Wide Band Gap Semiconductor Nanowires 2 E-Book

• Herausgeber: John Wiley & Sons
• Kategorie: Wissenschaft und neue Technologien
• Sprache: Englisch

Seitenzahl: 473

Veröffentlichungsjahr: 2014

Beschreibung

This book, the second of two volumes, describes heterostructures and optoelectronic devices made from GaN and ZnO nanowires. Over the last decade, the number of publications on GaN and ZnO nanowires has grown exponentially, in particular for their potential optical applications in LEDs, lasers, UV detectors or solar cells. So far, such applications are still in their infancy, which we analyze as being mostly due to a lack of understanding and control of the growth of nanowires and related heterostructures. Furthermore, dealing with two different but related semiconductors such as ZnO and GaN, but also with different chemical and physical synthesis methods, will bring valuable comparisons in order to gain a general approach for the growth of wide band gap nanowires applied to optical devices.

Leseprobe

Contents

Preface

PART 1: GaN and ZnO Nanowire Heterostructures

1 AlGaN/GaN Nanowire Heterostructures

1.1. A model system for AlGaN/GaN heterostructures

1.2. Axial AlGaN/GaN nanowire heterostructures

1.3. AlGaN/GaN core–shell nanowire heterostructures

1.4. Application examples

1.5. Conclusions

1.6. Bibliography

2 InGaN Nanowire Heterostructures

2.1. Introduction

2.2. Self-assembled InGaN nanowires

2.3. X-ray characterization of InGaN nanowires

2.4. InGaN nanodisks and nanoislands in GaN nanowires

2.5. Selective area growth (SAG) of InGaN nanowires

2.6. Conclusion

2.7. Bibliography

3 ZnO-Based Nanowire Heterostructures

3.1. Introduction

3.2. Designing ZnO-based nanowire heterostructures

3.3. Growth of ZnxMg1-xO/ZnO core–shell heterostructures by MOVPE

3.4. Misfit relaxation processes in ZnxMg1-xO/ZnO core–shell structures

3.5. Optical efficiency of core–shell oxide-based nanowire heterostructures

3.6. Axial nanowire heterostructures

3.7. Conclusions and perspectives

3.8. Bibliography

4: ZnO and GaN Nanowire-based Type II Heterostructures

4.1. Semiconductor heterostructures

4.2. Type II heterostructures

4.3. Optimal device architecture

4.4. Electronic structure of type II core–shell nanowires

4.5. Synthesis of the type II core–shell nanowires and their signatures

4.6. Demonstration of type II effects in ZnO–ZnSe core–shell nanowires and photovoltaic devices

4.7. Summary

4.8. Acknowledgments

4.9. Bibliography

PART 2: Integration of GaN and ZnO Nanowires in Optoelectronic Devices

5 Axial GaN Nanowire-based LEDs

5.1. Introduction

5.2. Top-down GaN-based axial nanowire LEDs

5.3. Bottom-up GaN-based axial nanowire LEDs

5.4. Carrier loss processes of axial nanowire LEDs

5.5. Controlling carrier loss of GaN-based nanowire LEDs

5.6. Conclusions

5.7. Bibliography

6 Radial GaN Nanowire-based LEDs

6.1. Radial GaN nanowire-based LED: an emerging device

6.2. Growth of GaN nanowires and radial nanowire-based devices

6.3. Radial GaN nanowire-based LED structure

6.4. Characteristics of radial nanowire devices

6.5. Further work and perspectives

6.6. Bibliography

7 GaN Nanowire-based Lasers

7.1. Introduction to nanowire lasers

7.2. Theoretical considerations and simulations

7.3. The first experimental observations of lasing in nanowires

7.4. GaN nanowire-based lasers

7.5. Toward wavelength tunability: nanowire lasers based on GaN/InxGa1-xN heterostructures

7.6. GaN nanowire lasers coupled with hybrid structures

7.7. Challenges and opportunities

7.8. Bibliography

8 GaN Nanowire-based Ultraviolet Photodetectors

8.1. Introduction

8.2. Growth and fabrication techniques

8.3. GaN nanowire photoconductive detectors

8.4. p–i–n junction-based GaN nanowire detectors

8.5. Single-wire GaN/AlN multiple quantum disk photodetectors

8.6. Single-wire InGaN/GaN core–shell photodetectors

8.7. Conclusions

8.8. Acknowledgments

8.9. Bibliography

9 ZnO Nanowire-based LEDs

9.1. Outline

9.2. Introduction

9.3. Growth of ZnO nanowires

9.4. White light emission from ZnO nanowires

9.5. ZnO nanowire white LEDs on solid crystalline substrates

9.6. ZnO nanowire white LEDs on flexible substrates

9.7. Enhancing the emission of ZnO nanowire-based LEDs

9.8. Conclusion and future prospective

9.9. Bibliography

10 ZnO Nanowire-based Solar Cells

10.1. Introduction

10.2. ZnO nanowire dye-sensitized solar cells

10.3. Quantum dot-sensitized nanowire solar cells

10.4. Extremely thin absorber solar cells

10.5. Nanowire arrays completely filled with inorganic absorbers

10.6. ZnO nanorod – organic hybrid solar cells

10.7. ZnO nanowire arrays for photoelectrochemical water splitting

10.8. Conclusions

10.9. Acknowledgments

10.10. Bibliography

List of Authors

First published 2014 in Great Britain and the United States by ISTE Ltd and John Wiley & Sons, Inc.

Apart from any fair dealing for the purposes of research or private study, or criticism or review, as permitted under the Copyright, Designs and Patents Act 1988, this publication may only be reproduced, stored or transmitted, in any form or by any means, with the prior permission in writing of the publishers, or in the case of reprographic reproduction in accordance with the terms and licenses issued by the CLA. Enquiries concerning reproduction outside these terms should be sent to the publishers at the undermentioned address:

ISTE Ltd27-37 St George’s RoadLondon SW19 4EUUKwww.iste.co.uk

John Wiley & Sons, Inc.111 River StreetHoboken, NJ 07030USAwww.wiley.com

© ISTE Ltd 2014The rights of Vincent Consonni and Guy Feuillet to be identified as the authors of this work have been asserted by them in accordance with the Copyright, Designs and Patents Act 1988.

