is expressed as the product of

breakdown voltage and cutoff frequency, which measures the suitability of a

semiconductor material for high frequency and high power transistor applications. The

theoretical calculation shows that the performance of the 2D/GaN HBT can be improved

by 5 times than that of the InP HBT devices. Here, the red triangles represent the case of

using p-GaN as the base for HBTs. Therefore, the purpose of this study is to demonstrate

2D/3D heterojunction by growing 2D materials on GaN substrates and growing GaN on

2D materials for base-collector and emitter-base junction respectively.

14

The major aim of this work is to demonstrate the 3D/2D/3D heterostructure by growing

various kinds of 2D materials on GaN substrates and vice versa. This thesis is organized

as follows. Chapter 2 of the thesis presents the growth of various 2D materials on GaN

substrate using MBE. The first section of the chapter investigates the growth mechanisms

of GaSe growth. A two-step growth method to grow coalesced crystalline GaSe on GaN

substrates is employed, and the structural properties are characterized by X-ray diffraction

(XRD), Raman spectroscopy, atomic force microscopy (AFM), and scanning transmission

electron microscopy (STEM). The microstructure of GaSe is comprehensively studied,

thus revealing different polytypes under growth condition, the preservation of surface

reconstruction, and randomly propagating defects in the GaSe film. In the following

section, the growth of SnSe2 and MoSe2 and their structural and electrical properties are

discussed. In particular, the band lineup of these 2D materials with GaN is estimated by

electrical and optical measurements, such as current-voltage (I-V), capacitance-voltage (C-

V), and internal photoemission (IPE) measurements.

Chapter 3 discusses the growth of wide bandgap semiconductors on MoS2. The first

part of the chapter discusses the growth mechanism of GaN using MBE. The difficulty of

GaN growth on MoS2 film is resolved by tuning the growth condition from Ga-rich to N-

rich regime. In this chapter, a two-step growth is employed to obtain a smooth GaN layer

on MoS2, and the study also compares the microstructure of the film with that of GaN

grown under N-rich condition. The second part of this chapter introduces the growth of

15

ZnO on MoS2 using ALD. The optimum growth condition for uniform coverage of the film

over the large area is obtained by measuring the structural and electrical properties of ZnO

films. The electrical properties of ZnO/MoS2 are discussed as well.

Chapter 4 discusses the device fabrication techniques that are be essential for

2D/3D heterojunction devices. One technique is to transfer 2D film to arbitrary substrate

in centimeter scale. Unlike conventional transfer methods, no chemicals or organic

materials are used for clean and sharp heterojunction interfaces. To evaluate the transfer

process, back gated FETs are fabricated and characterized. In the following section, the

layer by layer removal of MoS2 is presented. Due to chemical inertness of MoS2, dry

etching, which has etch rate fluctuation, is the only option to control the number of layers

in 2D materials. To overcome such an issue, repeating oxidation/oxidized layer etching

cycles was established to perform precise etch control. It also shows that this technology

is not limited to MoS2 only but can also be applied to other 2D materials such as WS2.

16

Chapter 2

Properties of 2D semiconductors on GaN heterojunctions

Two-dimensional (2D) metal chalcogenides are of great scientific interest for

electronic as well as optical devices due to their unique structural, electrical and mechanical

properties such as wide range of bandgaps [15, 52], valley-polarized carriers [53, 54],

strong spin–orbit coupling [55], and superconductivity [56]. Recently, artificial stacking of

these layered materials is being widely explored to create heterostructures for novel

applications. Most of these studies have been carried out by transferring layered flakes or

films from minerals [1, 2, 57] or synthesized materials obtained using chemical vapor

transport (CVT) [7, 58] or chemical vapor deposition (CVD) methods [59, 60]. In contrast

to such stacking methods, epitaxial techniques such as metal organic chemical vapor

deposition and molecular beam epitaxy (MBE) provide a more practical approach to

achieving large area epitaxial materials with precise control of layer thickness and doping.

