PAMBE Growth and Characterization of Superlattice Structures in Nitrides

THESIS

Presented in Partial Fulfillment of the Requirements for the Degree Doctor of Philosophy in the Graduate School of The Ohio State University

Jing Yang

Graduate Program in Materials Science and Engineering

The Ohio State University

2013

Dissertation Committee:

Dr. Roberto C Myers (advisor)

Dr. Siddharth Rajan

Dr. Wolfgang Windl

Dr. Jay Gupta

Jing Yang

2013

Abstract

Superlattice structures formed using the III- nitrides family of semiconductors have attracted a great deal of attention due to some unique properties. Within the III- nitrides, the large conduction band offset between GaN/AlN and InN/AlN provides very large electron confinement that could be useful for ultrafast intersubband–based photonics. Second, the large spontaneous polarization difference at any nitrides heterostructure interface coupled with the large lattice mismatch (3%-13%), shows the potential for polarization-driven applications. Finally, the epitaxial integration of rare earth pnictides (RE-Pn), such as ErAs and ErSb, in III-As semiconductors has been intensively studied due to the applications in novel high speed photodetectors and photoconductive switches. They motivate the parallel work of embedding GdN in a GaN matrix such as a superlattice structure.

In the first part of this thesis, the growth of InN/AlN multiple quantum well structures by plasma assisted molecular beam epitaxy is presented. The InN/AlN multiple quantum wells are grown on top of coalesced AlN nanocolumns on Si(111) substrates.

The structural and optical properties of InN/AlN quantum wells are thoroughly studied by scanning transmission electron microscopy, x-ray diffractometry and photoluminescence.

STEM confirms the formation of InN quantum wells between AlN barriers. XRD data indicate the successful tuning of the ultra-thin InN quantum well thicknesses from sample

II to sample. Photoluminescence measurements show emission in the visible that shifts as a function of quantum well thickness. Moreover, the time decay of PL from the different quantum wells is detected by time correlated single photon counting.

In order to study the effects of growth conditions on interface sharpness in

GaN/AlN superlattices, we investigate the possibility and demonstrate the success of using growth temperature above the GaN decomposition during the deposition of

GaN/AlN superlattices in the second part this work. This high temperature growth condition is compared to well-established low temperature conditions. For superlattices grown thinner than their critical thickness, the N-rich high-temperature sample displays slightly superior structural quality. Reciprocal space mapping has been conducted to study the different relaxation mechanisms in high temperature and low temperature growth conditions.

In the third part of the thesis, we report on the integration of the RE-Pn GdN as discrete particles in a GaN matrix, as well as the formation of a GdN/GaN superlattice structure. It is hypothesized that the growth of GdN particles proceeds in a similar fashion to that of Re-Pn in III-As zincblende systems, with initial Re-Pn island formation followed by overgrowth of the surrounding uncovered III-As matrix. Periodic structures of GdN nano-island layers spaced between GaN regions were prepared and subsequently characterized by a variety of methods. High resolution XRD shows that the cubic rock- salt GdN islands are epitaxially oriented to the hexagonal wurtzite GaN matrix with the relationship GdN [111]||GaN[0001] with 2.4 ML GdN deposit. Cross-sectional STEM

III combined with in-situ reflection high-energy electron diffraction allows for the study of island formation dynamics, which occurs after 1.2 monolayers of GdN coverage.

Dedication

This document is dedicated to my parents Jun Rao and Ziji Yang

Acknowledgments

In these four and a half years, I own my gratitude to so many people who help and support me en route to my Ph.D, as well as completing this dissertation. First I would like to thank my Ph.D. advisor Prof. Roberto Myers, for teaching me what is science, and his patience for putting up with my need-to-improve-a-lot writing skills. His dedication to science and pursuing the perfection impresses me all the time. I would also like to thank

Prof. Siddarth Rajan, for being on my committee and mostly, for giving me so many insightful suggestions for my research projects along the way. I admire his kindness and bright mind of unlimited ideas. Prof. Wolfgang Windl, the most “fun” instructor I’ve ever known, led me experiencing the computing world when I took the computational class. I am grateful for the discussion with him regarding the GdN/GaN band alignment class project which led to a great collaboration. Also, I want to thank Prof. Suliman Dregia.

I’ve learned so much materials science knowledge from his classes. He is always very helpful and offers great and long discussion regarding class, project and life.

More importantly, I want to thank all my labmates. I can’t say enough how fortunate I am to work with those guys. They are the most friendly, supportive labmates I ever dreamed of working with and it’s like a big family for this group. For some nights, when I was staying up to write this thesis, some of them helped me editing it at the same time due the limited time I had. Santino Carnevale, the first labmate after me who joined

VI the Myers group, is like a big brother to me. He is always there to help me out with MBE growth, class notes, draft proofreading, and explains the English idiom and expressions for me constantly. I have learned so much from him about the hardworking ethic and the organization skills. Thomas Kent, the most handy person I’ve met, has the best skills of putting things together such that the optical lab is very well equipped and maintained. I learned a lot from him regarding the scientific writing and hands-on skills. I would also like to thank Zihao Yang and ATM Golam Sarwar for many great discussions and suggestions on my thesis; Brandon Giles for his support for my English writing. The support from this group is one of the main reasons that I am getting close to the finish line. I will always be grateful.

Also, this work could not have been done without the support of the MBE team:

Digbijoy Nath, Sriram Krishnamoorthy from Dr. Rajan’s group whom I had learned so much from and shared great laughs with; Mark Brenner the staff member who maintains the system; Andrew Carlin, Chris Ratcliff, Krishna Swaminathan from Dr. Ringel’s group; formal postdoc Alessandro Giussani, Javier Grandal and Masihhur Laskar; along with the growers within the group (Santino, Thomas, Sarwar). This is the best team I ever worked with. I wish to keep the friendship with them for a lifetime.

I would also like to thank my Chinese friends within the MSE and ECE department, Meng Tong, Siwei Cao, Lin Li, Lang Qin, Fan Yang, Yufeng Zheng,

Rongpei Shi, Yipeng Gao, Zihao Yang, Yibin Gao, Zeng Zhang, Huimin Wang, Tengfei

Jiang and more. Their support in academia and life got me through some difficult times.

VII

Finally, I would like to thank my boyfriend Andy. Without his love and cooking within the last couple of month, I can’t imagine how I would have survived the intense time of rushing to finish my PhD. I would like to thank my Aunt Gang and uncle Cipeng, without their encouragement and support; I would not apply for graduate school in

United State at the first place. Also, I would like to thank my parents, my grandparents and my uncle Zihua back in China. Their love makes who I am today.

VIII

Vita

2008...... B.S. Optical Science and Engineering,

Fudan University, Shanghai, China

2008...... Graduate student, Department of Electrical

and Computer Engineering, Purdue

University

2009 to present ...... Graduate Research Associate, Department

of Material Science and Engineering, The

Ohio State University

2011 ...... Master of Science

Publications

1. C. M. Jaworski, J. Yang, S. Mack, D. D. Awschalom, J. P. Heremans, R.C. Myers,

“Observation of the Spin-Seebeck Effect in a Ferromagnetic Semiconductor”, Nature

Materials 9, 898-903 (2010)

2. S. D. Carnevale, J. Yang, P.J. Phillips, M. J. Mills, R.C. Myers, “Three-Dimensional

GaN/AlN nanowire heterostructures by separate nucleation and growth process,” Nano

Letters 11, 866-871 (2011)

3. C. M. Jaworski, J. Yang, S. Mack, D. D. Awschalom, R.C. Myers, J. P. Heremans, “

Phonon spin distribution due to the Spin-Seebeck effect”, Phys. Rev. Lett. 106, 186601

(2011)

4. Z. Zhang, C. A. Hurni, A. R. Arehart, J. Yang, R. C. Myers, J. S. Speck, and S. A.

Ringel, “Deep traps in nonpolar m-plane GaN grown by ammonia-based molecular beam epitaxy”, Appl. Phys. Lett. 100, 052114 (2012)

5. T. F. Kent, J. Yang, L. Yang, M. J. Mills, and R. C. Myers, “Epitaxial ferromagnetic nanoislands of cubic GdN in hexagonal GaN", Appl. Phys. lett. 100, 152111 (2012)

6. Sriram Krishnamoorthy , Thomas Kent, Jing Yang, Pil Sung Park, Roberto C. Myers, and Siddharth Rajan, "GdN Nanoisland-Based GaN Tunnel junctions", Nano Lett., 13, pp

2570–2575 (2013)

7. J. Yang, F. Yang, T. F. Kent, M. J. Mills and R. C. Myers, “Semipolar InN/AlN multiple quantum wells on {10-1 5} faceted AlN on silicon”, Appl. Phys. lett.

Submitted.

8. Jing Yang, Santino D. Carnevale, Roberto C. Myers,“Growth and structural characterization of AlN/GaN superlattice above decomposition temperature of GaN”,

Journal of Vacuum Science and Technology B, submitted.

Fields of Study

Major Field: Materials Science and Engineering

Table of Contents

Abstract ...... II

Dedication ...... V

Acknowledgments...... VI

Vita ...... IX

Publications ...... X

Fields of Study ...... XI

Table of Contents ...... XII

List of Tables ...... XVI

List of Figures ...... XVII

Chapter 1 : An Introduction to Molecular Beam Epitaxy and Characterization Techniques for III-nitrides...... 1

1.1 Molecular beam epitaxy system overview ...... 1

1.2 PAMBE growth of III-nitrides ...... 4

1.2.1 PAMBE vs MOCVD ...... 5

XII

1.2.2 GaN growth conditions and a history of the establishment of growth phase

diagram ...... 6

1.2.3 AlN growth conditions ...... 9

1.2.4 InN growth conditions ...... 10

1.3 Superlattice and multiple quantum well structures ...... 11

1.4 Methods of surface, interface and structural characterization ...... 14

1.4.1 Reflection high energy electron diffraction (RHEED) ...... 14

1.4.2 Atomic force microscopy, X-ray diffraction, and transmission electron

microscopy...... 14

Chapter 2 : InN/AlN Multiple Quantum Wells ...... 16

2.1 Introduction to the InN/AlN MQWs ...... 16

2.2 Growth and structural characterization of the InN/AlN MQWs ...... 18

2.2.1 Growth conditions ...... 18

2.2.1 STEM...... 21

2.2.2 Layer thickness and relaxation state ...... 23

2.2.3 Semipolar facet plane identification ...... 25

2.3 Polarization for semipolar planes ...... 30

2.4 Photoluminescence of MQWs with different InN thickness ...... 36

XIII

2.5 Growth window exploration ...... 41

Chapter 3 : GaN/AlN Multiple Quantum Wells ...... 45

3.1 Introduction to high temperature growth condition for GaN/AlN MQWs ...... 45

3.2 Reciprocal space mapping technique ...... 48

3.3 Growth of GaN/AlN MQWs ...... 52

3.4 Structural characterization by AFM, RHEED and XRD ...... 55

3.5 Relaxation mechanism ...... 58

3.6 An evaluation of high T growth conditions for both GaN and AlN ...... 64

Chapter 4 : GdN/GaN superlattice structure ...... 69

4.1 Introduction to rare earth incorporation into III-As ...... 69

4.2 Gd doping cell calibration ...... 75

4.3 Growth of GdN/GaN SL ...... 79

4.4 Structural characterization of GdN/GaN SL ...... 84

Chapter 5 : Conclusion...... 88

References ...... 91

Appendix A: Optical measurement for intersubband transition ...... 96

Appendix B: Gd:AlN/GaN 2DEG system ...... 100

Appendix C: Magnetic properties and device processing for spin Seebeck in GaMnAs 104

XIV

C.1 Magnetic properties of GaMnAs ...... 104

C.2 GaMnAs spin Seebeck device processing ...... 107

Appendix D: FTIR measurement for GaN/AlN nanowires ...... 110

D.1 MQWs on nanowires ...... 110

D.2 FTIR measurements on MQWs nanowire samples ...... 112

D.3 Absorption on MQWs drop-cast onto glasslides ...... 113

List of Tables

Table 1: Sample database of InN/AlN MQWs...... 20

Table 2: Inclined angles and Bragg’s angle of selective planes in wurtzite AlN...... 28

Table 3: Elastic constants, lattice parameters, piezoelectric coefficients and spontaneous polarization of AlN and InN...... 33

Table 4: Sample database of RSM and XRD analyses...... 59

XVI

List of Figures

Figure 1.1: Simplified representation of Plasma Assisted MBE growth chamber showing the effusion cells, the nitrogen plasma source, and the relative positions of the substrate stage, RHEED gun and phosphor fluorescent screen...... 3

Figure 1.2: Bandgap dependence vs. lattice constants of III-nitrides and other semiconductors [8]...... 5

Figure 1.3: Growth phase diagram of GaN [18]...... 9

Figure 1.4: Growth phase diagram of AlN [24]...... 10

Figure 1.5: Growth phase diagram of InN [25]...... 11

Figure 1.6: Schematic representation of 3 types of multiple quantum wells and superlattices [27]...... 13

Figure 2.1: Cross-sectional TEM dark field image for the InN/GaN SLs grown at 600 °C

[32]...... 17

Figure 2.2: in-situ RHEED patterns during different growth stages...... 19

Figure 2.3: Cross-sectional tilt and higher magnification (inset) SEM image of the

InN/AlN MQWs sample...... 20

Figure 2.4: HAADF STEM images of MQWs structure of the InN/AlN sample...... 22

Figure 2.5: An illustration of in plane lattice constant with imaged atomic spacing...... 23

XVII

Figure 2.6: HRSTEM image with atomic spacing averaged result and XRD simulation result...... 24

Figure 2.7: HRSTEM image and atomic model for image plane and facet orientation identification...... 26

Figure 2.8: The projection the of {10- 5} family of equivalent planes onto the (11- 0) plane...... 27

Figure 2.9: The inclined MQWs from basal plane with large scale STEM image...... 28

Figure 2.10: Off-axis XRD scan about AlN (10-15) peak and in plane phi scan...... 29

Figure 2.11: Schematic representation of “natural” and “inclined” coordinates in wurtize structure...... 31

Figure 2.12: The inclination angle dependence of strain components and polarization at the InN/AlN interface...... 35

Figure 2.13: XRD and PL spectra on 3,4,5 ML samples...... 38

Figure 2.14: Time-resolved PL measurements for 3 ML and 5 ML samples...... 40

Figure 2.15: XRD scans of samples grown at different temperatures...... 42

Figure 2.16: Cross-sectional SEM of the InN/AlN samples grown at different temperatures...... 43

Figure 3.1: Absorbance of doped (top left),undoped samples GaN/AlN MWQs (top right)

; homogeneous broadening of absorption peak (bottom left) and inhomogeneous broadening of absorption peak (bottom right)[1]...... 46

Figure 3.2: A geometric illustration of Bragg’s law [58]...... 49

XVIII

Figure 3.3: Powder diffraction geometry (a) and high resolution diffraction geometry (b) of a XRD system [58]...... 50

