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Optoelectronic device physics and technology of nitride semiconductors from the UV to the

terahertz

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2017 Rep. Prog. Phys. 80 106501

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10.1088/1361-6633/aa7bb2

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Reports on Progress in Physics

Optoelectronic device physics and technology of nitride semiconductors from the UV to the terahertz

Theodore D Moustakas and Roberto Paiella

Department of Electrical and Computer Engineering, Division of Materials Science and Engineering, Photonics Center, Boston University, Boston, MA 02215, United States of America

E-mail: tdm@bu.edu

Received 16 December 2016, revised 6 June 2017Accepted for publication 26 June 2017Published 1 September 2017

Corresponding Editor Professor Masud Mansuripur

AbstractThis paper reviews the device physics and technology of optoelectronic devices based on semiconductors of the GaN family, operating in the spectral regions from deep UV to Terahertz. Such devices include LEDs, lasers, detectors, electroabsorption modulators and devices based on intersubband transitions in AlGaN quantum wells (QWs). After a brief history of the development of the field, we describe how the unique crystal structure, chemical bonding, and resulting spontaneous and piezoelectric polarizations in heterostructures affect the design, fabrication and performance of devices based on these materials. The heteroepitaxial growth and the formation and role of extended defects are addressed. The role of the chemical bonding in the formation of metallic contacts to this class of materials is also addressed. A detailed discussion is then presented on potential origins of the high performance of blue LEDs and poorer performance of green LEDs (green gap), as well as of the efficiency reduction of both blue and green LEDs at high injection current (efficiency droop). The relatively poor performance of deep-UV LEDs based on AlGaN alloys and methods to address the materials issues responsible are similarly addressed. Other devices whose state-of-the-art performance and materials-related issues are reviewed include violet-blue lasers, ‘visible blind’ and ‘solar blind’ detectors based on photoconductive and photovoltaic designs, and electroabsorption modulators based on bulk GaN or GaN/AlGaN QWs. Finally, we describe the basic physics of intersubband transitions in AlGaN QWs, and their applications to near-infrared and terahertz devices.

Keywords: nitride semiconductors, LEDs, electroabsorption modulators, lasers, Intersubband devices, photodetectors

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Review

IOP

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1361-6633

Contents

1. Introduction and brief history of the field ......................2 2. Crystal structure and basic properties ............................3 3. Heteroepitaxial growth ...................................................5 4. Polarization effects .........................................................7 5. Device fabrication ..........................................................9

6. Light emitting diodes ...................................................11 6.1. Visible spectrum LEDs ........................................11 6.1.1. Potential origins of the high

IQE of blue LEDs ...................................12 6.1.2. Potential origins of the ‘green gap’

and solutions ...........................................12 6.1.3. Efficiency droop .. .........................................13

1361-6633/17/106501+41$33.00

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1. Introduction and brief history of the field

Over the past 40 years a new family of semiconductors, con-sisting of the binary compounds InN, GaN, AlN and their alloys, has emerged as one of the most important classes of semiconductor materials. All these materials are direct band gap semiconductors and, as shown schematically in figure 1, their energy gaps range from the near infrared for InN (Eg = 0.7 eV) to near UV for GaN (Eg = 3.4 eV) and to deep UV for AlN (Eg = 6.2 eV). As discussed in detail in section 3 these materials can be grown heteroepitaxially on a number of substrates and can also be doped n- and p-type. Furthermore, they were found to have some unique physical and chemical properties. Specifically, GaN has bonds considerably shorter (~20%) than most semiconductors, has large ionicity, has high hardness and thermal conductivity, it is radiation hard and exhibits chemical resistance to attack by acids and bases.

Thus, these materials are ideally suited for a broad range of applications as illustrated in table 1. Some of the device applica-tions described in this table have already been implemented in our current technology. For example, full color displays based on red, green and blue LEDs, are used extensively for outdoor and automotive applications as well as in displays such as TV and portable electronic devices. In parallel the introduction of

solid state lighting for general illumination is advancing rap-idly worldwide and its full implementation will have significant economic and environmental benefits to the society. Also, the development of violet-blue lasers has impacted very signifi-cantly the field of optical recording for information storage.

In addition to visible LEDs and lasers, which are based on InGaN alloys, significant research effort is currently being devoted to the development of UV optoelectronic devices based on AlGaN alloys. These alloys are well suited for such devices because their energy gap can be tuned by changing the alloy composition to cover the entire UV spectral region

Figure 1. Schematic of the energy band gap versus lattice constant of Nitride Semiconductors.

Table 1. Applications of nitride semiconductors.

Applications as optoelectronic materials

• Visible and UV lasers• Visible and UV LEDs• Solar-blind detectors• Optical modulators• Terahertz emitters and detectors based on intersubband transitions in QWs

Applications as energy materials

• Solid state lighting• Photovoltaic solar cells• Hydrogen production by water photolysis

Applications as electronic materials

• Transistors for high frequency• Transistors for high power• Transistors for high temperature

Applications as materials for sensors

• MEMS sensors due to their piezoelectric nature

6.1.4. White LEDs .............................................15 6.2. UV LEDs .............................................................15 6.2.1. Fundamental materials problems ............16 6.2.2. Growth methods of improving

the IQE of AlGaN alloys .........................16 7. Lasers ...........................................................................18 8. UV Photodetectors ...................................................... 20 8.1. Introduction ........................................................ 20 8.2. Photoconductive detectors .................................. 20 8.2.1. GaN ‘visible blind’ UV

photoconductive detectors ..................... 20 8.2.2. AlGaN ‘solar blind’ UV

photoconductive detectors ..................... 22 8.3. Photovoltaic UV detectors ................................. 23 9. UV electroabsorption modulators ............................... 23 9.1. Introduction ........................................................ 23 9.2. Electroabsorption modulators based

on bulk GaN films .............................................. 24 9.3. Electroabsorption modulators based

on GaN/AlGaN quantum wells .......................... 25 10. Intersubband transitions in AlGaN

quantum wells ............................................................. 27 10.1. Basic device physics ......................................... 27 10.2. Infrared device applications .............................. 28 10.3. Terahertz device applications ........................... 29 11. Recent developments .................................................. 31 11.1. GaN substrates .................................................. 31 11.2. LEDs on GaN substrates .................................. 31 11.3. LEDs on non-polar and semipolar

GaN substrates .................................................. 31 11.4. LEDs on silicon substrates ............................... 32 Acknowledgments ....................................................... 32 References ................................................................... 32

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from 360 to 210 nm. This research is motivated by a plethora of potential industrial and medical applications.