Library of Congress Control Number: 2014941789

British Library Cataloguing-in-Publication DataA CIP record for this book is available from the British LibraryISBN 978-1-84821-687-7

Preface

This book is devoted to the specific case of wires obtained from a given kind of semiconductors, namely the semiconducting materials with a direct and wide band gap (WBG). In short, semiconductors are considered as WBG semiconducting materials if their band gap energy is typically above 1.5/1.6 eV. The interest of these materials for optoelectronic devices lies in the fact that they are well-adapted for emission, detection or absorption processes in most of the visible range, and part of the UV range as well. From the more basic point of view, the large refractive index and high exciton binding energy as well as the strong photon/exciton interactions give rise to long sought effects such as polariton lasing at room temperature for instance. The two main materials composing the family of WBG semiconducting materials are GaN and ZnO. They have close band gap energy in the near UV region (i.e., around 3.3/3.4 eV), and have in common that their cationic alloys span the visible as well as the UV range (and also part of the near IR region for In-rich GaInN alloys). More importantly, they both crystallize, in standard conditions, in the strongly anisotropic wurtzite crystalline phase, leading to a large number of similar physical quantities such as lattice parameters and piezoelectric constants and of similar physical processes related for instance to polarity.

GaN and its alloys are now well-mastered and used in a flurry of industrial applications as optoelectronic devices. On the other hand, ZnO is less advanced in terms of industrial applications and its development is mainly hampered by the difficulty for controlling p-type doping. However, ZnO has a stronger exciton binding energy than GaN (60 meV vs. 25 meV) and also a stronger oscillator strength. GaN and related alloys are generally heteroepitaxially grown on foreign substrates since low-cost nitride substrates with large dimensions are still not available. In contrast, ZnO and related alloys can homoepitaxially be grown onto ZnO substrates with excellent structural properties but still with limited availability and sizes. Therefore, epitaxial growth is mostly carried out heteroepitaxially for both kinds of materials, typically yielding epitaxial planar layers with a high density of structural defects. If such WBG semiconducting materials with a rather poor structural quality are actually used for some optoelectronic devices such as commercial LEDs for the moment, the improvement of their overall structure would certainly be beneficial for additional potential optoelectronic devices but also for the understanding of the physical processes at stake in these devices.

The need for WBG semiconducting materials with better structural quality is one of the main reasons that propelled (nano)wires to their present day status in the field of semiconductor research: when grown onto foreign substrates, and as for the case of planar layers, wires can relax the elastic strain energy originating from large lattice mismatch by forming misfit dislocations. But these lie in the basal plane or bend towards the nearby lateral surfaces of wires, thus leaving defect-free materials in their core. This process whereby dislocations can bend towards the lateral growth front had been demonstrated beforehand in epitaxial lateral overgrowth (ELO).

The second reason behind the development of WBG semiconductor wires – considered for a long time as the unwanted result of wrong growth conditions when trying to synthesize 2-dimensional (2D) epitaxial layers– is related to the increasing interest for low-dimensionality objects, typically of sub-micron or nanometer size. The specific structural, optical, and electronic properties of these low-dimensionality objects open new opportunities for nanoscale optoelectronic devices, especially to fully exploit the strong photon/exciton interactions. As an example, wires allow for a full confinement of light in their section with free propagation along their axis. Such physics and the related optoelectronic applications are nonetheless limited by the large developed surfaces of the wires, for which surface passivation is for instance required in order to prevent light diffusion. Because of the presence of surface states, Fermi level pinning also leads to band bending affecting the carrier mobility along the wires and resulting in possible carrier trapping. In return, this specific property makes wires very invaluable objects to investigate surface effects in WBG semiconductors and can also be beneficial in photodetection applications.

Looking back in time, the first demonstration of semiconductor wire growth was achieved by the pioneering work of Wagner and Ellis in 1964 according to the vapor-liquid-solid (VLS) mechanism [WAG 64]. In the field of WBG semiconducting materials for optoelectronic devices, which are the materials that we are interested in in this book, one of the first “nano”objects that were looked into were dots, named quantum dots when the typical dimensions are smaller than the De Broglie’s wavelength, inserted as they were in 2D epitaxial layers. For instance, the dots can be grown according to the so-called Stransky-Krastanov mode owing to the elastic stress relaxation processes at play in lattice mismatched heteroepitaxial systems. This is nevertheless limited somehow to heteroepitaxial layers in a state of compressive strain, and of medium lattice mismatch range (typically a few percent). For one heteroepitaxial system, such dots have once and for all a fixed size given by the nature of the involved materials. Thus, one had to think of other possibilities for making sub-micron or nano objects with an easier control over their sizes and shapes. Instead of playing for instance with strain to form dots, the easier way to grow low-dimensionality structures is to try and depart from the 2D growth conditions, thereby changing the atomic diffusion and incorporation processes, hence using growth modes different from the usual 2D mode. This time, this leads to the controlled formation of 1D objects, now referred to as nanowires, microwires or more generally wires, depending on their lateral dimensions, or also as nanocolumns, nanorods or microrods.