Figure 3.4: Schematic drawing of the reciprocal lattice and diffraction condition (Ewald sphere construction) for the hexagonal nitride system...... 51

Figure 3.5: Different scan directions in reciprocal space [59] for the RSM...... 52

Figure 3.6: AFM image of a 1.5nm/20nm GaN/AlN SL structure...... 54

Figure 3.7: RHEED patterns of the test GaN/AlN MQWs structure...... 55

Figure 3.8: HRXRD scan about AlN (0002) diffraction peak for both high T and low T samples, with 9 periods (top) and 20 periods (bottom)...... 57

Figure 3.9：Reciprocal Space map around (11- 4) diffraction for 20 period sample, N- rich, high temperature ...... 61

Figure 3.10: Reciprocal Space map around (11- 4) diffraction for 20 period sample, Ga- rich, low temperature ...... 61

Figure 3.11: Reciprocal Space map around (11- 4) diffraction for 9 period sample, N- rich, high temperature ...... 62

Figure 3.12: Reciprocal Space map around (11- 4) diffraction for 9 period sample, Ga- rich, low temperature ...... 62

Figure 3.13: Reciprocal Space map around (11- 4) diffraction for sample 20 periods, Ga- rich, low temperature (low T buffer) ...... 64

Figure 3.14: RMS map of GaN growth condition probing above decomposition temperature under different conditions...... 65

XIX

Figure 3.15: Cross-sectional and tilted SEM images of selective Ga/N ratios at 820 oC. 66

Figure 3.16: AFM analyses of AlN homoepitaxial growth at 820 oC with different III/V ratios...... 67

Figure 4.1: STEM image [68] of ErAs incorporated into an InGaAs matrix and modeled structure [79] ...... 70

Figure 4.2: Time-resolved differential reﬂection traces on ErAs containing samples with different superlattice periods L [70]...... 71

Figure 4.3: Band diagrams of conventional (a) and ErAs inserted (a) tunneling junction

[71]...... 72

Figure 4.4: An illustration of epitaxy compatibility of GdN rock-salt on GaN wurtzite

[79]...... 74

Figure 4.5: In-plane relationship between GdN and GaN [77]...... 75

Figure 4.6: SIMS process illustration of surface material sputtering (from EAG’s website)...... 77

Figure 4.7: (Top) SIMS depth profile of the Gd:GaN calibration stack and (bottom) the doping concentration vs. cell temperature...... 78

Figure 4.8: The design of the TEM stack sample and the TEM image of the structure. .. 80

Figure 4.9: in-situ RHEED patterns for GdN deposition...... 81

Figure 4.10: Atomic resolution STEM images of GdN particles in GaN matrix...... 82

Figure 4.11: AFM image of the stack sample (GdN in GaN) surface...... 83

Figure 4.12: The illustration and TEM image of a 50 x GdN/GaN SL sample...... 85

Figure 4.13: HRXRD ω-2θ scan of the 50 x SL structure...... 86

Figure 4.14: AFM of the GdN/GaN SL surface...... 86

Figure A.1: Optical setup for transmission measurements of the ISB absorption...... 97

Figure A.2: Transmission spectrum of S and P-polarized light...... 98

Figure A.3: Absorption spectrum of Ga-rich and N-rich GaN/AlN MQWs samples...... 99

Figure A.4: An illustration of the Gd doped AlN/GaN HEMT structure...... 100

Figure A.5: Band diagram of the AlN/GaN HEMT...... 101

Figure A.6: SIMS depth profile of the Gd concentration...... 102

Figure A.7: Hysteresis loops of d=0.4nm and d=10nm samples at 5K ...... 103

Figure A.8: (a). Hysteresis loop raw data, (b) dominant diamagnetic signal at high field,

...... 105

Figure A.9: Hysteresis loop of GaMnAs (Mn%=15.8) along different applied magnetic field directions...... 106

Figure A.10: Magnetization vs. Temperature revels Tc of 150K...... 107

Figure A.11: A GaMnAs cleaved wafer and processed sample on a modified PPMS station...... 108

Figure A.12: The illustration of the Pt contacts geometry (first generation) and I-V characterization of the Pt contacts...... 109

Figure A.13: Second generation of the shadow mask...... 109

XXI

Figure A.14: Cross-sectional and top-view SEM images of the GaN/AlN MQW nanowires...... 110

Figure A.15: STEM images of MQW nanowires...... 111

Figure A.16: FTIR measurements on MQWs nanowire samples on reflectant mode. ... 112

Figure A.17: Absorption measurement on GaN/AlN MQWs NWs drop-casted onto glass slide...... 114

XXII

Chapter 1 : An Introduction to Molecular Beam Epitaxy and

Characterization Techniques for III-nitrides

1.1 Molecular beam epitaxy system overview

Molecular beam epitaxy (MBE) is one of several methods used to deposit thin films, which are grown in an ultra-high vacuum chamber at a high temperature, resulting in the epitaxial growth of atoms on a specific substrate. Due to its extensive application in fabrication of lasers and detectors because of its high accuracy in engineering a structure,

MBE is widely used to create single crystalline materials for semiconductors and oxide materials research. MBE was invented in the late 1960s by J. R. Arthur and Alfred Y.

Cho at Bell Telephone Laboratories, which was one of the most dynamic science research institutes in the world before they began concentrating less on basic science, material physics, and semiconductor research, and shifted its focus to more marketable products.

[2-5]

In theory, MBE is a relatively simple growth technique. It takes place in a high vacuum or ultra-high vacuum chamber, where the chamber maintains a very low base pressure between 10-9 to 10-11 Torr. A substrate is held at a controlled temperature and exposed to incident atoms or molecules beam that come out of separate cells containing different elements known as the source. Because of the ultra-high vacuum, MBE has a

1 very slow deposition rate, typically less than 1000 nm per hour, which allows the films to grow epitaxially. Since atoms grow layer by layer, the slow deposition rates require proportionally higher vacuum level to achieve the same impurity level as other deposition techniques. Simply put, the lower the deposition rate, the higher the vacuum level required. However, this is worthwhile because it guarantees precise fabrication in a target structure.

Typically MBE source materials are provided by one of the two following ways.

Either one can evaporate solid or liquid material from a furnace which is the cell in

Figure 1.1, or one can inject a gaseous species into the growth chamber. The high vacuum environment ensures that the mean free path of atoms and gaseous species is large, allowing the atom to travel a relatively long path before it collides with another atom or species. The high vacuum environment allows the beam to aim at the substrate surface directly.

Cryo pump

Cryo panel

for LN2 UHV chamber Car (substrate stage) Gate valve RHEED fluorescent RHEED screen gun Shutters

In Ga Al Effusion cells RF plasma source for nitrogen Figure 1.1: Simplified representation of Plasma Assisted MBE growth chamber showing the effusion cells, the nitrogen plasma source, and the relative positions of the substrate stage, RHEED gun and phosphor fluorescent screen.

The rate at which source material is deposited on the sample surface is controlled either by adjusting the temperature of a cell (Ga, Al, In as in Figure 1.1), or by controlling the rate at which a gaseous or vaporous source material is injected into the chamber, which is how the nitrogen plasma flow rate in Figure 1.1 is controlled.

Typically mechanical shutters are used to rapidly turn beam “on” and “off” by blocking the line of between the source of the beam and the substrate.

During operation, Reflection High Energy Electron Diffraction (RHEED) is used to monitor the growth of the materials. By observing of the RHEED pattern, one can tell the growth mode and the smoothness of the structure grown[6]. Moreover, a computer controlled shutter in front of every cell can open and close the shutter at the target times with very high accuracy, which helps controlling the thickness of the deposited structure with high precision. In this way, structures such as superlattice can be fabricated.

1.2 PAMBE growth of III-nitrides

The group III-nitrides have long been considered to be a promising material system for semiconductor device applications, because the bandgap of III-nitrides family spans a wide range and can be varying continuously from 0.7 eV (InN) to 3.4 eV (GaN) to 6.2 eV (AlN) [7] by changing the composition of an alloys as shown in Figure 1.2 [8].

As indicated in the figure, the InGaN and AlGaN alloys can span the entire visible spectrum while AlGaN and InAlN can cover a great part of the UV spectrum. Thus III- nitrides have great potential and versatility in optoelectronic applications. The large,

4 direct bandgap of these materials makes them advantageous for optoelectronic applications such as visible light emitters and solar-blind ultraviolet photodetectors [9].

Figure 1.2: Bandgap dependence vs. lattice constants of III-nitrides and other

semiconductors [8].

1.2.1 PAMBE vs MOCVD

Optimization of the materials for electronics applications can be achieved by RF plasma assisted molecular beam epitaxy (PAMBE). Early development of GaN growth was done primarily by utilizing metal-organic chemical vapor deposition (MOCVD) as the growth technique. MBE growth offers many advantages over MOCVD growth technology, such as low impurity incorporation due to the ultra-high vacuum environment, sharp interface control and in situ diagnostics such as RHEED, in order to monitor surfaces reconstruction during growth. In addition, PAMBE is a much lower

5 temperature process (e.g. 500 – 800 ºC) as compared to MOCVD (e.g. 900 - 1200 ºC).

The disadvantages of high temperature synthesis during MOCVD growth include thin film cracking due to thermal mismatch between substrate and layer grown on top, dopant redistribution in layers due to diffusion, and rough interfaces. However, MBE is limited to lower temperatures, which lead to the challenge of direct nucleation on non-lattice- matched substrates. Thus thin film grown by MBE prefers homoepitaxially on

‘templates’—lattice-matched substrates such as MOCVD-grown GaN (AlN) on c-plane sapphire. Only within the last few years, has PAMBE been successfully used for direct heteroepitaxy on substrates such as sapphire and SiC [10].

1.2.2 GaN growth conditions and a history of the establishment of growth phase diagram

It has been generally agreed in numerous publications that similar surface morphologies result in similar properties, thus establishing a growth diagram that can be used to identify the optimized growth conditions for thin films of desired morphology and properties [11, 12]. Although a great deal of work has been done in exploring different GaN growth conditions, namely Ga-stable and N-stable conditions [13-15],

Heying et al. developed the first growth diagram in 2000, defining Ga droplets, intermediate, and N-stable conditions in a substrate temperature vs. Ga flux diagram.

The three growth regimes are “Ga-droplets”—where Ga droplets are formed, “N- stable”—where low Ga flux is used, and “intermediate”—regime between the previous two. In early work, Ga-droplets, intermediate and N-stable region are reported to relate to

6 smooth, smooth with pits (or faceted edges), and rough surface morphologies respectively [11-13, 15, 16]. The boundary between the intermediate and N-rich region is easy to idenfy because it is a straight line of Ga flux corresponding to the stoichiometric point of GaN where active N-flux is equal to Ga flux. The key is to find the boundary between Ga-stable intermediate and droplets regimes. To do this, requires a series of growth at various Ga fluxes and substrate temperatures. The Arrhenius dependence of excess Ga flux vs 1/kT, which is fitted between the droplets samples and non-droplets samples, renders the activation energy for liquid desorption [11].

The growth map established by Heying is significant because it highlights the dependence of surface morphology on growth condition, namely, impinging Ga flux and temperature. However, there is no in situ characterization such as RHEED to monitor how the samples are grown. Moreover, the established growth phase diagram is at fairly low temperature (<700 oC) where GaN decomposition has not to be taken into consideration.

With the help of in situ characterization such as RHEED or QMS, further studies have completed more detailed work on the correlation of surface morphology and growth conditions within the growth phase diagram and extended to higher growth temperature.

A significant amount of work have been put into defining the growth mode within the growth phase diagram [17, 18], as well as 3-D and N-rich regimes at elevated temperatures [18-20]. First, Brown and Adelmann carried out Ga adsorption studies on the GaN surface. Their studies agree that Ga adlayer coverage increases with Ga flux to

7 form a 2.5 0.1 ML of Ga which defines the boundary of Ga-rich droplet region and Ga- rich intermediate region. Then, two different growth modes have been identified within the Ga-rich region, namely step-flow and layer-by-layer mode. Layer-by-layer growth mode can be identified by observing RHEED oscillations during growth. The transition from layer-by-layer to step flow mode happens at increased impinging flux to 1 ML [21,

22]. One of the most used conditions for electrical and optical quality considerations is with Ga fluxes just below the Ga-droplet formation, in other words, at the boundary of

GaN droplet and GaN-rich intermediate region. Not only can the condition produce high crystal quality materials, but also this cross-over point is achievable at all temperatures

[21, 23]. All the growth regions are shown in one map known as the GaN growth phase diagram, shown in Figure 1.3.

Figure 1.3: Growth phase diagram of GaN [18].

1.2.3 AlN growth conditions

The growth conditions for AlN have also been studied parallel to the establishment of the GaN growth phase diagram. The growth diagram of AlN is similar to the GaN growth diagram. With the variation of Al flux and substrate templates, the

AlN growth diagram also includes three regions, namely N-rich, Al-rich intermediate and

Al-rich droplet regions. In other words, the AlN homo-epitaxy growth diagram resembles the one for GaN, but with the temperature axis shifted to a higher temperatures range as shown in Figure 1.4 [24]. Although there has been more work done to study GaN growth kinetics and modes, it is at least known that in the Al-rich intermediate region for AlN growth, the smoothest AlN surface is achieved. The coalesced platelets render a surface

9 of monolayer smoothness enabling the 2-dimentianal growth in this region. In the more thorough study of GaN growth conditions, it has been generally agreed that pit density will decrease to zero with the increase of the adlayer thickness at or above 2.5ML [11,

18, 21, 22]. For AlN, on the contrary, pit density is zero at the intermediate regime, and is small at the droplet regime [24].

Figure 1.4: Growth phase diagram of AlN [24].

1.2.4 InN growth conditions

The growth of InN is the most complicated one among the III-nitrides family. InN decomposes at a lower temperature than the onset of In droplet desorbs from the surface.

Although there have been intensive studies regarding the PAMBE growth of InN, achieving highly crystalline InN thin films is still very challenging.

Figure 1.5: Growth phase diagram of InN [25].

1.3 Superlattice and multiple quantum well structures

The concept of superlattices (SLs) in semiconductors was first proposed by Esaki and Tsu in 1970 as a spatially modulated or artificial periodic structure consisting of alternating dissimilar layers A/B in the single crystalline materials setting [26]. When the period is small, the electrons are confined in a two-dimensional layer, thus the wave function of the structure can be calculated as in a one-dimensional periodic potential setting.

There is a very fine line between SLs structure and multiple quantum wells

(MQWs) structure. MQWs are similar in structure except that the quantum barrier thickness needs to be large enough to prevent electrons tunneling through. The SLs structure accentuates the alternating layers setting but is mostly assumed with a barrier width thin enough that wave functions do not decay to zero in the adjacent wells so the electrons can “see” the periodic potential as well.

The InN/AlN and GaN/AlN structures discussed in the following chapters are

MQWs because the AlN barrier (>5nm) is thick enough to prevent electron tunneling.

