Metal organic vapour phase epitaxy of AlN, GaN, InN and their alloys: A key chemical technology for advanced device applications

Accepted Manuscript Title: Metal organic vapour phase epitaxy of AlN, GaN, InN and their alloys: a key chemical technology for advanced device applications Author: Ian M. Watson PII: DOI: Reference:

S0010-8545(12)00264-0 doi:10.1016/j.ccr.2012.10.020 CCR 111654

To appear in:

Coordination Chemistry Reviews

Received date: Revised date: Accepted date:

26-8-2012 31-10-2012 31-10-2012

Please cite this article as: I.M. Watson, Metal organic vapour phase epitaxy of AlN, GaN, InN and their alloys: a key chemical technology for advanced device applications, Coordination Chemistry Reviews (2010), doi:10.1016/j.ccr.2012.10.020 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.

*Manuscript

Metal organic vapour phase epitaxy of AlN, GaN, InN and their alloys: a key chemical technology for advanced device applications

Author details: Ian M. Watson, Institute of Photonics, University of Strathclyde, Wolfson Centre, 106

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Rottenrow, Glasgow G4 0NW, UK. Tel. +44 (0)141 548 4120, fax +44 (0)141 552 1575

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e-mail: [email protected]

Contents

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1. Introduction

2. Basic properties of bulk materials and nanostructures

2.2 Alloy systems and bandgap engineering

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2.1 Crystal structures and physical properties of the binary compounds

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2.3 Dopants for conductivity control and advanced applications

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2.4 Quantum confined nanostructures and polarisation fields 3. Development of nitride MOVPE 3.1 Substrates and growth initiation 3.2 Conventional precursors

3.3 Reactor hardware and in situ monitoring

3.4 Growth conditions for different binaries and alloys 3.5 Mechanisms and growth simulations

3.6 Alternative precursors and innovative growth processes 3.7 Nanowire and nanorod growth 4. Summary and outlook

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Abstract This article reviews metal organic vapour phase epitaxy (MOVPE) processes developed for the group 13 nitrides AlN, GaN, InN and their alloys. The binaries are direct-gap semiconductors with respective bandgaps of 6.1 eV for AlN, 3.4 eV for GaN, and ~0.6 eV for InN, and adopt the hexagonal wurtzite crystal structure. The nitrides form continuous solid solutions, and are capable of being both n- and p-

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doped, thus making possible the growth of advanced heterostructure devices exemplified by GaN-based visible light-emitting diodes. Interest in nitride MOVPE from the late 1980s motivated significant work

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on single-source precursors, and also thermally labile nitrogen sources for use in two-source processes.

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The best developed of the former are azido compounds, which can have properties well tailored for lowtemperature film deposition. However, the nitride MOVPE processes that have come to dominate device

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manufacturing since the mid-1990s depend on the reaction between ammonia and metal alkyl sources, and deposit GaN at temperatures usually above 1000ºC. Most current nitride growth is performed

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heteroepitaxially on sapphire (0001) substrates, for which appropriate multistep growth initiation

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processes have been optimised. Current designs of nitride MOVPE reactor are engineered to avoid

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premature contact between the group 13 sources and ammonia, and feature in situ monitoring by optical means. The mechanisms of the growth chemistry are now understood to the extent that they are handled explicitly in multi-scale computational simulations of full processes. Particular recent advances in mechanistic understanding concern the role of nanoparticles that form in the gas phase, and which represent an important precursor loss channel. Methodologies for controlling the composition and properties of layers of GaN itself, and of ternary alloy layers with moderate (<25 mole %) contents of InN or AlN, are well established. However, greater challenges are posed by growth of layers InN, AlN, and of alloys close to these two binaries in composition. New variants of MOVPE continue to be explored as a consequence, and include processes with pulsed alternating precursor introduction to enhance lateral migration of adatoms on the surface of the growing film. A further important new emphasis in recent years is the controlled growth of nanowire and nanorod arrays, which already include core-shell heterostructures of considerable sophistication. Keywords 2

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Metal organic vapor phase epitaxy Metal organic chemical vapor deposition Gallium nitride Aluminum nitride Indium nitride

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Light-emitting diode

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1. Introduction

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This review focuses on metal organic vapour phase epitaxy (MOVPE) of the semiconducting group 13 nitrides, AlN, GaN, InN and their alloys, in the form of technologically useful thin films. MOVPE methods are a subset of

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the wider family of chemical vapour deposition or CVD processes, and the term refers specifically to processes in which layers with monocrystalline character are deposited on single-crystal substrates: a prerequisite for most

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advanced device applications. The alternative term metal organic chemical vapour deposition (MOCVD) is often used synonymously with MOVPE, but is in fact a less specific name encompassing processes producing

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polycrystalline and non-crystalline layers. MOVPE of the nitrides is a fundamentally chemical process, usually

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performed in continuous flow systems, and in most cases employing simple metal alkyl precursors reacted with ammonia. Despite its chemical roots, it has evolved to a state where the majority of research findings are now reported in the applied physics, materials science, nano-science or crystal growth literature. Consistent with this maturation well beyond the basics of depositing the desired solid phase from the vapour phase, most usage of MOVPE in both academic and industrial environments is for growth of complex multilayer structures for optoelectronic devices, as exemplified in Fig. 1. This is complemented by a continuing role in producing simpler multilayer samples for scientific investigation. A 2005 review of the materials chemistry of the group 13 nitrides gives particular emphasis to the mechanisms of conventional two-source MOVPE, and to the use of azido compounds and other single-source precursors for low-temperature deposition [1]. Gibart’s slightly earlier review of GaN MOVPE has a materials physics perspective, and concentrates particularly on means of improving the crystallographic perfection of heteroepitaxial GaN films [2]. The intention in this article is to first give a broad introduction to the valuable properties of the binary nitrides, their alloys and derived nanostructures, and to then survey the current status of MOVPE as a means to exploit these properties. The greatest emphasis will be placed on processes suitable for the production of device-quality multilayer structures in research and production settings,

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which have reached a high degree of sophistication. However, other topics in nitride MOVPE which are less mature, and still the subject of ongoing research efforts, will also be reviewed.

The devices made from nitride thin film structures impacting most on everyday life are light-emitting diodes (LEDs), either emitting directly in the blue-green spectral region, or used to generate white light via excitation of

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phosphors for solid sate lighting [3-5]. Nitride LEDs have improved steadily in performance, and fallen in unit cost, since their first introduction by the Nichia Corporation in 1993. The major driver for their use in solid-sate lighting

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is energy efficiency, and nitride LEDs some years ago exceeded the benchmark efficiency of ~70 lumens per Watt of electrical input power achievable with conventional fluorescent tubes [3]. GaN-based laser diodes pose much

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greater challenges than LEDs in terms of materials and device engineering, and achieved their first commercial impact in optical disk drives using devices emitting around 405 nm in the violet. They are now being developed

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further to provide the green and blue sources for full-colour projection displays [6]. Other important types of nitride device include LEDs emitting into the deep ultraviolet (UV) spectral region, which offer new opportunities in

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photocatalytic purification of air and water, sterilisation, and bioinstrumentation [7], and field effect transistors (FETs) for microwave power applications [8]. Additional applications which can be expected to develop fast

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include solar energy conversion, involving both photovoltaic and photochemical water splitting approaches, and

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chemical and biological sensing, in certain cases exploiting the properties of nitride nanowires [9-11].

Section 2 of this article surveys the properties of pure nitride materials, their doped counterparts and nanostructures underpinning the device applications just discussed. Most notable among the fundamental physical properties is the range of direct bandgaps, varying from 6.1 eV for AlN, through 3.4 eV for GaN, to ~0.6 eV for InN [9]. Bandgaps can be continuously varied between those of the binary compounds by the formation of ternary and quaternary alloys. Consequently the bandgaps of the technologically important material InxGa1-xN can span the entire visible spectrum (1.8 – 3.3 eV), and this justifies the choice of this alloy for active regions in visible LEDs and laser diodes, as well as the recent attention to solar energy conversion. The group 13 nitrides are usually synthesised as the hexagonal wurtzite phase, which contrasts with the cubic symmetry of many technologically mature semiconductors (Si, GaAs, InP etc.). The majority of nitride thin film samples are prepared with the hexagonal c-axis perpendicular to the substrate surface, such that only the (0001) basal plane is exposed on the film surface. However, different crystal faces of wurtzite crystals exhibit markedly different chemical reactivities,

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leading to significant anisotropy in crystal growth and etching processes. The wurtzite structure also lacks a centre of symmetry, which has profound consequences for the properties of nitride nanostructructures, owing to the presence of both a spontaneous (pyroelectric) polarisation and a piezoelectric polarisation in strained material [9,12]. These effects impact particularly on the properties of planar quantum wells (QWs) used in light-emitting devices, and provide motivations for epitaxy in new crystallographic orientations that mitigate the influence of

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electric fields. Recent progress in growth of bulk GaN crystals for use as substrates in homoepitxial growth is synergistic with these new directions in epitaxy [13]. Section 2 also covers doping issues, with discussions in turn

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of dopants used for conductivity control through introduction of itinerant charge carriers, and those where the

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inner-shell electron population of the dopant centre is of importance.

Section 3 provides detailed discussion of nitride MOVPE, which has achieved dominance for growth of most

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device structures over the main competitive methods hydride vapour phase epitaxy (HVPE) [2,12,13] and molecular beam epitaxy (MBE) [12,14]. Brief comparisons with these methods and a few others are made later in

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the current section. Section 3 begins with an overview of substrate properties, focussing especially on sapphire and silicon which have most relevance for mass produced devices. Conventional MOVPE using metal alkyl sources and

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ammonia has been refined to a highly sophisticated state over the past two decades by extensive interdisciplinary

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effort. Topics emphasised here include the essential role of in situ probes, the different configurations of reactor that have achieved success in large-area deposition, the use of computational modelling particularly to understand parasitic gas-phase reactions, and the significantly different growth regimes needed for optimised growth of the different binary compounds and alloys. There is a separate discussion of processes using non-conventional precursors or process conditions, for which motivations include growth at lower temperatures. Recent years have also seen an exciting extension of nitride MOVPE beyond its established role in preparing planar multilayer samples to the preparation of nanowire arrays. This topic has already advanced to a significant level of sophistication, for example, in demonstrator LEDs and radial core-shell structures [15,16].

Concerning alternatives to MOVPE for producing nitride thin film structures, various physical vapour deposition (PVD) techniques, including pulsed laser ablation, reactive sputtering, and reactive evaporation are capable of producing phase-pure textured polycrystalline layers that may suit less demanding applications. However, only the most sophisticated PVD technique MBE rivals MOVPE in its ability to produce highly complex epitaxial

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multilayers [12,14]. MBE invariably uses elemental metal sources evaporated from heated effusion cells. Transport of evaporated species in an ultra high vacuum (UHV) ambient takes place in a molecular flow (ie. collision-free) regime, and the name of the technique derives from this characteristic. The nitrogen component of nitride layers may be introduced from ammonia, as in most comparable MOVPE processes, although this approach requires special arrangements for pumping and cryogenic trapping of excess ammonia to maintain UHV conditions. It is

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more common to utilise plasma cells which generate active species from dinitrogen source gas. Several points of comparison between MOVPE and MBE for nitride materials apply equally to other inorganic semiconductors. For

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example, the compatibility of the UHV environment with electron beam propagation allows growing films to be interrogated by surface-specific diffraction methods and photoelectron spectroscopies, offering excellent insight

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into atomic-scale crystal growth mechanisms. This characteristic gives MBE advantages over MOVPE in materials research settings, although several factors favour MOVPE over MBE for commercial manufacturing. These include

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the relative lengths of downtime incurred by source changes and other hardware maintenance, the greater substrate capacity and scope for large-area deposition of modern MOVPE reactors, and somewhat slower growth rates in

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MBE of ~1m/h or lower. MBE also has difficulties in comparison to MOVPE with satisfactory nucleation of nitride growth on non-nitride substrates, for example by multi-stage processes discussed further in section 3.1 of

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produced by MOVPE or HVPE.

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this article. A common response is to perform MBE growth of device structures starting with nitride template layers

HVPE, in common with MOVPE, is a type of CVD process, and had considerable historical importance in producing samples for early physical and device-oriented studies of the group 13 nitrides [2,12,13]. The HVPE processes used for the nitrides evolved from those developed for other compound semiconductors, notably GaAs, and involve the transport of metals as volatile chlorides. Typically these are generated within the deposition system by passing hydrogen chloride over the molten metals. Ammonia is the usual nitrogen source. HVPE offers much faster growth rates for ntrides, typically in the range 10-100 m/h, than either MOVPE or MBE. However, the fast growth rates are incompatible with the abrupt interfaces needed in QWs and analogous advanced heterostructures. Practical HVPE systems use hot-wall tube furnaces with several independently controlled heating zones. Their construction limits the scope for rapidly switching different metal sources in and out of the deposition zone, as more easily achieved in MOVPE. Thus HVPE is poorly suited to direct deposition of multilayer device structures,

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but its fast growth rates make it well adapted to producing thick layers of material for use as templates in other epitaxy processes, or even cutting into bulk substrates [13].

The chemical reactivity and surface modification of the nitrides is outside the coverage of this article, but brief remarks are made here to guide the interested reader. The nitrides are well known as being resistant to traditional

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etching solutions, such that most etching in device fabrication uses energetic plasma chemistry converting the metal nitrides to volatile chlorides [17,18]. Nevertheless, considerable attention has been paid to developing wet etches

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for specialist purposes including controlled roughening, visualisation of dislocations, and selective removal of sacrificial layers [19,20]. Many of the more refined processes are photochemical or photoelectrochemical in nature,

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and anisotropic effects originating from the non-centrosymmetric crystal structure are often evident. Monolayer derivatisation of nitride surfaces with organic species has also attracted recent interest, and offers a route to novel

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chemical and biological sensors including enzyme-functionalised FETs [10]. Distinct routes used for organic functionalisation of GaN include surface hydroxylation followed by reaction with chlorosilane species [21], plasma

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hydrogenation of the surface followed by photochemical addition of an alkene [22], and surface chlorination followed by reaction with a Grignard reagent [23]. This field can be expected to develop significantly as part of the

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broadening scientific and technological interest in the properties of hybrid interfaces.