Interestingly, in terms of growth conditions, while most of the semiconductor (i.e., Si, Ge, arsenides, phosphides, …) wires can exclusively be grown by VLS or vapor-solid-solid mechanisms in the bottom-up approach, one of the most amazing properties of GaN and ZnO is their ability to grow in the form of wires following catalyst-free approaches (i.e., self-induced growth, spontaneous growth, …). These catalyst-free approaches are expected to reduce potential contamination into the wires and, more importantly, offer new valuable growth modes with great potentiality for optoelectronic devices. The first demonstrations of GaN and ZnO wire growth were shown in 1998 by molecular beam epitaxy [YOS 97, SAN 98] and in 2001 by vapor phase transport [HUA 01, PAN 01] and in solution [VAY 01], respectively. Basically, GaN wires can mainly be grown by molecular beam epitaxy and metal-organic chemical vapor deposition. In contrast, ZnO wires can additionally be deposited by vapor phase transport, pulsed-laser deposition or more specifically in solution via the low-cost and low-temperature chemical bath deposition technique for instance.

As discussed above, growing wires with dedicated properties in a reproducible way requires a good control of the growth conditions. When it comes to radial as well as axial heterostructures grown around or on top of the wires, things are somehow more complicated, since growth conditions very often have to be moved from the initial 1D case in order to stack the layers on top of each other. As in the case of any kind of heterostructures, managing the lattice mismatch issue may also be essential. This does depend upon the sizes involved and may potentially lead to the generation of misfit dislocations at the interfaces between the constituting layers. Moreover, owing to the specific geometry of the wires, other types of defects may also be introduced, such as stacking faults or inversion domain boundaries for instance, the origin of which has to be identified in order to better limit their occurrence. In return, identifying the right conditions for growing heterostructures with a good structural quality opens up a flurry of applications in the field of optoelectronics. These will benefit not only from the wave guiding properties of the wires (i.e., specific optical modes) but also from the control over the density of defects into the wires, leading to a decrease in the number of non-radiative recombination centers. These applications also take advantage of the larger surface to volume ratio at low-scale dimensions, leading for instance to much larger emitting or absorbing surfaces than in 2D layers or to efficient photodetectors.

The book has been organized along the lines of these introductory remarks.

Accordingly, it is the aim of the first part of volume 1 to focus on the specific properties of WBG semiconductor wires, in order to point out what differentiates these objects from their 2D counterparts. This appears as a necessary step in order to point out what these specificities could bring for the physics and applications of WBG semiconductors in the field of optoelectronics. It is nonetheless also the aim of this first part to try and pin-point the present day limitations associated with the use of WBG semiconductor wires, in order to draw possible solutions for a thorough use of these 1D objects. As for the second part of volume 1, it is dedicated to the different growth methods for the deposition of GaN and ZnO wires, stressing the mechanisms at play for the nucleation and growth of these 1D objects. The most interesting growth methods are discussed in detail with a special emphasis on the necessary ingredients to spontaneously grow GaN and ZnO wires. In volume 2, the first part aims at reviewing the different axial or radial heterostructures that can be integrated into GaN and ZnO wires. This is done to address relevant potential optoelectronic applications including LEDs, lasers, UV photodetectors and solar cells, which are presented and discussed in the second part of volume 2.

As revealed by the very numerous publications, the subject is far from being closed and new results emerge at a quick pace. With this in mind, this book is intended to give the reader a detailed overview of the current status of research in the field of WBG semiconductor wires for optoelectronic devices. As announced in the very title of this book, the choice was deliberately made to intermix chapters devoted to GaN and ZnO wires: the two materials have a lot in common, and the two communities will gain from mutual exchanges.

We hope that the reviews presented here by pioneering and world-leading scientists in the field, the discussion on the chemistry, physics, and applications of WBG semiconductor wires, together with the comparison between the two kinds of materials and between the different growth methods will be a useful source of information not only for the new comers in the field, but also for the already involved engineers and scientists who seek a detailed overview of the subject to give their work a new impulse.

Finally, we would like to warmly thank all our friends and colleagues who took part in this book project to create a lively, fruitful and high level place on the hot topic of WBG semiconductor wires.

Vincent CONSONNIGuy FEUILLETJune 2014

Bibliography

[WAG 64] WAGNER R.S., ELLIS W.C., Appl. Phys. Lett., 4, 89 (1964).

[YOS 97] YOSHIZAWA M., KIKUCHI A., MORI M., et al., Japanese J. Appl. Phys., 36, L459 (1997).

[SAN 98] SANCHEZ-GARCIA M.A., CALLEJA E., MONROY E., et al., J. Cryst. Growth, 183, 23 (1998).

[HUA 01] HUANG M.H., MAO S., FEICK H., et al., Science, 292, 1897 (2001).

[PAN 01] PAN Z.W., DAI Z.R., WANG Z.L., Science, 291, 1947 (2001).

[VAY 01] VAYSSIERES L., KEIS K., LINDQUIST S.E., et al., J. Phys. Chem., B 105, 3350 (2001).

PART 1GaN and ZnO Nanowire Heterostructures

1

AlGaN/GaN Nanowire Heterostructures

1.1. A model system for AlGaN/GaN heterostructures

In order to address real (optoelectronic) device applications based on GaN nanowires (NWs), the control of carrier confinement and of optical transition energies by alternating the optical band gap, either parallel or perpendicular to the growth direction, is of major importance. Within the group III-nitride (III-N) material system, this can be achieved either by the realization of AlGaN/GaN nanowire heterostructures (NWHs) that expand the energies of the involved optical transitions to the ultraviolet regime or by the realization of InGaN/GaN NWHs that open the way to the blue and green spectral region. Both types of NWHs impose different requirements in terms of growth conditions and – mainly due to the different growth temperature – exhibit different structural and morphological properties because the surface mobility of adatoms and the crystalline phase stability are strongly affected by the applied growth conditions. Therefore, AlGaN/GaN NWHs will be addressed in this chapter, while InGaN/GaN NWHs will be addressed in section 2.1.

The self-organized growth of GaN NWs results in nanostructures that, due to the possibility of lateral strain–relaxation during growth, exhibit very low densities of structural defects despite a large lattice mismatch with respect to the underlying substrate. Hence, they represent a perfect starting point for the growth of NWHs with optical properties that are only weakly influenced by recombination related to structural defects. Therefore, AlGaN/GaN NWHs are an ideal model system for the investigation of basic material properties because – although two-dimensional (2D) AlGaN/GaN heterostructures are well understood – the relations between structural characteristics on the one hand and optical as well as electrical characteristics on the other hand are still a topic of current research.