However, for the GdN/GaN alternating-layered structure, although GaN is not thin, we still call it a superlattice for convenience (quantum dot superlattice to be exact).

The MQWs or SLs can be categorized into three types due to the different band edge alignment thus the confinement energy schemes of the electrons and holes, as shown in Figure 1.6. In the type I setting, the electrons and holes are confined within the same layers known as quantum wells, thus the other layers are the barriers. Both InN/AlN and GaN/AlN MQWs have type I alignment. In type IIA setting, the electrons and holes are confined in the adjacent layers, thus this type is also called spatially indirect band gap or staggered type such as AlAs/GaSb SLs. For type IIB situation, not only the electrons and holes are confined in adjacent layers, but there is barely any energy gap between the confined electron and hole level. This is known as misaligned type such as InAs/GaSb

SLs. The natural band alignment is still under investigation for GdN/GaN SLs.

Figure 1.6: Schematic representation of 3 types of multiple quantum wells and

superlattices [27].

1.4 Methods of surface, interface and structural characterization

1.4.1 Reflection high energy electron diffraction (RHEED)

The RHEED technique is useful as an in situ surface monitoring method during epitaxial growth. It is highly surface sensitive that the pattern generated reflects the periodicity of the surface. Because of this fact, the intensity of the RHEED streaks is highly dependent on the order of the surface. For example, when the condition for the growth is layer by layer mode, as one complete mono-layer (ML) forms, the intensity of the streaks enhances. And when partial monolayers form, the streak intensity is weaker.

Thus by monitoring the streak intensity oscillations, the growth process and growth rate are directly monitored. Also, the streak geometry and spacing can render information such as the particular surface reconstruction, growth mode, and surface morphology

(surface corrugation) [6].

1.4.2 Atomic force microscopy, X-ray diffraction, and transmission electron microscopy

Other methods developed to study surfaces or interfaces have been used in this study, including atomic force microscope (AFM), X-ray diffraction (XRD), and transmission electron microscopy (TEM). AFM is one of the most common techniques to study the surface morphology post growth. It can be used for most of the surfaces because the force between an AFM tip and surface atoms drives deflection of a cantilever spring, causing changes in distance and thereby capacitance between the spring and an electrode. Thus surface morphology can be measured.

XRD is one of the most common techniques to characterize the structural quality based on Bragg’s law. It can be used for determining the arrangement and agreement of atoms. A lot of useful structure and interface information can be gathered. The more detailed explanation of XRD technique will be presented in Chapter 4.

TEM is a high resolution microscopy technique whereby a beam of electrons is transmitted through a sample when an image is formed from the interaction of the electrons transmitted through the specimen. It is one of the most direct and advance technique to probe the structure of the sample. In this study, STEM has been performed for most of the cases.

Chapter 2 : InN/AlN Multiple Quantum Wells

(This chapter is adapted from the Paper submitted to APL)

2.1 Introduction to the InN/AlN MQWs

InN-based MQWs are interesting to consider for optoelectronic applications because of the large confinement energy in conduction band using GaN or AlN as barrier layers [28-32]. In particular, it was recently shown that InN MQWs with just a few monolayers thick demonstrate visible wavelength interband photoluminescence (PL) even though the band gap of InN is in the near-infrared (0.65 eV) [32, 33]. Besides interband optoelectronics, ultrafast optoelectronics devices based on intersubband transitions could reach into the visible spectrum for InN/AlN MQWs due to the very large conduction band offset of 4 eV between InN and AlN [34, 35]. Moreover, the large polarization charges at the InN/GaN interface is predicted to lead to a spin-orbit- coupling-induced topological insulator state [36], which might also occur in InN/AlN due to the greater degree of polarization mismatch [37-39].

From the materials perspective, AlN and InN are the extreme end points of the

III-nitrides family. Though they both exhibit wurtzite crystal structure, their lattice mismatch is 13% and their melting points and bond energies are quite different as well

[7]. From this point of view, heterostructures of AlN and InN are considered unlikely. In

16 the case of InN/GaN (11% lattice mismatch), the possibility of epitaxial heterostructures is a bit more favorable. Yoshikawa et al. successfully realized the first InN MQWs embedded in a GaN matrix by PAMBE as shown in Figure 2.1. They demonstrated that for ultrathin layers (1 ML), InN layers form pseudomorphically (coherently strained) on

GaN [32].

Figure 2.1: Cross-sectional TEM dark field image for the InN/GaN SLs grown at 600 °C

[32].

Though InN/AlN heterostructures are expected to be more challenging to form than InN/GaN, it has been reported that In can be used as a surfactant for the growth of

AlN, providing a means to ensure high crystalline quality AlN [40]. Secondly, the intermixing effect observed in InN/GaN interfaces [41] can be largely circumvented in

InN/AlN heterostructures since the Al-N binding energy (2.88 eV) is much larger than the In-N binding energy (1.98 eV) (Ga-N binding energy is 2.20 eV) [42]. Here we explore the possibility of InN/AlN MQWs.

2.2 Growth and structural characterization of the InN/AlN MQWs

2.2.1 Growth conditions

In this study, multiple InN/AlN MQW structures were grown on n-type Si (111) wafers using a VEECO 930 PAMBE system with a base pressure of less than 9×10-11

Torr. Standard effusion cells are used for the In and Al sources with purity of 7N and 6N, respectively. Active nitrogen is supplied by a radio frequency plasma source operating at

350 W. The flow rate of 6N purity N2, is controlled by a mass flow controller operating at a flow rate (4.0-4.5 SCCM) adjusted to maintain an effective chamber pressure of 2×10-5

Torr during sample growths. The beam fluxes are measured using a nude ion gauge that can be rotated into the substrate’s growth position. RHEED, operating at 10 kV and 1.4

A, is used as an in situ surface technique to monitor the particular surface reconstruction of a material and the growth mode.

Before each growth, the Si (111) wafer is heated to 1000oC (measured by a calibrated infrared pyrometer) to remove the native oxide which is confirmed by emergence of a RHEED pattern reconstruction corresponding to the bare Si surface. After native oxide desorption, the AlN nano-columnar base is grown on the silicon substrate

18 held at 1000oC (pyrometer) under an Al/N flux ratio of 0.3. As shown in Figure 2.2a, from the distance between the primary diffraction lines in AlN [1 ̅00] (or [2 ̅ ̅0]) to Si

[110] in-plane lattice constant ratio is confirmed to be 4:5 as shown. Due to the in-plane lattice coincidence, the AlN and Si(111) interface energetically favors two dimensional growth rather than 3D islanding [43]. However, the ring and spot features during the high temperature growth of AlN (Figure 2.2b) are indicative of 3D growth mode [44], which has been known to occur under metal limited (III/N < 1) conditions [45]. Indeed, as shown in the cross-sectional SEM image (Figure 2.3 inset), nanocolumnar AlN forms under the conditions used here. The temperature of the substrate is then lowered and stabilized at 350oC (measured by a calibrated thermocouple at the substrate position) without any growth happening, and then followed by growth of InN/AlN MQWs (Figure

2.2c) using an In/N flux ratio of 0.15 and Al/N flux ratio of 0.3.

4:5

(a) (b) (c) Figure 2.2: in-situ RHEED patterns during different growth stages.

200nm

Figure 2.3: Cross-sectional tilt and higher magnification

(inset) SEM image of the InN/AlN MQWs sample.

Several InN/AlN multiple quantum well samples are successfully grown on top of coalesced AlN nanocolumns on Si(111). All SEM images in this work are obtained using a FEI Sirion scanning electron microscope operating at 10 kV. The details of each sample are showing in table 1.

Sample Growth InN thickness InN thickness number temperature (from XRD (from XRD simulation)* simulation)* 1 350oC 9.3 Å (44%) 62.7 Å (66%) 2 350oC 10.3 Å (98%) 56 Å (100%) 3 350oC 13.5 Å (100%) 65 Å (100%) 4 450oC 8.5 Å (76%) 67 Å (100%) 5 550oC N/A N/A Table 1: Sample database of InN/AlN MQWs.

How we measure the thickness and relaxation for each sample is explained in the following section. First, the detais of structural characterization (layer thickness and relaxation state) on sample 1 will be demonstrated. Scanning transmission electron microscopy (HAADF mode 80kV) and XRD are used to determine and confirm the thickness information and layer relaxation state. Then, the XRD analyses on sample 2 and 3 are conducted to demonstrate the success of tuning the InN thickness with the change of shutter open time.

2.2.1 STEM

The successful growth of InN/AlN multiple quantum wells are confirmed by

STEM (Figure 2.4). The STEM images reveal that the InN/AlN MQW structure forms in a zigzag pattern on top of the AlN base that was deposited at high temperature. This zigzag pattern is due to AlN columns coalescing into thin films with facets on top, and the InN/AlN MQWs following the orientation of facets. The InN QW interface is sharpest when close to the AlN base but AlN is getting progressively rougher and reduces the sharpness of the interfaces. This is due to the fact that the AlN spacer between the

InN layers is grown at a much lower temperature (350o C) than that of the typical AlN 2-

D growth temperature (800 o C) to prevent In desorption and InN decomposition. Thus the low Al adatom mobility is further reduced.

This can be modeled as statistical roughening of the AlN cladding layers due to the negligible Al adatom mobility at the substrate temperature used (350 oC). Here, the statistical roughening is not considered for InN because In has high enough mobility at

350 oC. Thus √ is the expected statistical roughening where n is the number of the spacer (n=1-20) under the InN layer grown at 350 oC and t is the AlN spacer thickness

(63Å for the sample shown in Figure 2.4). In the following section, we will discuss the other explanation for not so sharp InN/AlN interfaces.

Figure 2.4: HAADF STEM images of MQWs structure of the InN/AlN sample.

2.2.2 Layer thickness and relaxation state

In the literature, researchers do not usually estimate lattice parameters using

STEM images because in the scanning mode images can become stretched. However, the relative value among different layers can still provide information when assuming the stretching rate is uniform. Thus relaxation can be estimated.

The relaxation is estimated by assuming the initial AlN layer is relaxed because it forms nano-columnar/nanowire growth mode. Because this layer is relaxed, its lattice constant is used as a reference. The process is shown in Figure 2.5. By intensive image analysis that includes averaging 30+ measures in the units of 5 atoms from 3 different images (different region), the lattice constants for InN and AlN layer are estimated as

(1) √ where is the lattice consistent and is the single atomic spacing in the image plane.

Basal plane (0002)

Image plane {11 ̅0} Figure 2.5: An illustration of in plane lattice constant

with imaged atomic spacing.

From this, the relaxation is calculated as 44% for the InN layer and 63% for the AlN layers. Next, we input the relaxation state into an XRD simulator and dynamically simulate the structure, the result of which shows that each period contains 9.3 Å of InN and 63 Å of AlN. From the high resolution STEM images, the thickness of the first InN layer and AlN layer has been measured and averaged from more than 30 measures section from multiple images. The layer thickness is 9.2  0.8 Å and 60  6 Å for InN and AlN respectively. The fact that XRD simulation and STEM layer thickness are consistent with each other validates the method we use to estimate the relaxation state for

InN and AlN layers.

AlN XRD scan AlN(0002) InN 3 Simulation 10 O SL 63% (9.3 A InN/ 0 O relaxed 63 A AlN)

2 SL-1 10 SL

44% SL-2 1 Intensityc.p.s relaxed 1 2nm 10 0.30 0.33 -2 -1 0 1 a ' (nm)  (O) 0

Figure 2.6: HRSTEM image with atomic spacing averaged result and XRD simulation

result.

2.2.3 Semipolar facet plane identification

As shown in Figure 2.4, the quantum wells follow the zigzag orientation of the

AlN base top. Furthermore, the HAADF mode provides Z-contrast yielding bright In atoms and darker Al atoms in Figure 2.4, clearly revealing a single InN/AlN QW following the facet orientation of the underlying AlN nanocolumn. It would be very important to identify the orientation of the facet normal to study the polarization properties of the MQWs. In this section, the semipolar facet plane will be identified by

STEM analysis and will be further confirmed by off-axis XRD scan.

In order to identify the facet plane which the MQW grows on, the first step is to know which crystal plane is imaged. . The fast Fourier transform (FFT) of the STEM image (Figure 2.7a) is shown. By matching the modeled lattice (Figure 2.7c) to the imaged lattice (Figure 2.7a), the image plane is determined as {11 ̅0}, which is consistent with the angle and symmetry of the measured (Figure 2.7b) and modeled

(Figure 2.7b) reciprocal lattices. The facet direction is modeled in Figure 2.7c (dotted line) (generated by Crystal Maker) which allows identification of the facet plane as a

{10 ̅5} plane intersecting with the {11 ̅0} image plane.

c-direction

(a) 2nm (c)

61.74o

(b) (d)

Figure 2.7: HRSTEM image and atomic model for image plane and facet orientation

identification.

It is worth noting is that among the 6 of the {10 ̅5} family of equivalent planes, plane (10 ̅5) (01 ̅5) ( ̅015) (0 ̅15) are not perpendicular to the (11 ̅0) plane. As indicated in Figure 2.8 , only (1 ̅05) and ( ̅105) planes are perpendicular to the imaged plane. Thus the bright region in HRSTEM images can be either the cross-section of a perpendicular plane or the projection of an inclined plane. The latter case is one of the explanations for the not-so-sharp interfaces shown in Figure 2.4. Inter-diffusion, intermixing due to strain induced segregation can also account for the smeared out imaged interfaces.

(1 ̅05) ( ̅105)

(10 ̅5) (01 ̅5)

(0 ̅15) ( ̅015)

Figure 2.8: The projection the of {10 ̅5} family of equivalent planes onto the (11 ̅0)

plane.

Additionally, the inclined semipolar plane normal has been further confirmed by utilizing an off-axis ω-2θ XRD scan. The first hint for the (10 ̅5) plane is when the inclined angle from c-plane is measured as around 20o as shown in a relative large scale

STEM image of the MQWs structure. (Figure 2.9) The STEM image shows that the zigzag AlN and MQWs interfaces maintain the same incline angle from the Basal plane.

In Table 2, the planes with small incline angle ( <35o ) are listed, and the (10 ̅5) plane is most possible candidate.

Basal plane

Si-AlN growth interface

Figure 2.9: The inclined MQWs from basal plane with large scale STEM image.

hkl Angle to basal plane (o) Bragg angle (o) (10 ̅5) 20.29 55.54 (10 ̅4) 24.81 42.97 (10 ̅3) 31.65 33.03 (11 ̅5) 32.64 66.70 Table 2: Inclined angles and Bragg’s angle of selective planes in wurtzite AlN.

Then off-axis ω-2θ XRD scans are conducted on all small-angle inclined planes listed in Table 2 , except for the (11 ̅5) plane, due to its inaccessibility from the XRD geometry we use. Among the small angle inclination planes tested, only diffraction about

AlN (10 ̅5) peak exhibits the zeroth order superlattice peak (Figure 2.10a). Also, the off- axis in-plane Φ scan around (10 ̅5) zeroth order superlattice confirms the 6 fold symmetry expected for hexagonal {10 ̅5} facets (Figure 2.10).