2. Basic properties of bulk materials and nanostructures

2.1 Crystal structures and physical properties of the binary compounds The group 13 nitrides have high ionicities compared to other binary compound semiconductors, and their stable crystal structure under normal ambient conditions is the hexagonal wurtzite phase [24]. Discussion later in this review generally concerns only wurtzite-phase material. However, the cubic zinc blende phase can be obtained under special conditions, and the best understood example is the stabilisation of cubic films of GaN within a narrow process window achievable in MBE processes [25,26]. AlN, GaN and InN also all undergo phase transitions directly from the wurtzite phase to the rocksalt phase at sufficiently high pressures [27].

Table 1 compares selected structural and physical parameters for wurtzite-phase AlN, GaN and InN. These data are mostly taken from the review by Wu, which also summarises recent values for many other physical constants [9].

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Wu expands upon on the major revision of the accepted fundamental bandgap of InN, which caused much debate in among physicists around 2002. Older literature generally quotes an InN nitride bandgap of 1.9 eV, inferred from optical absorption measurements on low-quality thin films. The revised value as per Table 1 derives from measurements on higher-quality InN films grown by MBE, and, crucially, has been confirmed by observations of bandedge luminescence. More detailed band structure parameters, which are critical for the application of the

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nitrides and their alloys in devices, were reviewed by Vurgaftman and Meyer soon after the general acceptance of the sub-1 eV InN bandgap [28]. Other focussed compilations of physical properties concern the parameters

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required to analyse spontaneous and piezoelectric polarisation fields, which arise because of the lack of a centre of symmetry in the P63mc wurtzite space group. Analysis of these polarisation effects relies heavily on atomistic

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modelling, such that numerical values are prone to revision as computational techniques improve. Bernardini provides a recent compilation and critique of earlier literature [29]. Ambacher gives an introductory explanation of

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how spontaneous and strain-induced piezoelectric fields combine to generate predicted total fields with magnitudes

discussed further in 2.4.

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in the MV cm-1 range [12]. The impact of these fields on the behaviour of planar QWs and other nanostructures is

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The non-centrosymmetric character of the wurtzite structure also has chemical consequences involving the

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differential reactivity of crystal planes. The origin of these so-called polarity effects is illustrated in Fig. 2, which is a schematic of local bond configurations near the surface of a GaN film on a substrate. Epitaxial nitride films are most often grown with {0001} basal crystal planes aligned parallel to the substrate surface, in which case the two non-equivalent orientations shown are possible. The situation in which the crystal vector [0001] runs from the substrate to the film surface is known as the +c or Ga-face polarity, although the latter term has no implication for the actual atomic termination [12]. The opposite orientation is termed the –c or N-face polarity. Much discussion in the current literature concerns the control of polarity in GaN films with the two alternative {0001} orientations, and the exploitation of the different chemical reactivity of Ga- and N-face surfaces [30]. However, it is certain that the polarity effects will have much broader and diverse impacts as the growth of nitride epitaxial structures in new orientations becomes more important [31,32].

2.2 Alloy systems and bandgap engineering

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The isomorphous nature of the group 13 nitrides leads to the expectation that they form ternary and quaternary alloys maintaining the wurtzite structure, and allowing the application of ‘bandgap engineering’ concepts developed for other compound semiconductors. Fig. 3 illustrates the single most important aspect of alloy formation, which is the ability to adjust the bandgap of a ternary alloy to any chosen value between the bandgaps of the constituent binaries. The bandgap of GaN corresponds to a photon wavelength of 363 nm in the near UV, such

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that InxGa1-xN alloys of smaller bandgap are the basis of devices required to emit or absorb visible light, while AlyGa1-yN alloys of larger bandgap are used in most UV emitters and detectors. Very importantly, the direct nature

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of the bandgap, allowing efficient electron-hole recombination to emit photons without any need for momentum exchange with phonons, is maintained for all possible alloy compositions. In contrast to the near-linear dependence

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of lattice constants on composition in alloy systems (Vegard’s law), the tie-lines connecting the binary compounds in Fig. 3 are non-linear. Although simple theory predicts a parabolic form involving a single bowing parameter

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[28], the exact behaviour of InN-rich nitride alloys in particular is open to question. Thus for example, the bandgap-composition plot in ref. [9] differs considerably from Fig. 3 in the InN-rich region. Although Fig. 3 does

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tayloring of bandgap and lattice constant.

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not represent the properties of quaternary compositions directly, these give additional scope for independent

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The assumption that nitride alloys will always form single homogeneous phases with tailored properties is questionable on thermodynamic grounds, and is linked to the disparity in the ionic radii in Table 1. However, in practice alloys are grown only by thin film deposition methods operating far from equilibrium, and the durations of growth processes may be short compared with those required for metastable layers to decompose to equilibrium phase mixtures. Issues of phase separation have been studied in particular depth for the technologically important alloy InxGa1-xN. In an early computational study, Ho and Stringfellow calculated phase diagrams and predicted wide fields under which binodal or spinodal decomposition of metastable alloys was predicted under film growth conditions [33]. Karpov performed a modified analysis in which (0001)-oriented InxGa1-xN layers were compressively strained to match the in-plane lattice constant of GaN [34], which is a valid picture for the growth of thin InxGa1-xN QWs used in practical devices. The addition of strain modified the region of single-phase stability dramatically for GaN-rich alloys, suggesting that most device structures would be grown under conditions of single-phase stability. Despite this background of computational predictions, the existence or otherwise of nanoscale composition modulation in InxGa1-xN layers has been a subject of ongoing debate. Since many of the key

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experimental observations derive from transmission electron microscopy on thin foil specimens, the susceptibility of InxGa1-xN to electron beam damage is a pertinent factor [35-37].

Alloys incorporating group 13 or 15 elements outside the Al-Ga-In-N system have attracted steady scientific interest, but are not yet employed in commercial devices. Introduction of even a few atomic percent of a heavier

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group 15 element as a substituent on the nitrogen sublattice causes very a rapid reduction in bandgap, although crystal growth is complicated by immiscibility gaps resulting from atomic size mismatches. A significant recent

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achievement was the first growth of homogeneous GaNxBi1-x layers with bandgaps as low as ~ 1 eV by Novikov et al. [38]. This work used a plasma-assisted MBE process to deliberately target growth of amorphous layers at

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temperatures below 100ºC. The same group has advanced the growth of GaNxAs1-x alloys, which can retain crystallinity for x values up to ~0.1, and provide an promising electrode material for photoelectrochemical water

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splitting [39]. Novel boron-containing alloys have recently been grown by MOVPE, with the motivation of reducing the refractive index of the parent binary nitride to fabricate multilayer distributed Bragg reflectors. This

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approach has been demonstrated both in the BxAl1-xN [40] and BxGa1-xN [41] systems. There is also much unexplored scope for investigating alloys containing isoelectronic substituent metals outside group 13, and to

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which the development of better sources for MOVPE might contribute. A recent demonstration of this type

sources [42].

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involved growth of ScxGa1-xN alloy layers with x up to ~0.02 using ammonia-based MBE with elemental metal

2.3 Dopants for conductivity control and advanced applications Much of the device technology developed with the group 13 nitrides depends on the ability to grow epitaxial layers with either n- or p-type electrical conductivity as desired. This so-called amphoteric doping is discussed here for GaN as a prototypical system before consideration of ternary alloys. Silicon, which gives n-type conductivity, and magnesium, which gives rise to p-type conductivity, are essentially the only elements in use for standard conductivity control. The sources used to introduce these dopants in MOVPE will be discussed in section 3. The current sub-section will also cover the d- and f-block elements that have been studied as more exotic dopants, with motivations including spintronics, and exploitation of the unique optical properties of the lanthanide dopants. While discussed here in the context of relatively mature epitaxial thin films, all the doping principles discussed apply equally to bulk crystals and nanomaterials.

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The population of vacancies and light-element impurities resulting from most crystal growth methods gives native n-type conductivity in the nitrides [9,43]. Carbon and oxygen are impurity elements of particular concern in MOVPE, and can be introduced either from precursors or from out-gassing of reactor components. The most rigorous technique for quantifying concentrations of these elements at the low levels at which they still may affect

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important material properties is secondary ion mass spectrometry (SIMS), and examples of impurity concentrations measured by this technique in optimised samples are given in 3.4. Preliminary optimisation of MOVPE processes

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typically involves reduction of unwanted impurity incorporation to acceptable levels, after which intentional doping is essential to obtain material with the free electron concentrations required for devices such as LEDs. Silicon

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substitutes on the group 13 lattice site, and in GaN the resulting donor state is within <20 meV of the conduction band edge [12]. Thus silicon acts as an almost ideal dopant, and the free electron concentration is comparable to the

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concentration of silicon atoms. A typical silicon atom concentration in a GaN-based LED structure is 5 x 1018 cm-3,

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corresponding to substitution of about one in every 9000 gallium atoms.

In contrast to the relative ease of n-type doping, early difficulties in achieving p-type conductivity in GaN were a

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major reason that its promise as an optoelectronic material was not realised before the 1990s. Akasaki, himself the

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major contributor to solving this problem, has given an excellent summary of the key advances [44]. A prerequisite for p-type conductivity was the refinement of MOVPE to the point where compensation of the free electrons associated with adventitious impurities and crystal defects became possible. Magnesium, like silicon, substitutes on the group 13 lattice site, and was an obvious choice as a p-dopant. However, its incorporation during MOVPE as an electrically inactive Mg-H complex necessitates further post-growth processing to achieve electrically active centres. Akasaki’s group initially used electron beam irradiation to dissociate the Mg-H complexes [44], but Nakamura et al. demonstrated that simple thermal annealing sufficed [45], and this has become the standard technological method (often incorporated into the cooling sequence inside MOVPE reactors). The properties of ptype GaN remain fundamentally limited by the position of the magnesium acceptor state ~160 meV above the GaN valence band edge [46]. This results in only ~1% of the magnesium atoms generating mobile holes at room temperature. Consequently magnesium atom concentrations in device structures exceed 1019 cm-3, and may even approach the solubility limit in GaN of ~1020 cm-3 [47]. Competitive p-type GaN may have a typical hole concentration of 5 x 1017 cm-3, and a conductivity of 1-2 -1cm-1 [46]. This relatively low conductivity arises from

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the fact that room-temperature hole mobilities in p-GaN layers are rarely above 20 cm2V-1s-1, or about an order of magnitude lower than electron mobilities in the n-GaN layers they are combined with in devices.

Significant attention has been paid to p-type doping of AlyGa1-yN and InxGa1-xN alloys, in which the activation energy for the magnesium acceptor increases with increasing alloy bandgap [43]. As a consequence it becomes

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increasingly difficult to obtain useful p-type conductivity in more AlN-rich alloys, and practical device structures often employ AlyGa1-yN-based superlattices rather than uniformly doped layers in sections where hole conductivity

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is needed. In contrast, magnesium-doped InxGa1-xN layers of moderate InN fraction can exhibit much higher hole concentrations than obtainable in p-type GaN, into the 1018 cm-3 range. However, films of magnesium-doped InN

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and more InN-rich alloys exhibit the peculiarity of a surface electron accumulation layer, which prevents conventional electrical contact to the underlying p-type material [9]. From spectroscopic measurements, Wang et

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al. inferred a magnesium acceptor activation energy of ~61 meV in InN [48]. Conductivity optimisation of n-type AlyGa1-yN layers is a particular requirement for UV-LEDs, and becomes more challenging as the bandgap increases

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[6]. Contributory factors at a microscopic level include an increase in impurity incorporation, crystal defect density, and mid-bandgap trap states as the AlN fraction increases [49]. Nevertheless, there have been a few independent

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reports of growth of silicon-doped AlN layers showing useful n-type conductivity [49,50].

Several of the lanthanides have attracted interest as dopants in group 13 nitrides with the motivation of exploiting visible or infrared emission resulting from intra-4f electronic transitions. Examples of elements studied are listed in Table 2, and more detailed discussions are available elsewhere [51,52]. A simplistic model might visualise Ln3+ ions substituting only on group 13 lattice sites. However, there is often spectroscopic and other experimental evidence for more than one coordination environment and electronic state for the dopant, and these systems have also proved very challenging as regards theoretical interpretation. A practical issue contributing to the complexity is that the lanthanide is often introduced through energetic ion implantation after crystal growth, and this process needs to be followed by thermal annealing to remove the lattice vacancies, interstitials and other defects it introduces [53]. Success has also been achieved introducing lanthanides during MOVPE, and two recent reports are now discussed as markers of progress towards device applications. Nishikawa et al. described a red-emitting LED, using n- and p-doped GaN layers similar to the schematic in Fig. 1 of this article, but substituting a layer of europium-doped GaN for the conventional QW-based active region [54]. The thickness of the europium-doped

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GaN was 300 nm, and the dopant atom concentration was 9 x 1019 cm-3. Narrow-band red emission occurred at ~621 nm, corresponding to the 5D0 to 7F2 transition of Eu3+. Erbium-doped materials are meanwhile of interest for amplification of near-infrared signals used in fibre-optic communications, exploiting the 4I13/2 to 4I15/2 transition of Er3+ at ~1540 nm. Dahal et al. demonstrated amplification in planar waveguides fabricated from erbium-doped

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GaN, with a reported erbium atom concentration of order 1021 cm-3 [55].

The availability of volatile sources, including cyclopentadienyl and carbonyl complexes, makes doping of nitride

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thin films with d-transition metals in MOVPE processes very feasible. The major motivation, which also applies to some of the lanthanide doping studies, has been the prospect of obtaining long range magnetic ordering [56,57]. If

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achievable at ambient temperature, ferromagnetic ordering would open up many new device engineering opportunities exploiting spin-polarised currents, and the term spintronics has been coined for this field. Long range

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magnetic ordering can be mediated in semiconductors with relatively low concentrations of spin-active centres by a double exchange mechanism involving free charge carriers [56]. The system studied in most detail has been

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manganese-doped GaN, in which theoretical work from 2001 predicted ferromagnetic ordering to above 300 K via hole-mediated interactions between Mn2+ (d5) centres, assumed to substitute on 5% of the gallium sites [58].