Progress has also been made in developing electronic devices for high frequency, power, and temperature applica-tions. Furthermore, due to their piezoelectric nature these materials are being investigated for the development of sen-sors based on micro-electro-mechanical systems (MEMS). The growth and properties of these materials have been dis-cussed in a number of books and reviews [1–16].

Gallium nitride (GaN) was synthesized for the first time by Juza and Hahn [17] in 1938 by the reaction of ammonia with hot gallium. This method produced small needles and platelets, which were sufficient to study the crystal structure and determine the lattice constant of GaN. Grimmeiss and Koelmans [18] in 1959 used the same method of growth and produced small crystals of GaN and were able to measure their photoluminescence (PL) spectra. Maruska and Tietjen [19] in 1969 used the hydride vapor phase epitaxy (HVPE) method to produce large area thin film GaN crystals on sap-phire substrates. GaN films produced at that time were heav-ily doped n-type even though n-dopant impurities were not intentionally incorporated into the films. This n-type auto-doping was attributed by the authors to nitrogen vacancies. However, Seifert and co-workers [20] in 1983 have produced exper imental evidence that the n-type auto-doping is due to accidental incorporation of oxygen, which acts as a donor.

Following these initial developments, Pankove and co-workers [21] were able to incorporate zinc (Zn) during growth of GaN by HVPE and produced semi-insulating GaN films due to compensation of the n-dopants. This led to the fabrication of the first GaN LED having the device form M-i-n, where M is a metal. Subsequently, Maruska and co-workers [22] were able to form M-i-n GaN diodes, emitting violet light, by pro-ducing the semi-insulating i-layer using magnesium (Mg) as a compensating p-type dopant. These accomplishments led to a flurry of scientific activity in a number of laboratories world-wide. However, the inability to form conductive p-type GaN prohibited the exploitation of this new material for the fabrica-tion of efficient LEDs and other types of devices.

One major problem in the development of devices based on this class of materials is the lack of III-nitride substrates for homoepitaxial growth of devices. Thus, the great majority of work reported in the last 40 years dealt with the heteroepi-taxial growth of GaN on various foreign substrates. An impor-tant development in the late 1980s and early 1990s was the improvement in the heteroepitaxial growth of GaN which led to atomically smooth films, that were able to be doped intention-ally n- or p-type by incorporating Si or Mg respectively during film growth. This was accomplished by implementing novel nucleation steps prior to the growth of the epitaxial GaN films. Such nucleation steps included the nitridation of the sapphire substrate, followed by the growth of defective low temper-ature AlN or GaN buffer layers. The nitridation of sapphire was developed by the Boston University group and involved the conversion of the surface of sapphire from Al2O3 to AlN by exposing the surface to activated nitrogen produced either by a nitrogen plasma or by thermal cracking of ammonia [23, 24]. The AlN buffer was developed by Amano et al [25] and

the GaN buffer was developed independently by the Boston University group [23, 24, 26–28] and Nakamura [29] using the molecular beam epitaxy (MBE) and the metal organic chemi-cal vapor deposition (MOCVD) methods respectively. These nucleation steps facilitated the epitaxial growth and led to the annihilation of some of the threading defects during the growth of the device layers at higher temperatures.

The second major obstacle in the development of these materials was the p-type doping. In late 1980s Amano et al [30] were able to demonstrate p-type doping in GaN. Specifically, these authors were investigating the cathodoluminescence (CL) of GaN:Zn films and observed that the CL efficiency increased by two orders of magnitude upon exposure of the sample to low energy electron beam irradiation (LEEBI). Hall Effect measurements on GaN:Mg samples after the LEEBI treatment indicated that the films were p-type and conducting. It was also observed by Amano et al [31] that the high p- type conductivity extended only to 0.3 µm deep, which was the penetration depth of the electron beam.

This result was accounted for by van Vechten et al [32], who proposed that the shallow acceptor levels of Mg were compen-sated by hydrogen atoms through the formation of complexes with Mg. A similar model has earlier been proposed by Pankove et al [33] and Moustakas [34] to account for the de-activation of p-type dopants in crystalline and amorphous Si respectively upon incorporation of hydrogen into the lattice. The energy of the electron beam in the LEEBI experiment releases hydro-gen from the complex and the Mg becomes shallow acceptor approximately 0.16 eV above the valence band [35].

Nakamura et al [36] demonstrated that p- type conductivity can also be obtained by annealing the GaN:Mg films above 750 °C in nitrogen atmosphere or vacuum, which is consistent with thermal braking of the Mg-H bonds and hydrogen evo-lution out of the film. A verification of the role of hydrogen in de-activating p-type dopants was produced by Moustakas and Molnar [37], who demonstrated for the first time that GaN:Mg films, grown by MBE in a hydrogen free environment, have a hole concentration of 6 × 1018 cm−3 without requiring any post growth treatment. Post-hydrogenation of these films by Brand et al [38] with a hydrogen plasma demonstrated the de-activation of the p-type doping.

The development of the aforementioned novel nucleation steps together with the understanding of the p-type doping cleared the way for the fabrication and commercialization of various opti-cal and electronic devices. Indeed shortly after a group at Nichia Corporation in Japan, headed by Nakamura, demonstrated the fabrication of efficient blue LEDs [39] and blue–violet lasers [40].