In the following, the growth, structural, optical and electrical properties of different types of AlGaN/GaN NWHs are discussed. Here, we focus on those NWHs synthesized by a self-assembled bottom-up growth process. Resembling the chronology of the research work in the past two decades, we start the discussion with axial AlGaN/GaN NWHs grown along the polar growth direction (section 1.2) by plasma-assisted molecular beam epitaxy (PA-MBE). Here, the structural characteristics and the strain distribution of NWHs as a consequence of the different lattice parameters of AlGaN and GaN are discussed. Furthermore, the optical properties of GaN nanodisks (NDs) embedded in AlGaN/GaN NWHs as well as their dependence on the structural characteristics are summarized. The role of axial and lateral internal electric fields is discussed and the benefits of micro-photoluminescence spectroscopy (μ-PL) for the analysis of single NWs and even single NDs are demonstrated.

In section 1.3 radial or core–shell AlGaN/GaN NWHs are addressed. Also in this case, we start the discussion with a review of growth issues and structural properties before we summarize recent results of their optical characterization. Depending on the applied growth technique, core–shell NWHs can be grown with polar, semi-polar and non-polar side facets. In the first case, this gives rise to carrier accumulation at lateral hetero-interfaces and hence a higher degree of freedom for the design of electronic properties compared to PA-MBE grown axial NWHs. Such concepts are summarized in detail. The discussion in this section also includes the realization of one-dimensional (1D) GaN quantum wires (QWRs) realized by selective nucleation of GaN on the edges of AlN/GaN core–shell NWHs.

In section 1.4 we summarize two complementary application concepts for AlGaN/GaN NWHs, namely an optochemical sensor for the detection of oxygen and hydrogen based on GaN NDs in axial AlGaN/GaN NWHs and a resonant tunneling diode realized on the non-polar lateral surface of AlN/GaN double-barrier core–shell heterostructures. (The applications of NWHs as LEDs, lasers and UV photodetectors are addressed in Chapters 5–8). Finally, conclusions are presented in section 1.5.

1.2. Axial AlGaN/GaN nanowire heterostructures

With typical diameters of several 10–100 nm, GaN NWs cannot be considered 1D nanostructures and additional quantum effects due to a transition from quantum wells (QWs) to their counterparts embedded in NWHs, i.e. a further reduction in dimensionality, are not expected. Thus, the properties of embedded axial quantum structures, i.e. NDs, resemble the properties of 2D QWs rather than truly zero-dimensional (0D) behavior. Still, the morphological and optical properties of axial quantum structures embedded in NWHs are governed by their three-dimensional (3D) geometry that provides, e.g., the possibility for strain relaxation on the free lateral surface or strain management due to the presence of lateral shells. Due to the specific growth kinetics, PA-MBE has become the technique of choice for the growth of axial NWHs. The discussion in this section therefore focuses on this synthesis method. A detailed discussion of the self-induced growth by MBE can be found in Chapter 8 of Vol. 1 [CON 14].

1.2.1. Structural properties of axial AlGaN/GaN nanowire heterostructures

Homogeneous GaN NWs are widely considered as nanostructures exhibiting an extremely low density of structural defects. The possibility for strain relaxation on the lateral surfaces results in the absence of misfit dislocations when the NW diameter is small enough [YOS 97, CAL 00]. For the same reason it was found that homogeneously doped GaN NWs are free of strain for Si and Mg as dopants [FUR 08, RIC 08]. Whereas the incorporation of Si does not enhance the formation of structural defects even in high concentrations [FUR 08], it was reported in [ARB 09] that doping with Mg results in the formation of triple-twin domains which, in high concentrations, cause the formation of zinc blend atomic cells in the wurtzite stacking.

In axial AlGaN/GaN NWHs, the situation is different as the formation of coherent interfaces, due to the change in alloy composition along the growth axis, results in the generation of compressive (tensile) strain in the GaN NDs (the AlGaN barrier). In the 2D case, the deposition of a compressively strained GaN layer on an AlN substrate leads to the formation of islands which is an efficient mechanism for strain relaxation when a critical layer thickness of typically several monolayers (MLs) is exceeded. This relaxation mechanism is the driving force for the Stranski–Krastanov growth of GaN quantum dots (QDs) on AlN and AlGaN that has been intensely studied in recent years [DAU 08]. It remains an interesting question as to how the limited diameter of NWs and the corresponding possibility for strain relaxation influence the critical thickness for GaN grown on AlN embedded in NWs. In [GLA 06] strained layers on top of free-standing NWs were considered in a theoretical model showing that the critical layer thickness depends on the NW radius. It was further estimated that there exists a critical value of the radius, below which arbitrarily thick coherent layers should be obtainable (for a more profound discussion of stress relaxation see Chapter 2 in Vol. 1).

Regarding the general structure of GaN NDs, they have been found to appear as flat disks with sharp interfaces to AlN barriers [FUR 11]. ML fluctuations of the ND height have also been observed, particularly in the case of AlGaN barriers [RIG 10b, PIE 13]. It has been observed by several groups that the ND side walls are faceted and that the NDs are slightly bent downward on the outer edges [FUR 11, BOU 10, RIS 05a]. In [FUR 11], it was reported that this structure originates in faceting of the GaN base top surface, where -planes form the outer edges, rather than in strain relaxation of the embedded NDs. The top surface faceting is most probably a consequence of the N-face polarity of the NWs [MAT 12] in combination with the N-rich growth conditions. Subsequent overgrowth with ND and barrier material results in the typical shape of the ND stack.