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However, relatively soon afterwards, Graf and co-authors published a more realistic and extended analysis

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concluding that co-existence of the required concentrations of Mn2+ and mobile holes was not possible [59]. Attempts at co-doping GaN in MOVPE with manganese and magnesium (required to introduce holes) also showed that the two dopants competed for incorporation, thus posing an additional practical challenge [60]. A further complication in characterising samples with spin-active dopants is their tendency to undergo phase separation. This leads to the risk that magnetic properties measured by conventional techniques are dominated by minority nanophases, and Bonanni discusses the appropriate advanced techniques for rigorous characterisation [56]. A distinct application in which doping with a transition element has found practical utility concerns iron doping of GaN buffer layers used in the growth of heterojunction FETs. The buffer layers between the substrate and the active regions of such devices are required to have the lowest possible conductivity, ideally lower than that arising from native n-type conductivity in GaN grown without any intentional doping. It has been shown that iron doping, which introduces mid-bandgap electron trap states, is an effective and reproducible way of meeting this requirement, at atom concentrations of up to 3 x 1017 cm-3 [61].

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2.4 Quantum confined nanostructures and polarisation fields In crystal growth of inorganic semiconductors, quantum confined nanostructures result when a material is prepared in such a way that at least one dimension approaches the exciton radius, listed in Table 1 for the binary nitrides. The conventional classification of quantum confined nanostructures divides them into two-dimensional quantum wells or QWs, one-dimensional quantum wires, also frequently called nanowires or nanorods depending on their

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aspect ratio, and zero-dimensional quantum dots [62]. Planar QWs in particular, but also nanowire and nanorod arrays, are readily produced via MOVPE. Synthesis methods for zero-dimensional nitride nanocrystals are less

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mature, and progress with colloidal and related preparation routes has been recently reviewed elsewhere [63,64]. However, nitride quantum dot arrays have also been prepared by variants of MOVPE, and in such work the final

us

aim is often to produce an electrically driven device containing a buried layer of quantum dots as its active region

an

[65,66].

Planar QWs require interfaces of atomic-scale abruptness, and feature layers of a higher bandgap alloy on either

M

side of the QW itself. These are termed barrier or confinement layers. It is generally assumed that the material of lower bandgap grows pseudomorphically, and is elastically strained such that its in-plane unit cell constant matches

d

that of the underlying material. In light-emitting devices such as LEDs the confinement of injected charge carriers

Ac ce pt e

into the QW regions produces major enhancements in efficiency compared to devices with bulk-like active regions. However, single QWs rarely give competitive devices, and QWs are therefore employed in stacks containing a few repeat units as exemplified in Fig. 1. Barrier layers between individual QWs three or four times the QW thickness fully confine the electron and hole wavefunctions, and allow the QWs to behave in a non-coupled fashion. Future technologies are likely to make increased use of hierarchical nitride nanostructures [16] which are discussed further in 3.7. Conformal growth of a thin layer of a lower-bandgap alloy on the sidewalls of a nanorod, followed by a capping layer, can produce a radial QW structure, analogous to a colloidal core-shell-system. A contrasting situation arises if the composition of a nanorod is modulated during axial growth to give thin regions of lower bandgap material. A sufficiently small nanorod diameter will give rise to additional in-plane quantum confinement effects, in which case the low-bandgap regions are frequently termed ‘quantum disks’.

The properties of real QWs grown in wurtzite-phase nitride materials are affected significantly by the polarisation fields running in the [0001] crystal direction, and the associated quantum confined Stark effect (QCSE) [12, 32].

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Most QW and barrier combinations involve an in-plane lattice mismatch, such that a strain-related piezoelectric field is present in addition to the spontaneous polarisation. Fig. 4 (b) illustrates the situation in which the polarisation field is perpendicular to the plane of the QW, which applies in most thin film crystal growth of nitrides, using the (0001)- or c-plane orientation. Here the QCSE causes spatial separation of the electron and hole wavefunctions, which reduces the efficiency of radiative recombination in light-emitting devices. The emission is

ip t

also red-shifted in comparison to the field-free (or so-called flat-band) situation. The effects of the QCSE can be mitigated by switching to alternative growth orientations [32,67]. The favourable situation with no perpendicular

cr

component of the net polarisation field, because the polarisation vector lies in the plane of the QW, is also

illustrated in Fig. 4 (a). This requires crystal growth in one of the two orientations shown in Fig. 5 and designated

us

as non-polar. However, it has proved practically easier to grow nitrides in semi-polar orientations, in which the wurtzite c-axis and polarisation field is inclined to the surface normal. Three examples of such orientations are

an

shown in Fig. 5. The refinement of epitaxial film growth in the non-(0001) orientations discussed has been the subject of major efforts in MOVPE in recent years. A principal motivation has been the development of LEDs and

Ac ce pt e

3. Development of nitride MOVPE

d

orientations also offer benefits for laser diodes [6].

M

laser diodes emitting at longer wavelengths (green and yellow) within the visible spectrum [67], although these

3.1 Substrates and growth initiation

Monocrystalline substrates are a requisite for epitaxy, and some key characteristics of these are surveyed before considering MOVPE processes in detail. The properties of substrates also have a strong bearing on device fabrication and performance, such that different substrates may be preferred for different applications. The technological development of the group 13 nitrides has differed from that of nearly all other semiconductor families in that epitaxy developed faster than bulk crystal growth. Consequently MOVPE has been developed to an advanced state using heteroepitaxy, employing sapphire (Al2O3) substrates in the majority of cases. Sapphire substrates of 2 inch (50.8 mm) diameter were adopted as an early standard, and many large-capacity industrial epitaxy tools still use multiple 2-inch substrates, despite the availability of larger sizes [68]. An expedient to realise some of the advantages of homoepitaxial growth has been the use of pseudo-substrates or templates of relatively thick GaN layers pre-grown on foreign substrates, often using HVPE [2,13,69]. Several commercial vendors have developed GaN templates on either sapphire or silicon, in which the threading dislocation density (TDD) is reduced

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to relatively low levels by proprietary growth methods. However, future advances are likely to exploit the increased availability and affordability of bulk nitride substrates, enabled by advances in crystal growth discussed shortly.

Sapphire and silicon will undoubtedly remain important substrates for many years in situations where cost is critical, and parameters relevant to the growth of GaN films in the standard (0001) orientation are summarised in

ip t

Table 3. The in-plane lattice mismatch of ~14% between GaN and sapphire involves a 30º rotation between the conventional unit cell axes. The highly developed growth initiation sequences for MOVPE of GaN typically

cr

involve the following steps: pre-growth treatment of the sapphire often involving nitridation, growth of a GaN or AlN nucleation layer a few tens of nanometres thick at relatively low temperature, a temperature ramp during

us

which the nucleation layer restructures into discrete islands, a period of controlled transition from 3- to 2dimensional growth, and gradual structural improvement through mutual annihilation of dislocations as the film

an

thickness increases. Such growth sequences receive extensive coverage in other publications [2,44,74,75], and key stages are illustrated in Fig. 6 with examples form the author’s laboratory. Although most GaN layers grown by

M

MOVPE have the Ga-face or +c polarity, suitable manipulation of sapphire nitridation conditions can produce smooth layers of the inverted N-face polarity [76]. The TDDs obtained after growth of a few microns of standard

d

(0001)-oriented GaN are still typically in the low 109 cm-2 range, or some 5-6 orders of magnitude higher than

Ac ce pt e

TDDs present in semiconductor device structures grown by mature homoepitaxial techniques. Various dislocation filtering and lateral overgrowth techniques have been developed to further lower TDDs in GaN, and are reviewed elsewhere [2,74]. Lateral epitaxial overgrowth typically involves partly covering a GaN seed layer on sapphire by a thin mask of non-crystalline SiO2 or silicon nitride patterned into a stripe array, and then performing a second growth stage to obtain more structurally perfect material. The process is illustrated schematically in Fig. 7. Such processes have received renewed interest in the context of non-(0001) oriented growth [77]. A related but simplified process, that is being increasingly used in LED development and manufacturing, uses sapphire substrates etched into a submicron periodic relief pattern [78]. Growth on these substrates is performed in such a way as to leave voids at the GaN/sapphire interface. Light scattering from the voids has benefits for the performance of suitable designed LEDs, while the structurally perfect GaN that forms during lateral overgrowth of void regions contributes to an overall improvement in structural quality.

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Two further consequences of the dielectric nature of sapphire are significant. Its insulating character means that two-terminal devices such as LEDs must have both n- and p-contacts on the top side of the device, in contrast to situations where epitaxy can be performed on a conducting substrate. However, the transparency of sapphire into the deep UV is useful in two respects. Firstly, this property allows light extraction through the sapphire backside of nitride LEDs, and favours their mounting in the so-called flip-chip configuration advantageous for thermal

ip t

management [3,5]. Secondly, the transparency into the deep UV underpins the laser lift-off process which allows sapphire substrates to be fully removed from suitably supported epitaxial layers. Laser lift-off is usually achieved

cr

by irradiating the sample through the sapphire with a scanned excimer laser, the emission wavelength of which is

us

strongly absorbed by GaN, causing photo-induced dissociation of GaN close to the interface with sapphire [79].

The principal attractions of silicon as a substrate are its low cost, and availability in large wafer sizes. Growth

an

initiation steps in MOVPE are quite distinct from the corresponding processes summarised for sapphire substrates and have been extensively optimised [71-73]. The first material deposited onto the silicon is invariably AlN, and

M

this may be followed by a compositionally graded AlyGa1-yN buffer layer, a sequence of thin low-temperature AlN inerlayers, or other AlyGa1-yN-based multilayers before growth of a GaN-based device structure. A critical role for

d

these growth initiation layers is the management of strain both during growth and after cooling; a key challenge

Ac ce pt e

here is the thermal expansion mismatch between silicon and GaN, which will cause cracking in GaN layers grown without the appropriate measures. Strain management during MOVPE processes now often uses real time optical monitoring, as discussed in 3.3.

Other substrates potentially useful in nitride heteroepitaxy are discussed in a review by Liu and Edgar [70]. Hexagonal SiC was important in early developments, although its in-plane lattice mismatch of ~3.5% with GaN (0001) is large in absolute terms, and its cost has remained relatively high. The high thermal conductivity of SiC, 4.9 W cm-1 K-1 versus 1.3 W cm-1 K-1 for GaN itself [69], led to some continued emphasis on its use as a substrate for high-power FETs for which heat dissipation is a key issue. For applications where thermal management is particularly critical, a newer approach of growing GaN on diamond substrates may well become important, having recently been demonstrated using MOVPE [80]. Many of the other substrates used in heteroepitaxy of nitrides are metal oxides, which are required to remain chemically stable under the high-temperature growth conditions. The lower process temperatures and less reducing conditions used in MBE versus MOVPE have led to a wider range of

17

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substrates being acceptable in the former type of process. A historically important example was the first successful growth of nitride QWs in a non-polar (m-plane) orientation using plasma-source MBE and (100)-oriented tetragonal -LiAlO2 substrates [81]. ZnO has also attracted steady interest because of its wurtzite crystal structure, in-plane lattice mismatch of <2% with GaN (0001), and the mature hydrothermal technology developed for growth of large bulk crystals. Although ZnO undergoes chemical reduction under the conditions of conventional nitride

ip t

MOVPE using ammonia, Ougazzaden and co-workers successfully grew GaN directly onto ZnO buffer layers using 1,1-dimethylhydrazine [82]. This work will be discussed later in the context of other non-traditional MOVPE

cr

processes.

us

The application of homoepitxay in MOVPE and other process has until recently been hindered by the limited availability and high cost of bulk nitride substrates. These factors have meant that their initial impact has been in

an

high-cost, high-performance device applications, notably laser diodes. Freitas reviewed recent developments that are making true bulk GaN and AlN substrates more widely available and utilised [13]. Early effort on growth of

M

bulk GaN crystals concentrated on high-pressure growth from nitrogen-saturated gallium melts, which produces crystals with an unfavourable platelet morphology. The technological demands of this process can be appreciated

d

from the 60 kbar pressure of N2 in equilibrium with GaN at its melting temperature of ~2220ºC [83]. HVPE

Ac ce pt e

processes offer much faster growth rates, and are suitable for producing GaN crystal boules with a [0001] growth direction. These have sufficient thickness along the growth axis to be cut into large-area substrates with arbitrary orientations to support epitaxy in the non-polar and semi-polar orientations introduced in 2.4 and Fig. 5. Crystal growth involving transport of material from a GaN powder source onto a seed crystal in a medium of supercritical ammonia has also advanced greatly, and is termed ammonothermal growth by analogy with hydrothermal processes used for growth of ZnO, SiO2 and other materials. This method has the favourable characteristic of being able to produce equiaxed GaN crystal boules, which can be cut into a variety of orientations with minimal wastage. AlN bulk substrates can also be produced by sublimation routes, and are a superior choice to bulk GaN for growth of AlyGa1-yN-based structures with high AlN fractions, exemplified by LEDs for emission into the deep UV.

3.2 Conventional precursors The great majority of nitride MOVPE processes use a standardised chemistry, involving reaction between group 13 alkyl vapours and ammonia as they are passed over a heated substrate in a flow reactor. This is the classic route to

18

Page 18 of 73

deposition of binary compound semiconductors pioneered by Manasevit from the late 1960s, and employed in one of his original studies to growth of AlN and GaN on various substrates [84]. The overall reaction, exemplified by the use of trimethylgallium (TMGa) to deposit GaN, may be represented as: Me3Ga + NH3  GaN + 3 CH4

(1)

Actual mechanistic details are discussed further in 3.5, but an important difference between nitride growth and

ip t

MOVPE processes using arsine and phosphine sources arises from the stronger Lewis basicity of ammonia. This increases the likelihood of parasitic reaction sequences, starting with adduct formation in the cooler regions of

cr

reactors [85]. To mitigate this effect nitride reactors use specific hardware designs intended to avoid premature

us

mixing of group 13 and 15 sources, as discussed in 3.3.