2. Crystal structure and basic properties

An excellent review of the crystal structure of Nitride semi-conductors was presented by Trampert and co-workers [41]. In this section we present only a brief summary on the subject. Contrary to GaAs and other arsenide and phosphide III–V compounds, which have the cubic zincblende structure, the equilibrium structure of group-III nitride semiconductors is the wurtzite, belonging to P63Mc space group in which all the atoms are tetrahedrally coordinated with atoms of the opposite

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type. The wurtzite structure consists of two hexagonal close packed (hcp) lattices, displaced towards the c-direction (the direction perpendicular to the hexagonal base) by ± 3c/8, where c is the lattice parameter. The wurtzite crystal struc-ture is shown in figure 2(a) [42]. The equilibrium bulk lattice parameters for the three binary compounds as well as other relevant properties are listed in table 2. In all three compounds the ratio of c/a differs slightly from the ideal value of the wurtzite structure of 1.63, implying a certain degree of dist-ortion with respect to the ideal closed-packed arrangement. Figure 2(b) shows the important crystallographic planes of the wurtzite crystal structure.

Group III-nitrides can also crystallize in the rock-salt and zinc-blende polytypes. The transition to the rock-salt phase only occurs under high pressure [43–45], and so technologically it is not an important phase. The zinc-blende structure is metastable

and may be stabilized by heteroepitaxial growth on substrates with cubic symmetry [46, 47]. The in-plane symmetry of (0 0 0 1) wurtzite GaN is identical to the symmetry of (1 1 1) zinc-blend GaN. Furthermore, the cohesive energies of the wurtzite and cubic phases are very similar [48], allowing the conversion between the two phases to occur easily through the creation of stacking faults of the close packed (0 0 0 1) and (1 1 1) planes. Thus, basal plane stacking faults are abundant in III-nitrides.

The wurtzite crystal structure is not invariant with respect to inversion along the c-direction. Such an operation results in an exchange of group-III atoms with nitrogen atoms and vice versa. This lack of inversion symmetry is referred to as crys-tallographic polarity or more commonly just ‘polarity’ [41], and the c-direction is the polar direction. In GaN, the [0 0 0 1] direction is defined as the direction parallel to the Ga-N bonds pointing towards the nitrogen atoms, as illustrated in

Table 2. Properties of III-nitrides in Wurtzite structure.

Property AlN GaN InN

Lattice constant a = 3.11 Å a = 3.19 Å a = 3.55 Åc = 4.98 Å c = 5.185 Å c = 5.76 Å

Energy bandgap 6.2 eV 3.4 eV 0.65 eVCoefficient of thermal expansion (CTE) (×10−6 K−1)a

α − 4.2 α − 5.59 (300–900 K) α − 2.7–3.7α − 5.3 α − 3.17 (300–700 K) α − 3.4–5.7

Young modulus (Gpa)b C11 = 398 C11 = 396 C11 = 271C12 = 140 C12 = 144 C12 = 124C13 = 127 C13 = 100 C13 = 94C33 = 382 C33 = 392 C33 = 200

Bulk modulus (Gpa)c 199 176 122Index of refraction 1.9–2.2 2.5–2.6 —Thermal Conductivity (W cm−1 K−1) 2.8 1.3 0.8Melting point (K)d 3487 2791 2146

a [353].

b [354].

c [355].

d [356].

Figure 2. (a) GaN wurtzite crystal structure. Large yellow balls are nitrogen atoms and small grey balls are Ga atoms [42]; (b) important crystallographic planes of the wurtzite crystal structure. (a) Reproduced from (https://commons.wikimedia.org/wiki/File:Wurtzite_polyhedra.png). Image stated to be in the public domain.

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figure 2(a), and [0 0 01]-oriented films grown along this direc-tion are referred to as ‘Ga-polar’. Similarly, films grown along the [0 0 0 1] direction are termed ‘N-polar’. The two polarities exhibit several different chemical and physical properties, such as differences in resistance to chemical etching or the ability to incorporate impurities [49]. However, it should be stressed that polarity is not a surface property but instead is a bulk property.

The control of polarity during growth of III-nitride films is essential for the design and fabrication of electronic and optoelectronic devices. Thus, a number of analytical tools have been developed to determine the polarity of these films. Of these methods the convergent beam electron diffraction (CBED) in transmission electron microscopy (TEM) is the most direct and reliable technique [50]. An alternative method of polarity determination is the chemical etching using KOH or NaOH solutions. Etching of N-polar GaN with KOH leads to faceted surface morphology due to preferential attack of the domain boundaries. The Ga-polarity GaN etches slower due to its higher stability. The relative stability of the Ga-polarity is supported by the observation of faster growth rate in this polarity compared to the N-polarity [50].

The chemical bond in III-nitrides is covalent, but with a significant ionicity (group III cations and nitrogen anions). A consequence of the partly ionic nature of the chemical bonds between neighboring group-III and N atoms, which as shown in figure 2(a), are aligned along the c-axis of the crystal, together with the slight distortion from the ideal closed-packed arrange-ment results in a non-zero macroscopic polarization (‘spon-taneous’). The spontaneous polarization vector points along the [0 0 0 1] direction, which is parallel to the growth direction for N-polar films, and anti-parallel for Ga-polar films. Biaxial distortion of the crystal due to strain leads to a change in the magnitude of this macroscopic polarization, that is, wurtzite III-nitrides are also piezoelectric. Thus, the piezo electric polariza-tion is the induced change in the spontaneous polarization [51]. These issues are discussed in more detail in section 4 below.

3. Heteroepitaxial growth

Historically, one major problem in the development of elec-tronic and optoelectronic devices out of this class of materials has been the lack of III-nitride substrates to grow the devices homoepitaxially. Thus, the great majority of the reported devices were grown heteroepitaxially on foreign substrates such as sapphire, SiC and Si. Currently there is a significant effort to form bulk GaN and AlN crystals by a variety of meth-ods, and the development of quasi-GaN substrates by grow-ing thick GaN films on foreign substrates and then separating them from the substrate.

The heteroepitaxial growth of III-nitride films and devices has been investigated using several growth techniques. These includes hydride vapor phase epitaxy (HVPE), metalorganic chemical vapor deposition (MOCVD), molecular beam epi-taxy (MBE), pulsed laser deposition (PLD), and physical vapor deposition (PVD). Of these, the former three are the most popular methods.