100 AlN(10-15)

AlN(10-15) x10

SL (10-15) 10 102 0 Supperlattice 0 peak(10-15)

c.p.s Intensity Intensity c.p.s Intensity

1 50 51 52 53 54 55 56 57 -180 -120 -60 0o 60 120 180 (a)  (O) (b) ( )

Figure 2.10: Off-axis XRD scan about AlN (10-15) peak and in plane phi scan.

The stable facet for a crystal depends on the growth conditions, which alter the relative surface energy for particular sets of atomic planes. Hsu et al. observes

{11 ̅n}(n=1-3) facets for AlN grown by PAMBE on Si(111) at 800oC, with n increasing

29 as the Al/N flux ratio decreases [46]. By comparison, we grew AlN at higher temperature

(1150oC T/C, 1000oC Pyro) and lower Al/N flux ratio (0.3). Under these conditions, facets with a smaller inclination and lower Al/N dangling bond ratio are expected because of the reduction in surface energy [46].

2.3 Polarization for semipolar planes

Upon finding the semipolar planes MQWs grow on, it is very interesting to investigate the polarization state of the {10 ̅5} planes, which hasn’t been done in the

InN/AlN system. Thus in this section, the inclination angle from Basal plane dependence of polarization (P) will be calculated first assuming the coherently strained InN on AlN.

Then the specific case for the sample studied (44% relaxed InN and 63% AlN) is calculated for {10 ̅5} planes. The method presented in Refs [47, 48] has been adopted for the InN/AlN system.

A schematic representation of the wurtzite structure with an incline plane (blue shaded) and coordinate system relationship are shown in Figure 2.11. The xyz coordinate system (x║[11 ̅0], y║[1 ̅ 0], z║[0001]) represents the “natural” wurtzite coordinates.

The x’y’z’ coordinate system represents the “inclined” coordinates. It is rotated about the x (x’) axis with respect to the xyz system by an angle , which is the inclination angle of a semipolar plane from the Basal plane. The in-plane misfit parameters are defined along x’ and y’ directions. Different colors are used in Figure 2.11 for eye guiding.

[10 ̅5]

z z’

y y’ x (x’)

Figure 2.11: Schematic representation of “natural” and “inclined” coordinates in wurtize

structure.

First, we will deal with piezoelectric polarization. The strain components in the two coordinate systems are related by tensor transformation and results are as follows:

, (2)

After considering the symmetry of the wurtzite III-nitride structure, the piezoelectric polarization is calculated as:

( ) ( ). (3)

( )

As shown in Figure 2.11, the piezoelectric polarization component involves the projection of both in-plane piezoelectric polarization and out-of-plane piezoelectric

polarization onto the facet normal direction. Thus the total (along semipolar orientation) can be expressed as:

( ) . (4)

By inserting equations (2) and (3) into (4), the result for piezoelectric polarization in terms of strain is:

( ( ((

[( ] .

(5)

Since the spontaneous polarization component is straightforwardly calculated as the projection along the facet normal, the overall polarization at the InN/AlN interface is:

( ) (6)

The result calculated here corrects the mistakes in [47, 48].

In Table 3, the parameters used in the calculation for are listed. They are adopted from [47, 48] , however there is still some discrepancy between different reference sources.

AIN InN

( 396 223 ( 137 115 ( 108 92 ( 373 224 ( 116 48

a (Å) 3.112 3.54 c (Å) 4.982 5.705

2 (C/m ) 1.56 1.06 2 (C/m ) -0.57 -0.49 2 (C/m ) -0.48 -0.40

2 Psp (C/m ) -0.081 -0.032 Table 3: Elastic constants, lattice parameters, piezoelectric coefficients and spontaneous

polarization of AlN and InN.

To calculate the elastic strains in mismatched InN layers grown on inclined semipolar AlN, the well-established lattice parameter along with other parameters listed in Table 3 are used. Figure 2.12a presents the dependence of the strain components on the inclined angle from the basal plane in “inclined” coordinates. As shown in the figure, εz’z’ and εx’x’ are the opposite sign, which is expected because of Poisson’s effect.

The overall polarization at InN/AlN is calculated for both In-polar and N-polar cases. As shown in Figure 2.12b, the total polarization at the interface for a {10 ̅5} facet is 2/3 and 1/2 that of the c-plane, respectively. When assuming coherently strained layers, the structure will have non-polar effect happen at inclined plane in addition to non-polar plane (m/a-planes).



0.1 z'z'

)

i'j'

 (

0.0 

y'z' Strain  -0.1 x'x'  y'y' 0 20 40 60 80 Inclination angle (o) 0.3

piezo {10-15}

0.2 P

) 2 0.1

o C/m

( 51.1 total P 0.0 P

-0.1 0 20 40 60 80 Inclination angle (o)

Figure 2.12: The inclination angle dependence of strain components and polarization at

the InN/AlN interface.

In the case of partially relaxed InN and AlN layers, the pieozoelectric polarization at the interface contains contributions from both InN and AlN. For 44% relaxed InN on

63% relaxed AlN (3 ML sample), the piezoelectric polarizations are 0.077 C/m2 and

0.029 C/m2, respectively. The overall piezoelectric polarization at the InN/AlN interface is 0.106 C/m2, which is smaller than the fully strained InN layer case calculated in Figure

2.12. These result are similar to previous work on InGaN/GaN or AlGaN/GaN interfaces by Romanov et al. [48], showing that by engineering the orientation of a semipolar plane, any polarization state can be achieved. Not surprisingly, based on our calculations

(Figure 2.12b) we see the same versatility of semipolar planes for polarization engineering in the InN/AlN system.

2.4 Photoluminescence of MQWs with different InN thickness

Two additional MQW samples were prepared with larger InN layer thicknesses by increasing the In shutter open time. XRD measurements were used to determine the average InN and AlN layer thicknesses along the c-axis (Figure 2.13). Although the SL facets are inclined from the c-axis, the periodicity along the c-axis still leads to first and second order SL peaks due to interface scattering summation from all six {10 ̅5} facet planes. For all three samples, the AlN cladding thickness is between 23 and 26 ML and the average InN thickness varies from 3 to 5 ML projected along the c-direction. The dominant room temperature PL peak is 380 nm for the thinnest InN/AlN MQW sample (3

ML, 44% relaxed InN). The 4 ML MQW sample (fully strain relaxed InN) exhibits a PL

36 peak at 420 nm with a smaller shoulder close to 550 nm. Finally, the 5 ML MQW (fully strain relaxed InN) exhibits two distinct PL peaks: a dominant peak at 550 nm and a smaller peak at 420 nm, labeled (5-1) ML. This smaller peak could arise from 4 ML-thick regions occurring within the 5 ML sample due to interface thickness fluctuations.

Therefore, the PL spectra can be interpreted as arising from InN/AlN QWs with discrete shifts in wavelength with each integer ML change in average thickness. In the case of the

5 ML sample, distinct 5 ML and 4 ML (5-1 ML) PL peaks are clearly observed, whereas in the other two samples, secondary peaks appear only as smaller shoulders.

10 10 Simulation AlN(0002) 108 3 ML 6 SL SL SL 10 -2 -1 0 SL1 4 ML 104

2 Intensityc.p.s 10 5 ML 100 16 17 18  (o) 3 ML 4 ML 5 ML 1.0

5 - 1ML Normalized PL intensity PL Normalized 350 400 450 500 550 600 650 Wavelength (nm)

Figure 2.13: XRD and PL spectra on 3,4,5 ML samples.

Time-resolved PL measurements are carried out for both 3ML and 5 ML samples at each of the two distinct PL peaks. A short PL lifetime of 39 ps is observed for the 3

ML sample. Also a short PL lifetime of 30 ps is observed for the 4 ML (5-1) peak at 420 nm, whereas a relatively longer lifetime of 70 ps is observed for the 5 ML peak at 550 nm. The instrument response function (IRF) demonstrates that the measured PL decay curves are not below the time-resolution of the measurement. From this data set alone it is possible to conclude that the PL originates from exciton (band to band) recombination in InN rather than from traps. The sub-ns (and >ps) lifetimes are a strong indication of interband transitions [49, 50]. To our knowledge there are no previous reports of

InN/AlN MQW nor measurements of PL lifetime in InN/GaN MQW, thus it is unclear if the <100 ps lifetime is related to the high degree of quantum confinement in InN/AlN

MQW or due to a large increase in non-radiative recombination rate due to strain related defects at the interfaces. However, thicker InN QWs (16 ML) with lower energy barrier height In0.9Ga0.1N cladding layers have been measured by time-resolved pump-probe absorption yielding an interband relaxation time of 125 ps[50]. For these samples, the emission wavelength is 1650 nm as the quantum confinement is much lower than for the

InN/AlN samples studied here.

380nm

0.1

Normalizedintensity IRF 0 0.3 0.4 Time (ns) 1

550 nm

0.1 420 nm

IRF Normalized intensity Normalized

0 0.3 0.4 Time (ns)

Figure 2.14: Time-resolved PL measurements for 3 ML and 5 ML samples.

2.5 Growth window exploration

Finally, preliminary efforts of exploring the growth window for InN/AlN MQW samples have been done. Two more samples (#4 and #5) are grown at elevated temperatures and XRD data are shown in Figure 2.15. When the substrate temperature increases from 350 oC to 450 oC, the structural quality is improved indicated by the enhancement of the XRD SL peak sharpness. At the meantime, the InN quantum well thickness decreases, which is expected due to the above InN decomposition growth temperature. (Table 1) When the substrate temperature further increases to 550oC, a sharp peak correlated to 15% of InAlN alloy appears. There is literature showing the 550oC is the optimized growth temperature for InAlN alloy formation [51]. What happened in this sample is that InN decomposes into liquid In, parts of it incorporate into AlN when the Al shutter opens and forms InAlN due to a preferred temperature. Also, from cross-sectional

SEM images shown in Figure 2.16, the nanowire feature is getting more prominent as the temperature increases.

450oC InAlN 103 350oC AlN(0002) SL0 550oC 102 SL-1 SL SL 1 -2 1

Intensity c.p.s Intensity 10

100 16 17 18  (o)

Figure 2.15: XRD scans of samples grown at different temperatures.

350oC

450oC

550oC

Figure 2.16: Cross-sectional SEM ofthe InN/AlN samples grown at

different temperatures.

In summary, highly quantum confined InN/AlN MQWs are demonstrated by

PAMBE. Characterization by atomic resolution STEM and HRXRD confirm the formation of InN/AlN MQWs as well as identify {10 ̅5} semipolar facet planes. The polarization charge at these semipolar InN/AlN interfaces is calculated and found to be substantially reduced from that of an identical c-plane oriented interface. The room temperature PL wavelength shows discrete shifts due to discrete changes in quantum confinement as the InN QW thickness changes by integer ML values. This PL exhibits short lifetimes of less than 100 ps, indicative of exciton recombination in highly confined

QWs.

One direction for future work is to achieve 1 and 2 ML InN embedded in an AlN matrix, to investigate whether smoother interfaces are possible. Also, smaller temperature increment steps are useful to find the best temperature condition for InN/AlN MQWs.

Finally, exploration of the growth condition effects on formation semipolar plane orientation is very interesting. The InN/AlN MWQs will show more versatility when there is the freedom to choose which semipolar plane to grow, and thus modify the polarization at the interface.

Chapter 3 : GaN/AlN Multiple Quantum Wells

(This chapter is adapted from the paper submitted to JVSTB.)

3.1 Introduction to high temperature growth condition for GaN/AlN MQWs

Unlike the InN/AlN system, GaN/AlN MQWs have been studied intensively.

GaN/AlN MQWs are also of special interest due to the large conduction band offset

(CBO) of nearly 2eV, making GaN/AlN SLs very promising for intersubband (ISB) optoelectronic devices operating at telecommunication wavelengths of 1.55µm and

1.3µm. In most ISB transitions reported, both broad spectra (homogeneous broadening) and multiple shoulders (inhomogeneous broadening) exist since it is difficult to control the interface roughness and compositional profile within single monolayer precision in

GaN/AlN SLs as shown in Figure 3.1.[1, 52] Thus an important goal is achievement of the smoothest possible interface for engineering intersubband photonics. This chapter focuses on the development of a different GaN growth condition for the MQWs structure and studying the kinetics and the interfaces and comparing it with more widely used Ga- rich low temperature growth condition.

Figure 3 .1: Absorbance of doped (top left),undoped samples GaN/AlN MWQs (top right) ; homogeneous broadening of absorption peak (bottom left) and inhomogeneous broadening of absorption peak (bottom right)[1].

First, an overview of the established GaN/AlN MQWs growth condition is presented. To realize the optimized GaN/AlN structural quality, it has been agreed in a variety of publications that not only high quality GaN and AlN layers are required, but

46 also sharp GaN/AlN interfaces are necessary [53]. Previously, excess Ga growth front is used to stabilize the AlN surface as surfactant. In this case, not only GaN grown in the step-flow mode has very high crystalline quality, but also the 2D AlN growth can be achieved at stoichiometric flux of Al [54-56]. This condition will be referred to as the

“low T growth” (720o) moving forward.

However, there are some limitations of this low temperature growth condition.

First, because of the lower thermal-decomposition temperature of GaN compared to AlN,

GaN growth is the limiting factor to determine the growth temperature that is compatible for both GaN and AlN layers. Thus, most GaN/AlN SL samples are grown at temperatures between 700 oC and 750 oC where GaN decomposition can be neglected.

Second, there are reports showing intermixing effects at GaN/AlN interface when grown the structure above 730 oC, due to the segregation of Ga atoms at the interface. This makes the growth temperature window for GaN/AlN MQWs even narrower.

Furthermore, the AlN 2D growth window is quite narrow at this temperature as shown in

Figure 1.4 [24]. Thus, to keep the Al impinging flux just above the stoichiometric point is very critical at this temperature. In most studies, researchers are using an Al supply of

1.05 times the stoichiometric point [24, 56]. Another challenge for AlN growth is the low

Al adatom mobility at this low temperature.

Recently, a smooth 2D layer-by-layer growth mode for GaN at temperatures above the onset of GaN decomposition with an N-rich III/V ratio has been explored

(referred to as “high T growth” moving forward) [19, 57]. High temperature growth (820

47 oC) is an alternative way to improve the quality of the overall SL structure due to increased Al mobility leading to better crystalline quality of AlN layers compared to low

T growth [24]. Another potential advantage for growing GaN at N-rich condition is that there is no excess Ga at the GaN/AlN growth interfaces, thus sharper interfaces are expected.

3.2 Reciprocal space mapping technique

X-ray diffraction is one of the most frequently used techniques to study fine structure of materials non-destructively. It has been used in the previous section to confirm the structure for InN/AlN MQWs. In this chapter, XRD will be even more intensively used to study the interface as well as conducting the reciprocal space mapping. Thus it is necessary to have a detailed evaluation of XRD and how a Reciprocal

Space Map (RSM) is achieved theoretically.