The purities of all precursors are critical for growth of device quality layers. Specialised purification and analysis

an

methods have been developed for the group 13 alkyls, which vendors supply in stainless steel bubbler units that can be attached to growth systems [86]. The metal alkyl vapours are transported into the growth chamber in a stream of

M

highly purified hydrogen or nitrogen carrier gas, nitrogen being preferred for growth of InN-containing alloys as discussed later. Controlled transport is achieved by metering the carrier gas flow passed into bubbler, controlling

d

the bubbler temperature with an external thermostat bath, and finally having active control of the pressure inside

Ac ce pt e

the bubbler (often at one atmosphere for convenience). Under these conditions the molar transport rate of the alkyl vapour is proportional to its vapour pressure, which is in turn controlled by the bath temperature. The main group 13 sources employed in nitride MOVPE are TMGa, triethylgallium (TEGa), trimethylaluminium (TMAl), and trimethylindium (TMIn), for which vapour pressure relations are summarised in Table 4. The use of TEGa over TMGa is preferred in some situations on the grounds that a decomposition pathway involving –hydrogen elimination may reduce carbon incorporation into films. While the first three compounds discussed are all liquids in the relevant temperature ranges, TMIn is a solid. Entrainment of vapour from a solid source is less reproducible than can be achieved when bubbling the carrier gas through a liquid. TMIn is therefore often used as a so-called adduct source, where the bubbler contains solid TMIn in equilibrium with its saturated solution in the liquid 1:1 adduct it forms with N,N-dimethyldodecylamine. The vapour pressure of TMIn over this system is essentially identical to that over the solid [87].

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Page 19 of 73

Conventional nitride MOVPE processes always require an extremely large stoichiometric excess of ammonia, with values >103 being universal [1,2,12]. This excess of ammonia is larger than in comparable processes for depositing arsenide or phosphide semiconductors, and can be reconciled with the relative thermal stabilities of NH3, PH3 and AsH3, and the large H-NH2 dissociation energy of 435 kJ mol-1 [88]. Practical consequences in nitride growth include that fact large quantities of ammonia are used, and impurities in ammonia may become the dominant

ip t

sources of impurities in the deposited layers. Water vapour is a particularly troublesome impurity, contributing to unwanted oxygen incorporation into films, but a recent study showed that it could be reduced below detection

cr

limits of a few parts per billion levels by appropriate online purification [89]. The thermal stability issue also means that exhaust gas streams from reactors contain large quantities of uncracked ammonia, such that removal of this

us

before discharge to the atmosphere can be a significant issue in production environments.

an

As noted earlier in 2.3, silicon and magnesium are the principal dopants used for conductivity control in nitrides, while there is more specialised interest in doping with d-block and f-block elements. The commonest silicon

M

sources for producing n-type layers are the gases SiH4 and Si2H6, which are used a pre-diluted mixtures in a balance of hydrogen [43,90]. However, the use of less hazardous alkyl-substituted silanes is possible, and demonstrations

d

include the use of di(tertiarybutyl)silane [90]. The usual magnesium source employed for p-type doping is

Ac ce pt e

bis(cyclopentadienyl)magnesium. In common with TMIn, this is a solid at practical utilisation temperatures, and is therefore sometimes used as a solution source. One commercial embodiment involves addition of the involatile branched-chain alkane squalane to generate a solution phase inside the bubbler [47]. Cyclopentadienyl complexes afford a useful family of precursors with suitable vapour characteristics for doping with first-row transition metals [60,61]. To date, the majority of doping studies with lanthanides use -diketonate sources, which are not ideal in view of their oxygen content. However, this issue has not prevented the successful demonstration of a red electroluminescent device using such an approach [54].

3.3 Reactor hardware and in situ monitoring This discussion will emphasise the features of advanced commercial reactors designed to produce multi-layer nitride device structures such as LEDs and laser diodes, although valuable materials research has often utilised simpler locally built equipment. Most nitride MOVPE processes are operated at sub-atmospheric pressures down to ~50 mbar. In this regime computational fluid dynamics (CFD) is invaluable to optimise reactor designs [91], and

20

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the integration of CFD with simulations of chemical processes in multi-scale modelling is discussed further in 3.5. A modern commercial nitride MOVPE reactor will embody all or most of the following features: Computer control, whereby complex multilayer structures can be grown according to a pre-written recipe



Gas switching using the vent-run principle, which is essential to produce atomically abrupt interfaces



A temperature capability of at least 1100ºC



Ability to switch carrier gas flows from hydrogen to nitrogen



Capacity for multiple 2-inch substrates, or smaller numbers of larger substrates



Rotary motion of the substrate holder or individual substrate to maximise layer thickness and

cr

ip t



compositional uniformity

Non-contact heating of the substrate holder (susceptor) by radio-frequency induction or incandescent lamps



Physically separated introduction of the group 13 sources and ammonia to minimise parasitic reactions



In situ process monitoring by optical methods, using light beams transmitted through purged windows



Substrate loading and unloading via a glovebox or airlock, ensuring that the growth chamber is not

d

routinely exposed to the outside atmosphere

M

an

us



Ac ce pt e

Fig. 8 shows details of a popular design of single-substrate research reactor with horizontal gas flow, in which a simple splitter plate is used to avoid premature reactions between the metal alkyls and ammonia. This type of reactor has featured in a number of studies comparing experimental growth results with computational simulations [92-95]. Its basic flow dynamics and principle of separate injection channels have been extended to multiwafer reactors, with capacities exceeding forty 2-inch substrates [96, 97]. In such systems, substrates are arranged on a circular rotating susceptor over which the gas flow is radial. Gas inlets are near the centre of the susceptor, and exhaust pumping ports are arranged circumferentially around its edge. Individual substrates, or groups of substrates, can be given additional satellite rotation, and this style of reactor has become known as the planetary type. A competing type of multi-substrate reactor is known as the close coupled showerhead (CCS) design [2,96,97]. In this arrangement, group 13 sources and ammonia are injected vertically through a large number of inlet ports distributed across the susceptor. A third variant of high-capacity reactor is a type in which the susceptor is rotated at much higher speed, of order 1000 revolutions per minute, than in the designs described so far. This type of reactor features distributed reagent injection from ports above the susceptor, in common with the CCS type.

21

Page 21 of 73

However, the injection ports are a larger distance from the susceptor, and the high rotation speed dominates flow patterns in the reactor through viscous drag effects [98].

In situ monitoring techniques most routinely applied nitride MOVPE processes use light beams, and can assess film growth rates, surface roughness evolution, substrate curvature, and surface temperature, all with commercially

ip t

available instruments. Some more specialised techniques applied to monitor gas-phase chemistry, which contribute to the understanding of reaction mechanisms, are mentioned in 3.5 and 3.6. The information obtained from the

cr

standard optical monitoring techniques is extremely valuable for optimising the growth of complex device

structures, and improves the efficiency with which expensive growth systems are utilised in manufacturing [99].

us

The most universal method measures the sample reflectance at one or more wavelengths, using a light source outside the reactor. The growth rates of microscopically smooth nitride films can be measured by exploitation of

an

thin film interference effects, by exact analogy with many other deposition processes. Here the heteroepitxial nature of most nitride growth is turned to advantage, as the discontinuity in refractive index at the nitride-substrate

M

interface gives it a large reflection coefficient [100]. A special advantage of reflectance monitoring for nitride systems derives from the fact that growth initiation sequences in heteroepitaxy involve periods of island growth,

d

followed by gradual coalescence and development of an optically smooth surface. The specular reflectance of the

Ac ce pt e

sample in the early growth stages is greatly affected by scattering effects, and reflectance versus time measurements provide a powerful method of tracking the microscopic surface morphology [2,100]. Fig. 9 shows an example of the reflectance evolution during GaN growth on sapphire. Substrate curvature can also be measured by arranging for several parallel light beams to be incident on the sample. Analysis of the positions at which the reflected beams impinge on an array detector enables the curvature to be calculated, and this technique has become particularly important for strain management during growth of GaN-based structures on silicon substrates [71-73]. Attempts to measure sample surface temperatures by pyrometry have also become increasingly significant in recent years [99]. These are motivated by a desire to improve on established temperature control methods, where the measurement used for feedback control is taken from part of the susceptor, often well away from any substrate position. One basic challenge in applying pyrometry to nitride growth is the transparency of sapphire substrates into the mid-infrared, such that many pyrometric instruments measure the temperature of the susceptor surface below the substrate. A second issue is the continuous change in emissivity of the samples as film growth proceeds, which can, however, be handled by appropriately sophisticated instrumentation [99]. Creighton et al. successfully

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demonstrated a truly surface-specific instrument operating in a mid-infrared wavelength range (7-8 m) where sapphire becomes opaque but the reactant gases present in MOVPE are transparent [101].

Attempts to understand the microscopic surface processes in nitride MOVPE have been less numerous than studies using the optical monitoring techniques discussed so far, all of which are easily applied in manufacturing to multi-

ip t

substrate growth systems. The gas pressures used in MOVPE prevent the application of techniques requiring electron beam propagation, such as the surface electron diffraction methods routinely used in MBE. However, the

cr

powerful technique of spectroscopic ellipsometry (SE) has been utilised in nitride MOVPE, although optical access requirements are more demanding than for reflectometry. This is because SE requires incident and reflected beams

us

at oblique incidence (typically ~60º), whereas reflectometry is generally performed at normal incidence. Examples of SE applied to nitride MOVPE include an early study following the various stages involved in initiating GaN

an

growth on sapphire [102], and simultaneous composition and growth rate measurements on AlyGa1-yN structures [103]. The process ambients used in nitride MOVPE are also transmissive to X-ray photons, and there have been a

M

few studies using in situ X-ray diffraction to study growing layers. Richard et al., for example, used a synchrotron source to follow the strain and compositional evolution of InxGa1-xN layers grown on GaN buffer layers,

Ac ce pt e

d

specifically analysing the onset of structural relaxation [104].

3.4 Growth conditions for different binaries and alloys Many nitride device structures such as the LED shown in Fig. 1 have GaN initiating growth on the substrate, and making up most of their total thickness, while much thinner alloy layers form the active region. Often the desired performance can be achieved with InxGa1-xN and AlyGa1-yN compositions where x and y are kept below ~0.25. Nonetheless, significant changes in MOVPE conditions in addition to the obvious switching of precursor flows are needed to achieve successful multilayer growth. Table 5 compares parameters used for growth of GaN itself and the three ternary alloys in the author’s laboratory, in a research reactor configured as shown in Fig. 8. Although the detailed optimisation of growth conditions is reactor-specific, the tabulated parameters illustrate important general trends. Growth of AlN and InN layers pose challenges distinct from those encountered in GaN growth, which also extend to growth of AlN-rich AlyGa1-yN and InN-rich InxGa1-xN layers respectively. Examples of specialised MOVPE approaches successful in these situations are discussed below, after an initial survey of GaN growth.

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GaN itself may be grown under three or more different conditions during deposition of a complex multilayer device structure. The thickest layers are deposited at temperatures sufficiently high for growth rates to be limited by the flux of the gallium source, ie. growth in this regime is transport limited. Kaluza et al. experimentally verified the minimal temperature dependence of growth rate using in situ reflectometry [105]. Silicon and magnesium doping is

ip t

also straightforward under these conditions [47,90]. The exact GaN growth parameters used with a particular reactor will derive from a global optimisation to obtain the best compromises between growth rate, surface

cr

morphology, impurity incorporation, and structural quality, the latter principally relating to the population of threading dislocations. The very low levels of unwanted light-element impurities obtainable with optimised

us

processes are well illustrated by SIMS results on thick GaN layers without intentional doping from the author’s laboratory, indicating the presence of carbon at 2-3 x 1016 atoms/cm3 and oxygen at 4-5 x 1016 atoms/cm3; these

an

concentrations are about two orders of magnitude below those required for intentional dopants such as silicon. Concerning structural quality, the microstructures of GaN films are strongly determined by the processes in the

M

early stages of high-temperature GaN growth, during which 3-dimensional islands must coalesce to obtain 2dimensional growth, and many dislocations are removed through mutual annihilation [2,74]. In this regard the ratio

d

of lateral to vertical growth obtained at the 3-dimensional island stage is very critical even in conventional (0001)-

Ac ce pt e

oriented growth, and principles developed for controlling microstructural evolution are being increasingly applied to non-polar and semi-polar GaN orientations [106]. Initial growth of GaN nucleation layers on sapphire substrates is invariably performed under low-temperature conditions (eg. 540-550ºC in Table 5), under which reaction kinetics limit growth rates, but the apparent activation energy may still be hardware dependent [105]. The slow growth rate is actually advantageous to allow precise thickness control of these layers to an optimum value, typically in the 20-30 nm range [75,100]. The low surface mobilities obtained under these conditions are also advantageous in depositing nanocrystalline GaN uniformly over the substrate surface. Finally, many device structures require thin GaN layers to be grown at intermediate temperatures corresponding to those used to grow InN-containing alloys. A common example is the insertion of GaN barrier layers between In xGa1-xN QWs. In the author’s laboratory this was achieved using conditions similar to those for growth of In xGa1-xN itself as per Table 5, at growth rates of ~100 nm/h, and well within the kinetically limited regime [107].