Several researchers have used crystallographic models to explain the epitaxial relation between the nitride epilayers and

the (0 0 0 1) sapphire [52–55]. Lei and co-workers [52] pro-jected the (0 0 0 1) basal planes of sapphire, GaN and AlN (see figure 3) and have shown that GaN and AlN grow with 30° rotation with respect to the sapphire unit cell, with an orienta-tion relationship:

[1 1 − 2 0]GaN//[1 0 − 1 0]sapphire and [1 0 − 1 0]GaN//[1 1 − 2 0]sapphire.

(3.1)

The lattice mismatch between sapphire and GaN, as calcu-lated from this projection, is +16%, in close agreement with the experimentally measured value by x-ray diffraction [52] or electron microscopy [56]. It should be noted that without this rotation in the orientation the lattice mismatch between GaN and (0 0 0 1) sapphire would have been -34%.

This type of epitaxial relationship between all three binary nitrides and (0 0 0 1) sapphire is illustrated clearly in figure 4 [57]. If the GaN, AlN and InN were grown with the orien-tation [1 1 –2 0]sapphire // [1 1 –2 0]GaN, they would have been under tensile stress with the indicated in the figure lattice mis-match. On the other hand if the nitrides grow with 30° rota-tion with respect to the sapphire unit cell, the films will be under compression with significantly smaller lattice mismatch for the AlN and GaN. The lattice mismatch of the InN on the other hand is about the same with and without the 30º rota-tion. Thus, direct growth of InN on sapphire is likely to lead to domains of both orientations. The solution to this problem is to convert the surface of sapphire from Al2O3 to AlN, a pro-cess known as nitridation [23, 24], followed by the growth of a thin GaN or AlN template and then deposit the InN film. In this case the InN grows compressively with a lattice mis-match of 11% or 14% respectively. Growth of InN using this approach has led to high crystalline quality InN films [58].

Epitaxial thin film growth in general can be categorized into one of three modes [59]: (i) Frank–Van der Merwe mode or layer-by-layer growth mode, (ii) Volmer–Weber mode or island growth mode and (iii) Stranski–Krastanov mode or layer-by-layer followed by island growth mode. The mode of nucleation and growth is determined by the interfacial free energy between

Figure 3. Projection of bulk c-plane sapphire and GaN cation positions. The dots mark aluminum atom positions and the dashed lines show the sapphire (0 0 0 1) plane unit cells. The open squares mark gallium atom positions and the solid lines show the GaN (0 0 0 1) plane unit cells [8]. Reprinted from [8], Copyright (2002), with permission from Elsevier.

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the three surfaces: substrate (s), deposit (d), and vapor (v). Under equilibrium conditions the deposit forms a hemispheri-cal cap on the substrate with a contact angle θ given by [60],

cosθ = (σsv − σsd) /σdv (3.2)

where σij is the interfacial free energy between the two sur-faces i and j.

In the case of homoepitaxy, since the substrate and the deposit have the same structure, σsd = 0 and σsv = σdv. Thus, the contact angle is zero and growth occurs in the layer-by-layer or Frank–Van der Merwe mode. In the case of heteroepitaxy, the contact angle has a non-zero value, compli-cating the growth. Thus, in the presence of lattice mismatch between the substrate and the deposit, layer-by layer growth is never the equilibrium morphology, but is metastable with respect to cluster (island) formation. This is illustrated in fig-ure 5(a), which shows a scanning electron microscopy (SEM) image of the early stages of nucleation of GaN on (0 0 0 1)

sapphire [61]. It is apparent from these data that the nucleation of these films proceed through the formation and coalescence of hexagonal islands as schematically shown in figure 5(b). Continuing growth should lead to a material with columnar microstructure.

In general, the brittle to ductile transition in crystals occurs at about half of the melting temperature of the material. Above that temperature stress relaxation occurs by disloca-tion generation and glide. In GaN, for example, this transition is expected to occur at about 1400 K. Since GaN and its alloys are generally grown at lower temperatures (in particular by MBE), one expects that the hexagonal islands in figure 5(a) to be coherent to the substrate [61] at thicknesses well beyond the Mathews critical thickness [62]. Indeed Akasaki and Amano reported that AlGaN films (several thousand angstroms thick) and InGaN films (several hundred angstrom thick) are coher-ent to GaN templates [63]. Similarly, Okumura et al reported that 7000 Å thick AlN was grown coherently on 6H-SiC, even

Figure 4. Epitaxial relationship of III-nitride compounds and (0 0 0 1) sapphire [57]. Reproduced with permission from [57]. © 2004 American Vacuum Society.

Figure 5. (a) Heteronucleation of GaN on (0 0 0 1) sapphire; (b) schematic of incomplete coalescence of the hexagonal islands leading to the formation of threading dislocations at the domain boundaries [61]. [61] John Wiley & Sons. Copyright © 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.

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though the Mathews critical thickness is only 30 Å [64]. On the other hand due to incomplete coalescence of the hexagonal islands, as shown schematically in figure 5(b), one expects the formation of threading dislocations at the domain boundaries.

Thus, a key materials issue is the heteroepitaxial nucleation of the nitride films is the high dislocation density in the films. The epitaxial growth of these materials is further complicated by the existence of a metastable zincblende structure in addi-tion to the stable wurtzite structure. As discussed in section 2, the cohesive energies of these two structures are very similar, allowing conversion between the two phases to occur eas-ily, mainly by creation of stacking faults on the close packed (0 0 0 1) and (1 1 1) planes [52, 56]. Extrinsic stresses associ-ated with thermal mismatch between substrate and film as well as intrinsic stresses associated with island coalescence, grain growth and surface stresses are also important.

