Mechanisms in AlGaN-based Quantum Well Structures and Epitaxial Layers

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Mechanisms in AlGaN-based Quantum Well Structures and Epitaxial Layers

vorgelegt von M. Sc.

Christian Frankerl

an der Fakult¨ at II – Mathematik und Naturwissenschaften der Technischen Universit¨ at Berlin

zur Erlangung des akademischen Grades Doktor der Naturwissenschaften

- Dr. rer. nat. - genehmigte Dissertation

Promotionsausschuss:

Vorsitzender: Prof. Dr. Michael Lehmann Gutachter: Prof. Dr. Michael Kneissl Gutachter: Prof. Dr. Axel Hoffmann

Gutachter: Prof. Dr. Andreas Waag (TU Braunschweig) Tag der wissenschaftlichen Aussprache: 09. Juli 2021

Berlin 2021

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In this thesis, optical investigations of AlGaN-based single quantum well (SQW) and multiple quantum well (MQW) structures and epitaxial layers are performed and discussed, with a strong focus on time-integrated and time-resolved, temperature- and excitation power density-dependent photoluminescence (PL) spectroscopy. A profound understanding of the optical properties of these heterostructures is crucial for the com- mercialisation of light-emitting diodes (LEDs) operating in the deep ultraviolet (DUV) spectral region.

The internal quantum efficiency (IQE) of a light-emitting quantum well (QW) heterostructure is generally regarded as a reliable indicator of its growth quality. In this thesis, several approaches to improve the IQE of the AlGaN-based active region of a DUV LED are investigated, for instance by adjusting the growth morphology, reducing the threading dislocation density (TDD) and optimising the design of the active region itself. Furthermore, a variety of frequently employed methodologies for IQE determination are implemented and discussed, unveiling several potential sources of error which may potentially distort the IQE value obtained by simple optical measurements.

These include, among others, carrier transport effects, non-resonant optical excitation conditions, and growth morphology issues. As a result, a standard set of measurement conditions ensuring reliable IQE determination is derived. The results of this investigation considerably improve the comparability of the experimental results obtained throughout the global research community.

This thesis furthermore provides a detailed investigation of the phenomenon of carrier localisation in AlGaN-based QW heterostructures, covering a large number of samples with highly deviating QW compositions. Among these are AlGaN/AlN heterostructures with varying Al content and well width, both in SQW and MQW configuration.

The QW width is shown to be the main structural parameter determining the localisation strength, while the Al content has only a weak impact. It is demonstrated that the localisation effect inherently controls the carrier recombination dynamics of AlGaN-based QW heterostructures. At low temperatures and low carrier densities,

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the PL emission spectrum of the investigated structures is found to mainly originate from localised carriers, while the so-called ‘efficiency droop’ correlates with the saturation of localised states at increasing carrier densities. The anomalous temperature dependence of the PL peak emission energy is attributed to the thermal redistribution of carriers among a potential landscape of localised states and their delocalisation at elevated temperatures. Finally, the microscopic origins of the localisation effect in the AlGaN material system are investigated and compared to the InGaN sister system. A fundamentally different localisation process is observed, with the localisation strength directly coupled to the oscillator strength. As a consequence, a novel model resolving the phenomenon of ‘efficiency droop’ is developed, proposing delocalisation-induced Auger recombination as the responsible physical mechanism. Thus, this thesis can be considered as an important first step towards the unification of two competing but not necessarily contradicting theories.

One of the major challenges when optimising the external quantum efficiency (EQE) of an LED emitting in the DUV spectral region is to outcouple the light generated in the active region of the device. In the flip-chip architecture, the photons must pass both the AlN template layer and the n-side region of the LED, thus emphasising the importance of high crystal quality in the AlN template. In this thesis, the optical properties of AlN epitaxial layers deposited with a variety of distinct growth conditions are studied, including pulsed and continuous growth mode, varying growth temperatures and V/III ratios, as well as different polarities. Absorption and PL spectroscopy measurements reveal significant defect absorption bands in the DUV spectral region, the nature of which is identified by secondary-ion mass spectroscopy (SIMS) as related to the incorporation of carbon and oxygen impurities. It is demonstrated that high vapor supersaturation inhibits point defect incorporation, while pulsed growth conditions and N-polar domains strongly deteriorate the crystal purity. Similarities with other growth techniques such as physical vapour transport (PVT) and hydride vapour phase epitaxy (HVPE) are highlighted. The results of this thesis allow for a crucial choice of the AlN template growth parameters, inhibiting defect absorption bands at the desired emission wavelengths, in order to provide highly transparent templates for DUV LEDs processed in a flip-chip configuration.

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In der vorliegenden Arbeit werden optische Messungen an AlGaN-basierten Einzel- und Mehrfach-Quantentöpfen (QW) sowie Epitaxieschichten durchgeführt und diskutiert, wobei der Fokus auf zeitintegrierter und zeitaufgelöster, temperatur- und leistungs- dichteabhängiger Photolumineszenz-(PL)-Spektroskopie liegen soll. Ein fundiertes Ver- ständnis der optischen Eigenschaften solcher Heterostrukturen ist für die Kommerzial- isierung von im fernen Ultravioletten (DUV) emittierenden Leuchtdioden (LEDs) von entscheidener Bedeutung.

Als zuverlässiger Indikator für die Wachstumsqualität der QW-Heterostruktur wird zumeist deren interne Quanteneffizienz (IQE) bestimmt. In dieser Arbeit werden zahlreiche Ansätze zur qualitativen Verbesserung der AlGaN-basierten aktiven Re- gion einer DUV LED untersucht und diskutiert. Diese beinhalten die Anpassung der Wachstumsmorphologie, die Reduktion der Versetzungsdichte (TDD) und die Op- timierung der aktiven Region selbst. Weiterhin werden verschiedene weitverbreitete Methoden zur Bestimmung des IQEs auf potentielle Fehlerquellen untersucht, welche den durch optische Messungen bestimmten Zahlenwert verfälschen können. Mehrere solcher Fehlerquellen werden identifizert, einschließlich Ladungsträgertransporteffekte, nicht-resonante optische Anregungsbedingungen und Probleme mit der Wachstumsmor- phologie. Die Resultate dieser Untersuchung führen schließlich zu einer Zusammenstel- lung von standardisierten Messbedingungen, welche eine zuverlässige Bestimmung des IQEs ermöglichen und die Vergleichbarkeit der experimentellen Ergebnisse zwischen unterschiedlichen Arbeitsgruppen sicherstellen.

Die vorliegende Arbeit umfasst ferner eine detaillierte Studie der Ladungstr¨ager- lokalisierung in AlGaN-basierten QW-Heterostrukturen, wobei eine umfangreiche Aus- wahl an Proben mit h¨ochst unterschiedlichen Wachstumsbedingungen abgedeckt wird.

Beispielsweise werden für diese Studie der Al-Gehalt, die QW-Breite sowie die Anzahl der QWs der gewachsenen AlGaN/Al(Ga)N-Heterostrukturen variiert. Die Experi- mente zeigen, dass die QW-Breite die Stärke der Lokalisierung entscheidend beeinflusst, während der Al-Gehalt nur geringe Auswirkungen auf selbige hat. Weiterhin wird

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gezeigt, dass der Lokalisierungseffekt die Rekombinationsdynamik der Ladungsträger in AlGaN-basierten QW-Heterostrukturen grundlegend beeinflusst. Das PL-Emissions- spektrum wird bei niedrigen Temperaturen und geringen Ladungsträgerdichten haupt- sächlich von lokalisierten Ladungsträgern erzeugt und der sogenannte ‘Efficiency Droop’

korreliert mit der Sättigung von lokalisierten Zuständen bei hohen Ladungsträgerdich- ten. Die ungewöhnliche Temperaturabhängigkeit des PL-Emissionsmaximums kann der thermischen Umverteilung der Ladungsträger zwischen den lokalisierten Zuständen und ihrer Delokalisierung bei hohen Temperaturen zugeordnet werden. Schließlich werden die mikroskopischen Ursprünge des Lokalisierungseffekts im AlGaN-Materialsystem genauer beleuchtet und mit dem InGaN-Schwestersystem verglichen. Hierbei ergeben sich Hinweise auf einen Lokalisierungsprozess, der sich fundamental von den Vorgängen in InGaN-basierten Heterostrukturen unterscheidet und direkt an die Oszillatorstärke gekoppelt ist. Aus diesen Erkenntnissen wird ein neuartiges Modell zur Erklärung des

‘Efficiency Droops’ entwickelt. Der zugrundeliegende Mechanismus ist hierbei durch Delokalisierung induzierte Auger-Rekombination, womit diese Arbeit einen wichtigen ersten Schritt zur Vereinigung zweier konkurrierender, aber nicht notwendigerweise widerspr¨uchlicher Theorien darstellt.