The absence of out of plane chemical bonds in 2D layered materials enables flexibility

for epitaxy of on 3D materials[32, 61], and can therefore enable combinations of materials

for devices such as HBTs, vertical tunneling devices [62], and hot electron transistors

(HETs) [63]. More recently, renewed interest in 2D materials has led to exploration of

17

MBE growth of several materials including GaSe[37, 64, 65], MoSe2[8, 38, 66-68], WSe2[9] and

HfSe2[69]. Previous work on MBE growth of metal dichalcogenides on 3D substrates have

shown epitaxial registry between the 2D material and 3D bulk substrates.[8, 9, 69, 70].

In this chapter, the growth mechanisms of 2D materials and the transport of 2D/3D

heterojunction devices are investigated. Two-step growth of GaSe is introduced for

coalesced crystalline film. The microstructure of GaSe is carefully studied as well. In the

following part, 2D/3D transport is studied using SnSe2/GaN junction. The single-crystal

growth and transport mechanisms are presented. MBE growth of Nb-doped MoSe2 is also

attempted for demonstrating p-type 2D materials and the structural and electrical properties

are discussed.

For ideal HBT band alignment using 2D/GaN heterojunction, Mo or W based

dichalcogenides are the best candidates among other 2D semiconductors. However, the

growth study of 2D materials was initiated by growing GaSe due to higher adatom mobility

of Ga than that of refractory metals such as Mo and W, which provides better understanding

of growth kinetics of 2D materials.

GaSe has a layered crystal structure consisting of repeating units of covalently bonded

Se-Ga-Ga-Se held together by weak vdW force. Layered GaSe, however, occurs in several

polytypes displaying different stacking sequences, leading to 𝜀-, 𝛽- , 𝛾-, and 𝛿- phases of

18

the material shown in Figure 11.[71] Most common polytypes, 𝜀 (consists of two layers

per unit cell and has the space group, D13h) and 𝛽 (consist of two layers and has the space

group, D46h), have a 2H stacking sequence.[72] Bulk 𝜀-GaSe is a 2 eV direct bandgap

semiconductor and has been explored for applications in nonlinear optics, photovoltaics

and photodetectors.[33, 73]

Figure 11. Side and top view of GaSe unit cell with different polytypes.

Single crystal MBE growth of GaSe on GaAs(111)B substrates has been reported

by Keiji et al.[74] It has also been shown that GaSe and Ga2Se3 can be grown on GaAs

(001) substrates depending on the surface reconstruction.[75] Vinh et al. demonstrated the

growth of single crystal GaSe film on Si(111) substrate with 7 × 7 surface

reconstruction.[76] In addition, recent studies report that the growth of GaSe on sapphire

substrates produces crystalline films with random in-plane orientation of the domains.[37]

19

However, there have not been reports on GaSe growth on wide bandgap

semiconductors such as gallium nitride (GaN). Epitaxially grown high quality 2D materials

on GaN can enable vertical 2D/3D heterostructures[22, 77] that can enable vertical

tunneling devices33, HBTs, and HETs. We demonstrate the growth of highly crystalline

centimeter-scale few layer GaSe films on bulk 3D materials such as sapphire and GaN.

First, we have investigated the growth of continuous GaSe film on sapphire substrates at

various growth conditions and utilized the optimized condition to grow GaSe on a GaN

substrate. We report a two-step growth method to grow crystalline GaSe on GaN substrates

by employing a high temperature nucleation step for growth of single crystal domains

followed by the second step at lower growth temperature to achieve coalescence of the

film.

2.1.1. Experimental methods

MBE growth was performed in a Veeco Gen930 system with a standard thermal

effusion cell for gallium. While previous reports on GaSe growth use the standard

Knudson-type effusion cell to evaporate selenium, in this work, we use a valved cracker

source to supply Se. Se was evaporated using a valved cracker source with the cracker zone

at 950 oC in order to obtain Se2 species of Se.[78, 79] Growth was monitored in-situ using

reflection high-energy electron diffraction (RHEED). Prior to the growth, c-plane sapphire

and Fe-doped insulating GaN(0001)/sapphire substrates were solvent cleaned, annealed at

400 °C under ultra-high vacuum conditions (1×10-9 Torr) and loaded into the growth

20

chamber (base pressure 7×10-10 Torr). Sapphire substrates were then annealed at 850 oC in

the growth chamber for 30 minutes before ramping down the substrate temperature for