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Page 24 of 73

Growth of InN-containing alloys proved more challenging to optimise than that of GaN, and the low thermodynamic stability of InN is accepted as a fundamental issue. Attempts to determine the standard enthalpy of formation (Hf) of InN have given very scattered values, but a recent estimate of -36 kJ mol-1 [108] makes interesting comparison with a recent value of Hf for GaN of -165 kJ mol-1 [109]. The limited stability of InNcontaining alloys means that they must be grown at a comparatively low temperature, and with a large excess of

ip t

ammonia in order to maintain an adequate overpressure of active nitrogen species. It is also usual to minimise the amount of hydrogen present in the reactor by switching precursor carrier and auxiliary flows to purified nitrogen as

cr

per Table 5. This has the benefit of suppressing the reverse dissociation of ammonia, although there are additional reports that atomic and/or molecular hydrogen can etch InN-containing layers [110]. A study using grazing-

us

incidence X-ray fluorescence to study indium adsorption on GaN under realistic MOVPE conditions also showed introduction of molecular hydrogen increased indium desorption rates, and reduced its saturation coverage [111].

an

InxGa1-xN growth rates become kinetically limited under the conditions described, although this is not an intrinsic problem for QWs a few nanometres in thickness. The InN fraction in a deposited layer also decreases strongly with

M

increasing temperature for a given gas-phase composition. A common interpretation is that the kinetics of indium desorption govern the layer composition [94], although the possible additional role of gas-phase depletion

d

processes will be discussed in 3.5. The indium loss processes fortunately translate into a viable means of

Ac ce pt e

controlling InxGa1-xN compositions by variation of temperature alone, and representative data from the author’s laboratory are shown in Fig. 10. Further interpretation and computational simulation of these results is available in ref. [94], which also illustrates the similarity of trends observed in work from other groups. The methodology of using growth temperature to control InN fraction was successfully transferred to InzAl1-zN alloys, in which interest has grown in recent years, and has been used by multiple groups [95,112,113]. A lower growth pressure for InzAl1zN

compared to InxGa1-xN (eg. 75 versus 200 mbar in Table 5) is preferred to suppress the parasitic gas-phase

reactions which are an issue in growth of all AlN-rich alloys. The objective with InzAl1-zN-based structures is often to maintain z values of 0.17-0.18, corresponding to lattice matching with GaN on the (0001) basal plane. It is significant that growth of such AlN-rich alloys proves relatively straightforward, suggesting that indium atoms increase surface mobilities through surfactant behaviour, as also reported in growth of AlN/GaN superlattices [114].

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The binary InN has proved challenging to grow by MOVPE as smooth layers exhibiting intrinsic properties not dominated by defects, and the majority of InN samples used in scientific studies have been grown by MBE. By extrapolation of the conditions used for InxGa1-xN growth, the natural choice of growth condition for InN involves a combination of low and very large excess of ammonia, under which growth rates are inevitably somewhat low. Maleyre et al. reported early results on such conventional MOVPE growth in a temperature range of 550-625ºC and

ip t

with NH3/TMIn ratios of 5,000-36,000 [110]. Larger ammonia flows decreased growth rates, and the highest growth rate of 400 nm/h was obtained at the lowest NH3/TMIn ratio used. The same group has continued efforts to

cr

refine InN growth, and their observations include the superiority of an InN rather than a GaN nucleation layer for growth on sapphire (0001) substrates, a kinetically driven enhancement in growth rate by use of triethylindium as

us

an alternative to TMIn, and the effectiveness of CBrCl3 addition in obtaining smoother films [115]. The mechanism of the latter process is thought to involve formation of volatile indium halides, which competes with the

an

conventional MOVPE deposition chemistry in the sense of reducing vertical growth rates, but which enhances lateral growth rates. It was also shown that ammonia introduces hydrogen as an active n-dopant into InN layers, but

M

that it can be reversibly removed, in a manner analogous to the dissociation and re-formation of Mg-H complexes in magnesium-doped nitrides [116]. This factor has been one of the motivators for recent attention given to InN

Ac ce pt e

further in 3.6.

d

growth processes relying on generation of active nitrogen species using plasma sources [117,118], and discussed

Composition control inAlyGa1-yN layers is generally achieved by the classical MOVPE method of adjusting the ratios of gallium and aluminium sources in the gas phase. Rice and co-authors reported a functional dependence of the form

y = [TMAl]/[TMAl] + [TEGa]

(2)

where the parameter  was in the range 1.0 to 1.4, and depended only on ammonia partial pressure [119]. However, this study was made with a relatively low total system pressure (26 mbar) and ammonia partial pressures. A more general situation is one in which parasitic gas-phase reactions causing nanoparticle formation deplete the aluminium sources significantly, and result in a non-linear dependence of solid composition on precursor fluxes [85]. Other issues which become more important for alloys with y > 0.5 include low adatom surface mobilities, which require high growth temperatures of the in excess of 1200ºC, and increased incorporation of carbon and oxygen impurities. The latter issue is often linked to the oxophilicity of aluminium sources, and illustrative SIMS

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results from a set of Al0.52Ga0.48N layers grown under different conditions show oxygen concentrations from 3.8 x 1017 to 3 x 1018 atoms/cm3, and carbon concentrations from 6 x 1017 to 1.6 x 1018 atoms/cm3 [49]. The various challenges mentioned are taken to an extreme in growth of AlN itself, which is strongly preferred over GaN as a buffer layer for device structures using AlN-rich layers, because of its better lattice match and UV transparency. MOVPE processes which introduce the aluminium sources and ammonia in pulses of a few seconds duration have

ip t

had particular success in growth of high-quality AlN layers, example of which are now discussed. This approach was pioneered by Asif Khan’s group, which in early studies used non-overlapping TMAl and ammonia pulses, and

cr

regarded these processes as cases of atomic layer epitaxy, which would strictly imply self-limited adsorption of precursors and monolayer-by-monolayer growth [120,121]. Ref. [121] also emphasised an effective reduction in

us

the NH3/TMAl ratio from ~1,000 in conventional MOVPE to ~100 in the pulsed variant. Later work by the same group has featured more sophisticated precursor pulse sequences, for example allowing a degree of overlap

an

between ammonia and metal alkyl flows to adjust the relative contributions of different atomistic growth mechanisms [122]. The benefit of pulsed-flow techniques stressed recently by its originators has been enhancement

M

of surface migration, translating into growth of useful device material at lowered temperatures [7,122]. Similar mechanistic effects have been cited in pulsed MOVPE of AlN by other groups [123-125], and recent work has

d

combined pulsed precursor introduction with epitaxial lateral overgrowth methods to optimise the structural quality

Ac ce pt e

of thick AlN buffer layers for devices [126,127]. In this approach, a growth initiation layer is etched into an array of trenches, and the main growth is performed on this patterned template. Coalescence of features through lateral overgrowth typically leaves sealed void channels below the surface of the final film, as illustrated by Fig. 11. In contrast to related lateral overgrowth processes used in GaN growth, dielectric mask layers are avoided, as these would not be resistant to the higher process temperatures needed for AlN growth.

3.5 Mechanisms and growth simulations

The mechanisms of nitride MOVPE have intrinsic scientific interest, but the understanding obtained is increasingly used to support multi-scale simulations of production-scale reactors, which in turn can inform future process and hardware design. The current discussion will focus on recent developments in conventional nitride MOVPE using metal alkyl sources and ammonia, while some complementary investigations involving novel precursors are covered in 3.6. The review by Devi et al. inroduces much of the reaction chemistry and simulation principles applicable in nitride MOVPE [1], while Hirako et al. explicitly discuss the multiple reaction pathways that need to

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be considered [128]. Quantum chemical calculations continue to refine values for rate constants of relevant reactions [129-131], and so provide essential input parameters for process simulations. However, perhaps the most important advance in recent years has been an improved understanding of the so-called parasitic reactions that produce nanoparticles which are swept out of growth systems rather than contributing to useful film growth. This loss mechanism results from the large temperature gradients in the gas phase close the substrate position, such that

ip t

nanoparticles formed in the gas phase are repelled from the film surface by thermophoresis. Important experimental contributions have involved a stagnation flow reactor equipped with a laser scattering system, allowing the spatial

cr

distributions and characteristics of nanoparticles formed under different conditions to be studied in situ

[85,132,133]. Recent multi-scale computational simulations for production reactors have also begun to explicitly

us

include the formation of nanoparticles, which has particular impact on growth of AlN and AlyGa1-yN layers with high AlN fractions [134-136]. Insight into the superiority of pulsed MOVPE processes for AlN growth, as

an

discussed in 3.4, was also recently obtained by computational simulations including nanoparticle formation [137].

M

In the hypothetical absence of adduct formation, the important reaction pathways for group 13 trimethyl precursors would involve homolytic M-C bond cleavages, producing monomethyl radical species or ligand-free metal atoms

d

as the dominant species arriving at the surface of the growing films [1,128]. In the case of triethyl sources, -

Ac ce pt e

hydrogen elimination pathways producing metal hydride products can also contribute [1]. Subsequent reactions of species as mentioned with ammonia or NH3-x species are not rate-limiting in conventional models. In the more realistic scenario with adduct formation, a sequence of reactions as typified below for trimethyl sources must be considered:

Me3M + NH3  Me3M:NH3

(3)

Me3M:NH3  Me2M-NH2 + CH4

(4)

n Me2M-NH2  (Me2M-NH2)n

n = 2, 3....

(5)

This sequence of adduct formation (3), methane elimination (4) and oligomer formation (5) was demonstrated in a mass spectrometric studies on a flow-tube system using TMGa and NH3/ND3 [138], and the emphasis placed on it was also influence by the existence of the structurally characterised cyclic trimers such as (Me2GaNH2)3 [139]. Later studies using gas-phase infrared spectroscopy and metal alkyl concentrations more representative of MOVPE processes showed that both adducts Me3Ga:NH3 and Me3In:NH3 dissociate at temperatures above 200ºC, such that they could not play significant roles in film deposition or parasitic reaction pathways [140]. Enthalpies of formation

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for the adducts were estimated as -68 and -63 kJ mol-1 for Me3Ga:NH3 and Me3In:NH3 respectively. Me3Al:NH3 has a larger enthalpy of formation of ca. -113 kJ mol-1, but most critically undergoes fast methane elimination at temperatures around 250ºC, and undergoes facile dimerisation in the gas phase [141]. Further oligomerisation as per equation (4) is thus a plausible route to gas-phase clusters which act as nuclei for nanoparticles. Once formed, these nuclei may grow through a variety of microscopic processes, including direct depletion of non-associated

ip t

metal alkyls such as TMGa [132,135]. Precursor depletion through nanoparticle formation has most influence on growth of AlN and AlN-containing alloys, but also affects growth of GaN, InN and their alloys, in which

cr

nanoparticles form through radical chemistry. For example, a study of InN and InxGa1-xN deposition probed with in situ light scattering showed that nanoparticle formation could deplete almost 100% of the TMIn introduced within

us

experimental uncertainty at 800ºC, and the angular dependence of the scattering intensity was consistent with metallic character for the nanoparticles [133]. Fig. 12 illustrates the strong spatial localisation of the zone of

an

nanoparticle formation inside the reactor. In other work from the same group, nanoparticles formed under AlN, GaN and InN deposition conditions have been collected for ex situ transmission electron microscopy and

M

compositional analysis, verifying the presence of the relevant metal in each case [85].

d

Computational simulations of complete MOVPE processes can now be undertaken with commercial software

Ac ce pt e

packages, which must integrate handling of gas-phase reactions and key surface processes with macroscopic mass and heat transport phenomena specific to the detailed reactor geometry [91]. Published studies generally focus on comparisons between experimental measurements of the properties of deposited layers measured after growth with the predictions from simulations. Most of the approximations and simplifications necessary at present concern surface processes, which are very challenging to probe experimentally under MOVPE conditions [1]. Some appreciation of the increasing sophistication of simulation studies can be obtained by considering the scope of several publications already cited. The early studies focussed on single-substrate research reactors, initially calculating GaN growth rates and uniformity as a function of conditions and hardware geometry [92,93], and then considering compositions of InN-containing alloys [94,95]. Fig. 13 shows representative results on the predicted InN fraction as a function of growth setpoint temperature across epitaxial layers deposited on substrates of 50.8 mm diameter, using the methodologies from ref. [94]. Simulations of multi-substrate production reactors have similarly progressed from predictions of GaN growth rates and uniformity [96,97,134] to studies also predicting ternary alloy

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compositions and growth rates, specifically considering the gas-phase nucleation and growth of nanoparticles as a precursor loss mechanism in growth of AlyGa1-y N alloys [135].

3.6 Alternative precursors and innovative growth processes The deposition processes featured in the current and next sub-section do not necessarily produce epitaxial films, as

ip t

in true MOVPE, and the more general term MOCVD, as defined in section 1, will be employed henceforth where appropriate. Two undesirable features of conventional two-source MOVPE of nitrides are the high process

cr

temperatures, necessary to achieve adequate thermal cracking of ammonia, and the low efficiency with which ultrapure ammonia is utilised. Lower deposition temperatures are particularly desirable for growth of InN and InN-rich

us

alloy layers, as reviewed in 3.4, but would also allow the use of thermally sensitive substrates, and even embedding of thermally sensitive materials like colloidal quantum dots in epitaxial nitride structures [142]. The factors

an

mentioned have motivated investigations of more thermally labile nitrogen sources in two-source nitride MOVPE

M

since the 1990s, and also spurred studies of single-source precursors and plasma-based processes.

The various alternative nitrogen sources investigated for two-source growth include hydrazines, alkylamines,

d

nitrogen trifluoride, and hydrogen azide [1,2,86,143]. The only one of these sources to show realistic potential for

Ac ce pt e

producing device-quality layers routinely and safely has been 1,1-dimethylhydrazine (DMHy), which has acceptable properties for use on conventional MOVPE systems [144,145], and whose vapour pressure characteristics are given in Table 4. From experiments using mass spectrometry to study pyrolysis products formed in an MOVPE reactor, Bourret-Courchesne et al. concluded that thermal decomposition of DMHy started with cleavage of the N-N bond [145]. However, DMHy readily forms adducts with group 13 alkyls, such that some of the complications in ammonia-source MOVPE may not be avoided. One niche role for DMHy has been in the deposition of GaN directly onto ZnO buffer layers, which would not withstand the conditions of conventional growth [82]. This work builds on initial demonstrations of GaN growth using DMHy on GaN substrates, at temperatures as low as 550-690ºC, and using a DMHy/TMGa ratio of only ~10 [146]. Funato et al. performed a complementary study of growth of GaN and AlyGa1-yN on sapphire substrates [147]. Layers of GaN and AlyGa1-yN across the full composition range were deposited at 925ºC and a DMHy/Me3M ratio of 25. In the case of GaN, it was necessary to use hydrogen ambients to obtain smooth layers, and secondary ion mass spectrometry showed relatively high residual carbon concentrations at 1019 cm-3. InN growth was also attempted, but this material could

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not be obtained as phase-pure films. Another group succeeded in depositing InN films at temperature of 500-570ºC, and with DMHy/TMIn ratios of 100-200 [148]. However, this work used a special reactor design in which a capillary tube injected the DMHy close to the substrate position, which is unlikely allow uniformity of film properties across larger substrates.

ip t

By analogy with other compound semiconductors, single-source precursors for nitrides containing pre-formed M-N bonds have received considerable synthetic attention. The chemistry of candidate compounds is very rich [1,68,

cr

149-152], but viable precursors for MOVPE/MOCVD must possess a combination of volatility, strongly favouring mononuclear compounds with weak intermolecular interactions, and clean thermal decomposition pathways. These

us

factors have led to a focus on azido compounds as precursors to group 13 nitrides. The thermally labile character of the azido ligand offers an excellent source of active nitrogen species in the hot zone of a reactor. However, the

an

known explosive nature of many azido compounds and ionic azides (potentially present as impurities) means that the characteristics of new compounds must be assessed very stringently, particularly as they require heating to

M

obtain adequate vapour pressures for use [150]. Even in the absence of reported hazards, they may remain unattractive for industrial applications out of safety concerns. In view of the extensive coverage of nitride MOCVD

d

processes using single-molecule sources in previous reviews [1,64,143,150], the present survey will concentrate

Ac ce pt e

principally on ‘designed’ mononuclear azido compounds, which have most potential for specialist applications, and have been used in sophisticated film deposition studies in some cases.