To address some of the problems, related to the heter-oepitaxial growth, a number of nucleation steps have been developed for the growth of III-nitrides on various substrates. For example, growth on (0 0 0 1) sapphire by both MBE and MOCVD involves the nucleation steps of sapphire nitridation [23, 24, 65–72] and a low temperature GaN [23, 24, 26–29] or AlN buffer [25]. It is commonly agreed that a nitridated sapphire surface facilitates the growth of smooth and relaxed GaN films, leading to reduction in inversion domain bounda-ries and improved optoelectronic properties [71, 72]. The GaN or AlN low temperature buffer layers are generally polycrys-talline and rough with a very high density of defects including stacking faults and grain boundaries. In addition the buffer layer consists of a mixture of hexagonal and cubic phases of GaN [73, 74]. It is believed that the polycrystalline buffer accommodates the stress from the large lattice mismatch and reduces the propagation of defects generated at the interface by destructive interference and annihilation, thus reducing the dislocation density. Wu et al [75] reported that annealing of the buffer to the growth temperature results in partial con-version of zincblende GaN to wurtzite GaN. Keller et al [76] reported grain coarsening upon annealing and concluded that optimization of both the deposition and the re-crystallization

parameters are important. The two nucleation steps (nitrida-tion and buffer) play also an important role in determining the polarity of the nitride films [77]. The employment of a defec-tive low temperature GaN buffer facilitates the annihilation of threading defects as the film grows thicker. This is illustrated in figure 6, which shows the TEM cross-section image of a GaN film grown on a (0 0 0 1) sapphire substrate subjected first to the two nucleation steps of nitridation and low temper ature GaN buffer [61]. In recent years in addition to the aforemen-tioned nucleation steps patterning of the sapphire substrate has led to additional reduction in the concentration of the threading defects in heteroepitaxially grown III-nitride films on sapphire [78–80].

4. Polarization effects

A distinctive property of nitride semiconductors, which has important effects on the electronic, transport, and opti-cal properties of heterostructures, is the presence of strong spontaneous and piezoelectric polarizations [81–83]. These internal fields are a consequence of the partly ionic nature of the chemical bonds between neighboring group-III (cation) and N (anion) atoms, and are aligned along the c-axis of the crystal. In the case of zincblende materials (such as typical As-based semiconductors), the cation and anion sublattices are arranged in such a way that in the absence of strain there is no net internal polarization. In contrast, in nitride semi-conductors based on the wurtzite crystal structure, a relative shift of the sublattices with respect to the ideal closed-packed arrangement exists that produces a large spontaneous (pyro-electric) polarization. In both crystal types, a net piezo electric polarization is also introduced through the application of strain, due to the resulting relative displacement of neighbor-ing atoms.

The orientation of both spontaneous and piezoelectric polarizations along the c-axis depends on the atomic layers sequence in the crystal. As discussed in section 2, atoms in the wurtzite structure are arranged in alternating hexagonal lay-ers perpendicular to the c-axis, containing cations and anions, respectively. Assuming growth along the c-direction, if each group-III atom is stacked directly above a N atom, the crystal is said to have N-face polarity. In this case, the spontaneous polarization field PSP points away from the substrate towards the top surface of the epitaxial layer, in what will be taken here as the positive z-direction, i.e. PSP = PSPZ with PSP > 0. The opposite layer sequence is referred to as Ga-face in the case of GaN (and similarly for AlN and InN), and it results in a spontaneous polarization field in the negative z-direction. In general, both orientations can be synthesized depending on the growth conditions, although most nitride crystals nowa-days are formed with group-III-face polarity. Regarding the magnitude of PSP, it increases with increasing deviation of the unstrained crystal structure from the ideal closed-packed arrangement, i.e. in going from GaN to InN to AlN.

The piezoelectric polarization field PPE = PPEZ depends on the nonzero elements of the tensor describing the strain state of the crystal. Assuming pseudomorphic growth along the c-axis (i.e. the z-direction), these are [84]

Figure 6. Cross-section TEM of a GaN film showing the annihilation of threading defects due to the employment of the defective low temperature GaN buffer [61]. [61] John Wiley & Sons. Copyright © 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.

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εxx = εyy =a − a0

a0, εzz = −2

C13

C33εxx, (4.1)

where a0 and a are the equilibrium and strained values of the in-plane lattice constant, respectively (the latter being deter-mined by the composition of the growth template), and C13 and C33 are elastic constants. By definition of the piezoelectric coefficients eij, PPE is then given by [84]

PPE = ±2εxx

(e31 − e33

C13

C33

), (4.2)

where the plus and minus signs refer to group-III- and N-face polarity, respectively. It should be noted that in GaN, InN, AlN, and related alloys the quantity in parentheses in equa-tion (4.2) is negative, so that under group-III-face conditions PPE is negative for tensile and positive for compressive strain (and vice versa in the N-face case). The orientations and rela-tive magnitudes of PSP and PPE are illustrated in figure 7 for various combinations of layer and template compositions and polarity. All the material parameters required to compute these polarizations fields can be found, e.g. in [85].

The composition dependence and large magnitude of PSP and PPE play an important role in the optoelectronic device physics of nitride semiconductors, due to the associated polarization charge distributions. In general, any discontinu-ous change ΔP in the normal component of the polarization across an interface is equivalent to a surface charge distribu-tion of density σP = –ΔP, and therefore results in the forma-tion of an electrostatic field perpendicular to the interface. In the case of bulk crystals with uniform composition, such polarization charges only exist on the free surfaces, where they can be effectively neutralized by the mobile charges present in the atmosphere. In contrast, in heterostructures polarization charges are also induced on each heterojunction, where they can produce large internal electric fields which in turn can strongly bend the conduction- and valence-band edges near the interface. Referring to the geometry of figure 7, σP is posi-tive if the net polarization density P = PSP + PPE decreases (in absolute value and sign) upon crossing the interface in the positive z-direction, and vice versa; some examples are shown in the figure. When the polarization charge distribution is pos-itive, it can be partly compensated by available conduction-band electrons in the crystal, leading to the formation of a 2D

electron gas. Similarly, a 2D hole gas can be created when σP < 0. These effects are widely exploited in the development of GaN-based high electron mobility transistors (HEMTs) providing excellent high-speed and high-power performance.