GaSe growth (400-500 oC). The Gallium sub-oxides on GaN substrates were removed in-

situ prior to the growth by using the following Ga polish technique. GaN substrates were

exposed to a Ga flux of ~1×10-8 Torr until the RHEED showed an amorphous pattern at

400 oC. The substrate was then heated to 700 oC for 30 minutes, followed by a ramp down

to the growth temperature. Streaky RHEED patterns with Kikuchi line patterns were

obtained prior to initiation of GaSe growths on sapphire and GaN substrates. The substrate

temperature was measured using the thermocouple attached with the continuous azimuthal

rotation (CAR) heater. The BEP of Se was fixed at 1×10-6 and 1×10-5 Torr during growth

and was measured using a nude ion gauge with a tungsten filament. Samples were grown

at different substrate temperatures (350-600 oC) and Ga:Se flux ratios. The growth was

initiated by opening the Se shutter for 2 minutes followed by opening of the Ga source at

the growth temperature.

The crystalline quality of the GaSe films was evaluated through X-ray diffraction

(Bruker, D8 Discover) and Raman spectra (Renishaw) with a 1 mW laser at 514 nm. The

surface morphology of the samples was examined by atomic force microscopy (AFM)

(Veeco Instrument, DI 3000). The microstructure of GaSe was examined by cross-sectional

scanning transmission electron microscopy (STEM). Due to the oxidation of GaSe in

ambient conditions [80], AFM scans were performed immediately after the growth. XRD

was measured after covering the GaSe surface with SPR955 photoresist. For the STEM

measurements, the photoresist was removed using solvents and the surface was capped

21

with Au metal immediately to prevent oxidation. Graphical illustrations of GaSe crystal

structure was generated using VESTA software [81].

2.1.2. Growth of mechanisms of GaSe on GaN

Growth of GaSe was explored on c-plane sapphire substrates by varying the

substrate temperature and the Ga:Se flux ratio. C-plane sapphire was chosen due to the

hexagonal symmetry of the basal plane, which is similar to that of GaSe, and the high

chemical and thermal stability of sapphire. The substrate temperature was varied from

350oC to 500o C, while changing the Ga:Se ratio from 1:50 to 1:200, and holding the Se

flux at 1×10-6 Torr. Growth was performed for one hour. The Se shutter was opened for

two minutes and the streaky RHEED pattern of sapphire substrates (Figure 12(a)) remained

before the opening of the Ga shutter indicating that the sticking of Se adatoms is very poor

in the absence of Ga flux. Upon opening the Ga shutter, the RHEED pattern corresponding

to m- (101̅0) and a- (112̅0) planes of GaSe was observed, and the RHEED pattern did not

change along the different azimuths (i.e. in-plane rotation of the substrate). The coexistence

of the RHEED patterns corresponding to the m- and a-planes of GaSe was also reported

earlier.[37] This indicates that GaSe nucleated with random in-plane orientation. However,

no polycrystalline rings were observed in the RHEED. The inverse of the ratio of spacing

between m-plane and a-plane streaks in the RHEED image (Figure 12(b)) was measured

to be 1.72, which is very close to the theoretical value of √3. XRD spectra of the samples

grown in the range of conditions mentioned above, with the exception of the extremely Se-

22

rich condition (Tsub = 400 oC, Ga:Se = 1:200) showed diffraction peaks corresponding to

the (002) family of planes in layered-GaSe. However, with extremely Se-rich condition,

the Ga2Se3 phase was observed in XRD and a spotty RHEED pattern was observed. The

growth window for GaSe in order to maintain a streaky RHEED pattern was found to be

very narrow at a given substrate temperature. The RHEED pattern remained streaky and

the intensity remained constant only at a certain Ga flux at a given substrate temperature;

Higher Ga flux resulted in complete RHEED dimming and lower Ga flux resulted in a

spotty RHEED pattern.

Figure 12. RHEED patterns observed along [101̅0] azimuth for (a) sapphire substrate and

(b) GaSe film. (c) XRD spectra of GaSe grown at 400 oC with Ga:Se of 1:200 (black) and

1:100 (red). Se flux was at maintained at 1×10-6 Torr.