The single-source precursor to GaN studied most extensively has been biasazido(dimethylaminopropyl)gallium, Me2N(CH2)3Ga(N3)2 , often abbreviated to BAZIGA. This compound provides a prime example of the success of intermolecular adduct formation in stabilising a mononuclear species [153], and its schematic structure is shown in Fig. 14 (a). The compound melts at ~35ºC, and can be distilled under vacuum at 150ºC [1]. Although the dimethylaminopropyl ligand offers the opportunity for Ga-C bond cleavage through -elimination, mechanistic investigations discussed shortly indicate this pathway is not significant for BAZIGA under MOCVD conditions. A comprehensive study was made of the deposition of GaN films from BAZIGA at a range of temperatures in a simple cold-wall reactor, both with and without the addition of ammonia [154]. Continuous crystalline layers showing the same triaxial orientation relationship to nitrided sapphire (0001) substrates as GaN layers grown by conventional two-source MOVPE were obtained without ammonia addition at 900ºC, and showed the expected

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band-edge photoluminescence, while non-crystalline layers were obtained below 700ºC. The issue of carbon incorporation from the precursor was addressed through the use of the highly sensitive techniques X-ray photoelectron spectroscopy (XPS). Although no quantitative estimates of residual carbon concentrations were made, there was a striking elimination of a detectable carbon signal from material deposited at 900ºC when even small additions of ammonia were made. Appropriate modification of the MOCVD conditions were able to produce

ip t

dense, oriented GaN nanopillar arrays, in which growth was catalysed by a gallium metal particle at the tip of each pillar [155]. In the same study, production of entangled networks of non-oriented GaN nanowires was

cr

demonstrated with the addition of hydrogen. Complementary work directed towards mechanistic understanding of the pyrolysis pathways of BAZIGA included the matrix isolation of Ga(N3) [156], and the identification of various

us

nitrogen-rich species of the general type HxGaNy, including GaN6, by a mass spectrometric technique sampling species formed near a hot substrate [157]. The matrix isolation studies provided evidence that the -elimination

an

pathway is insignificant in the pyrolysis of BAZIGA under MOCVD conditions, and it was further suggested that the rate of this process was affected by the other ligands present. Most unusually for a single-source ‘designed’

M

precursor, a full multiscale simulation of the type discussed in 3.5 has been performed for deposition of GaN from BAZIGA, and compared with experimental growth results [158]. This analysis gave an overall activation energy of

d

321 kJ mol-1 for GaN deposition, and suggested the rate limiting step was Ga-C bond cleavage. The precursor

Ac ce pt e

Et2Ga(N3)(MeHNNH2) used by Sung et al. provides an interesting example of incorporating a second type of thermally labile nitrogen-containing ligand [159]. This compound was used to deposit triaxially oriented layers of GaN on silicon (111) substrates at 800ºC. Analysis of volatile pyrolysis products by gas chromatography suggested hydrogen transfer from methylhydrazine to the ethyl ligands to form ethane and the imine MeNNH as products, thus indicating that the hydrazine N-N bond was not cleaved.

Paralleling the detailed investigation of BAZIGA as a precursor to GaN films and nanostructures, InN deposition from [Me2N(CH2)3]2In(N3) has been the subject of several reports. In the solid state, the compound is a coordination polymer with bridging azido groups [160]. However, chemical ionisation mass spectrometry indicates a monomeric gas-phase structure, with the presumed indium coordination environment shown in Fig. 14 (b) [161]. Film deposition studies showed that phase-pure continuous InN films could be obtained on nitrided sapphire substrates at 450ºC, but that clusters of InN nanowires (termed whiskers by the original authors) radiating from specific nucleation sites on the substrate, formed at lower temperatures [162]. Ammonia addition was found to be

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deleterious to film deposition, providing an interesting concurrence with observations in conventional two-source MOVPE of InN [110]. More systematic preparation of InN nanowires from [Me2N(CH2)3]2In(N3) used nonnitrided sapphire substrates [163]. A study of the pyrolysis products of [Me2N(CH2)3]2In(N3) under MOCVD conditions using mass spectroscopic sampling identified InN, HInN, HInN2 and InN3 as important species [161], in

ip t

general concurrence with studies by the same technique on the pyrolysis of BAZIGA.

Single-source MOCVD of AlN using azido precursors has a longer history than work on GaN and InN, and early

cr

studies include the use of the trimeric [Et2Al(N3)]3 [164]. Subsequently Fischer and co-workers prepared

monomeric [Me2N(CH2)3]2Al(N3) [165,166], an analogue of the indium source used in studies just discussed.

us

However, the crystal structure of Me2(N3)Al(H2NtBu) reported at the same time showed that a chelating ligand was not strictly necessary to obtain a monomeric compound [165,166]. Little appears to have been reported on AlN

an

films obtained from these azido compounds, although their volatilities are comparable to those of the analogous precursors to GaN and InN. In related work by the Fischer group, ligand-stabilised dialkyl aluminium amides of the

M

type Me2Al[NR´(CH2)2NR2] were prepared and used in film deposition experiments [167]. However, selfdecomposition of the compound with R´ = Et and R = Me gave only metallic aluminium films, showing that this

Ac ce pt e

d

class of compound is not an effective single-source precursor for AlN.

The relatively high stability of the adducts formed between aluminium alkyls and ammonia was remarked upon in 3.5, and leads to the possibility of using such adducts themselves, or oligomers formed according to reactions (3) to (5), as single-molecule precursors for AlN. Although this approach does not appear to being pursued actively now, past examples include the use of Me3Al:NH3 [168] and [Et2AlNH2]3 [164] for film deposition. A number of formally similar uses of adduct precursors formed from alane, AlH3, for nitride growth were reviewed by Jeiger and Gladfelter [169]. Derivatives of alane have attracted interest in MOVPE/MOCVD with the expectation that the absence of Al-C bonds will lead to reduced carbon incorporation in films compared to processes using alkyl precursors. The specific adducts used in nitride growth were H3Al:NMe3 and H3Al:NMe2Et, but these required cointroduction of ammonia in thin film growth, and so cannot be regarded as single-source precursors. A contrasting example of a single-source precursor used for AlN deposition is the hydrazidoalane [Me2Al--N(H)NMe2]2 which is dimeric in the solid sate with bridging nitrogens [170]. Polycrystalline AlN films with strong preferred orientation were deposited on silicon (111) substrates at 800ºC using this precursor [171].

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The final class of alternative nitride MOCVD processes now considered use physical means to generate active species from dinitrogen introduced into the growth system. This approach avoids the high cost and other drawbacks inherent in the use of ammonia, and has usually been applied in situations in which lowered deposition temperatures. Danielsson and Janzén made computational studies on the generation of active nitrogen species by

ip t

purely thermal means, and proposed a novel reactor configuration in which a N2/H2 mixture was preheated to 1700ºC [172]. However, the approach taken to dissociate N2 in practice has been plasma activation, and there are

cr

thus analogies with the routine use of plasma sources to generate active nitrogen species in nitride MBE [14]. For application in MOCVD processes, a so-called remote plasma configuration is common, in which the plasma is

us

generated in a flow tube upstream of the main reactor chamber. Corr et al. reported results using optical emission spectroscopy to characterise plasmas generated in such a system, and correlated the results with GaN film growth

an

characteristics [173]. This process used the conventional gallium source TMGa, but operated at a pressure of only ~4 mbar, lower than conventional nitride MOVPE processes, and a substrate temperature of 750ºC. In these

M

authors’ proposal concerning mechanisms, the activated diatomic species N2+ and N2* (where the * symbol indicates an electronically excited state) contributed only to GaN deposition, whereas the monoatomic species N,

d

N+ and N* could also cause etching. The addition of hydrogen to the plasma gas mixture was shown to have

Ac ce pt e

strongly beneficial effects on microstructure of deposited GaN layers, and also suppressed the intensity of emission observed from CN species. Two earlier reports discuss more diverse characterisation of GaN and InN layers obtained with a remote plasma system of the same type [174,175]. Significantly, these included the demonstration of triaxial orientation in GaN layers grown on ZnO-buffered glass substrates at temperatures at 650-670ºC [174]. A GaN thickness variation of less than 3% over glass substrates of 150 mm diameter was also noted [175]. A refined plasma-assisted process recently developed for InN growth allows rapid pulsing of both precursor introduction, and the plasma source [118]. Rapid pulsing of the plasma source allows active species to be delivered in the afterglow mode, under which ions and more damaging energetic species are removed by collisions before reaching the substrate, and N2* may dominate [176]. The pulsing of precursor introduction is meanwhile thought to enhance surface migration of adatoms as discussed earlier in the context of AlN growth.

3.7 Nanowire and nanorod growth

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The hexagonal crystal structures of the group 13 nitrides give the possibility of strongly anisotropic crystal growth, frequently involving fast growth along the unique c-axis to give nanostructures with a regular hexagonal cross section. In combination with situations where nucleation sites on a substrate are sparse, this can lead to the spontaneous formation of nanostructures rather than continuous films in simple MOCVD experiments. Examples include the studies on GaN and InN deposition from single-source precursors discussed in 3.6 [155,163]. However,

ip t

more systematic preparation of single-crystal nitride nanowire and nanorod arrays has developed into a substantial research area, motivated in particular by applications envisaged in novel forms of nano-structured LED [15]. The

cr

present survey will mainly draw on examples of growth of GaN and GaN-rich alloys, although there is parallel interest in nanostructures based on InN [9,11] and AlN [177]. Importantly, the application of sophisticated

us

MOCVD (or MBE) hardware to the growth of elongated nanostructures allows abrupt interfaces analogous to those used in planar device structures to be introduced, but in two distinct modes. Axial heterostructures feature

an

modulation of the composition along the growth axis, and can, for example, produce quantum disks as mentioned in 2.4, or distributed Bragg reflectors exploiting refractive index differences between successive layers. Radial

M

nanostructures are produced by a change in growth mode to deposit a material of different composition on the sidewalls of an already formed nanostructure, and are termed core-shell structures by analogy with colloidal

d

quantum dot systems. Growth of radial nanostructures in nitrides by MOVPE has evolved to a high degree of

Ac ce pt e

sophistication, such that core-shell structures exactly analogous to the planar LED structure shown in Fig. 1 were produced several years ago [178]. QWs in such core-shell structures also usually have the advantages of non-polar or semi-polar growth orientations as explained in 2.4.

Historical preparations of semiconductor nanowires (termed whiskers in the older literature) from the vapour phase were mediated by a molten nanoparticle at the tip of the growing nanowire, and this growth mode became known as the vapour-liquid-solid (VLS) mechanism [15,16]. Analogous MOCVD processes for forming GaN nanowires have been successfully developed by seeding a substrate surface with catalytic metal nanoparticles, most frequently nickel or gold [15]. In other situations, molten droplets of gallium may act as self-catalysts [155,179]. Process optimisation has been to a large extent empirical, although a recent study applied computational simulations to a nickel-catalysed MOCVD process for GaN nanowire growth in a tubular reactor [180]. Nanowire deposition was observed at low NH3/TMGa ratios of 33-67, in the zone where the calculated gas-phase temperature was in the range 580 to 790ºC, and concentrations of adduct-derived cluster species were relatively high. Further mechanistic

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insight into the catalyst’s role has been obtained from experiments using electron beam lithography to pattern a nickel-iron seed layer into features of controlled size [181]. This work gave proved quantitatively that the diameter of a catalyst particle governed the diameter of the nanowire whose growth it mediates. Post-process analysis of the catalyst particles involved in nickel-assisted MOCVD of GaN nanowires also indicated that the ordered solid phase Ni3Ga formed under MOCVD conditions, such that the growth process occurred through a vapour-solid-solid

ip t

mechanism similar to the classical VLS mechanism [182]. Disadvantages of catalyst-initiated nanowire growth include the possible need to remove the catalyst particles by chemical treatment before nanowires arrays can be

cr

processed into devices. In addition, the elements present in the catalyst particles may diffuse into the nanowires as unwanted impurities, and structural defects such as stacking faults may arise as a result of the growth mechanisms

us

supported by particular catalysts [183]. The application for which catalyst-initiated growth is perhaps most suitable is the production of arrays of vertical nanowires with relatively large spatial separations, which are subsequently

an

removed from the substrate, and manipulated by fluidic or similar positioning methods [16]. Superposition of horizontal nanowires running in two non-parallel directions allows the fabrication of ‘crossbar’ nano-devices, of

d

simple prototype [184].