It is well known that for any of the diffraction techniques, (XRD, RHEED, TEM), in order to have constructive interference, one must satisfy the condition called Bragg’s law:

, (3.1) where d is the distance between the lattice planes, is the angle measured between the incident beam and the plane of interest. For X-ray diffraction, the Cu kα line is used as the source beam with of 0.154 nm for high resolution diffraction. A geometric illustration is shown in Figure 3.2.

Figure 3.2: A geometric illustration of Bragg’s law [58].

There are two major categories of x-ray diffraction geometries: one is powder diffraction which does not employ a monochromator for the incident beam; the other one is high resolution diffraction. The two examples of geometry are illustrated in Figure 3.3.

Powder diffraction proceeds such that the incident angle is maintained while the detector rotates about the sample as shown in Figure 3.3a. In this way, all powder particle orientations can be detected. For reciprocal space mapping (RSM), a high resolution diffraction geometry is required. As shown in Figure 3.3b, the x-ray source is monochromated by diffraction on four Gemanium crystal mirrors. Additionally, an analyzer crystal can be added in front of the detector (triple-axis set up) to increase the resolution at the expense of very low count rate. In the Bruker D8 HRXRD system used in this study, 2 can reach as high as 115o due to a camera placed to observe the precise position of small size samples. However, it limits the option for accessible RSM. This is

49 the reason why RSM around the (11 ̅4) diffraction peak was performed instead of around the mostly reported (10 ̅5) diffraction peak.

Figure 3.3: Powder diffraction geometry (a) and high resolution diffraction geometry (b)

of a XRD system [58].

Then, let’s do the Ewald sphere construction to see the target diffraction in reciprocal space. The theoretical maximum 2θ angle in a diffractometer is 180o, which defines the boundary of a semicircle as a total accessible area in reciprocal space with a radius of 2/λ, shown as diffractometer circle in Figure 3.4. The two Laue-range semicircles cover the spots not accessible, where the sample will be blocked by either incident or outgoing beam. The Ewald circle, with the radius of 1/λ, represents all the possible points in reciprocal space that could satisfy the Bragg equation. The Ewald circle needs to pass through (0000) and the point of target diffraction. And the incident (ki) and

50 the diffraction (kdifr) beam vector, with the length of 1/λ, indicate the angles with respect to the basal plane.

[0001]

(0006) (11 ̅5) (0005) (22 ̅4) (11 ̅4) (0004) kdifr Diffractometer circle Laue-range Laue-range (2θ < ω) ( ω < 0) k [ ̅ ̅20] Ewald i sphere

Figure 3.4: Schematic drawing of the reciprocal lattice and diffraction condition (Ewald

sphere construction) for the hexagonal nitride system.

There are three possible scan directions in reciprocal space. The first one, usually known as ω-2θ, is illustrated in Figure 3.5. Under this scan mode, the simultaneous rotation of the sample (θo/s) and the detector (2θo/s) leads the scan direction along the scattered vector (kdifr – ki). The second direction, as indicated in Figure 3.5, is perpendicular to the scattered vector. This is the ω scan also known as rocking curves, because the sample rotates (“rockes”) with the detector angle fixed. The third direction is

51 along 2θ. This is when the detector rotates around the sample while the sample keeps still. In reciprocal space, the movement is a long the Ewald sphere as in Figure 3.5. With the combination of two of these three directions, the scans can cover an area in reciprocal space, thus RSM can be produced. In this work, ω-2θ scan and ω scan are used to generate the RSMs.

Q y Qy ω- 2θ scan ω=αi 2θ=αi +αf Qx 2θ scan ω scan

αi +αf αi

Figure 3.5: Different scan directions in reciprocal space [59] for the RSM.

3.3 Growth of GaN/AlN MQWs

In order to investigate the impact of changing growth kinetics on interface roughening and structure quality of SLs, several short period highly confined GaN/AlN

52 multiple quantum wells are prepared under both low T and high T growth conditions. The

GaN/AlN SLs are grown using a VEECO 930 PAMBE system with a base pressure of less than 9×10-11 Torr. All samples are grown on c-plane 50-nm-thick AlN on sapphire templates commercially available through Kyma Technologies. Prior to the growth, 500 nm of Ti was deposited on the back of the sapphire by an E-beam evaporator for heat absorption during growth. Standard effusion cells are used for the Ga and Al sources with 7N pure source material. Active nitrogen is supplied by a Veeco radio frequency plasma source operating at 350W with N2 source purity of 6N. The flow rate of N2 is controlled by a Mass Flow Controller adjusted to maintain an effective chamber pressure of 2×10-5 Torr during growth. The substrate temperature is monitored by an infrared pyrometer calibrated using the melting point of Al (660oC). Before each growth, the AlN template is heated to 1100 oC to remove the oxides on the surface. The growth rate is calibrated by the growth of multiple AlN/GaN SL samples with different AlN and GaN growth times, followed by XRD analysis in order to find the thickness of each layer. The growth rate is further confirmed and stoichiometric flux (metal/N = 1) is calibrated by preparing multiple samples with different fluxes of Al (Ga) with identical N-plasma conditions. Cross-sectional SEM images are used to determine the effective 2D growth rate versus Al (Ga) beam effective pressure and find the intersection of the N-limited and metal-limited growth region. Upon finding the stoichiometric point, the III/V ratio is known for any Ga or Al flux used, assuming active N is consistent with the calibration

53 growths. RHEED is used as an in situ surface sensitive technique to monitor particular surface reconstruction and the growth mode.

The target structure of the sample consists of 9 period SLs of 1.5 nm GaN quantum wells embedded in 5 nm AlN thick barriers on top of a 50 nm AlN buffer layer.

The low T growth samples are grown at 720oC and measured by a pyrometer with a Ga/N ratio of 2.5 (Ga-rich) and an Al/N ratio of 1.05. The high temperature samples are grown at 820oC with a Ga/N ratio of 0.7 (N-rich) and Al/N ratio of 1.05. Additional SL samples with 20 periods are prepared under both low T and high T growth conditions to study the relaxation mechanism.

1.10 nm

0.00 nm Figure 3.6: AFM image of a 1.5nm/20nm GaN/AlN SL

structure.

3.4 Structural characterization by AFM, RHEED and XRD

To confirm that the N-rich (0.7 Ga/N ratio) high T (820 oC) condition is applicable to the SL structure, a 9 period SL sample is grown with thick AlN (20 nm) spacers and thin GaN (1.5 nm) inserts. As shown in Figure 3.6, the surface is very smooth

(0.3nm RMS) with hillocks and step features for the AlN layer. When comparing the

RHEED patterns of the first GaN layer and the first AlN spacer as in Figure 3.7a and

Figure 3.7b, they are both streaky and have the same line spacing ([1 ̅00]), indicating the formation of a strained 2D GaN layer. Similar RHEED patterns of the last GaN and AlN layer (9th) are shown in Figure 3.7c and Figure 3.7d. Closer to the end of growth the GaN

RHEED pattern turns spotty, indicating a rougher layer than the first GaN layer deposited. However, the spacing is still the same as in the first layers indicative of a strained structure.

(a) (b)

Figure 3.7: RHEED patterns of the test GaN/AlN MQWs structure.

Following the analysis of this initial SL sample, additional samples with the same structure are grown under both high T and low T conditions. Both of the samples have smooth surfaces, with RMS roughnesses of 0.4 nm for the high T and 0.2 nm for the low

T sample (5 µm by 5 µm range). Sharp satellite peaks are observed in the XRD patterns for both the N-rich, high T sample and Ga-rich, low T sample, as shown in Figure 3.8. N- rich, high T SLs render stronger SL fringes which is consistent with more well defined, sharper, interfaces. This result is not a complete surprise. Since the GaN layer is grown under N-rich condition, no Ga segregation should occur at the GaN/AlN interfaces, thus resulting in sharper interface. In addition, due to the higher temperature, the enhancement of Al atom mobility should lead to the higher crystal quality AlN, confirmed by the XRD

ω scans. A 9 period SL structure was chosen because the overall combined thickness of the SL is below the Matthews-Blackeslee predicted critical thickness, approximated as: h

2 = as / Δa = as /ε, where as is the lattice parameter of the substrate (AlN in this case) and ε is the strain of the epitaxial film. Here the epitaxial SL layers are approximated as one uniform AlGaN layer with composition equal to the weighted average composition of the total GaN/AlN structure.

9 10 9 x N-rich high temperature 9 x Ga-rich low temperature AlN 7 SL 10 0 (0002) SL SL -1 5 -2 SL1

10 SL-3

3 x 104 SL0 10 SL-1 SL-2 SL1 SL-3

Intensity(c.p.s.) 1 SL 10 -4

16 17 o 18

10  ( ) 10 20x N-rich high temperature 8 20x Ga-rich high temperature 10 AlN(0002) SL0 SL-1 6 SL-2 10 SL1 SL-4 SL-3 4 10 SL0 SL Intensityc.p.s 2 -1 SL-2 SL 10 SL-3 1 100 16 18  (o)

Figure 3.8: HRXRD scan about AlN (0002) diffraction peak for both high T and low T

samples, with 9 periods (top) and 20 periods (bottom).

Next, SL structures with 20 periods instead of 9 are formed using the same low T and high T conditions. The XRD results (Figure 3.8.) suggest that the Ga-rich, low T SL exhibits sharper interfaces than the N-rich, high T SL based on the relative SL satellite peak intensity.

3.5 Relaxation mechanism

The sizable lattice mismatch of c-plane GaN and AlN makes the study of relaxation properties extremely important. Not only will the piezoelectric field affect the overall internal electric field which we will discuss in the following chapter, but also the electrical and optical properties of the devices are largely dependent on defects. Both defects and piezoelectric properties are partially shaped by the strain relaxation mechanism [54]. Due to the lack of a main slip system that is not parallel to the basal plane in hexagonal symmetry, plastic relaxation mechanisms are still under investigation.

In this study, XRD reciprocal space mapping is the core technique to study the in- plane relaxation state of the film. The reciprocal space maps around the (11 ̅4) reflection on both low T and high T SL samples are conducted under double axis mode to allow for higher intensity counts from XRD scan and triple mode.

Sample 20 x N-rich 20 x Ga-rich 9 x N-rich high 9 x Ga-rich high low temperature low temperature temperature temperature Average lattice 3.136  0.002 3.133  0.001 3.115  0.001 3.124  0.001 parameter a (Å) Average lattice 5.018  0.001 5.013  0.003 5.004  0.001 5.022  0.001 parameter c (Å) GaN layer 10.6  0.2 19.8  0.2 20.0  0.2 18.1  0.2 thickness (Å) (100%) (22.1%) (100%) (4.4%) (relaxation) AlN layer 55.0  0.5 48.3  0.5 43.8  0.5 49.5  0.5 thickness (Å) (99.9%) (14.0%) (66.1%) (0.5%) (relaxation) Table 4: Sample database of RSM and XRD analyses.

Reciprocal space mapping (RSM) around the (11 ̅4) reflection is performed on both the low T and high T SL, 20 period samples. The detailed information for each sample from RSM is combined with XRD analysis and shown in Table 4 . As shown in

Figure 3.9 and Figure 3.10, the RSM of the high T, 20 period sample exhibits a larger lateral correlation for SL peaks. This is especially true for the higher order SL peak, which is a strong indication of higher dislocation density in the high T condition sample.

Unlike the low T sample, the SL peaks for the high T sample are not aligned in-plane, representing that a gradual lattice relaxation occurs in each GaN/AlN interfaces. On the other hand, the 1st, 0th and -1st SL peaks for the low T, 20 period sample have the same in-plane lattice parameter, indicating the in-plane lattice relaxation steady state condition has been reached. Kandaswamy et al. have reported that the strain state of SLs reaches

59 the steady state under Ga-rich conditions after some periods of growth and that the final relaxation state is independent of the substrate [60]. Because excess Ga at the interface will decrease the (0001) surface free energy, Ga-rich conditions minimize strain relaxation [61]. This explains why Ga-rich low temperature conditions can produce smooth interfaces even beyond the critical thickness.

5.2 20x N-rich high temperature

)

5.1 SL1 -1

A (

z z AlN Q 5.0 SL0

4.9

SL-1

3.9 4.0 o 4.1 4.2 Qx(A-1)

Figure 3.9：Reciprocal Space map around (11 ̅4) diffraction for 20 period sample, N-

rich, high temperature

Figure 3.10: Reciprocal Space map around (11 ̅4) diffraction for 20 period sample, Ga-

rich, low temperature.

5.2 9x N-rich high temperature

5.1 ) SL

-1 1

Å (

z AlN

Q 5.0

SL0

4.9

SL-1 3.9 4.0 o 4.1 4.2 Qx(A-1) Figure 3. 11: Reciprocal Space map around (11 ̅4) diffraction for 9 period sample, N-rich, high

temperature.

Figure 3.12: Reciprocal Space map around (11 ̅4) diffraction for 9 period sample, Ga-rich, low

temperature.

The RSM analysis for the 9 period sample at the high T condition (Figure 3.12) shows that the gradual in-plane lattice expansion has already begun within 9 periods, but there is no excessive lateral correlation compared to the 20 period sample. From the RSM for a low T 9 period sample (Figure 3.13), it is clear that the structure has already reached steady-state within 9 periods. Previous work has shown that strain accommodation through the formation of misfit dislocations at a heterostructure interface is more favorable at higher temperatures, resulting in a smaller critical thickness than that at lower temperatures [62, 63]. The RSM shows that for 9 periods, when the theoretical critical thickness has not been reached, the sample grown at 820 oC is fully relaxed, while the sample grown at 720 oC is strained. Combining the RSM results from both low T and high T conditions for 9 and 20 period samples, we conclude that the N-rich high T condition gives rise to a sharp GaN/AlN interface. The advantage of a dry interface and the high crystalline quality AlN are negated upon growth beyond the critical thickness because strain relaxation occurs more rapidly at higher substrate temperature. Counter intuitively, the Ga-rich, low T condition, in spite of having excess Ga on the surface during growth, produces smooth interfaces in SL structures grown beyond the critical thickness.

One thing worth mentioning is that the samples studied above were all grown at

820 oC for the AlN buffer. The substrate temperature was brought down to 720 oC for low temperature samples later on before the growth of MWQs. However, one sample grown at 720 oC for the full structure has also been studied. The RSM is shown in Figure 3.14.

Compared to the RSM on the same structure sample with buffer grown at 820 oC (Figure

2.9), the reciprocal peak of AlN has bigger in-plane broadening correlated to AlN buffer degradation. This is evidence that higher temperature renders better crystal quality of

AlN.

5.2

20 X Ga-rich low temperature )

-1 5.1

(

Qz 5.0

4.9

3.9 4.0 o 4.1 4.2 -1 Qx (A )

Figure 3.13: Reciprocal Space map around (11 ̅4) diffraction for sample 20 periods, Ga-

rich, low temperature (low T buffer).