23

Surface morphology of GaSe films as a function of growth conditions is shown in

Figure 13(a). At 350 oC, a Ga:Se ratio of 1:50 resulted in a relatively smooth surface

morphology, while a reduced Ga flux (Ga:Se = 1:100) was required at a growth temperature

of 400 oC. With an increase in substrate temperature from 350 oC to 450 oC, relatively

smooth surface morphology could be maintained only with a reduction of Ga flux.

Figure 13. (a) AFM images of GaSe as a function of growth temperature and Ga:Se BEP

flux ratio. RMS surface roughness is marked in the image. Red boxes indicate growth

conditions to obtain relatively smooth surface morphology. (b) Surface morphology of

GaSe film grown at the optimized condition (400 oC, Ga:Se=1:100) showing atomic

steps. (c) Step height of GaSe film taken from the red line in (b).

24

This can be explained as follows. With an increase in substrate temperature, the sticking

coefficient of Se is expected to reduce exponentially and hence the Ga flux that is required

to maintain stoichiometry at the surface is lower at higher substrate temperatures. This is

also expected to result in a reduction in the growth rate of GaSe with an increase in the

substrate temperature, assuming unity sticking coefficient for Ga adatoms at the growth

temperature used. This observation is in agreement with the RHEED patterns observed

during the growth. At the optimized conditions, where the adsorbed Ga and Se adatoms are

close to stoichiometry, the RHEED pattern remained streaky throughout the growth.

However, when the Ga flux is higher than the stoichiometry (Ga:Se =1:50, Tsub = 400 oC,

450 oC) the RHEED showed an amorphous pattern indicating the presence of excess Ga on

the surface during the growth. With Se-rich conditions, a spotty (i.e. rough) RHEED pattern

as observed. At higher substrate temperatures (>500 oC), the Se sticking coefficient is very

low and no growth was observed. At the optimized conditions with streaky RHEED pattern

and bright RHEED intensity, atomic steps were clearly observed (Figure 13 (b)). The step

height measured from AFM (0.8 nm) matches closely with the thickness of monolayer

GaSe. The growth rate was found to be 1.25nm/min with a total film thickness of 75 nm.

In order to estimate sticking coefficient of Se, we use the measured Ga flux (BEP)

needed to obtain a relatively smooth surface at various temperatures (350-450 oC). The

sticking coefficient for Ga was assumed to be 1, which is expected at such low temperatures

(

25

340 360 380 400 420 440 460

0.004

0.006

0.008

0.010

0.012

0.014

0.016

0.018

0.020

0.022

Temperature (oC)

stickin

g c

oe

ffic

ien

t

1.0

1.2

1.4

1.6

1.8

2.0

Gro

wth

ra

te (

nm

/min

)

Figure 14. Se sticking coefficient and growth rate at different growth temperatures.

The Ga flux required for a relatively smooth surface morphology (Marked in red

boxes in Figure 13(a)) at different temperature of 350, 400, 450 oC is 2E-8, 1E-8 and 5E-9

Torr respectively. The Se flux that incorporates can be calculated as 1.1E-8, 5.8E-9, and

2.9E-9 (accounting for mass difference and cell temperature between Ga and Se using

Hertz-Knudsen equation). Incoming Se flux was fixed at 1E-6 Torr, and sticking

coefficient is the ratio of the Se flux that incorporates into the crystal to the incoming flux.

Hence the estimated sticking coefficient at 350, 400 and 450 oC is 1.1E-2, 5.8E-3 and 2.9E-

3 respectively, and is plotted in Figure 14.

26

Figure 15. AFM scan image showing particles removed on GaSe film after HCl etching.

On the other hand, some particles were observed on the surface. It is suspected that

the particles to be Ga droplets, since pure Se dose not stick on the surface higher than 100

oC, confirmed experimentally shown in Figure 15. HCl treatment was performed, which

would nominally remove excess Ga. An etching test on a specimen with many particles on

the surface (Ga:Se=1:50, 400 oC) was deliberately conducted. The images shown below

are before and after etching using HCl, it can be confirmed that most of the larger particles

have disappeared.