M

which nano-LEDs made by the superposition of individual n- and p-doped GaN nanowires can be viewed as a

Ac ce pt e

There are two basic variants of catalyst-free growth of elongated nanostructures, respectively termed selective-area growth (SAG) and self-organised growth. The former technique will be discussed first, and gives the great advantage of allowing lithographic control over the positions where nitride growth is initiated. This enables for example, the fabrication of periodic functional structures such as photonic crystals [15]. SAG is also typically used to grow nanostructures with smaller length/width ratios than usual in catalyst-assisted processes, and these will be termed nanorods in the current discussion. Many embodiments of SAG of GaN nanorod arrays involve initial epitaxial growth of a GaN seed layer, which is coated by a thin dielectric mask, which is in turn patterned to create openings where further growth is to be initiated. Such a process is illustrated schematically in Fig. 15 (a). There are close parallels with the lateral overgrowth methods mentioned in 3.1, but with the significant difference that the GaN growth mode must be controlled to give fast vertical growth, combined with minimal lateral growth over the mask surface. Two-source MOVPE processes featuring pulsed alternating introduction of the metal alkyls and ammonia have proven effective in obtaining the required anisotropy. An initial report by Hersee’s group highlighted vertical growth rates over 2 m/h along the [0001] axis for GaN nanorods of diameter 220 nm [185]. It

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was further noted that the pulsed precursor introduction gave a low effective NH3/TMGa ratio, consistent with the observed requirements for successful nanowire growth under catalytic seeding conditions. Recent work by Yeh et al. using a similar pulsed MOVPE process produced highly regular GaN nanorod arrays as illustrated in Fig. 15, which were used to template the growth of non-polar InxGa1-xN/GaN QWs in core-shell mode [186]. Non-pulsed MOVPE has also been successful in SAG of GaN nanorod arrays, and the approach of Bergbauer and co-workers

ip t

produced N-polar nanorods with a [000] growth direction, initiating the GaN growth on nitrided sapphire without any pre-grown GaN seed layer [187, 188]. These studies investigated the effects of H2/N2 ratio in the

cr

reactor, and also the pitch of the mask openings. Below a critical H2/N2 ratio, pyramidal structures grew in conjunction with lateral overgrowth of the mask. However, with H2/N2 ratios above 1, nanorods with vertical side

us

facets formed [187]. Nanorods with the highest aspect ratio of ~8 were obtained with a H2/N2 ratio of 5, and analysis of the total volume of GaN grown with different masking patterns indicated that precursors adsorbed over

an

the entire mask surface contributed [188]. Successful core-shell growth of InxGa1-xN/GaN QWs was again achieved

M

[187].

For many device applications envisaged for nanorod and nanowire arrays, the mask patterning step needed for SAG

d

is unnecessary complication. Those growth methods characterised as self-organised involve controlled nucleation

Ac ce pt e

on a substrate surface without the use of a patterned mask, and have been developed in nitride MBE extensively, but much less widely in nitride MOVPE/MOCVD processes. One important issue is the directional supply of atom fluxes to the substrate in MBE that can favour anisotropic vertical growth [15]. A major recent advance in the sophistication of results obtained with self-assembled MOVPE of GaN nanorods was achieved by in situ deposition of an ultrathin (~2 nm) SiNx layer on sapphire (0001) substrates before commencement of the GaN growth [189]. Examples of the nanostructures obtained are shown in Fig. 16. The GaN core parts of these were produced using a three-stage process: initial SiNx deposition from a SiH4/NH3 mixture, which is a self-limiting process, a short GaN growth initiation step to promote some lateral growth from nucleation sites on the SiNx-coated sapphire, and the main GaN growth performed under conditions strongly favouring vertical growth [189]. The last step was performed with an NH3/TMGa ratio of ~15, and also with the addition of the dopant SiH4, which accelerates vertical growth. The GaN growth initiation time was found to control the structural quality and mean diameter of the nanostructures formed, and the latter was varied between 200 nm and 1.5 m. Aspect ratios exceeding 100 could be obtained. Following the initial growth and characterisation of core-shell non-polar InxGa1-xN/GaN QWs

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close to the tips of GaN nanorods [190], complete core-shell nano-LEDs with n- and p-doped regions have been demonstrated [191]. After detachment from the growth substrate, individual nano-LEDs were contacted at opposite ends to observe electroluminescence, and it was found that this could be selectively excited either from the polar or non-polar QWs grown on different facets.

ip t

4. Summary and outlook MOVPE of group 13 nitrides has developed since the late 1980s as the principal technique enabling the unique

cr

properties of these semiconductors to be exploited in mass-produced optoelectronic and electronic devices. Great success has been achieved in both manufacturing and research laboratories by adopting the classical MOVPE

us

chemistry applied to other compound semiconductor systems, involving the reaction of metal alkyls with ammonia. A major motivation for research into alternative processes has been the desire for reduced deposition temperatures,

an

most successfully addressed in research studies by employing single-source azido precursors, or by the use of 1,1dimethylhdrazine as a replacement for ammonia. However, other processes involving plasma activation of

M

dinitrogen and conventional metal alkyl sources are under development, and these may prove more attractive for

d

industrial uptake for applications such as nitride deposition on temperature-sensitive substrates.

Ac ce pt e

Particular challenges in conventional MOVPE of the nitrides include the need for heteroepitaxial growth in most cases, the high thermal stability of ammonia, and the occurrence of parasitic reactions that result in depletion of precursors through gas-phase nucleation and growth of nanoparticles. In response to the first of these issues, multistep sequences have been developed to initiate GaN growth on sapphire and silicon substrates, and real-time process control is greatly assisted by the now-routine use of in situ optical reflectometry. Homoepitaxial growth of nitrides is likely to become more widespread in the future, as a result of recent advances in bulk crystal growth of GaN and AlN. The mechanistic aspects of conventional two-source nitride MOVPE have been investigated by a variety of experimental means, and important recent advances include the use of light scattering methods to observe zones of nanoparticle formation. Mechanistic models have reached the state of maturity where they can be incorporated into multi-scale computational simulations of complete growth reactors, and simulations of this type will be increasingly used to optimise hardware and process design. They can already accurately predict, for example, the variations in growth rate and alloy composition as a function of position obtained in multi-substrate production reactors.

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The nitride devices already established in mass production, exemplified best by LEDs emitting in the blue-green part of the visible spectrum, use multilayer planar structures in which the thickest layers are comprised of GaN itself. These GaN layers are routinely doped with silicon, to obtain useful n-type conductivity, and with magnesium, to obtain p-type conductivity. Lanthanides and d-block transition elements may become increasingly

ip t

used as dopants for specialist purposes in the future. In the relatively thin functional layers of In xGa1-xN and AlyGa1-yN incorporated into current devices, x and y may be kept below ~0.25 in many cases. The established

cr

methodologies for compositional control in MOVPE involve growth temperature variation in the case of In xGa1-xN, and the more classical variation of gas-phase composition in the case of AlyGa1-yN. Layers of GaN and these alloys

us

can be grown with atomically abrupt interfaces as required to produce advanced quantum confined heterostructures. MOVPE of more InN- and AlN-rich ternary alloys, as well as the parent binaries InN and AlN, is

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the subject of much ongoing research motivated by new device applications. Advanced techniques applied to AlN

M

growth include pulsed precursor introduction to enhance surface migration, and maskless lateral overgrowth.

Direct growth of nanowire and nanorod arrays as an alternative to continuous monocrystalline layers has also

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become important in recent years. Suitable variation in growth conditions during an MOVPE process can produce

Ac ce pt e

radial or core-shell heterostructures with comparable sophistication to conventional planar devices. Such core-shell nanostructures provide an important means of producing QWs in non-polar and semi-polar crystal orientations, which mitigation the effects of polarisation fields. Future development of growth processes for elongated nanostructures can be expected to focus on the so-called self-organised variants, which avoid the use either of catalytic nanoparticles or lithographic patterning of a masking layer.

Acknowledgements

The author thanks the EPSRC for funding, and colleagues at Strathclyde University Institute of Photonics and Department of Physics for helpful discussions. Former PhD students Christopher Deatcher, Chaowang Liu and Katarzyna Bejtka are thanked particularly for their contributions to growth studies and film characterisation. SIMS analyses on the author’s samples were made by Dr Irwin Sproule of the Institute for Microstructural Sciences, NRC Ottawa.

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[153] A. Devi, H. Sussek, H. Pritkow, M. Winter, R.A. Fischer, Eur. J. Inorg. Chem. (1999) 2127. [154] A. Devi, W. Rogge, A. Wohlfart, F. Hipler, H.W. Becker, R.A. Fischer, Chem. Vap. Dep. 6 (2000) 245.

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[155] A. Wohlfart, A. Devi, E. Maile, R.A. Fischer, Chem. Commun. (2002) 998. [156] J. Müller, S. Bendix, Chem. Commun. (2001) 911.

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[159] M.M. Sung, C. Kim, S.H. Yoo, C.G. Kim, Y. Kim, Chem. Vap. Dep. 8 (2002) 50. [160] R.A. Fischer, H. Sussek, A. Miehr, H. Pritzkow, E. Herdtweck, J. Organomet. Chem. 548 (1997) 73. [161] J. Schäfer, J. Wolfrum, R.A. Fischer, H. Sussek, Chem. Vap. Dep. 5 (1999) 205. [162] R.A. Fischer, A. Miehr, T. Metzger, E. Born, O. Ambacher, H. Angerer, R. Dimitrov, Chem. Mater. 8 (1996) 1356.

[163] H. Parala, A. Devi, F. Hipler, E. Maile, A. Birkner, H.W. Becker, R. A. Fischer, J. Cryst. Growth 231 (2001) 68. [164] D.C. Boyd, R.T. Haasch, D.R. Mantell, R.K. Schulze, J.F. Evans, W.L. Gladfelter, Chem. Mater. 1 (1989) 119. [165] R.A. Fischer, A. Miehr, H. Sussek, H. Pritzkow, E. Herbtweck, J. Müller, O. Ambacher, T. Metzger, Chem. Commun. (1996) 2685.

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[166] R.A. Fischer, A. Miehr, O. Ambacher, T. Metzger, E. Born, J. Cryst. Growth 170 (1997) 139. [167] J. Khanderi, D. Rische, H.W. Becker, R.A. Fischer, J. Mater. Chem. 14 (2004) 3210. [168] A.C. Jones, S.A. Rushworth, D.J. Houlton, J.S. Roberts, V. Roberts, C.R. Whitehouse, G.W. Critchlow, Chem. Vap. Dep. 2 (1996) 5. [169] J.A. Jeigler, W.L. Gladfelter, Coord. Chem. Rev. 206-207 (2000) 631.

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[170] Y. Kim, J.H. Kim, J.E. Park, H. Song, J.T. Park, J. Organomet. Chem. 545-546 (1997) 99. [171] Y. Kim, J.H. Kim, J.E. Park, B.J. Bae, B. Kim, J.T. Park, K.-Y. Yu, Y. Kim, Thin Solid Films 339 (1999)

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[175] C. Martin, K.S.A. Butcher, M. Wintrebert-Fouquet, A. Fernandes, T. Dabbs, P. P-T. Chen, R. Carman, Proc. SPIE 6894 (2008) 689407.

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[176] P. Choi, Y. Kaufman, R. Aliaga, Appl. Phys. Lett. 57 (1990) 440. [177] Kendry, K.-T. Yong, S.F. Yu, J. Mater. Sci. 47 (2012) 5341.

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[179] V. Purushothaman, V. Ramakrishnan, K. Jeganathan, RSC Advances 2 (2012) 4802. [180] R.A. Burke, D.R. Lambourn, X. Weng, J.M. Redwing, J. Cryst. Growth 311 (2009) 3409. [181] B.S. Simpkins, P.E. Pehrsson, M.L. Taheri, R.M. Stroud, J. Appl. Phys. 101 (2007) 094305. [182] X. Weng, R.A. Burke, J.M. Redwing, Nanotechnology 20 (2009) 085610. [183] X. Zhou, J. Chesin, S. Crawford, S. Gradecak, Nanotechnology 23 (2012) 285603. [184] Z. Zhong, F. Qian, D. Wang, C.M. Lieber, Nano Lett. 3 (2003) 343. [185] S.D. Hersee, X. Sun, X. Wang, Nano Lett. 6 (2006) 1808. [186] T.-W. Yeh, Y.-T. Li, L.S. Stewart, P.D. Dapkus, R. Sarkissian, J.D. O’Brien, B. Ahn, S.R. Nutt, Nano Lett. 12 (2012) 3257. [187] W. Bergbauer, M. Strassburg, Ch. Kölper, N. Linder, C. Roder, J. Lähnemann, A. Trampert, S. Fündling, S.F. Li, H.-H. Wehmann, A. Waag, Nanotechnology 21 (2010) 305201. [188] W. Bergbauer, M. Strassburg, Ch. Kölper, N. Linder, C. Roder, J. Lähnemann, A. Trampert, S. Fündling, S.F. Li, H.-H. Wehmann, A. Waag, J. Cryst. Growth 315 (2011) 164.

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[189] R. Koester, J.S. Hwang, C. Durand, D. Le Si Dang, J. Eymery, Nanotechnology, 21 (2010) 015602. [190] R. Koester, J.-S. Hwang, D. Salomon, X. Chen, C. Bougerol, J.-P. Barnes, D. Le Si. Dang, L. Rigutti, A. de Luna Bugallo, G. Jacopin, M. Tchernycheva, C. Durand, J. Eymery, Nano Lett. 11 (2011) 4839. [191] G. Jacopin, A. de Luna Bugallo, P. Lavenus, L. Rigutti, F.H. Julien, L.F. Zagonel, M. Kociak, C. Durand, D.

ip t

Salomon, X.J. Chen, J. Eymery, M. Tchernycheva, Appl. Phys. Exp. 5 (2012) 014101.

Figure captions

cr

Fig. 1 – (a) Schematic cross section showing the functional layers required in a GaN-based LED; (b) cross section of a processed LED showing metal contacts and etching down to the n-type layer; (c) illuminated bare LED with

us

voltage applied via wire bonds to n- and p-contacts.

Fig. 2 – Bond orientations at the film-substrate interface for Ga- and N-face polarities of GaN films. Reproduced

an

with permission from O. Ambacher, J. Phys. D: Appl. Phys. 31 (1998). Copyright © 1998 Institute of Physics Publishing.