Eine der größten Herausforderungen bei der Optimierung der externen Quanten- effizienz (EQE) von LEDs im DUV-Spektralbereich ist es, das in der aktiven Re- gion erzeugte Licht aus dem Bauelement auszukoppeln. In der Flip-Chip-Architektur müssen die Photonen sowohl das AlN-Template als auch die n-Seite der LED passieren, weshalb der Kristallqualität des AlN-Templates eine besondere Bedeutung zukommt.

In der vorliegenden Arbeit werden die optischen Eigenschaften von AlN-Epitaxieschich- ten untersucht, welche mit Hilfe einer Vielzahl von verschiedenen Wachstumsbedin- gungen abgeschieden werden. Diese behinhalten etwa gepulstes und kontinuierliches Wachstum, unterschiedliche Wachstumstemperaturen und V/III-Verhältnisse, sowie verschiedene Polaritäten. Absorptions- und PL-Spektroskopiemessungen offenbaren beträchtliche, von Punktdefekten hervorgerufene Absorptionsbänder, deren zugrunde liegenden Defektatome mittels Sekundärionen-Massenspektrometrie (SIMS) als Sauer- stoff und Kohlenstoff identifiziert werden. Es wird gezeigt, dass ein hoher Grad der Dampfübersättigung den Einbau von Defektatomen reduziert, während gepulstes Wach- stum und das Vorhandensein von N-polaren Domänen die Kristallqualität verschlech- tert. Gemeinsamkeiten mit weiteren Wachstumstechnologien, etwa physikalischer Gas- phasenabscheidung (PVT) und Hydridgasphasenepitaxie (HVPE) werden ebenfalls her- ausgearbeitet. Die Ergebnisse dieser Arbeit erlauben es, die Wachstumsparameter der AlN-Epitaxieschicht derart einzustellen, sodass Absorptionsbänder im gewünschten Emissionwellenlängenbereich der DUV LED unterdrückt werden und hochtransparente AlN-Templates für eine Flip-Chip-Architektur hergestellt werden können.

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Authored Articles

This cumulative thesis is based on the following publications:

[P1] C. Frankerl, M. P. Hoffmann, F. Nippert, H. Wang, C. Brandl, N. Tillner, H.-J.

Lugauer, R. Zeisel, A. Hoffmann, and M. J. Davies,

“Challenges for reliable internal quantum efficiency determination in AlGaN- based multi-quantum-well structures posed by carrier transport effects and morphology issues”,

Journal of Applied Physics 126, 075703 (2019), DOI: 10.1063/1.5100498.

Contribution: Design and execution of all experiments except for atomic force microscopy (AFM) measurements (performed by N. Tillner). Analysis of data and preparation of the manuscript. Numerical calculations were performed by H.

Wang and samples were grown by M. P. Hoffmann.

[P2] C. Frankerl, F. Nippert, M. P. Hoffmann, H. Wang, C. Brandl, H.-J. Lugauer, R.

Zeisel, A. Hoffmann, and M. J. Davies,

“Strongly localized carriers in Al-rich AlGaN/AlN single quantum wells grown on sapphire substrates”,

Journal of Applied Physics 127, 095701 (2020), DOI: 10.1063/1.5144152.

Contribution: Design and execution of all experiments, analysis of data and preparation of the manuscript. Samples were grown by C. Brandl.

[P3] C. Frankerl, F. Nippert, M. P. Hoffmann, C. Brandl, H.-J. Lugauer, R. Zeisel, A.

Hoffmann, and M. J. Davies,

“Carrier dynamics in Al-rich AlGaN/AlN quantum well structures governed by carrier localization”,

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physica status solidi (b) 257, 12 (2020).

DOI: 10.1002/pssb.202000242.

Contribution: Design and execution of all experiments, analysis of data and preparation of the manuscript. Samples were grown by C. Brandl.

[P4] C. Frankerl, F. Nippert, A. Gomez-Iglesias, M. P. Hoffmann, C. Brandl, H.-J.

“Origin of carrier localization in AlGaN-based quantum well structures and im- plications for efficiency droop”,

Applied Physics Letters 117, 102107 (2020), DOI: 10.1063/5.0018885.

Contribution: Design and execution of all experiments, analysis of data and preparation of the manuscript. Numerical calculations were performed by A. Gomez- Iglesias and samples were grown by C. Brandl.

[P5] N. Tillner,* C. Frankerl,* F. Nippert, M. J. Davies, C. Brandl, R. L¨osing, M.

Mandl, H.-J. Lugauer, R. Zeisel, A. Hoffmann, A. Waag, and M. P. Hoffmann,

“Point defect-induced UV-C absorption in aluminum nitride epitaxial layers grown on sapphire substrates by metal-organic chemical vapor deposition”,

physica status solidi (b) 257, 12 (2020), DOI: 10.1002/pssb.202000278.

* Both authors contributed equally to this work.

Contribution: Design of all experiments in collaboration with N. Tillner. Ex- ecution of all experiments except for photoluminescence excitation (PLE) spectroscopy (performed by F. Nippert) and secondary-ion mass spectroscopy (SIMS) measurements (performed by R. L¨osing). Analysis of data and preparation of the manuscript in collaboration with N. Tillner. Samples were grown by N. Tillner.

Co-Authored Articles

The following publications were co-authored alongside this thesis. The relevance of these works for the presented thesis is also provided.

[P6] G. Jacopin, C. Frankerl, N. Tillner, M. J. Davies, G. Rossbach, C. Brandl, M. P.

Hoffmann, R. Zeisel, A. Hoffmann, and H.-J. Lugauer,

“Influence of the growth substrate on the internal quantum efficiency of AlGaN/

AlN multi quantum wells governed by carrier localization”, physica status solidi (b), 2000464 (2020),

DOI: 10.1002/pssb.202000464.

Relevance: This work focusses on the optical properties of AlGaN-based multiple

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quantum well (MQW) structures grown either on sapphire or on native AlN substrates. It is demonstrated that the dislocation-mediated spiral growth in the MQW region of structures deposited on sapphire substrates generally leads to a more efficient localisation of carriers compared to the samples grown on native AlN substrates.

[P7] S. Gelfert, C. Frankerl, C. Reichl, D. Schuh, G. Salis, W. Wegscheider, D. Bougeard, T. Korn, and C. Sch¨uller,

“Magneto-Raman spectroscopy of spin-density excitations in (001)-grown GaAs- AlGaAs quantum wells in the regime of the persistent spin helix”,

Spintronics XI 10732, 118 - 123 (2018), DOI: 10.1117/12.2320205.

Relevance: Unrelated to this work.

[P8] S. Gelfert, C. Frankerl, C. Reichl, D. Schuh, G. Salis, W. Wegscheider, D. Bouge- ard, T. Korn, and C. Sch¨uller,

“Inelastic light scattering by intrasubband spin-density excitations in GaAs- AlGaAs quantum wells with balanced Bychkov-Rashba and Dresselhaus spin-orbit interaction: Quantitative determination of the spin-orbit field”,

Physical Review B 101, 035427 (2020), DOI: 10.1103/PhysRevB.101.035427.