Using the optimized growth conditions obtained from growth studies on sapphire

(Tsub = 400 oC, Ga:Se 1:100), we next explored the growth of GaSe on GaN templates (2

µm GaN/sapphire). The films were grown for one hour and the growth rate was found to

be 0.75 nm/min with a total film thickness of 45 nm. While the lattice mismatch between

GaN and GaSe (18%) is high, GaN provides a direct route for device design using 2D/GaN

heterostructure based devices.

27

Figure 16. (a) RHEED pattern of GaSe showing coexistence of a- and m-planes. (b) XRD

pattern, and (c) Surface morphology of GaSe film grown on GaN substrate.

X-ray diffraction of the films (Figure 16(b)) showed diffraction peaks corresponding to

(002), (004), (006) and (008) planes of GaSe. Complete surface coverage with spiral

hillocks and atomic steps (RMS roughness = 0.85 nm, Figure 16(c)) was obtained.

However, the sample showed in-plane disorder (Figure 16(a)) showing both m-plane and

a-plane spacing, and the RHEED pattern was insensitive to substrate rotation.

2.1.3. Two-step growth of GaSe on GaN substrate

Control of in-plane orientation of the crystal domains during nucleation is very

critical to obtaining single crystalline GaSe films. GaSe growth temperature was hence

increased from 400 oC to 575 oC to control the in-plane orientation. Higher growth

temperature necessitates higher Se flux due to reduction in sticking coefficient of Se with

increase in the substrate temperature. This is qualitatively similar to the effect of flux ratios

28

observed at lower growth temperature of 400 oC. The Se beam flux was increased to 1×10-

5 Torr. Figure 17(a) shows XRD spectra of GaSe films grown at different growth conditions

on GaN substrates. The sample grown at 500 oC with at a ratio of 1:400 showed (111)

oriented Ga2Se3 phase due to excess Se.

Figure 17. (a) The XRD spectra for GaSe film grown at different condition with Se flux

at 1x10-5 Torr. The asterisks indicate the substrate peaks of GaN (002) and Sapphire

(006) at 34.5 and 42 degree, respectively. (b-e) RHEED patterns of GaN and GaSe along

the [112̅0] and [101̅0] azimuth showing basal plane alignment. (f) AFM image of the

GaSe film showing aligned triangular domains.

With an increase in Ga:Se ratio from 1:400 to 1:100 diffraction peaks corresponding to

both GaSe(002) and Ga2Se3(111) planes were measured. With further increase in Ga:Se

ratio to 1:100, only GaSe(002) was detected at higher growth temperature of 575 oC.

Figure 17(b)-(e) show the RHEED patterns of GaN substrate and GaSe film grown at 575

oC with Ga:Se = 1:100 along the [112̅0] and [101̅0] directions of the GaN substrate. The

29

RHEED patterns corresponding to m- and a-planes of GaSe (Figure 17(d) and (e),

respectively) were observed along the same azimuth as GaN. The basal planes of GaSe was

found to be perfectly aligned with the GaN substrate ([11 2̅0]GaSe//[11 2̅0]GaN and

[101̅0]GaSe//[101̅0]GaN) and six-fold symmetry of GaSe was clearly observed. Unlike

the film grown at 400 oC with in-plane disorder, GaSe streaks corresponding to m- and a-

planes of GaSe appeared only at every 60o azimuthal rotation spacing. The inverse of the

RHEED spacing ratio between GaN and GaSe was found to be 1.170, which is very close

to the ratio (1.173) of bulk lattice constants of GaSe (3.74 nm) and GaN (3.189 nm). This

clearly suggests that the epilayer is fully relaxed and the growth proceeds by van der Waals

epitaxy. While the higher temperature growths led to single phase films, surface coverage

was found to be incomplete. A step height corresponding to 1 ML of GaSe (0.8 nm) was

measured at the edge of a triangular domain that grew on top of another triangular domain.

Large area (10 μ m × 10 μ m) AFM scan and STEM measurements confirmed the

observation of incomplete surface coverage from AFM scans. More details regarding the

microstructure of the film is discussed later in the manuscript.