M

Fig. 3 – Bandgaps as a function of composition for wurtzite (hexagonal) and zinc-blende (cubic) nitride semiconductor alloys. Reproduced with permission from I. Vurgaftman, J.R. Meyer, J. Appl. Phys. 94 (2003) 3675.

d

Copyright © 2003 American Institute of Physics

Ac ce pt e

Fig. 4 – Schematic representation of band profiles and electron and hole wavefunctions in (a) non-polar and (b) polar quantum wells.

Fig. 5 – Schematic of important polar, non-polar and semi-polar crystal planes in hexagonal GaN. Reproduced with permission from T. Paskova, phys. stat. sol. (b) 245 (2008) 1011. Copyright © 2008 John Wiley and Sons. Fig. 6 – Micrographs illustrating the typical morphology evolution of GaN grown on sapphire (0001): (a) obliqueview scanning electron micrograph showing the nucleation layer at the point of maximum roughness after annealing; (b) topographic atomic force microscope image showing the smooth surface obtained after hightemperature growth and island coalescence (step height between atomically smooth terraces c/2, ~0.3 nm); (c) cross-sectional transmission electron micrograph illustrating the propagation of some threading dislocations into the main part of the film, and the more defective zone near the GaN-sapphire interface. Part (c) courtesy of Dr D. Zhu, University of Cambridge. Fig. 7 – Schematic sequence of steps in a typical lateral epitaxial overgrowth process for GaN, shown as crosssections across the width of mask stripes, whose height is greatly exaggerated for clarity. Threading dislocations

49

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approaching an inclined free surface are bent towards the horizontal (steps 3 and 4), and can also undergo mutual annihilation (step 5). The end result is a planar GaN template with stripe regions of greatly reduced threading dislocation density compared to the initial seed layer. Fig. 8 – Schematic longitudinal section of the deposition cell of the widely used Aixtron 200/4 RF-S nitride MOVPE system. Process gases are confined within a quartz inner cell of rectangular cross-section, itself housed

a glovebox beyond the bulkhead seen at the right-hand side of the photograph.

ip t

inside the metal outer casing (length ~0.5 m) shown in the photograph inset. Loading and unloading is conducted in

cr

Fig. 9 – Annotated plot of the normal-incidence reflectance versus time during growth of a GaN layer on sapphire (0001) in the author’s laboratory. The period of near-zero reflectance around 2900 s corresponds to the point of

us

maximum roughness shown in Fig. 6 (a). The GaN growth rate during the main growth, extracted from the period of the reflectance oscillations, was 1.85 m/h.

an

Fig. 10 – Plot of InN mole fraction x in InxGa1-xN layers versus setpoint temperature for the conditions summarised in Table 5, derived from analyses by two methods unaffected by the strain state of the alloy layer. The dashed line

M

is a guide to the eye.

Fig. 11 – Cross-sectional scanning electron micrographs of coalesced AlN grown on a stripe-patterned seed layer

d

using pulsed MOVPE. Part (a) shows the result of single-stage overgrowth at 1400ºC, while part (b) shows the

Ac ce pt e

result of two-stage overgrowth started at 1300ºC. Reproduced with permission from V. Kueller, A. Knauer, C. Reich, A. Mogilatenko, M. Weyers, J. Stellmach, T. Wernicke, M. Kneissl, Z. Yang, C.L. Chua, N.M. Johnson, Photon. Technol. Lett. 24 (2012) 1603. Copyright © 2012 IEEE. Fig. 12 – Laser light scattering from localised layers of particles formed in InN MOVPE in a stagnation point flow geometry (a) in a H2/NH3 ambient, and (b) in a N2/NH3 ambient, both near 800ºC and at 263 mbar pressure. The ambient gas strongly affects the distance of the nanoparticle layer relative to the substrate heater (outlined). Reproduced with permission from J.R. Creighton, M.E. Coltrin, J.J. Figiel, Appl. Phys. Lett. 93 (2008) 171906. Copyright (c) 2008 American Institute of Physics. Fig. 13 – Calculated InN mole fraction in InxGa1-xN layers across a rotating 50-mm substrate as a function of growth setpoint temperature. Courtesy of Dr E. V. Yakovlev, Soft Impact Ltd, St. Petersburg. Fig. 14 – Schematic molecular structures of (a) Me2N(CH2)3Ga(N3)2 , BAZIGA, and (b) [Me2N(CH2)3]2In(N3). Fig. 15 – (a) Schematic of the formation of vertical GaN nanorod arrays by SAG; (b) oblique- and plan-view scanning electron micrographs of a typical nanorod array labelled to show different crystal facets; (c)-(f) oblique-

50

Page 50 of 73

view scanning electron micrographs of nanorod arrays grown on respective pitches of 250, 500, 750 and 1000 nm. All scale markers are 500 nm. Reproduced with permission from T.-W. Yeh, Y.-T. Li, L.S. Stewart, P.D. Dapkus, R. Sarkissian, J.D. O’Brien, B. Ahn, S.R. Nutt, Nano Lett. 12 (2012) 3257. Copyright © 2012 American Chemical Society. Fig. 16 – (a) Oblique-view scanning electron micrograph of GaN-based nanorods produced by a self-organised

ip t

process; (b) schematic of the core-shell multiple QW structure grown near the end of the nanorods; (c) highermagnification scanning electron micrograph of an individual nanorod viewed perpendicular to the growth axis.

cr

Reproduced with permission from R. Koester, J.-S. Hwang, D. Salomon, X. Chen, C. Bougerol, J.-P. Barnes, D. Le

Ac ce pt e

d

M

an

(2011) 4839. Copyright © 2011 American Chemical Society.

us

Si. Dang, L. Rigutti, A. de Luna Bugallo, G. Jacopin, M. Tchernycheva, C. Durand, J. Eymery, Nano Lett. 11

51

Page 51 of 73

*Highlights (for review)

   

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ce pt

ed

M

an

us

cr

ip t



GaN-based semiconductor alloys are the key components of many modern optoelectronic devices Metal organic vapour phase epitaxy has become the dominant method for preparing these materials Two-source processes using metal alkyls and ammonia have proven highly successful Ongoing research challenges with planar films include InN and AlN deposition There is growing interest in growth of nanostructures, such as nanorod arrays

Page 52 of 73

cr

i

Figure(s)

(a) Active region (see expansion)

n-type GaN

InxGa1-xN quantum well (typical 3 nm)

ed

Sapphire substrate

GaN barrier (typical 7-10 nm)

M an

Typical 4-6 m

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p-type GaN (typical 130-200 nm)

p-contact pad

Semitransparent electrode over p-type GaN

(c)

Ac

(b)

ce pt

High bandgap electron blocking layer (typically p-type AlyGa1-yN)

n-contact pad

350 m

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Ac

ce

pt

ed

M

an

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cr

i

Figure(s)

Page 54 of 73

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ce

pt

ed

M

an

us

cr

i

Figure(s)

Page 55 of 73

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cr

i

Figure(s)

M an

(a) Non-polar QW e-wavefunction

(b) Polar QW

e-wavefunction

ce pt

ed

Energy

Conduction band profile

Valence band profile

h-wavefunction

h-wavefunction

Ac

Distance along growth direction

Page 56 of 73

Ac

ce

pt

ed

M

an

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cr

i

Figure(s)

Page 57 of 73

cr

i

Figure(s)

(a)

M an

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(b)

ce pt

ed

1 mm

Threading dislocations running vertically

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(c)

GaN/sapphire interface Page 58 of 73

Figure(s)

2. Selective area growth

1. Mask layer (normally

only in mask openings

a stripe array)

ip t

Dislocation-rich GaN seed layer

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cr

Sapphire substrate

4. Continued lateral growth

3. Lateral overgrowth

results in coalescence

pt

ed

M

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of mask begins

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5. New defects may form

6. Surface almost planarised

Ac

at coalescence fronts

Page 59 of 73

Metal alkyl flow

ed

Ammonia flow

Water-cooled quartz Hole for susceptor plate, fixed reflectometer to loading door beams Substrate (50 mm f) Inner quartz cell, ceiling part

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Graphite susceptor disks, with gas-foil rotation

M an

Reagent separation plate

cr

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Figure(s)

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Frosted quartz bowl supporting graphite disks

Flat induction heating coil

Exhaust gas to rotary pump

Pyrometer lightpipe, sighting through hole in quartz liner

Page 60 of 73

cr

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Figure(s)

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TMGa flow restarted for main GaN growth

0.3

an

0.35

Optically smooth film supports successive constructive and destructive interference conditions as thickness increases

TMGa flow stopped; temperature ramp started

M

0.25

ed

0.2

pt

0.15

0.05

ce

0.1 Start of nucleation layer growth

0 2000

Ac

Calibrated reflectance at 600 nm

0.4

3000

TMGa flow stopped to end growth

4000

5000

6000

7000

Elapsed process time (seconds)

Page 61 of 73

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Figure(s)

cr us an

0.25

Electron probe microanalysis

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0.2

Rutherford backscattering

ed

0.15

ce

pt

0.1

0.05

0 740

Ac

Measured InN mole fraction x

0.3

760

780

800

820

840

860

880

900

920

940

Growth setpoint temperature (deg C)

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ce

pt

ed

M

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Figure(s)

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ce

pt

ed

M

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Figure(s)

Page 64 of 73

-0.03

ed

M an

o

T=760 C

o

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0.23 0.22 0.21 0.20 0.19 0.18 0.17 0.16 0.15 0.14 0.13 0.12 0.11

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InN mole fraction x

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Figure(s)

-0.02

-0.01

T=800 C o

T=820 C o

T= 832 C 0.00

0.01

0.02

0.03

Distance from centre (m)

Page 65 of 73

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cr

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Figure(s)

CH3 H 3C

(b)

M an

(a)

N3

ce pt

Ga

ed

N

In N N

CH3 CH3 CH3

CH3

Ac

N3

N3

Page 66 of 73

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ce

pt

ed

M

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Figure(s)

Page 67 of 73

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te

d

M

an

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cr

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Figure(s)

Page 68 of 73

Table(s)

Table 1 – Properties of the binary group 13 nitrides at 300 K, after ref. [9] unless otherwise noted GaN

InN

Lattice constant a (nm)

0.3112

0.3189

0.3533

Lattice constant c (nm)

0.4982

0.5185

0.5693

Ratio c/a (cf. 1.633 for ideal

1.6009

1.6259

Fractional atomic coordinate u (cf. 0.375 for ideal wurtzite) M3+ ionic radius (nm)

0.3821a

0.3770a

0.039c

0.047c

Density (g cm-3)

3.23

6.15

6.81

Fundamental bandgap (eV)

6.14

3.43

0.64

Exciton binding energy (meV)

60

34

9

Exciton Bohr radius (nm)

1.4

2.4

8

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cr

0.3787b 0.079c

Experimental value from ref. [29]

b

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a

1.6114

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wurtzite)

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AlN

Calculated value from ref. [29]

ce pt

From ref. [12]

Ac

c

Page 69 of 73

Table(s)

Table 2 – Examples of lanthanides used to dope group 13 nitrides Principal motivation

Thulium, Tm

Blue emission from intra-4f transitions

Praseodymium, Pr

Green emission from intra-4f transitions

Europium, Eu

Red emission from intra-4f transitions

Erbium, Er

Amplification at ~1540 nm from intra-4f transitions

Gadolinium, Gd

High local magnetic moment expected from f7 configuration of Gd3+

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ce pt

ed

M

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cr

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Element

Page 70 of 73

Table(s)

Table 3 – Comparative properties of sapphire and silicon substrates for growth of (0001)oriented GaN; data from refs. [69] and [70] unless stated

Silicon

Substrate plane considered

(0001)

(111)

In-plane orientation relationship

GaN

[Al2

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Sapphire

O3 GaN [ Si [a

cr

[ -16.9a

In-plane lattice mismatch with +13.9 GaN (%)b In-plane thermal expansion 1.34 coefficient as a fraction of that of GaNc Melting point (ºC) 2030

Electrical resistivity ( cm)

>1011

us an

8.1-8.6

1414 1.1

Up to 5 x 104 if undoped

M

Bandgap (eV)

0.46

a

From ref. [72]

b

150 mme

ed

Wafer sizes demonstrated in 100 mmd GaN growth

c

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Positive for GaN repeat larger than substrate repeat

Value >1 inducing compressive film strain during cool down

d

Example from ref. [68]

Example from ref. [73]

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e

Page 71 of 73

Table(s)

Table 4 – Vapour pressure relations for important precursors, of the form log10 P [Torr] = A – B/T[K]

A

B (K)

P at 20º (Torr)

Reference

Trimethylgallium

8.07

1703

181

86

Triethylgallium

9.165

2503

4.2

Trimethylaluminium

8.224

2135

8.7

Trimethylindium

10.52

3014

1.7

1,1-dimethylhydrazine

8.19

1780

130

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Compound

86

86

145

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ce pt

ed

M

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cr

86

Page 72 of 73

Table(s)

Table 5 – Comparison of GaN and ternary alloy growth conditions in work at Strathclyde University Institute of Photonics, with TMGa, TMAl, TMIn and NH3 sources

Molar ratio NH3/total group 13 2050

GaN nucleation Layer Standard undoped GaN InxGa1-xN

540-550

1100-1170

H2

200

1700

760-920a

N2

200

6200

AlyGa1-yNb

1170

H2

75

960

InzAl1-zN

760-860c

N2

75

2500

Typical growth rate (nm/h) 350

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H2

Reactor pressure (mbar) 200

2000

cr

Carrier gas

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Temperature range (ºC)

100

1800 60

an

Material

a

Composition variation in x from 0.02 to 0.24 achieved through temperature variation with constant [TMIn]/[TMIn]+[TMGa] ratio of 0.4 b

M

Composition variation in y from 0.025 to 0.22 achieved by varying [TMAl]/[TMAl]+[TMGa] ratio up to 0.45 c

Ac

ce pt

ed

Composition variation in z from 0.11 to 0.24 achieved through temperature variation with constant [TMIn]/[TMIn]+[TMAl] ratio of 0.4

Page 73 of 73