The electric fields produced by the polarization charges in heterostructures can be computed using simple electro-static arguments (i.e. by requiring conservation of the elec-tric displacement flux across each interface), together with an appropriate set of boundary conditions. Here we consider the important case of multiple-quantum-well (MQW) structures, which are used routinely in most optoelectronic device appli-cations considered in this article. The well and barrier mat-erials could be for example GaN and AlGaN (or InGaN and GaN), and their properties will be denoted by the subscripts w and b, respectively. Continuity of the displacement flux across each interface requires that εwFw + Pw = εbFb + Pb, where the symbols F and ε indicate the built-in electric fields and the permittivities, respectively. A simple analytical solution for Fw and Fb can be obtained in the limit of an infinitely periodic structure, where the intrinsic voltage drop across each well/barrier pair must be zero, i.e. FwLw + FbLb = 0 (L being the layer thickness). Combining these two equations gives

Fw =(Pb − Pw) Lb

εbLw + εwLb (4.3)

for the polarization-induced built-in electric field in the well layers. The same expression with the subscripts w and b inter-changed applies to Fb.

A schematic illustration of the resulting conduction- and valence-band lineups of a generic Ga-face GaN/AlGaN MQW structure is shown in figure 8. As clearly seen in this figure, the band edges in the well and barrier layers are strongly tilted (in opposite directions) due to the polarization-induced built-in electric fields. Very large values of up to a few MV cm−1 are typically computed for these fields in c-plane nitride heterostructures. As a result, the energy minima and enve-lope functions of all subbands supported by the QWs are strongly modified relative to the ideal flat-band case (i.e. for Fw = Fb = 0) through the quantum-confined Stark effect. In particular, the energy separation between the ground-state conduction and valence subbands, which determines the operating wavelength of interband optoelectronic devices, is

Figure 7. Direction of the spontaneous and piezoelectric polarization fields and sign of the interface polarization charge distribution, for different combinations of layer and template compositions and crystal polarity.

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substantially red-shifted by Fw. Furthermore, the envelope functions of these conduction and valence subbands become more tightly localized near opposite sidewalls of each well. This phenomenon is illustrated in figure 8, where the squared envelope functions of the ground-state conduction and valence subbands are also plotted (referenced to their respective energy levels). As a result, the mutual overlap between these electron and hole wavefunctions is decreased (again compared to the ideal flat-band case), leading to reduced interband oscillator strengths and hence reduced absorption, spontaneous-emis-sion, and gain coefficients.

Finally, it should be noted that a more accurate descrip-tion of polarization phenomena in nitride heterostructures requires more realistic boundary conditions (to account for the finite number of QWs in practical structures), and the inclusion of space charge effects due to the available mobile electrons and holes. The latter effects can be modeled through the self-consistent solution of the Schrödinger and Poisson equations including the polarization-induced interface charge distributions. The main conclusions described in the previous paragraph remain valid even with this more realistic descrip-tion, although the built-in electric fields can be partly screened in the presence of high carrier densities.

5. Device fabrication

The fabrication of optoelectronic devices requires the forma-tion of Ohmic contacts with low contact resistivity. The choice of such contacts depends sensitively on the position of the intrinsic surface states in the energy gap of the semiconduc-tor. As first was pointed out by Shockley, the position of such states in semiconductors and insulators depends strongly on the ionicity of the chemical bond between their atoms [86]. For covalent materials, with large overlap between adjacent atoms, the surface states occur in a narrow half-filled band near the center of the energy gap. However, in ionic mat erials, with small overlap between adjacent atoms, the surface states have moved towards the band edges, and therefore the surface Fermi level is unpinned. The chemical bond in III-nitrides is covalent, but with a significant ionicity. Thus, one expects

that the intrinsic surface states due to dangling bonds to have moved towards the band edges [61]. Such result has been con-firmed for GaN experimentally by Dhesi et al [87] and Ryan et al [88], using angle resolved photoemission spectroscopy.

Since the surface Fermi level of clean GaN surface is unpinned, Ohmic contacts and Schottky barriers can be formed to these materials by proper choice of the metal work function. For example, to form Ohmic contacts to n-GaN and its alloys one has to use metals with low work function [89]. Also to form Ohmic contacts to p-GaN and its alloys one has to use metals with high work function. Specifically, for ionic semiconduc-tors the barrier height (Φ) of a metal in intimate contact with the surface of the semiconductor should be proportional to the work function (W) of the metal. This should be contrasted with covalent semiconductors in which the barrier height of metals in intimate contact with the surface of the semiconductor is independent of the workfunction of the metal.

In a series of papers the Caltech group had verified these predictions by measuring the barrier heights of metals depos-ited in ultrahigh vacuum on cleaved surfaces of a number of n-type semiconductors [90–94]. Kurtin and co-workers [94] established for the first time that if the crystalline solids are ordered by some measure of their bond ionicity, then many electronic properties follow a universal abrupt transition between the covalently and ionically bonded materials. In their discussion these authors have adapted the Pauling ionic-ity scale [95], and have considered the difference in electron-egativity (ΔX) of the two constituents of a binary compound as an approximate measure of the ionicity of the compound. Figure 9 shows that the dependence of the derivative dΦ/dW as a function of the electronegativity difference (ΔX) of the investigated compounds. The results indicate an abrupt trans-ition between covalent and ionic semiconductors.

Foresi and Moustakas [89] have investigated metal contacts to GaN and found that there is a direct correlation between barrier height and work function of the metal. Figure 9, which was first reported in [94], was modified by Moustakas [61] to include the Nitride semiconductors, using their electron-egativity difference in the Pauling scale. This puts the nitride semiconductors above the knee of the curve, implying that

z

E

z

E

Figure 8. Conduction- and valence-band lineups and squared envelope functions of the ground-state electron and hole subbands in a Ga-face GaN/AlGaN MQW structure.

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Schottky barriers on these semiconductors should have barrier heights which depend directly on the work function difference between the metal and the semiconductor.