Relevance: Unrelated to this work.

Conference Contributions

[C1] C. Frankerl, F. Nippert, M. P. Hoffmann, H. Wang, C. Brandl, H.-J. Lugauer, R.

“Carrier localization in AlGaN quantum wells”,

International Symposium on Semiconductor Light Emitting Devices (ISSLED), Magdeburg, Germany (2020) [postponed due to the COVID-19 pandemic].

[C2] N. Tillner, C. Brandl, C. Frankerl, H.-J. Lugauer, A. Waag, and M. P. Hoffmann,

“Epitaxial Growth of UV-C LEDs on nano-PSS”,

International Symposium on Semiconductor Light Emitting Devices (ISSLED), Magdeburg, Germany (2020) [postponed due to the COVID-19 pandemic].

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“Carrier redistribution and recombination dynamics in Al-rich AlGaN/AlN quantum well structures governed by carrier localization”,

International Workshop on Nitride Semiconductors (IWN), Berlin, Germany (2020) [postponed due to the COVID-19 pandemic].

[C4] G. Jacopin, C. Frankerl, N. Tillner, M. J. Davies, G. Rossbach, C. Brandl, M. P.

Hoffmann, R. Zeisel, A. Hoffmann, and H.-J. Lugauer,

“Internal quantum efficiency of AlGaN-based quantum well structures - The role of the growth substrate”,

International Workshop on Nitride Semiconductors (IWN), Berlin, Germany (2020) [postponed due to the COVID-19 pandemic].

[C5] C. Frankerl, M. P. Hoffmann, F. Nippert, H. Wang, C. Brandl, N. Tillner, H.-J.

“Optical internal quantum efficiency determination of UVC LEDs - towards a standardization of experimental conditions”,

International Conference on UV LED Technologies Applications (ICULTA), Berlin, Germany (2021).

“Improvement of the internal quantum efficiency by strong localization of carriers in AlGaN-based heterostructures”,

International Conference on UV LED Technologies Applications (ICULTA), Berlin, Germany (2021)

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AFM atomic force microscopy CCD charge-coupled device CL cathodoluminescence DAP donor-acceptor pair DH double heterostructure DUV deep ultraviolet EBL electron blocking layer EL electroluminescence

EQE external quantum efficiency FWHM full width at half maximum HAADF high-angle annular dark field HT/LT high temperature/low temperature HVPE hydride vapour phase epitaxy IQE internal quantum efficiency LED light-emitting diode

LEE light extraction efficiency MBE molecular beam epitaxy

MOCVD metal-organic vapour phase epitaxy

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MQW multiple quantum well ND neutral density

PL photoluminescence

PLE photoluminescence excitation PVT physical vapour transport QCSE quantum-confined Stark effect QW quantum well

RMS root mean square

SIMS secondary-ion mass spectroscopy SQW single quantum well

SRH Shockley-Read-Hall

STEM scanning transmission electron microscopy TCSPC time-correlated single photon counting TD threading dislocation

TDD threading dislocation density UV ultraviolet

XRD X-ray diffraction

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Abstract iii

Zusammenfassung v

List of Publications vii

List of Acronyms xi

1 Introduction 1

1.1 Applications of DUV Emitters . . . 3

1.2 Historical Development of Blue and DUV Emitters . . . 5

1.2.1 ZnSe versus GaN . . . 5

1.2.2 Development of (In)GaN . . . 6

1.2.3 Towards AlGaN-based DUV LEDs . . . 8

1.3 Thesis Outline . . . 10

2 Fundamentals of III-Nitride Heterostructures 11 2.1 Crystal Structure of III-Nitrides . . . 12

2.2 Ternary III-Nitride Alloys . . . 14

2.3 Polarisation Fields . . . 15

2.3.1 Spontaneous and Piezo-Electric Polarisation . . . 15

2.3.2 The Quantum-Confined Stark Effect . . . 16

2.4 Carrier Dynamics in AlGaN-based QW Heterostructures . . . 19

2.4.1 Non-Radiative Recombination Processes . . . 19

2.4.2 The ABC Model . . . 20

2.4.3 Determination of the IQE . . . 24

2.5 Carrier Localisation . . . 27

2.5.1 ‘S’-shape Temperature Dependence of the PL Peak Energy . . . 27

2.5.2 Carrier Decay Times and Decay Shape . . . 29

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CONTENTS

2.5.3 Potential Mechanisms Responsible for Carrier Localisation . . . 32

2.6 Further Loss Mechanisms in AlGaN-based LEDs . . . 36

2.7 Efficiency Droop . . . 38

2.7.1 Auger Effect . . . 38

2.7.2 Saturation of Localised States . . . 39

2.8 Design of Commercial AlGaN-based LEDs . . . 42

2.9 Summary . . . 45

3 Growth and Characterisation 47 3.1 MOCVD Growth of III-Nitride Heterostructures . . . 48

3.2 Experimental Methods and Setups . . . 49

3.2.1 Optical Spectroscopy . . . 49

3.2.2 Further Characterisation Techniques . . . 53

3.3 Calculation of the Effective Carrier Density . . . 56

3.3.1 Neglecting Carrier Decay . . . 56

3.3.2 Including Carrier Decay . . . 58

3.4 Investigated Sample Structures . . . 60

4 Design, Evaluation and Optimisation of AlGaN-based QWs 61 4.1 The Role of the Growth Morphology . . . 62

4.1.1 Time-Integrated and Time-Resolved PL Spectroscopy . . . 62

4.1.2 STEM measurements . . . 67

4.1.3 AFM measurements . . . 69

4.1.4 Conclusion . . . 70

4.2 Challenges for Reliable IQE Determination . . . 71

4.2.1 Publication P1 . . . 71

4.2.2 Summary . . . 79

4.3 Impact of the Threading Dislocation Density . . . 80

4.4 Optimisation of the QW Design . . . 83

4.4.1 Barrier Thickness Variation . . . 83

4.4.2 Cap Layer Thickness Variation . . . 84

4.5 Conclusion . . . 86

5 Carrier Localisation in AlGaN/AlN QW Heterostructures 87 5.1 Carrier Localisation in AlGaN/AlN SQWs . . . 88

5.1.2 Additional Measurements . . . 97

5.1.3 Summary . . . 99

5.2 Carrier Localisation in AlGaN/AlN MQWs . . . 101

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5.2.1 Sample Structure and Measurement Methods . . . 101

5.2.2 Time-Integrated PL measurements . . . 101

5.2.3 Time-Resolved PL Measurements . . . 104

5.2.4 Summary . . . 105

5.3 Carrier Recombination Dynamics Governed by Carrier Localisation . . 106

5.3.2 Additional Measurements . . . 114

5.3.3 Summary and Discussion . . . 115

5.4 Microscopic Origin of Carrier Localisation and Efficiency Droop . . . . 117

5.4.2 Summary and Discussion . . . 133

6 Absorption and Light Outcoupling Efficiency of DUV LEDs 139 6.1 Point Defect Incorporation in III-Nitrides . . . 141

6.1.1 ‘Blue’ and ‘Yellow’ Luminescence in Al(Ga)N Epitaxial Layers . 141 6.1.2 Absorption in Al(Ga)N Epitaxial Layers . . . 145

6.2 Defect Absorption in AlN Template Layers . . . 148

6.2.2 Summary . . . 155

7 Summary and Outlook 157 7.1 Thesis Summary . . . 158

7.2 Outlook . . . 162

List of Figures 165

List of Tables 168

Bibliography 169

Acknowledgements 231

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Introduction

For the fabrication of green, blue and ultraviolet (UV) light-emitting diodes (LEDs) and other optoelectronic devices, the III-nitride material system is most commonly employed, comprising gallium nitride (GaN), indium nitride (InN) and aluminium nitride (AlN). The emission energies (wavelengths) which can be achieved by alloying these constituents range from 0.7 eV (1800 nm) in the infrared (IR) for InN up to 6.2 eV (200 nm) in the deep ultraviolet (DUV) for AlN [1]. This is depicted in Figure 1.0.1, along with the lattice parameter a.