It has also been observed by different groups that, while the GaN base and the GaN NDs grow in a 1D growth mode, the growth of the AlGaN barrier often exhibits a lateral growth rate that results in the formation of a lateral shell that increases in thickness during growth of each barrier [RIS 05a, TCH 08, BOU 10, ZAG 11]. The lateral growth rate of the barrier material, which depends on the growth temperature and the total metal flux during barrier growth, strongly impacts the morphology and the resulting strain distribution of the GaN/AlGaN ND stack. In [LAN 10], no lateral growth of the AlN barrier was found (see Figure 1.1(a)), while in [BOU 10] the presence of a lateral shell was mentioned (see Figure 1.4). Also in [RIS 05a], the presence of a lateral AlGaN shell is reported. In [TCH 08], a lateral growth rate for AlN of 35% compared to the axial growth rate was found, and in [FUR 11], it was found to be 11% for AlN barriers and linearly decreasing with the Al concentration in the barriers [Al]bar (see Figure 1.2).

Figure 1.1.a) High-resolution transmission electron microscopy (HRTEM) image of a GaN/AlN NWH. Five periods of AlN/GaN grown on a GaN base are visible. The arrow indicates the growth direction; b) high-angle annular dark field (HAADF) image showing that no significant inter-diffusion occurs between AlN and GaN; and c) profile of the c lattice-parameter, along the growth axis taken in the central part of an NW, obtained from the geometrical phase analysis of the HRTEM image. The arrow indicates the growth direction. For convenience, the x-axis origin has been taken as the top of the GaN NW base before the growth of the first AlN layer. (Reprinted with permission from [LAN 10]. Copyright © 2010, American Physical Society)

In a first systematic experimental study on elastic strain relaxation in AlN/GaN NW super-lattices, Landré et al. found by in situ high-resolution X-ray diffraction experiments that AlN/GaN (2.3 nm/2 nm) super-lattices (Figure 1.1a, b) on a GaN NW base are in elastic equilibrium, i.e. the strain is distributed between the GaN and AlN layers according to their layer thickness and an averaged in-plane lattice parameter corresponding to an Al0.55Ga0.45N alloy is adopted. By HRTEM analysis, an increase in the c-parameter in the GaN NDs compared to the GaN NW base was found that corresponds to a decrease in the in-plane lattice parameter according to the transition from relaxed GaN to that of Al0.55Ga0.45N (Figure 1.1(c)). The c-parameter in the AlN barriers (lower value in Figure 1.1(c)) corresponds to the value that is obtained when the in-plane lattice parameter for Al0.55Ga0.45N and the Poisson ratio for AlN are considered. No evidence for strain relaxation due to the presence of misfit dislocations was found [LAN 10].

A different situation arises when the growth parameters are changed and an enhanced lateral growth rate during the growth of the AlN barriers is obtained. In [FUR 11], a lateral growth rate of 11% from the axial growth rate was observed during barrier growth in a nine-fold AlN/GaN ND structure. As is depicted in Figures 1.2(a) and (b), this results in an increasing ND diameter and a decrease in the shell thickness along the growth direction as well as full encapsulation of the GaN NDs with AlN. Hence, free relaxation of NDs at the NW periphery can no longer be assumed and compressive stress along the c-direction is also exerted in that region. This effect remains but becomes less pronounced with decreasing Al-concentration in the barriers [Al]bar since a linear dependence of the lateral growth rate on [Al]bar was observed.

As a consequence of these boundary conditions, the strain state of the NDs is altered compared to that described in [LAN 10] (Figure 1.1). In particular, a different strain state of each individual ND within a multi-ND structure has to be considered, which complicates the discussion of optical properties. In [FUR 11], such AlN/GaN ND structures with different ND heights were investigated by high-angle annular dark field STEM with respect to the formation of misfit dislocations.

Strain relaxation due to the formation of dislocations in AlN/GaN ND structures has also been found by HRTEM analysis reported in [BOU 10]. Here, the authors found the formation of dislocations at the interface between the GaN ND and the AlN barrier (depicted in Figure 1.4). The presence of dislocations could explain the observed deviation of results from geometric phase analysis (GPA) and strain simulations based on a valence-force model [KEA 66]. Using the latter, a strain maximum in the ND center and in the edges of the basal plane was estimated but not observed in GPA.

Whereas the composition of AlGaN barriers has mainly been regarded as being constant throughout the barrier, a detailed analysis by Pierret et al. has revealed that, depending on the intentional Al-concentration, significant fluctuations of the Al-concentrations can occur in AlGaN NWs grown on a GaN base if the growth temperature is too high [PIE 13]. This might also result in the destabilization of AlGaN/GaN interfaces of GaN NDs and influence the emission properties of GaN NDs [RIG 10b].

In summary, the structural properties of GaN NDs embedded in AlGaN/GaN NWHs strongly depend on the impact of mechanical strain by the AlGaN barriers and due to the presence of a lateral AlGaN shell. Both factors influence the possible formation of dislocations on the one hand and the radial and axial strain distribution on the other hand. They are also expected to affect the optical characteristics of NWHs via polarization-induced internal electric fields.

1.2.2. Optical properties of axial AlGaN/GaN nanowire heterostructures

After the first reports on the self-assembled catalyst-free synthesis of GaN NWs and particularly of their intriguing optical characteristics [YOS 97, CAL 00], only a few reports dealt with an in-depth optical analysis of such structures [RIS 03, CAL 00]. In 2005, Ristic et al. extended this field of optical properties of NWs to those of NWHs, realized by embedding GaN NDs between AlxGa1 – xN barriers in a GaN NW [RIS 05a].