While high temperature growth of GaSe at 575oC resulted in (002)-oriented single

crystal domains, the layers did not coalesce to form a continuous layer. Growth at 400 oC

with a Ga:Se ratio of 1:100 resulted in coalesced (002)-oriented GaSe layers with in-plane

disorder. To obtain single crystalline GaSe with complete surface coverage, we designed a

two-step growth method illustrated in Figure 18(a). After forming the nucleation layer at

575 oC with 1×10-5 Torr of Se beam-equivalent pressure (BEP) flux, the growth

30

temperature reduced to 400 oC with a reduced Se flux of 1×10-6 Torr followed by 30

minutes of GaSe growth with 1:100 of Ga:Se ratio.

Figure 18. (a) Schematic of the two-step growth of GaSe on GaN substrates. (b) RHEED

patterns of GaSe after the two-step growth. (c) XRD scan of GaSe after first nucleation

step (black) and second (red) low temperature growth step. (d) XRD phi scan at GaSe

(103) and GaN (102) planes confirming basal plane alignment.

Figure 18(b) shows the RHEED patterns along the [112̅0] and [101̅0] azimuthal

orientations. Six-fold symmetry was maintained after the second low temperature step,

indicating that the basal planes are aligned with the GaN substrate and there is no in-plane

disorder. Figure 18(c) displays XRD spectra of grown GaSe films after the first nucleation

31

step (black) and the second low temperature growth step (red). The GaSe layers grew along

the (002) orientation, and a higher order peak (006) was observed after second step growth

mainly due to the increased thickness of the film. No additional phase such as Ga2Se3 was

observed after the second step growth. An off-axis ϕ scan of the GaSe (103) plane was

performed, and six-peaks with 60o spacing were observed. The ϕ scan was repeated along

the (102) plane of GaN and six peaks were found at the identical azimuth angles as GaSe,

confirming the observation of basal plane alignment from RHEED.

Figure 19. (a) AFM image of the GaSe after two-step growth. (b) Raman spectra of the

GaSe film grown after first (red) and second (blue) step. Substrate is also shown for

comparison. (c) Raman intensity mapping of the A11g peak over 20 μm by 20 μm.

Figure 19(a) shows the surface morphology of GaSe after the two-step growth

process with a rms roughness of 1.1 nm. Surface coverage was found to be complete. Figure

19(b) shows the Raman spectra for GaSe grown after the first nucleation step (red), and the

second low temperature step (blue). The Raman mode corresponding to a shift of 143 cm−1

32

comes from the GaN/sapphire substrate. After the two-step growth, the Raman spectra

matches the typical spectra expected from bulk GaSe with Raman modes at 134.3 cm-1

(A11g), 211.7 cm-1 (E12g), 250.2 cm

-1 (E21g), and 307.6 cm-1 (A21g).[82] The A

11g and A

21g

modes correspond to the out-of-plane vibration modes, while the E21g and the E2

2g modes

are associated with the in-plane vibrational modes of GaSe. In contrast to the enhanced

intensity of these Raman peaks with the film thickness, no significant peak shift of A21g

mode due to the change in thickness[83] was observed because of sufficiently thick GaSe

film after the first step growth. The appearance of E21g peak in GaSe has been reported in

the literature.[33, 84] Nevertheless, at present the assignment of the new mode remains

unclear. In addition, it is difficult to differentiate the polytypes from the Raman spectra as

they show similar vibration modes.[85] Contour plot in Figure 19(c) shows the intensity

map of the dominant A11g Raman mode over a 20 μm × 20 μm area indicating complete

surface coverage. Thus, this two-step growth method enables formation of coalesced

multilayer GaSe films.

Photoluminescence spectrum for two-step grown GaSe film is shown in Figure 20.

Pronounced emission is observed at 1.99 eV which is attributed to the band edge emission

from ε-GaSe. We hypothesize that the peak at 1.61 eV could be arising from defects within

the band gap of GaSe. The other peak observed at 1.93 eV could be originating from the

mixed polytypes found in the sample.

33

Figure 20. Photoluminescence spectra of two-step grown GaSe film. PL spectra for

Substrate (GaN/sapphire) is also included for reference.

2.1.4. Microstructure of GaSe on GaN substrate

The microstructure of MBE-grown GaSe films were investigated in detail using

STEM measurements. STEM images from two regions of the GaSe nucleation layer grown

at 575 oC is shown in Figure 21(a) and (c). An abrupt GaSe/GaN interface and 5-8 GaSe

monolayers separated by van der Waals gaps could be clearly resolved in the STEM

images. Ball-and-stick model generated using VESTA is superimposed on the atomic

resolution image to identify the stacking sequence. The stacking sequence indicates that

the films grown are of the 𝜀-GaSe polytype, in Figure 21(b).