Figure 1.0.1: Room temperature band gap energies and lattice parameter afor AlN, InN, GaN and their composite ternary alloys. Band gap values of the binary alloys and bowing parameters, which are used in the calculation of the band gap energies and lattice parameters of the ternary alloys, are taken from refs. [1–3].

Binary and ternary compounds only containing Ga and Al as metal atoms are classified as wide band gap semiconductors, with band gaps varying from the near-UV (GaN,

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3.4 eV) into the DUV spectral range (AlN, 6.2 eV). Therefore, by alloying GaN and AlN, emission wavelengths from 200 – 350 nm can be realised, rendering DUV LEDs suitable for a large variety of potential applications. These are briefly discussed in the following section.

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1.1 Applications of DUV Emitters

UV radiation is commonly classified into three spectral bands: UV-A (400 – 320 nm), UV-B (320 – 280 nm) and UV-C (280 – 100 nm) [4]. This classification is motivated by the consequences of UV radiation for materials and organisms depending on the specific photon wavelength [5, 6]. This thesis focusses on the UV-C emission region, henceforth referred to as DUV.

The spectrum of the Sun’s solar radiation closely matches the theoretical emission spectrum of a black body with a temperature of approximately 6000 K [7]. For living organisms on earth, the DUV share of this spectrum is strongly attenuated by the atmosphere of the Earth [8]. Hence, evolution has not given rise to an efficient defence mechanism against highly-energetic photon exposure [4]. It has been demonstrated by microbiological studies that exposure to DUV radiation renders many pathological bac- teria, spores and viruses inactive [9–11]. Consequently, DUV emitters are particularly suitable for sterilisation [12, 13], disinfection [14, 15] and water purification purposes [14, 16, 17]. These applications become even more important in light of the 2020 Coronavirus pandemic, as recent ground-breaking studies report on the successful inactivation of SARS-CoV-2 using intense DUV radiation [18, 19]. Further applications of DUV devices include, for instance, gas sensing [20, 21], UV curing and printing [22, 23].

Figure 1.1.1: Comparison of a typical emission spectrum of an AlGaN-based DUV multiple quantum well (MQW), similar to those investigated in this thesis, with the emission of a standard low-pressure mercury vapour lamp and the germicidal efficacy curve of E. Coli. The last two curves are taken from ref. [24].

The efficiency of DUV-stimulated inactivation processes is a strongly varying function of the photon wavelength. This dependence can be described with the germicidal

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1.1. APPLICATIONS OF DUV EMITTERS

efficacy curve, which peaks around 265 nm for several pathogenic microorganisms such as E. coli [9], as illustrated in Figure 1.1.1. The figure also depicts the emission spectra of a typical DUV MQW as well as a standard low-pressure mercury vapour lamp [24].

For a long time, mercury-based devices have found widespread application for most UV-related applications. However, there are several serious drawbacks limiting the applications of mercury vapour lamps, such as their bulky and fragile design, long warm-up times, strong heat radiation and limited range of emission wavelengths [4].

Furthermore, the Minamata Convention, which entered into force in 2017, intends to gradually ban products containing mercury, further highlighting the need for an environmentally friendly alternative [25]. Hence, many researchers see the future in III-nitride-based DUV LED devices as they are considerably smaller, operate at lower voltages, have longer lifetimes, provide a tunable emission wavelength based on the material composition and do not contain any toxic materials [4].

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1.2 Historical Development of Blue and DUV Emitters

The very first reports on the generation of light by current (called electroluminescence (EL)) date back to as early as the 1950s, with SiC as the material of choice [26, 27].

Shortly after, the 1960s saw the realisation of the first LEDs, which were initially based on p-n-junctions in GaAs- [28] and GaP-based [29] homoepitaxial layers. However, for a long time, achieving efficient blue and green light emission posed an exceptionally challenging task. In the 1970s, the external quantum efficiency (EQE) of early devices based on ZnSe [30], ZnS [31] and SiC [32] was only around 0.1 %. At the same time, GaP-based red-emitting LEDs were reported to achieve EQEs of more than 10 %, a value two orders of magnitude higher [33].

1.2.1 ZnSe versus GaN

In the 1980s, two material systems were generally considered as possible candidates for the fabrication of efficient blue-emitting LEDs: ZnSe and GaN [34]. Most research groups believed that ZnSe was the most promising material as it could be grown with high structural quality on widely available single-crystal GaAs substrates due to the very small lattice mismatch between ZnSe and GaAs, which amounts to only 0.3 % [34, 35]. In contrast, GaN had to be grown on sapphire given that a suitable, lattice- matched substrate material was lacking [34]. Using this approach, GaN epitaxial layers could only be grown with extremely high threading dislocation densities (TDDs) [36], owing to the large lattice mismatch of ∼16 % between GaN and sapphire. A vivid example highlighting the popularity of ZnSe in comparison to GaN at this point in time is provided by Nobel Prize laureate Shuji Nakamura during his Nobel Lecture in 2015 [34]:

“At the Japan Society of Applied Physics conference in 1992, there were approximately 500 individuals attending the ZnSe sessions, whereas for GaN, there were around five, including the chair Professor Isamu Akasaki, speaker Hiroshi Amano and myself, as a member of the audience. Not only was ZnSe more popular at the time, GaN was actively discouraged with researchers stating ‘GaN has no future’ and ‘GaN people have to move to ZnSe material’.”

(Shuji Nakamura, Nobel Lecture 2015) During that time, ZnSe could be grown on GaAs substrates with TDDs less than 10³cm⁻², while research efforts on GaN were stuck at TDDs of the order of 10⁹cm⁻². In addition, the 1990s saw several fundamental breakthroughs in heterostructure design

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1.2. HISTORICAL DEVELOPMENT OF BLUE AND DUV EMITTERS

and crystal growth of the ZnSe material system, leading many researchers to the belief that widespread commercial implementation was imminent [37–39]. However, all these hopes were eventually shattered as, despite immense research efforts, severe issues regarding material degradation during standard device operation could never be solved [40, 41]. Instead, to the surprise of the scientific community, it was found that efficient LED operation could, in fact, be achieved in GaN-based devices despite the high amount of threading dislocations (TDs) incorporated in the deposition process [42]. To this day, the failure to develop highly efficient ZnSe-based LED devices is one of the most unexpected disappointments in the history of solid state physics.

1.2.2 Development of (In)GaN

Johnson et al. were the first to report on the synthesis of GaN as early as 1932, employing a direct reaction between hot metallic Ga and NH₃ gas [43]. The first synthesisation of crystalline GaN, however, was only achieved almost four decades later in 1969 by Maruska and Tietjen using hydride vapour phase epitaxy (HVPE) [44]. These early GaN layers were found to be unintentionally n-doped, possibly due to nitrogen vacancies and interstitial defects by Ga atoms. In addition, the wafers suffered from large cracks and pits [44]. In 1972, the first GaN-based light-emitting device, a simple i-n structure, was demonstrated by Pankove et al., producing a blue luminescence at room temperature [45]. In this design, an i-GaN layer was fabricated by doping with Zn, compensating the electron population of the inherently n-type GaN.

Later, it was shown by Maruska et al. that violet luminescence can also be achieved when Zn atoms are substituted with Mg atoms [46]. At this point, progress in device and material development slowed down for several years due to insufficient material quality as a result of inherent limitations in the HVPE growth process.

Research efforts started to make progress only with the introduction of more sophisti- cated growth techniques, primarily molecular beam epitaxy (MBE) and metal-organic vapour phase epitaxy (MOCVD) technology. As a consequence, two significant steps were made towards the realisation of effective III-nitride emitters in the early 1980s.