The electron affinity of GaN was measure to be 4.1 eV [96]. Therefore, any metal with a work function equal to or lower than 4.1 eV should form essentially Ohmic contacts to n-GaN and any metal with a work function higher should form a rectifying contact to n-GaN. In the original paper Foresi and Moustakas [89] have chosen Al to form an Ohmic contact to n-GaN, since its work function is 4.08 eV [97]. However, these authors have observed that the contact resistivity increases upon annealing, a result attributed to the formation of a thin interfacial layer of AlN. To address this problem these authors had proposed the incorporation of a thin transition metal layer (e.g. Ti or Cr) between Al and GaN to act as a diffusion bar-rier. Today a multi-stack of Ti/Al/Ti (or Ni)/Au, annealed to 800–900 °C is used universally to form Ohmic contact to n-GaN [98].

Since the valence band of GaN is about 7.0 eV below the vacuum, to contact p-GaN films one has to use high work function metals. Such metals are, for example, Au, Pd and Pt [97]. However, because Au has the tendency to diffuse and introduce states in the gap of GaN it was recommended to insert again a thin transition metal film between Au and GaN to act as a diffusion barrier [89]. Today Ni/Au annealed at about 500 °C is used universally to form an Ohmic contact to p-GaN.

The same combination of metal stacks used to form Ohmic contacts to GaN can also be used to form Ohmic contacts to n- and p-InGaN alloys, since these alloys have a smaller energy gap than GaN. However, for AlGaN alloys the electron affin-ity is expected to decrease as the AlN mole fraction increase. Thus, Ohmic contacts to n-AlGaN require metals with work function less than 4.0 eV. Due to the lack of conducting met-als with low work function, the problem of forming Ohmic

contacts to n-AlGaN becomes challenging. Indeed, it was shown that using the traditional metal contact schemes of Ti/Al/Ti/Au, one can form good Ohmic contacts to n-AlGaN only up to about 20% AlN mole fraction [99]. Other work-ers have demonstrated that vanadium-based contacts can form good Ohmic contacts to n-GaN [100, 101] and also to n-AlGaN [102–104].

France et al [103] demonstrated the formation of Ohmic contacts to n-AlGaN over the entire alloy composition uti-lizing vanadium-based contacts in the form V(15 nm)/Al(80 nm)/V(20 nm)/Au(100 nm). As shown in figure 10, these authors have found, that the formation of Ohmic con-tacts using this metallization scheme, requires progressively higher annealing temperature as the AlN mole fraction in the AlGaN alloys increases. The data of figure 10 suggest that the mechanism by which V-based schemes form Ohmic contact to the III-nitride semiconductors is through reaction between vanadium and the nitride film, which leads to the formation of vanadium nitride. Since the Al–N bonds are stronger than the Ga–N bonds, one expects that reaction between vanadium and nitrogen should occur at higher temperatures as the AlN mole fraction increases. Schweitz and Mohney studied the phase equilibria in transition metal-Al–N system and suggested that vanadium nitride (VN) is a thermodynamically favored form of V contacts on AlGaN at 600 °C [105]. It is also known that VN is a low work function metal (3.56 eV) [106]. Song et al [100] have shown that upon annealing at temperatures higher than 600 °C, other phases of vanadium nitride are also formed, which potentially may have even smaller work functions. Also, further doping of the sample by the nitrogen vacancies created upon VN formation may also contribute to the low contact resistivity. Pookpanratana et al have investi-gated the microstructure, intermixing and chemical structure of such vanadium-based contacts to n-GaN and n-AlN before and after rapid thermal annealing [107–109].

The specific contact resistivity of such vanadium-based contacts to n-AlGaN films versus AlN mole fraction was reported by France et al [103] and the results are shown in figure 11. According to these data the specific contact resis-tivity has a value of 10−6 Ω · cm2, for AlGaN films with AlN

Figure 10. Optimum annealing temperature of V/Al/V/Au to form Ohmic contacts to n-AlGaN alloys versus AlN % mole fraction [103]. Reprinted with permission from [103]. Copyright (2007), AIP Publishing LLC.

Figure 9. Dependence of index S, (change in barrier height over the change in metal work function) on the electronegativity difference between the components of the compound [61]. [61] John Wiley & Sons. Copyright © 2013 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.

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mole fraction up to 70% and increases to 1 Ω · cm2 in pure AlN. This predictive facility of making metal contacts to nitride semiconductors has accelerated the development of their optoelectronic devices.

6. Light emitting diodes

Following the discovery of the novel nucleation steps for the heteroepitaxial growth of GaN in late 1980s and early 1990s, which led to atomically smooth GaN films capable of inten-tional n- and p-type doping, Nakamura and co-workers [39, 110, 111] reported the development of bright blue–green LEDs followed by the commercialization of these devices by Nichia Corp.

As discussed previously the III-nitride materials are direct band gap semiconductors and thus, are suitable to form LEDs from the near IR to deep UV. Such devices are usually grown heteroepitaxially on (0 0 0 1) sapphire or 6H-SiC substrates. Recently there are also commercialization efforts of develop-ing such LEDs on (1 1 1) Si and bulk GaN wafers. The vis-ible spectrum LEDs are based on InGaN alloys, while the UV LEDs are based on AlGaN alloys.

6.1. Visible spectrum LEDs

Excellent reviews on the physics, design, performance and applications of visible spectrum LEDs have been published by Krames et al [112] and by Harbers et al [113]. The semiconduc-tors used in LEDs for solid state lighting are III–V compounds based on InGaN and AlGaInP. The external quantum efficiency (EQE) versus peak wavelength of LEDs, based on these two classes of materials, is shown in figure 12 [112]. These data clearly indicate that the efficiency of LEDs based on both fami-lies of materials is reduced from above 50% to less than 20% in the green-yellow part of the electro magnetic spectrum, a phenomenon known as ‘green gap’. However, even though the

efficiency of green LEDs is poor, such devices are used widely in displays and other applications due to the high light sensitiv-ity of the eye in this spectral region, as indicated by the lumi-nous eye response curve, V(λ), in figure 12. In this article we are addressing only LEDs based on the InGaN materials system.