1.4 1.6 1.8 2.0 2.2 2.4

PL

in

ten

sity (

a.u

.)

Energy (eV)

Substrate (GaN/Sapphire)

GaSe two-step

34

Figure 21. (a) Cross-sectional TEM image of GaSe film growth after first step at 575 oC.

(b) Magnified image from the boxed area in (a). (c) GaSe TEM image taken from the

same sample but different area. (d) and (e) Magnified images from (c). (f) Ball-and-stick

model of 𝜀- and 𝛽-GaSe types. 60o rotation of every other layer in GaSe structure in 𝜀-

type turns out to be 𝛽-type.

However, a 60o rotation of the Se-Ga-Ga-Se tetralayer is observed in the region highlighted

in Figure 21(d), in which the Ga atoms sit on top of Se atom. Figure 21(f) shows the

simulated crystal structure of 𝜀-GaSe with a 60o rotation of every other layer resulting in

𝛽-GaSe polytype crystal structure. Such a rotation of the basal plane would not be captured

in the RHEED or XRD measurements due to the six-fold symmetry of both the 𝛽 and 𝜀

polytypes of GaSe. In spite of the rotation of the first tetralayer, subsequent GaSe stacking

is pure 𝜀-type. This may be attributed to the fact that the 𝜀 polytype is energetically more

stable than the 𝛽 type.[86] Similar lattice rotations and the resultant formation of grain

boundaries have been reported in the case of MoS2.[30, 87, 88] Dumcenco et al.[30] has

reported simulated data on the binding energies for MoS2 and sapphire substrate as a

35

function of orientation angle of MoS2 grains. It was pointed out that only 0o or 60o

orientations of the lattice were energetically favorable and stable.

Figure 22. (a) and (b) Cross sectional STEM images of GaSe after two-step growth taken

from different region. (c) Magnified image from boxed in (b). Ball-and-stick models of

GaSe and GaN are also presented. (c) Defects formed in the middle of GaSe film marked

with an arrow.

The microstructure of the coalesced GaSe films grown using the two-step method

was also investigated using cross-section STEM. Total number of layers after two-step

growth was found to be 25-27 from STEM measurements, and 20-22 layers were grown in

the second step. This implies a growth rate of 0.7 nm/min, which is similar to the low

growth temperature (Tsub = 400 oC) sample. The first five layers are identical to the

nucleation sample discussed in the previous section. A region with 60o rotation of first

layer was also observed in the two-step sample and is shown in Figure 22(a). However,

inclusions of 𝛽-type is observed along with the dominant 𝜀-type GaSe. Figure 22(b) shows

36

a magnified image of a region cropped from the boxed region in Figure 22(a). Surface

reconstruction of the GaN surface can be clearly observed in the image. Ga atoms (red

arrow) at the surface are bonded directly to a Ga atom below it, suggesting a 1×1

reconstruction of Ga atoms. On top of the surface Ga atoms, two atoms (green arrows)

were observed above every second Ga atom. We hypothesize that these could be Se atoms

passivating the GaN surface. This suggests that van der Waals epitaxy can be used to

maintain surface reconstructions on the GaN surface, and which could have important

implications for Fermi level pinning and dangling bond termination at heterostructure

interfaces. The electronic properties of these artificial two-dimensional interfacial layers

could be of great interest, but are outside the scope of the present work. We also observed

that defects formed in one area of GaSe film did not propagate along c-axis towards surface

due to the absence of bonding between individual 2D layers (Figure 22(c)). However,

certain amount of defect propagation is indeed observed and further careful study is

required to understand extended defects in 2D crystals. The GaSe growth study has

provided an overall understanding of 2D material growth. The growth rate is predominantly

determined by the amount of Ga flux. However, unlike Ga, migration-enhanced epitaxy

may be more effective in the case of TMD growth using refractory metal such as Mo, W,

or Nb. In addition, the transport study using GaSe could not followed up due to poor

stability of the film in air ambient. The oxidation of GaSe and change into a-Se was

reported earlier.