Firstly, the crystalline quality could be improved significantly and secondly, an effective p-doping technique for GaN epitaxial layers was found.

In 1983, Yoshidaet al. were the first to report on the MBE growth of crack-free GaN layers with improved electrical and optical properties, achieved by inserting a thin AlN underlayer between the sapphire substrate and the GaN film [47]. The increased material quality was a result of the reduction in lattice mismatch and the closer match of the thermal expansion coefficients of GaN and AlN (in contrast to GaN and sapphire).

Later, Amano et al. pursued the same approach but employed the MOCVD growth technique, which enabled even more significant improvements in crystalline quality in

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comparison to MBE [48]. Independently, Nakamura reported on similar GaN crystal quality by replacing the AlN buffer layer with a low temperature GaN buffer layer [49].

The high density of unintentionally incorporated n-type dopants during growth ren- dered the p-type doping of GaN epitaxial layers extraordinarily difficult, with early p-type layers revealing very high resistivities of ∼10⁶Ω cm, basically equivalent to an insulating material [50]. It was suggested by Van Vechtenet al. that acceptor-hydrogen complexes are formed in the crystal, leading to self-compensation and explaining the low density of available acceptors in p-type GaN [51]. Thus, a method had to be developed to dissociate these complexes, in turn increasing the density of active acceptors.

In 1988, Amano et al. had already reported on a persistent increase in the luminescence intensity of Zn-doped GaN epitaxial layers after cathodoluminescence (CL) measurements, demonstrating a first possibility to decrease the resistivity [52]. A similar observation was made by Nakamura et al. in a Mg-doped GaN layer after thermal annealing, also significantly lowering its resultant resistivity [53]. This proved that thermal annealing was a feasible approach to dissociate the acceptor-hydrogen complexes and reduce the undesired self-compensation effect. Following these advances, the first LED emitting in the UV spectral region was demonstrated by Amano et al.

in 1989, based on a classical pn-junction [54].

The last critical step towards achieving a commercialisable, i.e. sufficiently efficient LED was the development of a double heterostructure (DH). InGaN was quickly identified as an ideal candidate material for the active region [34]. In atoms implemented into the GaN crystal can be utilised as means to confine the charge carriers and varying their concentration in the alloy enables the colour of the emitted photons to be delib- erately tuned. Given the highly defective nature of the material, high quality layers of InGaN could not be realised until the 1990s [34]. Moreover, the fabrication of a DH LED required precise and immediate control over the various growth parameters to en- sure sufficient interface quality between GaN and InGaN. In addition, the introduction of In into the GaN lattice leads to significant strain since In atoms are ∼20 % bigger in size than Ga atoms, resulting in the incorporation of a large number of defects [34].

Early reports on InGaN alloy growth were provided by Osamura et al. in 1975 [55], Nagatomo et al. in 1989 [56] and Yoshimoto et al. in 1991 [57], with all groups struggling to achieve good crystal quality. The first high-quality InGaN layer was grown by Nakamura and Mukai in 1992, employing a self-built two-flow MOCVD reactor [58].

Nakamura et al. then continued to work on the inclusion of these layers into an LED structure. Their efforts paid off and in 1993, Nakamura et al. demonstrated the first blue-emitting DH LED [59]. Soon after, they further succeeded in growing the first high- brightness blue LED in 1994, employing a Zn-doped InGaN/AlGaN DH [60]. Later on and until the present day, device development has shifted towards utilising a stack of quantum wells (QWs) to form the active region, with Nakamuraet al. once more being

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1.2. HISTORICAL DEVELOPMENT OF BLUE AND DUV EMITTERS

the first group to report on the successful development of high-brightness InGaN single quantum well (SQW) LEDs emitting in the blue and violet spectral range, exhibiting much narrower emission linewidths compared to devices comprising DHs [61].

1.2.3 Towards AlGaN-based DUV LEDs

The late 1990s and early 2000s saw the success of research efforts on AlGaN-based UV-emitting devices with wavelengths below 360 nm [62–64]. However, it took until 2004 before the first sub-300 nm device was successfully fabricated, an LED emitting at a wavelength of 280 nm, developed by Sunet al. [65]. The efficiencies of AlGaN and AlGaInN QW heterostructures could be substantially increased by utilising a pulse-flow growth method, which allowed for the growth of a high-quality AlN buffer layer [66–

68]. The first commercially available LEDs emitting in the DUV spectral region were fabricated by Sensor Electronic Technology, with a maximum EQE of 11 % achieved for a 278 nm-emitting LED in 2012 [69, 70].

Figure 1.2.1: Overview of the reported EQEs of UV-emitting LEDs developed in the last decades. The figure is reproduced from ref. [4] and the datapoints are taken from refs. [70–84].

An overview of state-of-the-art UV LED devices is provided by Kneissl et al. [4]

and discussed in the following. In Figure 1.2.1 are shown the EQEs of UV LEDs reported throughout almost two decades of development. The strong dependence of the performance of the LEDs on their respective emission wavelengths is striking. A substantial decrease in EQE is observed for emission wavelengths shorter than 360 nm, which indicates the transition from (well-established) InGaN- to (comparatively im- mature) AlGaN-based LED technology. Near-UV emitters are based on low-In InGaN heterostructures and as such employ the same materials and technologies as blue LEDs,

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thus being ahead of their AlGaN-based counterparts by many years of development.

As a consequence, near-UV LEDs can be fabricated with efficiencies close to those of blue LEDs, achieving EQEs mostly above 50 %. On the contrary, AlGaN-based DUV LEDs exhibit EQEs primarily below 10 % - with the exception of recently achieved peak EQEs of 20 % near 275 nm - and are observed to quench rapidly for wavelengths below 250 nm, corresponding to Al-rich active regions.

It has become evident that DUV LED technology still offers a great potential for improvement, without being able to identify a single cause for the current limitation of the achievable EQEs. For example, in case of blue LEDs, important key performance parameters such as injection, radiative recombination and light extraction efficiencies are all in the 90 % range [4]. All of these values are generally much lower for DUV LEDs [4, 85–87]. As a consequence, to this day, the full potential of AlGaN-based devices has not yet been fully exploited. The understanding of the fundamental physics of this material system is still very limited, impeding a more commercialised application of the devices fabricated from it. Many of the remaining technological hurdles are elaborated in the course of this thesis.

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1.3. THESIS OUTLINE

1.3 Thesis Outline

In Chapter 2 of this thesis the fundamental properties of III-nitride QW heterostructures and epitaxial layers are discussed. This includes the crystal structure and the formation of ternary alloys, spontaneous and piezo-electric polarisation fields, as well as the quantum-confined Stark effect (QCSE). Subsequently, recombination dynamics in AlGaN-based QWs are described and key parameters and models such as the internal quantum efficiency (IQE), the ABC model, the efficiency droop and the localisation of carriers are introduced. Finally, the design of a commonly employed AlGaN-based DUV LEDs structure is presented.

In Chapter 3 the experimental techniques employed in this thesis are described. This includes the MOCVD growth of III-nitride heterostructures, time-integrated and time- resolved photoluminescence (PL) spectroscopy, photoluminescence excitation (PLE) and absorption spectroscopy, as well as additional material characterisation methods (atomic force microscopy (AFM), secondary-ion mass spectroscopy (SIMS), scanning transmission electron microscopy (STEM) and X-ray diffraction (XRD)). Finally, an estimation of the optically induced effective carrier density in the grown QW structures and an overview of the sample structures investigated in this thesis is provided.

Chapter 4 focusses on the determination and the improvement of the IQE in AlGaN- based QW heterostructures. The appropriate choice of excitation conditions and extraction methodology is discussed and detrimental effects such as carrier transport and smeared growth morphology are identified. After demonstrating a reliable methodology to determine the IQE, a variety of experiments is carried out which aims to improve the quality of the active region of a DUV LED, including a brief study on the impact of the TDD and the QW design.