Figure 1.4.a) HRTEM image showing a dislocation at the AlN/GaN interface for three successive inclusions; and b) enlargement showing the insertion of an extra (0002)-plane in AlN. ([BOU 10]. Copyright © 2010, IOP Publishing. Reproduced by permission of IOP Publishing. All rights reserved)

In contrast to 2D systems, where the lateral extension of the QWs can be regarded as infinite, the NW geometry imposes specific geometric boundary conditions with zero stress along the lateral surface, which results in a strain gradient along the NW diameter. Hence, to calculate the potential profile along the NW diameter, these strain variations and the corresponding interplay between strain-dependent contributions of the deformation potential and the piezoelectric polarization have to be analyzed. This mechanism gives rise to the so-called “strain confinement” that was first mentioned in [RIS 05a, RIS 05b] and systematically elaborated in [RIV 07]. For compressively strained GaN NDs between AlGaN barriers, the resulting strain-induced piezoelectric polarization fields cause a carrier confinement in the NW center, whereas the deformation potential leads to a confinement of carriers at the NW surface, as schematically shown in Figure 1.5 I, II. As the deformation potentials for conduction band and valence band differ for GaN [VUR 03], situations with the confinement for holes (electrons) determined by piezoelectric polarization and that for electrons (holes) by the deformation potential can occur, depending on the Al concentration in the barriers, [Al]bar, and the ND thickness, dND (see Figure 1.5 III (IV)). From mechanical considerations it was concluded that for thick (thin) NDs, piezoelectric (I) (deformation potential (II)) confinement prevails, while for intermediate thicknesses the cases III and IV can occur.

By calculation of the corresponding wave functions, it was demonstrated that the spatial separation of electrons and holes can be controlled by the extrinsic parameters [Al]bar and dND of the NWH. The results obtained applying this “strain confinement model“ were used to explain the observed decrease in PL intensity with decreasing ND thickness as a reduced influence of the piezoelectric polarization and a corresponding spatial separation of electrons (confined in the ND center) and holes (confined at the periphery) leading to a reduced oscillator strength for radiative recombination and an enhanced non-radiative recombination via surface states [RIV 07]. Thus, the presence of a strained NW core and the strain relaxation at the lateral surfaces gave rise to interesting physics in AlGaN/GaN NWHs and stimulated further research activities in this direction.

In contrast to the ND material, i.e. GaN, the barrier material, AlGaN or AlN, also exhibits a lateral growth rate under some growth conditions and hence leads to the formation of a lateral shell consisting of the barrier material [TCH 08, FUR 11]. In that case, PA-MBE growth of a multi-ND NW heterostructure results in a 3D confinement due to the presence of a lateral AlGaN or AlN shell that decreases in thickness along the growth direction as shown in Figure 1.2(a). The presence of this lateral shell was demonstrated to severely impact the optical emission properties of AlGaN/GaN NWHs in various respects: the compressive stress on the ND lateral surface exerted by the AlN shell can significantly alter the band profile in the ND and, depending on its thickness, cause transitions from case (I) to case (II) of Figure 1.4 along a single multi-ND NWH shown in Figure 1.2a.

Figure 1.5.Schematics of the conduction and valence bands for the in-plane directions of a GaN/AlGaN QD. For I (II), the piezoelectric (deformation) potential determines the net potential of both the conduction and valence band. For III (IV), the piezoelectric potential determines only the net potential of the valence (conduction) band, whereas the deformation potential determines the net potential of the conduction (valence) band. (Reprinted with permission from [RIV 07]. Copyright © 2007, American Physical Society)

Furthermore, as the lateral growth rate of the shell strongly depends on the Al-concentration in the barrier, the thickness of the lateral shell and therefore its strain impact on the GaN NDs increases with increasing [Al]bar. The compressive axial strain exerted by the AlGaN shell results in a reduction of the piezoelectric contribution to the internal electric fields and thus to an attenuation of the quantum-confined Stark effect (QCSE), i.e. to a blue-shift in the transition energies compared to the case when no shell is present, as shown in Figure 1.6(b) by a comparison of experimental results for the transition energies in nine-fold AlGaN/GaN ND samples and 3D simulations of the entire structure that were carried out using the software package nextnano3 [NEX].

The interplay between improved carrier confinement and increasing influence of polarization-induced internal electric fields is also reflected in the temperature stability of the ND-related PL emission, shown in Figure 1.7. Here, an increasing temperature stability is observed for low and medium [Al]bar, i.e. [Al]bar < 0.30, as the height of the potential barriers increases and the 3D confinement is improved due to the formation of a lateral shell. At [Al]bar ≈ 0.30, the highest temperature stability was observed, comparable to that of GaN quantum dots in an AlN matrix [GUI 06]. For further increase in [Al]bar, the increasing internal electric fields lead to a decreasing exciton binding energy in the NDs [GRA 99a, LEA 90] and hence a faster quenching of the PL intensity (not shown).

1.2.3. Lateral internal electric fields

Figure 1.8(a) depicts the evolution of the lateral conduction band profiles with increasing [Al]bar (the valence band profiles follow the same trend and are not shown for clarity) as extracted from the set of calculations discussed above and presented in Figure 1.6. It is evident that the lateral internal electric fields increase with increasing [Al]bar, exceeding the critical field of 80 kV/cm for the separation of excitons in GaN [SHO 03] for high Al concentrations. As a result, the effect of “strain confinement” proposed by Rivera and Ristic et al. [RIS 05a, RIV 07] is enhanced by the presence of a lateral AlGaN shell as in this case compressive stress is also applied in the axial direction on the ND periphery. The evolution of the lateral electric fields with increasing [Al]bar was proved by time-resolved PL analysis in [FUR 11] as depicted in Figure 1.8(b). Here, it is shown that for an ND thickness of 1.7 nm, which is below the exciton Bohr radius in GaN of 2.7 nm [STE 99, GAL 00], the PL decay time of an ND ensemble (as shown in Figure 1.2), significantly increases when [Al]bar exceeds 30%, i.e. when the lateral band profiles turn from flat band to strongly U-shaped (see Figure 1.8(a)). The long decay times of 2 ns and more demonstrate the transition from an exciton to spatially (laterally) separated electron and hole wave functions. The inset in Figure 1.8(b) demonstrates that this effect is not caused by the presence of axial internal fields (discussed below), as an increase in the ND thickness in AlN/GaN NWHs results only in a weak increase in the PL decay time. In comparison, AlN/GaN QWs show a strong increase in the PL decay time up to more than 100 ns when the QW thickness increases. The presence of lateral electric fields in NDs and their absence in QWs are evidenced by the significantly higher PL decay time for NDs with small heights. The evolution for large ND or QW heights also demonstrates that the apparent axial internal electric fields in AlN/GaN ND samples are smaller than those which are observed for AlN/GaN QWs (also discussed below).