37

The growth mechanisms of 2D material was explored using GaSe in the previous

sections. However, due to the poor stability of the film was a major obstacle to perform

electrical characterization. Here, we report SnSe2/GaN diode to study the transport of

2D/3D heterojunction via electrical and optical measurements.

To date, band lineups for various heterojunctions between 2D and 3D materials have

been proposed. For instance, type-I band alignment was demonstrated in n-MoS2/p-Si [89],

p-MoS2/n-SiC [90] and p-MoS2/n-GaN [22, 91]. Unlike transition metal dichalcogenides,

Sn has two oxidation states (Sn2+ and Sn4+) which gives two stoichiometric phases, SnSe

and SnSe2. SnSe is an orthorhombic layered structure [92] with p-type conductivity [93],

while SnSe2 is intrinsically an n-type semiconductor [94] and is known to have two crystal

structures. One is the 2H phase with D6h (P63/mmc) symmetry and the other is the 1T phase

with D3d (P3̅m1) symmetry. The bulk 1T phase of SnSe2 has been reported to have a direct

energy band gap of 1 eV [94, 95] with electron affinity of 5 eV [70]. This high electron

affinity has been exploited to form type-III heterojunctions with black phosphorus [96] and

WSe2 [70].

In this paper, we report on the growth and electronic properties of SnSe2/GaN

heterojunctions. The combination of such a high electron affinity low band gap material

such as SnSe2 with a wide band gap material such as GaN presents a unique heterojunction

combination that is not possible with the III-Nitride system alone. While the band gap of

38

InGaN can be tuned to be as low as 1 eV, lattice mismatch between InN and GaN (11%)

makes it very challenging to grow high composition InGaN on GaN.

2.2.1. Experimental methods

The epitaxial growth of SnSe2 on GaN was performed in a Veeco GEN930 MBE

system with a standard thermal effusion cell for Ga and Sn. A valved cracker source (with

the cracker zone at 950°C) was used to evaporate Se. The sample surfaces were monitored

in-situ by reflection high-energy electron diffraction (RHEED) operated at 15 keV. The

structural quality of the SnSe2 films were evaluated through X-ray diffractometry (XRD)

(Bruker, D8 Discover) and Raman spectroscopy (Renishaw) equipped with a 514 nm laser.

The thickness of SnSe2 film was measured by X-ray reflectometry (XRR) (Bruker, D8

Discover). Atomic force microscopy (AFM) (Bruker Icon 3) was used to examine the

surface morphology of the film. VESTA software [81] was used to generate graphical

illustrations of the SnSe2 crystal structure.

Semi-insulating and n-type (0001) oriented GaN/sapphire substrates were used for

the study. Pre-growth surface preparation included solvent cleaning followed by a 1 hour

400°C anneal under ultra-high vacuum conditions (~1×10-9 Torr). Samples were then

loaded into the growth chamber (base pressure ~7×10-10 Torr) and exposed to Ga polish

procedure to remove gallium sub-oxides on GaN surface prior to the growth. The procedure

used is as follows. GaN surface was exposed to Ga flux of ~1×10-8 Torr at 400°C until the

RHEED intensity dropped. The substrates were then heated to 700°C for 30 minutes to

recover GaN RHEED pattern, followed by a ramp down to the growth temperature of

39

210°C. The substrate temperature was measured using a thermocouple attached to the

continuous azimuthal rotation (CAR) substrate heater.

For growth of SnSe2, the Se:Sn beam equivalent pressure (BEP) flux ratio

(measured using a nude ion gauge with a tungsten filament) was maintained at ~250. The

surface covered with Se by opening the Se shutter for two minutes. Growth was then

initiated by opening the Sn shutter. This procedure is qualitatively similar to that described

previously for the growth of GaSe on GaN[97]. Growth was carried out for 1 hour, and

terminated by closing all shutters, and immediately cooling down the sample to room

temperature.

2.2.2. Structural properties of SnSe2/GaN junction

Figure 23(a)-(d) show the RHEED patterns for the GaN substrate before growth,

and the SnSe2 film after growth was completed, along the [112̅0] and [101̅0] directions.

The streaky RHEED patt