Chapter 5 presents an in-depth investigation of the role of carrier localisation in AlGaN/AlN QWs. The impact of the localisation and delocalisation of carriers on the recombination dynamics in a large set of samples is discussed, including AlGaN/AlN SQWs and MQWs of varying QW width and Al content. Furthermore, the microscopic origins of carrier localisation are explored and a novel model describing the efficiency droop in III-nitride heterostructures is introduced, based on the obtained experimental evidence.

In Chapter 6 the incorporation of point defects in AlN and AlGaN epitaxial layers is studied and several detrimental absorption bands are identified. Concepts to reduce absorption and improve the light outcoupling efficiency of DUV LEDs are discussed and appropriate growth parameters are identified.

Chapter 7 summarises the thesis and presents an outlook on further experiments.

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Fundamentals of III-Nitride Heterostructures

This chapter discusses the fundamental properties of semiconductor materials based on AlGaN/Al(Ga)N heterostructures, including epitaxial layers and quantum well (QW) structures, with the focus being on the latter. The crystal structure of III-nitride alloys is described and the inherent polarisation fields, giving rise to the quantum- confined Stark effect (QCSE), are discussed. Basic models describing the carrier recombination dynamics in AlGaN/Al(Ga)N QWs are introduced and central concepts such as the internal quantum efficiency (IQE), the ABC model, carrier localisation and the efficiency droop are presented. Finally, the design of a commercial light-emitting diode (LED) emitting in the deep ultraviolet (DUV) spectrum is illustrated, with remaining challenges outlined.

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2.1. CRYSTAL STRUCTURE OF III-NITRIDES

2.1 Crystal Structure of III-Nitrides

Depending on the symmetry of the growth substrate and the employed growth conditions, III-nitride semiconductors crystallise either in wurtzite or zinc-blende crystal structure [88–90]. The wurtzite crystal structure is the thermodynamically most stable configuration and all samples investigated in this work are grown in the wurtzite phase in{0001}-orientation. In this configuration, N atoms are arranged in hexagonal units, with alternating layers of group-III elements and further layers of N atoms stacked on top. The wurtzite unit cell is defined by two lattice constants, labelled a and c, whereby a represents the spacing between adjacent atoms in the hexagonal unit cell and cdenotes the separation of the atomic planes. III-nitride semiconductor growth is primarily performed on foreign substrates. Among the most popular substrate materials are sapphire (Al2O2), Si and SiC. For reasons of availability and cost, the most commonly used substrate for wurtzite AlN growth - serving as the basis for all structures grown during this thesis - is sapphire, in which the Al atoms of the nitride alloy align themselves with the O atoms of the upper sapphire interface.

Figure 2.1.1: Schematic diagram depicting the relative positions of the lattice sites of AlN grown on sapphire.

Similar to GaN growth, the lattice of AlN grown on sapphire is rotated by 30° with respect to the sapphire lattice, resulting in a reduced lattice mismatch of ∼13 % [91].

This is sketched in Figure 2.1.1. However, the lattice mismatch of the rotated AlN lattice is still significant. In addition, Yimet al. determined a∼37 % difference in thermal expansion coefficients of sapphire and AlN, altogether resulting in severe compressive strain in the epitaxial growth process [92]. This promotes the formation of a high density of crystalline defects, including atomic misfit dislocations (vacancies and interstitial atoms) and extended defects (threading dislocations (TDs)). Despite intense efforts to reduce the threading dislocation density (TDD) of Al(Ga)N epitaxial layers grown on sapphire substrates, this value remains generally high in the order of 10⁹cm⁻²

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[93, 94]. The lowest values reported by metal-organic vapour phase epitaxy (MOCVD) growth are of the order of 10⁸cm⁻², which can be achieved by means of high temperature annealing and sputtering [95–97]. This is in stark contrast to lattice-matched III-V semiconductors, such as GaAs, where TDDs have been routinely achieved below 10³cm⁻² for decades [98, 99].

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2.2. TERNARY III-NITRIDE ALLOYS

2.2 Ternary III-Nitride Alloys

Room temperature lattice parameters of the binary III-nitride alloys are summarised in Table 2.1. The variation in the lattice parameters a andc is a result of the different degrees of strain in the growth process, as previously discussed in Chapter 2.1.

Parameter InN GaN AlN

Lattice parametera [A] 3.545 3.189 3.112 Lattice parameterc [A] 5.703 5.185 4.982 Electron effective massm^ke [m₀] 0.07 0.20 0.32 Electron effective massm^⊥_e [m₀] 0.07 0.20 0.30 Hole effective massm^k_h [m₀] 1.67 1.96 3.53 Hole effective massm^⊥_h [m₀] 1.61 1.87 11.14

Table 2.1: Room temperature (300 K) lattice parameters of InN, GaN and AlN. Here, m^ke indicates the effective mass within the growth plane, while m^⊥_e denotes the effective mass perpendicular to the growth plane. The values are taken from refs. [1, 100, 101].

LEDs fabricated to emit light in the DUV emission spectrum employ the ternary alloy AlxGa1−xN both in the active region and as host material forming the n- and p-side of the device. The lattice parameters for such a ternary alloy are commonly determined by using Vegard’s law, which consists of a linear interpolation of the binary constituents’ lattice parameters [102]:

a(Al_xGa1−xN) =x·a_AlN+ (1−x)·a_GaN, (2.1) wherexdenotes the fraction of Al atoms contained in the AlGaN alloy anda_AlN and a_GaN are the respective lattice parameters which can be found in Table 2.1. The band gap energy of ternary III-nitride alloys, on the other hand, cannot be determined by simple linear interpolation. Instead, a quadratic bowing parameterb must be included.

Hence, the band gap energy of a Al_xGa1−xN alloy is calculated as:

E_G^AlGaN =x·E_G^Al+ (1−x)·E_G^Ga−x·(1−x)·b (2.2) The bowing parameter of Al_xGa1−xN was experimentally determined to around 0.6 – 0.9 eV [103–105]. No clear consensus has been reached so far to narrow this range down.

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2.3 Polarisation Fields

Epitaxial layers grown on c-plane sapphire in the wurtzite crystal phase exhibit strong macroscopic polarisation fields as a result of their inherent polarity. These fields are oriented along the {0001}-axis and can be divided into spontaneous and (strain-induced) piezo-electric polarisation fields [106, 107], significantly affecting the recombination dynamics in the QWs as they provoke the QCSE.

2.3.1 Spontaneous and Piezo-Electric Polarisation

.

Figure 2.3.1: Schematic orientation of the atomic bonds in an AlN crystal lattice. An Al atom (blue sphere) is surrounded by a tetrahedron of N atoms (red spheres). A and B denote the atomic bonds andθand φindicate the angles between them, respectively (see text).

In the wurtzite crystal orientation, group-III atoms, represented by an Al atom in Figure 2.3.1, are surrounded by a tetrahedron of N atoms. A certain net polarisation is associated with each atomic bond as a result of the large electronegativity of N compared to Al. The bond between the group-III atom and the N atom at the apex of the tetrahedron is labelled A, while each individual bond between the bottom three N atoms and the group-III atoms are labelled B. θ denotes the angle betweenA and B, while φ indicates the angle between twoB, respectively, with their values compiled by Schulz and Thiemann [108].