Figure 1.7.Arrhenius-plots of the ND PL intensity for different Al concentrations in the barrier [Al]bar. An increasing temperature stability with increasing [Al]bar is observed for values up to [Al]bar ≈ 0.30. For higher values, the temperature stability decreases due to the increasing influence of polarization-induced internal electric fields (not shown). For a color version of this figure, see www.iste.co.uk/consonni/nanowires2.zip

1.2.4. Axial internal electric fields

1.2.5. Optical characterization of single-AlGaN/GaN nanowires containing GaN nanodisks

In the preceding sections, it has been demonstrated that control of the growth process and hence the structural properties of AlGaN/GaN NWHs provides nanostructures that facilitate the analysis of the basic optical properties of confined structures within an NW and even help in advanced studies of more complex optical characteristics. However, wire-to-wire fluctuations of the structural properties also affect the optical properties of individual NWHs. Variation of the ND diameter, the ND height or the Al concentration in the barriers results in broadened PL emission lines that strongly limit the possibilities of assigning the optical characteristics to specific microscopic properties. Hence, μ-PL analysis of single NWs (that might contain several NDs) with reasonable statistics is a key technique for gaining information that would not be possible to be obtained from ensemble measurements. This was demonstrated by Rigutti et al. in [RIG 10b], where the origin of the energy dispersion in AlGaN/GaN NWH for [Al]bar < 0.14 was studied. In Figure 1.10(a), the PL spectrum of an ensemble of NWHs with nine-fold GaN NDs in Al0.16Ga0.84N barriers is compared to the μ-PL spectrum of two single NWHs of the same sample and to the μ-PL spectrum of a single NWH containing a single NDs embedded in Al0.16Ga0.84N barriers.

Basically, the spectra are characterized by emission features in two different energy regimes: the emission around 3.47 eV related to the GaN NW base and an emission band between 3.6 and 3.7 eV that is related to the nine-fold GaN ND structure. While ND-related emission of the ensemble exhibits a single broad peak between 3.62 and 3.63 eV with a FWHM of 45 meV the single-NW μ-PL spectrum is characterized by multiple narrow peaks with a spectral width of around 3 meV, which the authors assigned to emission from the individual NDs. On the one hand, this assignment is confirmed by comparison to the μ-PL spectrum of an NW with a single ND shown in the same diagram, and on the other hand, this comparison shows that the NDs in one NWH show a considerable dispersion of the PL emission energy. Within this report it was found that this dispersion is caused by two different effects: the presence of a (thin) lateral shell with decreasing thickness along the growth direction that gives rise to a different strain state in each of the NDs and the presence of ML fluctuations in the ND thickness that were observed by HRTEM analysis. The influence of both effects has been modeled using the software nextnano3 [NEX]. As displayed in Figure 1.10(b), both effects comparably contribute to the energy dispersion along the NW, and for both mechanisms the resulting variation in transition energies increase with increasing [Al]bar, in accordance with the results reported for higher Al-contents in [FUR 11].

However, it should be noted that the number of emission peaks exceeds the number of NDs in the wire and also the single-ND emission spectrum shows a double-peak structure which was attributed to different exciton localization energies within a single ND due to alloy disorder in the barriers, also confirmed by the results reported in [PIE 13].

The μ-PL analysis of single NWHs containing GaN NDs was also applied for the investigation of the PL polarization properties in [RIG 10a]. In this work, it was demonstrated that the near-band-edge excitonic emission from the GaN NW base was polarized parallel to the NW axis (π-polarized), while the ND emission in different multi-ND samples showed polarization perpendicular to that axis (σ-polarization), as displayed in Figure 1.11. The authors have assigned the π-polarization of the GaN-base to the specific NW geometry and to the dielectric contrast between the NW and the surrounding medium. The polarization of the ND emission was explained by the selection rules of the dipole-matrix elements. Due to the confinement-induced increase in the splitting between XA- and XB-transition energies, the ND emission is of pure, σ-polarized XA-character at low temperatures. Hence, the investigated AlGaN/GaN NWHs present a nanophotonic system in which the two main emission lines are perpendicularly polarized.

The first μ-PL analysis of single GaN NDs in AlN/GaN NWHs even facilitated the identification of biexciton transitions as reported by Renard et al. [REN 08]. In that work, GaN inclusions with a height of 1 nm between AlN barriers of 13 nm (bottom) and 8 nm (top) were realized. μ-PL analysis at low-excitation energies revealed emission in the range between 3.8 and 4.2 eV with line widths down to 1 meV, i.e. similar characteristics as observed for single-GaN/AlN quantum dots grown in the Stranski–Krastanov growth mode [KAK 04, BAR 06]. By power-dependent analysis of the single-ND PL intensity, the authors have identified excitonic and biexcitonic characteristics for different NDs (Figure 1.12). The latter is reflected by the quadratic power dependence of the PL intensity in Figure 1.12, that reflects a transition from a biexcitonic to an excitonic state with a biexciton binding energy of 20 meV.

As a similar experimental approach, μ-PL studies of single-Al0.14Ga0.86N/GaN NWHs containing one ND were reported by Jacopin et al. [JAC 13]. In that work, the ND height was systematically varied and the resulting effect on the transition energies was analyzed (see Figure 1.9(b)). Here, it has been shown that 3D simulations including the presence of polarization-induced internal electric fields and the impact of the lateral AlGaN shell have to be performed in order to achieve agreement with the experimental results. In contrast, the pronounced influence of the NW diameter on the interband transition energy that was estimated by the simulations could not be verified as under the experimental conditions in the single-NW experiments excitonic transition in the NW core rather than recombination processes at the NW periphery dominate the PL emission.