3.6 An evaluation of high T growth conditions for both GaN and AlN

Before the high T growth condition was applied to SL growth, a series of growths were conducted to reproduce the growth condition in literatures reported by Koblmuller et al [18]. The growth list spans different high temperatures and III/V ratios. However,

64 the very smooth layer by layer growth mode has never been reproduced. Here, a Root

Mean Square (RMS) roughness map of all the GaN samples grown on AlN template with different growth condition is shown in Figure 3.14. Several SEM images of GaN grown at 820 oC with different III/V ratio are shown in Figure 3.15. III/V ratio up to eight has been tested out under this temperature due to the extensive absorption. However, the surface of the thin firm GaN was still very rough with RMS of 9 nm. Error bars in the left column of Figure 3.15 is 200 nm while is 500 nm in the right column images.

Figure 3.14: RMS map of GaN growth condition probing above decomposition

temperature under different conditions.

Ga/N (III/V): 0.9

1.0

1.1

1.3

2.0

Figure 3.15: Cross-sectional and tilted SEM images of selective Ga/N ratios at 820 oC.

III/V: 1.05x III/V: 1.15x III/V: 1.3x RMS:2.35nm RMS: 0.477nm RMS: 0.564nm

Figure 3.16: AFM analyses of AlN homoepitaxial growth at 820 oC with different III/V

ratios.

The growth of AlN at high T shows a very smooth surface and step flow features.

When III/V is 1.15 at 820 oC, AlN surface is the smoothest. However, with excess Al at the growth interface, the SL growth needs periodic pause with nitrogen plasma shutter open to nitride the surface and consume the Al. Otherwise, the AlGaN quantum wells would form. The periodic pause is a degradation factor for optical properties of the SLs.

In summary, we have demonstrated strained GaN/AlN SLs on AlN grown under

N-rich, high temperature and Ga-rich, low temperature conditions by PAMBE. By comparing the structural properties of these SLs, we explore the different strain relaxation mechanisms by conducting XRD and RSM measurements. We find that the N- rich, high T conditions result in smoother interfaces in strained heterostructures grown below the critical thickness, but in relaxed structures the higher temperature leads to an increase in strain relaxation, which has a deleterious effect on interface sharpness. The

67 complete story of this part of growth study would be developing an AlGaN buffer with the same composition of the MQWs layer. Without strain relaxation, the N-rich high T condition might be able to apply for sharp interfaces. Another route is to simply grow more periods. This is the way to study whether and when N-rich high T condition can also render the steady state.

Chapter 4 : GdN/GaN superlattice structure

4.1 Introduction to rare earth incorporation into III-As

A couple of decades ago, people become interested in incorporating ErAs into III-

V systems [64, 65]. Semimetal ErAs grown epitaxially on GaAs was first demonstrated by Palmstrom in 1988 using MBE [66]. Interest in incorporating rare-earth ErAs into a

GaAs matrix started because of the structural compatibility between the two. ErAs has rock-salt crystal structure with a small lattice mismatch (1.6%) to the zincblende structure of GaAs. Thus ErAs can be lattice matched to InGaAs with low In composition [67]. The

ErAs/GaAs interface is shown in Figure 4.1. As modeled in the figure, the rock-salt

ErArs sit perfectly on top of InGaAs zincblende structure, with a continuous arsenic sublattice. From the STEM image in Figure 4.1, it is confirmed that ErAs nanoparticles are compatible to the InGaAs matrix [68].

Next, an intensive study of optical and electrical properties of ErAs nanoparticles in GaAs was conducted and researchers found great potential for high speed photodetector and THz source applications [69, 70]. Kadow et al demonstrated the ultrafast absorption of ErAs/GaAs superlattices (Figure 4.2)

Group III

Figure 4.1: STEM image [68] of ErAs incorporated into an InGaAs matrix and modeled

structure [79].

Figure 4.2: Time-resolved differential reﬂection traces on ErAs containing samples with

different superlattice periods L [70].

Last but not least, ErAs nanoparticles are beneficial for band engineering for tunnel junctions, due to the gap states created by ErAs particles. When ErAs is inserted into the tunneling region, it will bring down the barrier height and barrier distance as illustrated in Figure 4.3 [71].

Figure 4.3: Band diagrams of conventional (a) and ErAs inserted (a) tunneling junction

[71].

In parallel, the epitaxial integration of rare earth pnictides with the nitride matrix is also very interesting. GdN is especially interesting, because it has been confirmed as a ferromagnetic materials with a Currie temperature of 70K [72]. By incorporating GdN into GaN, one can manipulate the magnetic properties of the system in addition to some of the unique properties brought in by GaN. Also, because of the strong spin correlation, the Gd integrated system is a good candidate for a spin injection source. Moreover, several groups have shown that Gd doped GaN is ferromagnetic at room temperature [72-

75]. Finally, GdN could potentially function as a carrier recombination center and assist in the formation of tunneling. It has been demonstrated by Krishnamoorthy et al [76] based on this growth study.

However, unlike ErAs on GaAs, GdN has a cubic based rock-salt unit cell while

GaN has a hcp based wurtzite unit cell. The incompatibility of four-fold symmetry of rocksalt (001) on top of six-fold symmetry of wurtzite (0002) makes the stacking of GdN on GaN more complex. However, the rocksalt (111) GdN is three-fold symmetry and thus is compatible with wurtzite (0002) as shown in Figure 4.4. It has been demonstrated by Scarpulla et al, when 3 fold symmetry GdN[111] is put on top of 6 fold symmetry

GaN[0001], there are two rotational variants confirmed by XRD off-axis φ scan as shown in Figure 4.5. Scarpulla grew less than 100 nm thick of GdN thin film, which remains epitaxial respect to the GaN base, however the GaN capping layer on top is polycrystalline [77].

GdN (Rock Salt) [111]

Gd GaN (Wurtzite) N [0001] Ga

Figure 4.4: An illustration of epitaxy compatibility of GdN rock-salt on GaN wurtzite

[79].

Figure 4.5: In-plane relationship between GdN and GaN [77].

4.2 Gd doping cell calibration

When the Gd doping cell was first installed into the MBE system, a doping calibration was required. Thus a Gd doped stack structure was designed. The structure contained 30 mins growth of Gd:GaN layer with the cell stabilized at 900 oC, followed by

15 mins growth of Gd:GaN at cell temperatures of 950 oC,1000 oC, 1050 oC, 1100 oC with a separation time of 10 mins to let the Gd cell reach equilibrium. Because very little material inside the Gd cell, it stabilized at the target temperature very fast. The GaN matrix was grown at metal rich condition with a growth rate of 278 nm/hr.

The doping level of the sample was probed by SIMS analysis. Conducted by the

EAG group, secondary ion mass spectrometry (SIMS) analysis is to gain the doping profile information by bombarding the sample surface with a primary ion beam and detected by mass spectrometry illustrated in Figure 4.6.

The second ion mass beam contains resputtered primary ions, electrons and photons. The secondary particles can have kinetic energies varying from zero to several hundred eV. The mass spectrometry can identify the target species determined by these atomic mass. Usually another known species in the structure is collected as well to use its location information as a marker.

Figure 4.6: SIMS process illustration of surface material sputtering (from EAG’s website).

During SIMS analysis, the sputter rate off a sample surface can be controlled. The rate is determined by the structure and orientation of the sample, primary beam intensity and the target resolution. A depth profile is rendered by sputtering the sample continuously and produces the mass spectra as a function of depth, as shown in Figure

4.7 top. To depth information is gathered by measuring the depth using a profilometer and then converting the time axis into depth due to a known sputtered rate.

In Figure 4.7 bottom, the Gd doping level vs. cell temperature is plotted. The unit of y axis was related to the real flux (cm-2/s) by the dividing the calibrated growth rate

(nm/s). This is the Gd doping calibration result that for future reference, where any target doping concentration and thus the cell temperature can be calculated based on it. For

GdN thin layer growth, it is necessary to convert the Gd doping into a 2D growth rate by

77 comparing the GaN and GdN area density. Due to that Gd cell is a doping cell, a very low growth rate of GdN is expected.

1E+21

) 1E21 3 1E+20

1E+19

1E+18 1E18 1E+17

1E+16

1E+15 Gd (average) 1E141E+14

00 500500 10001000 15001500 Concentration (Atoms/ cm (Atoms/ Concentration Depth (nm)

1E20 )

-3 1E19

(

Gd Gd

1E18

0.0009 0.0010 0.0011 -1 -1 T (C ) Figure 4.7: (Top) SIMS depth profile of the Gd:GaN calibration stack and (bottom) the

doping concentration vs. cell temperature.

4.3 Growth of GdN/GaN SL

It has been show that a 100 nm layer of GaN on GdN has very poor crystalline quality [77]. Maintaining a high crystalline of GaN is very important for the SL structure.

Determining how to do this is the purpose of this study.

First a calibration stack sample was prepared and grown (shown in Figure 4.8).

The purpose of gradually increasing the GdN thickness by 0.2 ML and separating adjacent layers with a 10nm GaN spacer is to determin if a GaN layer can grow epitaxially on GdN when it’s ultra-thin; if there is a critical layer thickness for a GdN layer that GaN can remain eptixially on top and whether the GaN capping layer will remain in the wurtzite single crystalline structure.

GdN layers were deposited at extremely N-rich condition (Gd/N ratio of 1/45) with a deposition rate of 0.4ML/min. The nominal thickness of the GdN increases from

0.2 ML to 2.4 ML in increments of 0.2 ML. GaN spacers were grown at Ga-rich condition (Ga/N ratio of 2) with a deposition rate of 268 nm/hr. The structure was grown on a n-GaN buffer at a substrate temperature of 730oC. An AlN template (on sapphire) was used. As shown in the cross-sectional z-contrast HAADF STEM image of the sample

(Figure 4.8), for up to 1.2 ML GdN, no change in the structure of the heavily Gd doped discrete GdN region is observed. However, upon reaching a thickness of 1.2ML, Gd precipitates out of the GaN matrix and forms discrete clusters.

One benefit of having an extremely low GdN growth rate is that the RHEED transition can be seen very clearly. As shown in Figure 4.9, for the ﬁrst 0.4ML of GdN

79 coverage, the pattern is representative of the wurtzite Ga-face 1x2 reconstruction.

Between 0.5ML and 1.2 ML, the 2x4 reconstruction happens. Beyond 1.2 ML of coverage, the wurtzite1x2 pattern again is visible, indicating a temporary change in the surface reconstruction during growth of the GdN layer between 0.5 ML and 1.2 ML.

Figure 4.8: The design of the TEM stack sample and the TEM image of the structure.

Figure 4.9: in-situ RHEED patterns for GdN deposition.

When comparing the cross-sectional STEM image and in-situ RHEED pattern, it is of great interest to correlate the change of structure with the surface reconstruction.

Thus high resolution STEM is very useful to study the onset of the GdN structure change by zooming in a bright region of 0.6 ML and 1.2 ML. From the atomic resolution data shown in Figure 4.10 , distinct cubic particles of GdN are observed in the 1.4ML layer, but no second phase exists in 0.6 ML layer. Thus we can conclude that when the GdN deposited is thinner than 1.2 ML, Gd is soluble in the GaN matrix and thus maintains the hexagonal structure. GdN cubic precipitation takes place when there is more than 1.2 ML of GdN deposited. After 1.2 ML and up to 2.4ML, GdN particles with a clearly cubic

81 structure surrounded by a hexagonal GaN matrix can be seen. Additional Gd deposition leads to increased lateral growth, suggesting that the height of the nanoparticle is self limited and further growth will proceed by lateral expansion of the GdN islands.

Figure 4.10: Atomic resolution STEM images of GdN particles in GaN matrix.

After having clear evidence to correlate the RHEED transition to the formation of the cubic GdN, a model is proposed to explain the island formation dynamics. When the

Gd shutters open, Gd atoms start impinging the surface and wetting it further. This explains the dim RHEED pattern at the early age of growth. When the deposited GdN approaches 0.4ML, impinging Gd atoms are still soluble in overgrowth GaN only with less Ga droplets on the surface, which is the reason for the RHEED pattern brightening up. When Gd is more than 0.5 ML, it is speculated that wurtzite GdN forms. This is when

82 the RHEED transition happens. When the GdN exceeds 1.2 ML, due to the huge strain

(10%) energy, GdN starts to precipitate and island formation is favored to minimize the surface energy. However, the area ratio of GaN to imbedded GdN nano-particles is still high. Thus the RHEED pattern is dominated by the regular GaN surface reconstruction.

Later on, due to the overgrowth between the islands, the following Ga spacers maintain the single crystalline wurtzite structure. A smooth surface of the sample is confirmed by

AFM as shown in Figure 4.11, with 1.2 nm RMS, indicating the highly crystalline GaN matrix.

Figure 4.11: AFM image of the stack sample (GdN in GaN) surface.

4.4 Structural characterization of GdN/GaN SL

After calibration of the GdN precipitation threshold, a SL structure of 50 periods of 2.4ML GdN with 10 nm GaN spacers is designed (illustrated in Figure 4.12). Although

GdN forms nano-particles after exceeding 1.2 ML, it is more favorable to incorporate greater amounts of GdN to study the magnetic properties. A similar growth condition to the TEM stack structure is used with higher Gd/N ratio, and thus a higher deposition rate

(2 ML/min). The cross-sectional STEM image shown in Figure 4.12 provides atomic number contrast, showing the expected discrete GdN particles in a GaN matrix with a

GaN/GdN period thickness of 11.8  0.4 nm and GdN layer thickness of 5.6  0.3 nm obtained from image analysis.

Then the high resolution XRD ω-2θ scan is presented in Figure 4.13. The superlattice peaks indicate defined SL interfaces and good periodicity of the structure. In addition, the fact that there is only one single GdN peak existing and that the angle is consistent with the rock-salt (111) peak, indicates clear epitaxial orientation of the GdN

[111] to the wurtzite [0001]. When compare this result to the STEM image where a cubic precipitate was observed at 1.4 ML Gd nominal deposition, a conclusion can be draw that the cubic in 1.4 ML was a metamorphic state where (111) is not aligned to hexagonal

(0002) direction. GdN (111) align to GaN (0002) happens later in the growth when approaching thicker layers. Finally, the GaN/GdN period thickness can be calculated from the analysis of the superlattice fringe angle spacing, and it is determined to be 10.98

84 nm which is consistent with the value estimated from STEM image analysis. The AFM confirms a very smooth surface (0.8 nm RMS) as expected. (Figure 4.14)

Figure 4.12: The illustration and TEM image of a 50 x GdN/GaN SL sample.

10M GaN (0002)

) 1M

AlN (0002) 100k Al O

+1 2 3

-1

counts/s

( +2

10k SL

GdN (111) -2 1k SL

Intensity 100

10 15 16 17 18 19 20 21 22 o  )

Figure 4.13: HRXRD ω-2θ scan of the 50 x SL structure.

Figure 4.14: AFM of the GdN/GaN SL surface.