When an AlN crystal is grown in wurtzite crystal geometry, the weighted charge distribution centres of subsequently grown layers do not overlap. As a consequence, even in a fully relaxed material without strain, a residual net polarisation is present, called spontaneous polarisation. In a ternary alloy such as AlGaN, the resulting spontaneous

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2.3. POLARISATION FIELDS

polarisation is determined by the spontaneous polarisations of its binary constituents together with a quadratic bowing parameter. This is called the ‘bowed Vergard rule’

[106, 109]:

P_SPÂB =x·P_SPÂ + (1−x)·P_SP^B −x·(1−x)·bÂB_SP (2.3) Here,xrepresents the composition of the constituentA,P_SPÂ (P_SP^B ) is the spontaneous polarisation of the constituent A(B) and bÂB_SP is the bowing parameter of the resulting ternary alloy. P_SP amounts to −0.034 C m⁻² and −0.090 C m⁻² for GaN and AlN [1], while bÂB_SP was determined as 0.019 C m⁻² for Al_xGa1−xN [106, 107].

Even more importantly, as the lattice parameters of the binary III-nitride deviate significantly, strain is introduced within epitaxial layers and QW heterostructures depending on alloy composition. Of particular interest for this thesis are Al_xGa_1−xN/AlN QW heterostructures grown on c-plane sapphire. An increase of the Ga content in a ternary Al_xGa1−xN alloy results in an increase of the a lattice parameter. Hence, the Al_xGa1−xN layer is compressively strained when grown pseudomorphically on an AlN layer or an AlGaN layer with higher Al content. In turn, the angles θ and φ depicted in Figure 2.3.1 are modified, resulting in a piezo-electric polarisation PP Z. In the important case of biaxial strain in a QW heterostructure,P_{P Z} is oriented along thez-axis and its strength can be calculated using the following equation [110]:

P_{P Z} = 2

a_B−a_QW a_QW

×

e₃₁−e₃₃C₁₃ C₃₃

(2.4) Here,a_B (a_QW) represent the lattice parameters of the barrier (QW) material,e₃₁and e₃₃ denote the piezo-electric matrix elements and C₁₃ and C₃₃ are the matrix elements of the elasticity. These values can be found e.g. in ref. [109].

2.3.2 The Quantum-Confined Stark Effect

In the preceding section, it was determined that a non-zero net electric field across a QW heterostructure grown in polar crystal orientation (e.g. AlGaN deposited on c- plane sapphire) exists. In the active region of a typical DUV LED, AlGaN-based QWs with well widths around 1 – 3 nm are grown, separated by slightly thicker barriers with higher Al content. For optical experiments, binary AlN is also a popular barrier material. The large barrier band gap and the thin QW layer provide a means to confine the carriers into a small volume in the active region, while the emission energy may also be tuned by a variation of the alloy composition of the confinement layer. However, the strong spontaneous and piezo-electric polarisation fields result in large electric fields

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of up to 2 – 3 MV cm⁻¹ [111, 112]. These fields are found to increase rapidly with barrier thickness and Al content [112], significantly tilting the potential profile of the QW and barrier region compared to non-polar QW growth without net polarisation, as schematically depicted in Figure 2.3.2.

CB

AlGaN AlN

AlN

VB

𝑬_𝟏

AlGaN AlN AlN

𝑬_𝟐

non-polar growth direction polar growth direction

Figure 2.3.2: Schematic diagram of the QCSE. Left: An AlN/AlGaN/AlN QW heterostructure grown in non-polar crystal orientation without net polarisation fields.

Right: An equivalent QW system grown in polar direction.

The tilting of the band profile reduces the ground state recombination energy and leads to a spatial separation of the electron and hole wave functions. This effect is commonly referred to as the QCSE [113, 114], illustrated in Figure 2.3.2. The electric fields in the QWs push the electron and hole populations towards opposing interfaces of the QWs, reducing the electron-hole wave function overlap as well as the oscillator strength of the ground-state transition. This decreases the radiative recombination probability of the carriers as well as their recombination energy. In general, the QCSE is enhanced with increasing QW width [115] due to the increasing spatial separation of the electron and hole wave functions. Given its severe impact on the carrier recombination dynamics, the QCSE plays a decisive role in the optical emission properties of polar semiconducting structures.

At very high carrier densities, the dipole moments of the spatially separated electrons and holes contribute significantly to the net electric field across the QW, however with opposing polarity to that of the polarisation field produced by piezo-electric and spontaneous polarisation. Hence, high carrier densities effectively reduce the net polarisation strength inside a QW, partially restoring the nominal band gap and recombination efficiencies of the ground state transition for a structure without polarisation fields.

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2.3. POLARISATION FIELDS

Experimentally, the carrier density-induced screening of the electric field across the QW can be observed as a blueshift of the recombination energy and an increasing efficiency of the radiative recombination due to the increasing wave function overlap, altogether suggesting a reduction of the QCSE strength [116–119].

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2.4 Carrier Dynamics in AlGaN-based QW Heterostruc- tures

Light is produced inside a QW heterostructure by the radiative recombination of a conduction band electron with a valence band hole, releasing energy in form of a photon.

The photon wavelength corresponds to the energetic separation of the original electron- hole pair [120]. This process has been theoretically described already in 1927 through Fermi’s Golden Rule [121] and is the desired recombination channel of any light-emitting device. No momentum conversation is necessary in case of a direct semiconductor such as AlGaN, meaning that the process is not phonon-assisted [122]. In an ideal active region, each electron-hole pair is converted into a photon, and there are no losses through non-radiative recombination channels. However, any real QW heterostructure is far from ideal, with a non-negligible fraction of carriers decaying via a non-radiative decay channel, reducing the light emission efficiency.

This chapter aims to introduce the most important non-radiative recombination processes as well as the decisive structural parameter labelled IQE and several methodologies for the determination thereof. Furthermore, two intrinsic effects significantly affecting the carrier recombination dynamics are discussed, namely the efficiency droop and carrier localisation. A significant fraction of the experimental work in this thesis is spent investigating the microscopic origin of these effects, hence they are explained in greater detail.

2.4.1 Non-Radiative Recombination Processes

Figure 2.4.1 provides an overview of the recombination channels in AlGaN-based QW heterostructures. Bimolecular electron-hole recombination constitutes the main radiative recombination channel, producing a photon with a wavelength corresponding to the effective band gap of the alloy, modified by quantum confinement effects and the QCSE [123]. Bimolecular recombination is a two-particle process, thus its recombination rate can be expressed as Rbimolecular = Bn², where B is the bimolecular recombination coefficient.

SRH recombination involves a carrier, which can be either an electron or a hole, get- ting captured by a trap state in the forbidden zone of the band gap before recombining with the opposing type of carrier [124, 125]. These traps are mainly formed by defects in the crystal lattice, such as (native or foreign) point defects or TDs. SRH recombination can either occur non-radiatively, emitting a phonon, or radiatively, emitting a photon of a wavelength much higher than the band gap, and hence not contributing to the (desired) band gap emission. The recombination rate of SRH processes is hence a function of the density of trap states and their carrier capture cross section. Carriers

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2.4. CARRIER DYNAMICS IN ALGAN-BASED QW HETEROSTRUCTURES

Shockley-Read-Hall Bimolecular Auger

𝐸

𝑘

𝐸

𝑘

𝐸

𝑘

photon defect

defect luminescence or phonon

Figure 2.4.1: Schematic overview of the main recombination channels in AlGaN- based QWs: Shockley-Read-Hall (SRH), bimolecular, and Auger recombination.

trapped at deep defect states are strongly localised, resulting in comparatively long carrier lifetimes [126]. With only one particle involved, the SRH recombination rate is a linear function of the charge carrier density n, i.e. RSRH =An, where A is the SRH recombination coefficient.

At sufficiently high carrier densities, an additional three-particle process gains importance. The Auger effect is a non-radiative process, where the released energy of an electron-hole recombination event is transferred to a third carrier (either an electron or a hole), exciting it into a higher energy state. This is exemplarily illustrated for an electron as the third, excited particle in Figure 2.4.1. While first studied in individual atoms [127], the Auger effect has been experimentally observed and described theoretically in a variety of solid state material systems [128–131]. As Auger recombination is a three-particle process, the Auger recombination rate can be expressed asR_Auger =Cn³, where C is the Auger recombination coefficient.