1.2.6. Electrical transport properties

The electrical transport properties of axial AlGaN/GaN NWHs have only rarely been investigated. One reason for this is the high resistivity of non-intentionally low-doped NWs due to lateral depletion effects [CAL 05], and although Si-doping has been demonstrated [RIC 08, FUR 08], the precise control of the transport properties has not yet been reported. A possible approach to overcome this problem is Ge-doping which facilitates achieving high-impurity concentrations without influencing the NW morphology [SCH 13]. In addition, the degrees of freedom for the design of nanoelectronic devices based on axial NWHs are limited, as transport properties parallel to the hetero-interfaces, which are of high interest in III-N devices (see section 1.3), are not accessible. In contrast, electronic transport across tunneling barriers is possible as it has been reported for InAs/InP double-barrier heterostructures in [BJO 02]. In [SON 10], a low-temperature transport study of an AlN/GaN double-barrier NWH (shown in Figure 1.13(a)) has been reported. NDR behavior was observed for voltages between 0.1 and 1 V, as it is shown in Figure 1.13(b). This behavior was assigned to the presence of discrete electronic levels as a consequence of quantum confinement between the AlN barriers.

Figure 1.12.Dependence of the energy-integrated photoluminescence (PL) intensity of the exciton-emission (X) and biexciton-emission (XX) as a function of the excitation power on a log-log scale. The solid lines represent a linear and a quadratic dependence. (Reprinted with permission from [REN 08]. Copyright © 2008, American Chemical Society)

In addition, also single-electron tunneling was observed at significantly lower voltages (Figure 1.14), pointing to another quantization mechanism. This latter effect was assigned to quantization of electronic levels perpendicular to the NW axis. The authors have estimated a level spacing of 1-10 mV from the data shown in Figure 1.14. In general, these results indicate that axial NWHs can serve as a tool for exploring nanoscale quantum transport effects.

1.3. AlGaN/GaN core–shell nanowire heterostructures

Group III-nitride core-shell NWHs are widely considered as offering a higher application potential in optoelectronic light emitters [MAN 13] or in nano-photovoltaics compared to the axial NWHs described above [QIA 08, DON 09]. As the main benefits of this type of nanostructures, exemplarily shown in Figure 1.15, the large active area between core and shell as well as the non-polar surface of the NW core are often put forward [MAN 13]. In order to realize core–shell structures that show quantum confinement, the insertion of a QW between the inner NW core and the outer NW shell has to be achieved. Thus, for optoelectronic applications, InGaN/GaN core–shell NWHs are a topic of intense research [MAN 13] and are discussed in detail in Chapter 2.

Figure 1.13.a) Schematic illustrations of a single NW-based device with metal contacts on an oxidized Si substrate. The lower left (lower right) scheme shows an n–i–n NW without (with) AlN double tunnel barriers; and b) ISD – VSD characteristic showing the evolution of the negative differential resistance (NDR) appearing at negative VSD for different VG. The same qualitative gate voltage dependence has been observed in most of the devices fabricated from the same sample. (Reprinted with permission from [SON 10]. Copyright © 2010, American Chemical Society). For a color version of this figure, see www.iste.co.uk/consonni/nanowires2.zip

In contrast to InGaN/GaN core–shell NWHs, which are widely discussed with respect to their potential application as light emitter structures, AlGaN/GaN core–shell NWHs have been mainly addressed with respect to their basic structural and optical properties. In the following, the state of the art in AlGaN/GaN core–shell NWHs is summarized and specific structural, optical and electronic properties are discussed.

1.3.1. Structural properties

Due to the enhancement of the lateral growth rate for AlGaN compared to GaN, AlGaN/GaN core–shell NWHs have also been grown by molecular beam epitaxy, in contrast to InGaN/GaN core–shell NWHs, which are mainly synthesized by vapor deposition processes [LI 12].

Figure 1.14.a) Differential conductance of the double barrier NWH shown in Figure 1.13a versus VG, revealing Coulomb blockade peaks. The measurement was done by using the lock-in technique with an alternate frequency of 13.305 Hz and an excitation amplitude of 500 μV; b) color scale plot of dISD/dVSD versus VG and VSD. All measurements were taken at 4.2 K; c) magnification of (b) at the region indicated by a black rectangle. Peaks in dISD/dVSD denoting the onset of tunneling via ground and excited states have been highlighted by dotted and dashed lines, respectively. (Reprinted with permission from [SON 10]. Copyright © 2010, American Chemical Society). For a color version of this figure, see www.iste.co.uk/consonni/nanowires2.zip

In 2008, Tchernycheva et al. have investigated the lateral growth rate of GaN during NW growth using thin AlN marker layers (visible in Figure 1.16(a) [TCH 08]) and found an increase from 0.8–3.5% to 7–11% of the axial growth rate when the Ga-BEP was increased from 7.2 × 10–8Torr to 1.25 × 10–7Torr.

The finite lateral growth rate of GaN made possible the realization of GaN-confined areas grown in the non-polar 〈100〉-direction on the lateral surfaces as also visible in Figure 1.16(a). Narrow emission lines with an FWHM of 1 meV, observed above the band gap of GaN in μ-PL spectroscopy of single NWHs (Figure 1.16b), were assigned to the presence of non-polar QW structures.

Due to the lateral growth rate of AlGaN, the core–shell geometry was also observed during MBE growth of axial NWHs. As already mentioned above, Furtmayr et al. reported a lateral growth rate of 11% of the axial growth rate for AlN, that was shown to decrease with decreasing [Al]bar and gave rise to the formation of a lateral shell around the AlN/GaN NWH core as depicted in Figures 1.2 and 1.3