In summary, GdN/GaN SL was achieved with 50 periods of 2.4 ML GdN embedded in a single crystalline GaN matrix. The structural characterization has been done to confirm to GdN(111) parallel to GaN(0002) epitaxial relationship. There are a lot of potential to continue the growth study of the GdN in GaN matrix. First, one can deposit intermediate nominal GdN thickness such as the1.4 ML to study the metamorphic state of GdN, in order to study whether that phase is accountable for some unique magnetic phase. Also, the TEM calibration sample was stopped at the GdN nominal thickness of 2.4 ML due to the limited time with a super slow growth rate of GdN. When growth rate of GdN is increased, thicker GdN layer can be tried and further islanding dynamics can be studied. On the other hand, we can study how many GaN monolayers are needed to bury the islands and maintain the single crystalline matrix. The systematical way for this part of the study can reveal the role of GaN to GdN ratio as to the magnetic phase.

Chapter 5 : Conclusion

In summary, three kinds of SL structures have been demonstrated with the in- plane strain varying from 3% to 13% grown by PAMBE. For the smallest strain energy system in this study (GaN/AlN SLs), a novel growth condition for GaN was tested and the different relaxation mechanisms were probed. The results indicate that for high T condition, it won’t be too useful unless there are few repeats of SL such that the structure is fully strained. Otherwise, this condition can be applied in a heterostructure with only one interface. In fact, the latter has been achieved in a HEMT structure which produces high 2DEG mobility. In this part, structural characterizations were the center of the study. Further investigations of this new condition for SLs involve optical and electrical measurements.

Then, for the 9% in-plane lattice mismatch system (GdN/GaN SLs), pioneer work has been shown to incorporate rare-earth nanoparticles into single crystalline GaN matrix. The close study of the growth dynamics for island formation has been built a foundation for the demonstration of several devices achieved by collaborators. First, Kent et al. [77] have demonstrated very sharp 4f Gd emission from a PINLED structure. Also,

Krishnamorthy et al. [76] have demonstrated that by incorporate GdN islands into a GaN based tunneling junction, it achieves very low tunneling resistance. The future direction

88 for this study involves identify different magnetic phases and the origins of those phases.

From growth prospective, one can test out the possibility of matching GdN with InN matrix due to a very good lattice match. Also, the effects of growth condition on island formation mechanism are also of significant interest.

Lastly, this work demonstrate the first highly quantum confined InN/AlN MQWs.

The lattice mismatch in this system, 13%, makes it the extreme case in III-nitrides family.

The intensive structural characterizations by STEM, HRXRD identifies the {10 ̅5} semipolar facet planes. Also, the polarization at these semipolar InN/AlN interfaces is calculated and is substantially reduced from that of an identical c-plane oriented interface. The calculation provides a roadmap for polarization manipulation in the

InN/AlN system, and to prove that InN/AlN system has the similar versatile for polarization engineering. One direction for future work is to achieve 1 and 2 ML InN embedded in AlN matrix to investigate whether smoother interfaces are possible. Also, smaller temperature increment steps are useful to realize the best temperature condition for the InN/AlN MQWs. Furthermore, exploration the growth condition effects on semipolar plane is very interesting. Because the InN/AlN MWQs will show more versatility when there is the freedom to choose which semipolar plane to grow on, thus modify the polarization charges at the interface. Finally, when one want to utilize the extreme polarization at InN/AlN interfaces, c-direction is the way to go, thus to achieve the InN/AlN MQWs in nanowires are very important. The possible starting point includes

89 growing AlN on Si (001) substrate and on GaN template. The author has already observed the nanowire formation on top of the GaN template when at low III/V ratio.

References

1. Tchernycheva, M., et al., Systematic experimental and theoretical investigation of intersubband absorption in GaN/AlN quantum wells. Physical Review B, 2006. 73(12): p. 11. 2. Smith, D.L. and V.Y. Pickhardt, Molecular-Beam Epitaxy of Ii-Vi Compounds. Journal of Applied Physics, 1975. 46(6): p. 2366-2374. 3. McCray, W.P., MBE deserves a place in the history books. Nature Nanotechnology, 2007. 2(5): p. 259-261. 4. Kometani, T.Y. and W. Wiegmann, Measurement of Ga and Al in a Molecular- Beam Epitaxy Chamber by Atomic-Absorption Spectrometry (Aas). Journal of Vacuum Science & Technology, 1975. 12(4): p. 933-936. 5. Cho, A.Y. and W.C. Ballamy, Gaas Planar Technology by Molecular-Beam Epitaxy (Mbe). Journal of Applied Physics, 1975. 46(2): p. 783-785. 6. Dabrowska-Szata, M., Analysis of RHEED pattern from semiconductor surfaces. Materials Chemistry and Physics, 2003. 81(2-3): p. 257-259. 7. Strite, S. and H. Morkoc, Gan, Ain, and Inn - a Review. Journal of Vacuum Science & Technology B, 1992. 10(4): p. 1237-1266. 8. Lafont, U., H. van Zeijl, and S. van der Zwaag, Increasing the reliability of solid state lighting systems via self-healing approaches: A review. Microelectronics Reliability, 2012. 52(1): p. 71-89. 9. Razeghi, M. and R. McClintock, A review of III-nitride research at the Center for Quantum Devices. Journal of Crystal Growth, 2009. 311(10): p. 3067-3074. 10. Johnson, M.A.L., et al., A critical comparison between MOVPE and MBE growth of III-V nitride semiconductor materials for opto-electronic device applications. Mrs Internet Journal of Nitride Semiconductor Research, 1999. 4: p. art. no.- G5.10. 11. Heying, B., et al., Control of GaN surface morphologies using plasma-assisted molecular beam epitaxy. Journal of Applied Physics, 2000. 88(4): p. 1855-1860. 12. Heying, B., et al., Optimization of the electron mobilities in GaN grown by plasma-assisted molecular beam epitaxy, in Proceedings of the International Workshop on Nitride Semiconductors. 2000, Inst Pure Applied Physics: Tokyo. p. 158-161. 13. Tarsa, E.J., et al., Homoepitaxial growth of GaN under Ga-stable and N-stable conditions by plasma-assisted molecular beam epitaxy. Journal of Applied Physics, 1997. 82(11): p. 5472-5479.

14. Okumura, H., et al., Analysis of MBE growth mode for GaN epilayers by RHEED. Journal of Crystal Growth, 1998. 190: p. 364-369. 15. Held, R., et al., N-limited versus Ga-limited growth on GaN(0001) by MBE using NH3. Surface Review and Letters, 1998. 5(3-4): p. 913-934. 16. Okumura, H., et al., Analysis of MBE growth mode for GaN epilayers by RHEED. Journal of Crystal Growth, 1998. 189: p. 364-369. 17. Mula, G., et al., Surfactant effect of gallium during molecular-beam epitaxy of GaN on AlN (0001). Physical Review B, 2001. 64(19): p. 12. 18. Koblmuller, G., et al., In situ investigation of growth modes during plasma- assisted molecular beam epitaxy of (0001) GaN. Applied Physics Letters, 2007. 91(16): p. -. 19. Koblmuller, G., et al., High electron mobility GaN grown under N-rich conditions by plasma-assisted molecular beam epitaxy. Applied Physics Letters, 2007. 91(22): p. -. 20. Koblmuller, G., et al., Dislocation reduction in AlGaN/GaN heterostructures on 4H-SiC by molecular beam epitaxy in the thermal decomposition regime. Applied Physics Express, 2008. 1(6): p. -. 21. Adelmann, C., et al., Gallium adsorption on (0001) GaN surfaces. Physical Review B, 2003. 67(16): p. 9. 22. Brown, J.S., et al., Ga adsorbate on (0001) GaN: In situ characterization with quadrupole mass spectrometry and reflection high-energy electron diffraction. Journal of Applied Physics, 2006. 99(7): p. -. 23. Adelmann, C., et al., Dynamically stable gallium surface coverages during plasma-assisted molecular-beam epitaxy of (0001) GaN. Journal of Applied Physics, 2002. 91(12): p. 9638-9645. 24. Koblmueller, G., et al., Growth diagram and morphologies of AlN thin films grown by molecular beam epitaxy. Journal of Applied Physics, 2003. 93(12): p. 9591-9596. 25. Gallinat, C.S., et al., A growth diagram for plasma-assisted molecular beam epitaxy of In-face InN. Journal of Applied Physics, 2007. 102(6): p. -. 26. Esaki, L. and R. Tsu, Superlattice and Negative Differential Conductivity in Semiconductors. IBM Journal of Research and Development, 1970. 14(1): p. 61- 65. 27. Yu, P.Y. and M. Cardona, Fundamentals of semiconductors : physics and materials properties. 3rd, rev. and enlarged ed. 2001, Berlin ; New York: Springer. xviii, 639 p. 28. Dimakis, E., et al., Growth optimization of an electron confining InN/GaN quantum well heterostructure. Journal of Electronic Materials, 2007. 36(4): p. 373-378. 29. Jogai, B., Three-dimensional strain field calculations in multiple InN/AlN wurtzite quantum dots. Journal of Applied Physics, 2001. 90(2): p. 699-704.

30. Lin, W., et al., Band Engineering in Strained GaN/ultrathin InN/GaN Quantum Wells. Crystal Growth & Design, 2009. 9(4): p. 1698-1701. 31. Lin, W., S.P. Li, and J.Y. Kang, Near-ultraviolet light emitting diodes using strained ultrathin InN/GaN quantum well grown by metal organic vapor phase epitaxy. Applied Physics Letters, 2010. 96(10). 32. Yoshikawa, A., et al., Proposal and achievement of novel structure InN/GaN multiple quantum wells consisting of 1 ML and fractional monolayer InN wells inserted in GaN matrix. Applied Physics Letters, 2007. 90(7). 33. Wu, J. and W. Walukiewicz, Band gaps of InN and group III nitride alloys. Superlattices and Microstructures, 2003. 34(1-2): p. 63-75. 34. Kuo, C.T., et al., Is electron accumulation universal at InN polar surfaces? Applied Physics Letters, 2011. 98(5). 35. Wu, C.L., C.H. Shen, and S. Gwo, Valence band offset of wurtzite InN/AlN heterojunction determined by photoelectron spectroscopy. Applied Physics Letters, 2006. 88(3). 36. Miao, M.S., et al., Polarization-Driven Topological Insulator Transition in a GaN/InN/GaN Quantum Well. Physical Review Letters, 2012. 109(18). 37. Yu, E.T., et al., Spontaneous and piezoelectric polarization effects in III-V nitride heterostructures. Journal of Vacuum Science & Technology B, 1999. 17(4): p. 1742-1749. 38. Bernardini, F., V. Fiorentini, and D. Vanderbilt, Spontaneous polarization and piezoelectric constants of III-V nitrides. Physical Review B, 1997. 56(16): p. 10024-10027. 39. Bernardini, F. and V. Fiorentini, Spontaneous versus piezoelectric polarization in III-V nitrides: Conceptual aspects and practical consequences. Physica Status Solidi B-Basic Research, 1999. 216(1): p. 391-398. 40. Gogneau, N., et al., Surfactant effect of gallium during the growth of GaN on AlN(000(1)over-bar) by plasma-assisted molecular beam epitaxy. Applied Physics Letters, 2004. 85(8): p. 1421-1423. 41. Kadir, A., et al., MOVPE growth and characterization of InN/GaN single and multi-quantum well structures. Journal of Crystal Growth, 2008. 311(1): p. 95-98. 42. Dvorak, M., S.H. Wei, and Z.G. Wu, Origin of the Variation of Exciton Binding Energy in Semiconductors (vol 110, 016402, 2013). Physical Review Letters, 2013. 110(16). 43. Bourret, A., et al., Growth of aluminum nitride on (111) silicon: Microstructure and interface structure. Journal of Applied Physics, 1998. 83(4): p. 2003-2009. 44. Cheze, C., et al., In situ investigation of self-induced GaN nanowire nucleation on Si. Applied Physics Letters, 2010. 97(4). 45. Landre, O., et al., Molecular beam epitaxy growth and optical properties of AlN nanowires. Applied Physics Letters, 2010. 96(6). 46. Hsu, K.Y., et al., Molecular beam epitaxy growth of wurtzite AlN nanotips. Applied Physics Letters, 2008. 93(18).

47. Xie, M.Y., et al., Elastic constants, composition, and piezolectric polarization in InxAl1-xN: From ab initio calculations to experimental implications for the applicability of Vegard's rule. Physical Review B, 2012. 86(15). 48. Romanov, A.E., et al., Strain-induced polarization in wurtzite III-nitride semipolar layers. Journal of Applied Physics, 2006. 100(2). 49. Nam, K.B., et al., Growth and deep ultraviolet picosecond time-resolved photoluminescence studies of AlN/GaN multiple quantum wells. Applied Physics Letters, 2001. 78(23): p. 3690-3692. 50. Valdueza-Felip, S., et al., Carrier localization in InN/InGaN multiple-quantum wells with high In-content. Applied Physics Letters, 2012. 101(6). 51. Fernandez-Garrido, S., Z. Gacevic, and E. Calleja, A comprehensive diagram to grow InAlN alloys by plasma-assisted molecular beam epitaxy. Applied Physics Letters, 2008. 93(19). 52. Wang, Z., et al., Ultrafast hole burning in intersubband absorption lines of GaN/AlN superlattices. Applied Physics Letters, 2006. 89(15): p. 3. 53. Monroy, E., et al., Observation of hot luminescence and slow inter-sub-band relaxation in Si-doped GaN/AlxGa1-xN (x=0.11, 0.25) multi-quantum-well structures. Journal of Applied Physics, 2006. 99(9): p. -. 54. Kandaswamy, P.K., et al., Strain relaxation in short-period polar GaN/AlN superlattices. Journal of Applied Physics, 2009. 106(1): p. 9. 55. Kandaswamy, P.K., et al., GaN/AlN short-period superlattices for intersubband optoelectronics: A systematic study of their epitaxial growth, design, and performance. Journal of Applied Physics, 2008. 104(9): p. 16. 56. Monroy, E., et al., MBE growth of nitride-based photovoltaic intersubband detectors. Superlattices and Microstructures, 2006. 40(4-6): p. 418-425. 57. Koblmuller, G., et al., Influence of Ga/N ratio on morphology, vacancies, and electrical transport in GaN grown by molecular beam epitaxy at high temperature. Applied Physics Letters. 97(19): p. 3. 58. Moram, M.A. and M.E. Vickers, X-ray diffraction of III-nitrides. Reports on Progress in Physics, 2009. 72(3): p. -. 59. Holý, V., U. Pietsch, and T. Baumbach, High-resolution X-ray scattering from thin films and multilayers. Springer tracts in modern physics,. 1999, Berlin ; New York: Springer. xi, 256 p. 60. Kandaswamy, P.K., et al., Strain relaxation in short-period polar GaN/AlN superlattices. Journal of Applied Physics, 2009. 106(1). 61. Northrup, J.E., et al., Structure of GaN(0001): The laterally contracted Ga bilayer model. Physical Review B, 2000. 61(15): p. 9932-9935. 62. Miles, R.H., et al., Accommodation of Lattice Mismatch in Gexsi1-X/Si Superlattices. Journal of Vacuum Science & Technology B, 1988. 6(4): p. 1382- 1385. 63. Miles, R.H., et al., Dependence of Critical Thickness on Growth Temperature in Gexsi1-X/Si Superlattices. Applied Physics Letters, 1988. 52(11): p. 916-918.