2.4.2 The ABC Model

In order to quantitatively study the radiative emission efficiency of AlGaN-based heterostructures, an empirical rate equation model originally proposed by Shen et al. is employed [132]. Steady state conditions for the rate of change in carrier density with time are assumed. Defining n as the carrier density, G as the carrier generation rate and R as the recombination rate yields

dn

dt =G−R = 0. (2.5)

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In thermal equilibrium, the probability of bimolecular radiative recombination is directly proportional to the density of electrons (ne) and holes (pe), i.e. gt=R =Bnepe. In an intrinsic semiconductor, n_e =p_e =n_i holds, with n_i denoting the intrinsic carrier density, thus simplifying the recombination rate to R =Bn²_i. After optical excitation, the system reverts to the equilibrium state and ^dn_dt 6= 0.

This thesis focusses on purely optical experiments. Hereby, an excess carrier density is induced by a laser source directly into the QW region without the need for electrical contacts. To account for such conditions, the equilibrium carrier densitiesneandpe are expanded to n(t) =n_e+ ∆n(t) and p(t) =p_e+ ∆p(t), respectively, as the laser source induces an excess population of electrons ∆n(t) and holes ∆p(t). In case of resonant, optical excitation, ∆n(t) = ∆p(t) holds without further assumptions. In general, the optically induced carrier density significantly exceeds the thermal equilibrium density, i.e. n_e+ ∆n(t)≈∆n(t) holds in good approximation. Finally, non-radiative SRH and Auger recombination are taken into consideration. Eventually, Equation (2.5) reads

−d∆n(t)

dt ≈A∆n(t) +B∆n(t)²+C∆n(t)³, (2.6) where A, B and C represent the previously introduced non-radiative SRH, bimolecular and Auger recombination coefficients, respectively.

Equation (2.6) is employed to experimentally determine the IQE, which is defined as the ratio of radiatively recombining charge carriers to the full population of carriers present in the active region, thus quantifying the radiative emission efficiency of a QW heterostructure. The IQE is a key system parameter of crucial importance as it directly affects the achievable external quantum efficiency (EQE) of an LED via the relation EQE = η_injη_extIQE, whereη_inj andη_extdenote the carrier injection efficiency and light extraction efficiency, respectively. In the majority of the optical experiments presented in this thesis, η_inj is unity, as carriers are excited directly into the QWs, while η_ext is generally a function of the emission photon wavelength and strongly dependent on the crystal quality of the outcoupling AlN cap layer.

The IQE is calculated as the ratio of the recombination processes yielding a photon (N_rad) to the total number of recombination processes occurring in the active region, including non-radiative recombination channels (N_rad+N_nonrad):

IQE(n) = N_rad

N_rad+N_nonrad = Bn²

An+Bn²+Cn³, (2.7) whereBn² represents the radiative recombination rate of carriers within the QW and An+Bn²+Cn³ denotes the total carrier recombination rate via SRH, radiative and Auger recombination processes. Experiments determining absolute values for theA, B

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2.4. CARRIER DYNAMICS IN ALGAN-BASED QW HETEROSTRUCTURES

andC recombination coefficients have been performed primarily in the InGaN material system. Typically extracted room temperature values are of the order of 1×10⁶ – 5×10⁷s⁻¹(A), 1×10⁻¹²– 2×10⁻¹¹cm³s⁻¹ (B) and 3×10⁻³¹– 2×10⁻³⁰cm⁶s⁻¹(C) [132–137]. While the room temperature efficiency of AlGaN-based QW heterostructures remained very low until recently, in 2018 Nippertet al. were the first to experimentally determine these parameters by means of time-resolved optical spectroscopy in AlGaN- based QWs as 1.1×10⁷s⁻¹ (A), 2.4×10⁻¹¹cm³s⁻¹ (B) and 3.4×10⁻³⁰cm⁶s⁻¹ (C) [138], revealing comparable values.

Interestingly, in both cases - but especially in case of AlGaN - the C parameter is significantly larger than expected for direct semiconductors with (ultra-)wide band gaps, which is theoretically predicted to decrease exponentially with band gap energy [131]. Hence, for constant temperature, according to Equation (2.7) the behaviour of the IQE with varying excitation power density is as follows: At low carrier densities, non-radiative SRH recombination is the dominating decay channel, resulting in a very low light emission efficiency. With increasing carrier density, SRH recombination centres gradually saturate and the IQE increases until reaching its maximum value IQE_max at n_max = p

A/C. At even higher carrier densities, Auger recombination dominates, decreasing the IQE again.

Of particular importance is the strong temperature dependence of the recombination coefficientsA,B andC, which is known from both experimental studies and theoretical considerations [137, 139–145]. As a result, the IQE is both a function of temperature and carrier density:

IQE(n, T) = B(T)n²

A(T)n+B(T)n²+C(T)n³ (2.8) It is commonly assumed that the temperature dependences of the radiative and Auger recombination coefficients are of the formsB(T) = ^B_T⁰ andC(T) = C₀exp−_k^E^α

BT, where B₀ and C₀ are constant factors and E_α is the activation energy related to Auger recombination [144, 145]. The precise form of the temperature dependence of the SRH recombination coefficient is disputed, but an activation process is generally assumed, as A(T) strongly quenches with decreasing temperature [146]. As a result, at sufficiently low temperatures, the initial increase of the IQE with carrier density is not observed in high-quality sample structures. Instead, the emission efficiency is found to be approximately constant before finally decreasing at high carrier densities. Furthermore, as the recombination rates follow different functions of temperature, the position of the photoluminescence (PL) efficiency maximum shifts with temperature [138].

Given the IQE is a strongly varying function of both carrier density and temperature, it is important to clarify what exactly is meant when stating the IQE value of a specific

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semiconductor structure. In the remainder of this thesis, whenever an IQE value is stated without further elaboration, IQE_maxcorresponding to the carrier densitynmax = pA/C at room temperature is provided.

Limitations of the ABC Model

Although the ABC model is widely used to determine the IQEs of both InGaN- and AlGaN-based heterostructures, its derivation is based on several assumptions, not all of which are inherently justified:

(1) The ABC model does not consider any non-radiative effects other than SRH and Auger, therefore neglecting losses due to poor hole injection, carrier leakage and carrier spillover (c.f. Chapter 2.6). In this thesis, as the majority of experiments is conducted employing purely optical excitation, the first two effects can be ruled out, while the usage of AlN or high-Al content AlGaN barriers in all samples and the absence of external electric fields renders carrier spill-over equally unlikely. In addition, no typical fingerprint of carrier spill-over, such as an additional luminescence of a photon energy corresponding to the band gap of the barrier, could be observed.

(2) The carrier concentrations of both electrons and holes are equal. This is very difficult to achieve in an electrically-driven LED device due to the difficulty in p-doping of (Al)GaN as a consequence of the high Mg acceptor ionisation energy (c.f. Chapter 2.8) and the known issues regarding hole injection (c.f. Chapter 2.6).

However, in purely optical experiments employing resonant excitation conditions, each absorbed photon generates exactly one electron-hole pair in the active region, fulfilling the assumption in this case.

(3) The ABC model neglects possible excitonic effects. Excitons are believed to significantly impact the recombination properties of non-polar Al(Ga)N structures [147–150]. However, as the polarisation fields in polar Al-rich AlGaN/Al(Ga)N QW heterostructures strongly exceed the typically reported exciton binding energies [112, 151–153], these effects are considered unlikely to strongly influence the recombination dynamics of AlGaN/AlN QWs. A more detailed justification can be found in Chapter 5.3.

(4) The recombination coefficients (more specifically, their respective ratios) do not depend on the carrier density. In general, this assumption does not hold as the carrier density - e.g. by screening of the QCSE - modifies the charge distribution and hence wave function overlap, alteringB and C but not necessarily A. This is a serious disadvantage of the oversimplified ABC model.