64. Griebel, M., et al., Tunable subpicosecond optoelectronic transduction in superlattices of self-assembled ErAs nanoislands. Nature Materials, 2003. 2(2): p. 122-126. 65. Lambrecht, W.R.L., Electronic structure and optical spectra of the semimetal ScAs and of the indirect-band-gap semiconductors ScN and GdN. Physical Review B, 2000. 62(20): p. 13538-13545. 66. Palmstrom, C.J., N. Tabatabaie, and S.J. Allen, Epitaxial-Growth of Eras on (100)Gaas. Applied Physics Letters, 1988. 53(26): p. 2608-2610. 67. Klenov, D.O., et al., Interface atomic structure of epitaxial ErAs layers on (001)In0.53Ga0.47As and GaAs. Applied Physics Letters, 2005. 86(24). 68. Klenov, D.O., et al., Scanning transmission electron microscopy of ErAs nanoparticles embedded in epitaxial In0.53Ga0.47As layers. Applied Physics Letters, 2005. 86(11). 69. Kadow, C., et al., Subpicosecond carrier dynamics in low-temperature grown GaAs on Si substrates. Applied Physics Letters, 1999. 75(17): p. 2575-2577. 70. Kadow, C., et al., Self-assembled ErAs islands in GaAs: Growth and subpicosecond carrier dynamics. Applied Physics Letters, 1999. 75(22): p. 3548- 3550. 71. Nair, H.P., A.M. Crook, and S.R. Bank, Enhanced conductivity of tunnel junctions employing semimetallic nanoparticles through variation in growth temperature and deposition. Applied Physics Letters, 2010. 96(22). 72. Yoshitomi, H., et al., Optical and magnetic properties in epitaxial GdN thin films. Physical Review B. 83(15): p. 155202. 73. Dhar, S., et al., Colossal Magnetic Moment of Gd in GaN. Physical Review Letters, 2005. 94(3): p. 037205. 74. Natali, F., et al., Epitaxial growth of GdN on silicon substrate using an AlN buffer layer. Journal of Crystal Growth. 312(24): p. 3583-3587. 75. Roever, M., et al., Tracking defect-induced ferromagnetism in GaN:Gd. Physical Review B. 84(8): p. 081201. 76. Krishnamoorthy, S., et al., GdN Nanoisland-Based GaN Tunnel Junctions. Nano Letters, 2013. 13(6): p. 2570-2575. 77. Scarpulla, M.A., et al., GdN (1,1,1) heteroepitaxy on GaN (0,0,0,2) by N2 plasma and NH3 molecular beam epitaxy. Journal of Crystal Growth, 2009. 311(5): p. 1239-1244. 78. Kent, T. F., et al. Atomically Sharp 318nm Gd:AlGaN Ultraviolet Light Emitting Diodes on Si with Low Threshold Voltage, Appl. Phys. Lett. 102, 201114 (2013) . 79. Kent, T. F., et al. Room Temperature Ferrmagnetism in GaN-AlN quantum Confined Heterostructures, APS march meeting, Session H15.

Appendix A: Optical measurement for intersubband transition

Due to the selection rule, all samples (discussed in Chapter 3) and reference which is bare AlN substrate were prepared in multi-bounce 45 o optical waveguides as shown in Figure A.1.The experiment consists of four measurements of transmission spectra with the arc lamp monochromator set to 0 degree and the IR grating set. IR achromatic collimation lenses are used after the light source, before and after the waveguide and before the spectrometer which is connected to an InGaAs IR detector.

Glan-Thompson polarizer is in the path and transmission spectra for either the SL sample or the reference sample are collected for both s-polar and p-polar relative to the sample normal. The experiment consists of four measurements, a sample at both s and p polarizations and a reference at both polarizations. The transmittance spectrum is then calculated as the quotient of the sample over reference spectrum. When doing low temperature measurement, samples were mounted in the close cycle He cryostat, with the redesigned expander mount, which allows vibration free operation, as well as low thermal loss to touching of the cold finger to the side wall. When doing room temperature measurement, the front and the back side of the cryostat windows are removed to prevent the minimum scatting of the light.

As shown in the Figure A.3, the absorption spectra are very noisy. This is due to the process of quotient of the sample/reference spectrum. When the size of the sample

96 and reference are not identical, it will change the optics pass. This is the reason all groups doing ISB measurement are using a FTIR with a special attachment to detect the signal.

Figure A.1: Optical setup for transmission measurements of the ISB absorption.

P-polarized light S-polarized light

10k

Transmission(a.u.)

10 1200 1600 2000 wavelength (nm)

Figure A.2: Transmission spectrum of S and P-polarized light.

110511A IR Transmission Spectra for different polarization 1.5 1.4 p-polarized 1.3 s-polarized 1.2 1.1 1 0.9 0.8 0.7

0.6 tansmittance Units)(Arb. 0.5 0.6 0.8 1.0 1.2 Energy (eV) 110519A IR Transmission Spectra for different polarization 1.6 1.4 p pol N rich 1.2 s pol N rich 1

0.8

0.6

tansmittance (Arb. Units) tansmittance 0.4

0.6 0.8 1.0 1.2 Energy (eV)

Figure A.3: Absorption spectrum of Ga-rich and N-rich GaN/AlN MQWs samples.

Appendix B: Gd:AlN/GaN 2DEG system

All samples were grown by PAMBE on a Veeco 930 MBE system on semi- insulating GaN/sapphire templates available from Lumilog. Samples were grown under the Ga-rich (III/V=2) with growth temperature of 720 oC. Sample structures is shown in

Figure A.4. The 150 nm of GaN buffer layer was first deposited on AlN substrate, followed by the 5 nm of AlN layer on top. This is a regular HEMT structure with the high electron density due to the polarization at the interface. A thin layer of GdN was inserted very close to the AlN/GaN interface with the distance of d away from the interfaces. One sample was grown with d=0.4nm which is aligned with the center of 2DEG and another with d=10 nm which should not have any overlap with 2DEG. Gd dopant concentrations for the doping were calibrated by SIMS and RBS analysis at 2.0x1014 cm2, corresponding to 0.2 monolayer of GdN (Figure A.6). The simulated band diagram is in Figure A.5 and measured magnetic properties are in Figure A.7

Figure A.4: An illustration of the Gd doped AlN/GaN HEMT structure.

100

Figure A.5: Band diagram of the AlN/GaN HEMT.

101

20 9.0x10 2500000 Gd Al marker 2000000 6.0x1020

) 1500000 3

1000000 3.0x1020

500000

0.0 0

4.00E+020

20 Al units) (arb.

6.0x10 3.00E+020 concentration (atoms/cm concentration

2.00E+020

Gd 3.0x1020

1.00E+020

0.00E+000 0 10 20 30 Depth (nm)

Figure A.6: SIMS depth profile of the Gd concentration.

102

300 T = 5 K d = 10nm 200

100

/ Gd 0 d = 0.4 nm

 -100

-200

-300 -4 -3 -2 -1 0 1 2 3 4 H (kOe)

Figure A.7: Hysteresis loops of d=0.4nm and d=10nm samples at 5K

103

Appendix C: Magnetic properties and device processing for spin Seebeck in GaMnAs C.1 Magnetic properties of GaMnAs

All samples were grown in the method described in Mack et al. 2008. Because the semi-insulating (001) GaAs substrate does not rotate during growth, As:Ga flux ratio varies continuously thus changing the Mn concentration. For the Spin-Seebeck measurement, we took the center piece along [-110] where As:Ga is precisely balanced to achieve stoichiometry, when there is no suppression of ferromagnetism. Therefore the center piece gives the highest Mn concentration, result in high Currie temperature. An example of the magnetic data analysis is done on sample 080326B. (Figure A.8)

Magnetic characterizations up to 300 K were carried out in a Quantum Design superconducting quantum interference device (SQUID) magnetometer. The sample held in a plastic straw with a small piece of straw and tape around the sample piece to keep it in place. The magnetic field was applied along the Spin-Seebeck measurement direction and other magnetic axis directions for anisotropy study. Hysteresis loops were recorded at same temperatures where the Spin-Seebeck measurements were done. Due to the diamagnetic background brought by GaAs substrate, the ferromagnetic data analysis required to have diamagnetic background subtraction. The process to correct the diamagnetic background subtraction and sample volume normalization is shown in

Figure A.8. The anisotropy of GaMnAs magnetic properties is shown in Figure A.9, and

104

M vs. T measurement is shown in Figure A.10. The sample was first warm it up to 300K without field; then the dipoles were aligned at 1T when cooling to 4 K. We measure the magnetization vs. temperature curve when warming up the system at 300 Oe field. Currie temperature Tc is measured at 150K.

-5 6.0x10 -2.0x10-4 Data point

-5 4.0x10 -4 Linear Fit of Long Moment (emu) -3.0x10

2.0x10-5 -4.0x10-4

0.0

-5.0x10-4

-2.0x10-5 -4

Long Moment (emu) Moment Long -6.0x10 Long Moment (emu) Moment Long -4.0x10-5 -7.0x10-4 -6.0x10-5 -1000 -500 0 500 1000 10000 12000 14000 16000 18000 20000 Field (Oe) (a) (b) Field (Oe) 1.0x10-4

8.0x10-5 1.0x102 080326B hysteresis loop[110]

6.0x10-5 50K

(emu) -5 1 back back 4.0x10 5.0x10

2.0x10-5 0.0 0.0

-2.0x10-5

-4.0x10-5 -5.0x101

-6.0x10-5

Long Moment (emu/cm^3) Moment Long

Long moment with moment Long ground subtraction ground subtraction -8.0x10-5 -1.0x102

-1.0x10-4 -1000 -500 0 500 1000 -1000 -500 0 500 1000 (c) Field (Oe) (d) Field (Oe)

Figure A.8: (a). Hysteresis loop raw data, (b) dominant diamagnetic signal at high field,

105

120 75 080326B hysteresis loop 50K 080326B hysteresis loop 50K 80 [100] 50 [100] [110] [110] 40 [-110] 25 [-110] 0 0

-40 -25 Long Moment (emu/cm^3) Moment Long -80 (emu/cm^3) Moment Long -50

-120 -75 -1000 -500 0 500 1000 -100 0 100 200 Field (Oe) Field (Oe)

Figure A.9: Hysteresis loop of GaMnAs (Mn%=15.8) along different applied magnetic

field directions.

106

70 [100] 60 [110] 50 40 30 20

Long Moment (emu/cm^3) Moment Long 10 0 0 100 200 300 Temperature (K)

Figure A.10: Magnetization vs. Temperature revels Tc of 150K.

C.2 GaMnAs spin Seebeck device processing

The samples were cleaved into 3-5 mm by 10-15 mm rectangular pieces. A thin layer of Ti (less than 1 nm) was deposited on GaMnAs for adhesion purpose in Electron

Beam Evaporator, followed by a 10 nm of Pt layer which is the same thickness as sample prepared by Uchida et al. The example of a processed device is shown in Figure A.11.

107

Figure A.11: An GaMnAs cleaved wafer and processed sample on a modified PPMS

station.

The contacts of first generation were evaporated using a hand-taped Al foil as the mask. The contacts of second generation were evaporatd using a machined Al alloy shadow mask (Figure A. 13) thus the width of the Pt contacts is uniform and better controlled. To exam the contact properties of Pt strip, I-V measurements are conducted, shown in Figure A.12. The results indicate Pt has Ohmic contact to GaMnAs surface.

108

080326B 0.2 Pt 1 Pt 2 Pt 1 Pt 3 Pt 1 Pt 4 0.1 Pt 2 Pt 3 Pt 2 Pt 4 Pt1 Pt2 Pt 3 Pt 4

0.0 Pt3 Pt4 V (V) -0.1

-0.2

-8.0x10-4 0.0 8.0x10-4 Pt strips I (A)

Figure A.12: The illustration of the Pt contacts geometry (first generation) and I-V

characterization of the Pt contacts.

~ 20 mm

Figure A.13: Second generation of the shadow mask.

109

Appendix D: FTIR measurement for GaN/AlN nanowires D.1 MQWs on nanowires

Figure A.14: Cross-sectional and top-view SEM images of the GaN/AlN MQW

nanowires.

110

The GaN/AlN SL structure is grown on top of the GaN nanowire using a two-step method and is kept the growth temperature at the second step (high temperature condition). As shown in Figure A.14, the nanowire diameter increases along the growth.

This is due to the radio/vertical growth rate ratio is relatively high for AlN as the result of the low Al atom mobility. The STEM images of the nanowire SL structure are shown in

Figure A.15.

Figure A.15: STEM images of MQW nanowires.

111

D.2 FTIR measurements on MQWs nanowire samples

FTIR on SL NWs on reflectant mode MQW sample 1 SQW sample 1 200 MQW sample 2

GaN nanowire sample

T% a.u. T% 100

1500 2000 2500 Wavelength (nm)

Figure A.16: FTIR measurements on MQWs nanowire samples on reflectant mode.

The preliminary FTIR measurement results are showing in Figure A.16. The absorption peak shifts as the change of the quantum thickness. Moreover, a reference sample with only GaN nanowire grown under the identical condition was shown in the

112 same figure as well. The flat spectrum for the region shown here confirms that the absorption features are not from the nanowire plasmatic absorption.

D.3 Absorption on MQWs drop-cast onto glasslides

In order to develop the method to gather ISB absorption, drop-casted on the glass slide has been developed. First, a MQW nanowires (NWs) sample grown on Si substrate was in ultrasonic acetone solution for 2 hrs. Upon removal of the Si substrate from the solution that full of nanowires, the solution was transferred into a drop-cast tube with a rate control valve. Then, the valve was opened to the minimum to produce the slowest drop rate such that before the next drop of NWs solution fall onto the glass slide, the previous drop of acetone has evaporated and left the NWs behind. In order to enhance the evaporation of acetone solution, the glass slide was put onto a heat station. In order to confirm the feasibility of this method, the identical set of drop-cast was done onto a Si substrate for SEM measurement. As shown in Figure A.17 bottom, there are cluster or single nanowires laying on the Si substrate, demonstrating the success of drop-cast NWs onto glass slide. The absorption of MQWs band-to-band transition was shown in the

Figure A.17 as well. The consistence of the absorbed peak center and simulation result reveals that the drop-cast method is one feasible way to do absorption for NWs. Another dipping feature at 290 nm arises from the glass absorption.

113

glassslide

NWs on glassslide 1

0.1 3 2

0 transmittance -1 3.2346

(eV) -2 0.01 c eV E -3 -4 -5 -6 0 50 100 150 z(A) 290 300 310 320 330 340 350 360 370 380 390 400 410 420 430 wavelength (nm)

Figure A.17: Absorption measurement on GaN/AlN MQWs NWs drop-casted onto glass

slide.

114