The EQE of an LED is defined as the product of the inter-nal quantum efficiency (IQE- ratio of generated photons to electrons injected into the LED’s contacts) and the extraction efficiency (EXE—the number of photons extracted from the LED versus photons generated in the junction) [112]. The IQE depends sensitively on extended and point defects, which act as non-radiative recombination centers, as well as on the device design, which prevents carriers from escaping from the active region prior to recombination. The EXE depends sensitively on the ability to extract the light from the active region of the device to the free space. Light extraction from an LED chip is generally a challenging problem due to the relatively high index of refraction of these materials and the reader is referred to specialized articles addressing this issue [112–114]. In the following we address primarily the material issues, which affect the internal quantum efficiency of LEDs based on InGaN alloys.

As discussed by Chichibu et al [115] the IQE efficiency at room temperature can be estimated by the ratio of the photo-luminescence intensity at room temperature and that at low temperatures (~10 K). This of course assumes that recombi-nation at low temperatures is 100 % radiative, which is a rea-sonable assumption. These authors have used this method to estimate the IQE of various AlInGaN alloys and QWs and the results are shown in figure 13. As seen in these data the IQE of c-plane InGaN QWs increases with the InN mole fraction to a value of about 90% for 15–20 % In and then decreases. Since these materials are highly defective, this extremely high IQE is difficult to rationalize. In the following we dis-cuss the potential sources of this high IQE and the observed decrease of the IQE at InN mole fraction more than 15–20%.

Figure 11. Specific contact resistivity versus AlN mole fraction for V-based contacts on n-AlGaN in the entire alloy composition [103]. Reprinted with permission from [103]. Copyright (2007), AIP Publishing LLC.

Figure 12. State-of-art external quantum efficiencies for high-power visible-spectrum light-emitting diodes. V(λ) is the luminous eye response curve [112]. Reproduced with permission from [112]. © Copyright 2007 IEEE.

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6.1.1. Potential origins of the high IQE of blue LEDs. As dis-cussed in section 3, nitride semiconductors suffer from high concentration of dislocations and other extended defects due to heteroepitaxial growth on foreign substrates. However, con-trary to LEDs based on traditional III–V compounds, whose performance decreases drastically if the concentration of dislocations exceeds 105 cm−2 [116], Ponce and Bour [117] reported early on that InGaN based LEDs can be relatively very efficient even if the concentration of edge dislocations is as high as 1010 cm−2. This result is consistent with the ionic bonding in nitride semiconductors as discussed previously. In ionic semiconductors the energy states due to dangling bonds at free surfaces as well as at edge dislocations, which are the plurality of dislocations in GaN, have move towards the band edges, where they act as traps rather than recombination states [61]. The evidence suggest that the screw dislocations, whose density is several orders of magnitude smaller than the edge dislocations, are non-radiative recombination centers [118].

Other potential sources responsible for the high IQE of InGaN based blue LEDs are band structure potential fluctua-tion present in these materials, leading to exciton localiza-tion, which prevents them from migrating and recombining non-radiatively in point and line defects. A number of authors have attributed the high efficiency of blue LEDs to compo-sitional inhomogeneities in InGaN/GaN QWs due to local phase separation [119]. Also, another source of band struc-ture potential fluctuations is the partial alloy ordering in InGaN alloys [120, 121].

It has also been suggested that basal plane stacking faults introduce band structure potential fluctuations [61]. As dis-cussed previously III-nitride semiconductors can exist in the stable wurtzite and metastable zincblende structures. The enthalpy of formations of these two allotropic forms differs by only ~10 meV [48]. Since these materials are grown at 1000 K or higher, conversion between the two phases can occur easily

by creation of stacking faults on the close packed (0 0 0 1) and (1 1 1) planes. Thus, basal plane stacking faults are abun-dant in these semiconductors. However, since the basal plane stacking fault in the wurtzite structure is the equivalent of a monolayer of a cubic domain and since the energy band gap of the cubic structure is 0.2 eV smaller than the wurtzite struc-ture [46], the stacking faults introduce band structure potential fluctuations. Therefore such defects are expected to be benefi-cial in the performance of an LED, since electron–hole pairs injected in the active region of the LED are localized in such potential fluctuations. This is exactly the opposite in cubic III–V compounds, where a stacking fault is the equivalent of a monolayer of a wurtzite structure with larger energy gap than the matrix.

6.1.2. Potential origins of the ‘green gap’ and solutions. The physical origin in the reduction of the IQE as the InN mole fraction in the InGaN/GaN MQWs increases to more than 15–20% is the subject of intense investigation over the past several years [112, 122–129]. Some authors have attribute this phenomenon to the formation of dislocations in the QWs due to the larger lattice mismatch in the InGaN/GaN QWs as the InN mole fraction in the InGaN increases [125] and others to point defects due to the lower growth temperature required to incorporate higher In percentage [126]. Yet another potential source of the observed reduction in IQE with the InN mole fraction is the increase of the internal electric field in the InGaN/GaN QWs due to spontaneous and piezoelectric polar-ization [127]. As discussed in section 4, these internal electric fields lead to a distortion of the quantum wells, which reduce the overlap of the wavefunctions of the injected electron–hole pairs (quantum-confined Stark effect) [128, 129].

Alternative sources for this reduction in IQE are alloy-ing phenomena related to phase separation and partial alloy ordering. At low InN mole fractions these phenomena may be beneficial by introducing band structure potential fluctuations, which lead to exciton localization and thus efficient radiative recombination. However, at high InN mole fraction the InGaN alloys may undergo complete phase separation to its binary compounds. Alternatively, the strain in the InGaN alloys may be relieved through alloy ordering. However, such ordering is partial consisting of domains having an ordered or a random structure with different energy gaps [120, 121, 130]. Based on experimental results and theoretical predictions in AlGaN alloys [131, 132] the energy gaps of the two types of domains may form a type-I heterostructure at low InN mole fraction and convert to type-II heterostructure for high InN mole frac-tion. In the latter case the injected electron–hole pairs become physically separated resulting in reduction of radiative recom-bination rate [131].

A number of approaches have been propose in order to improve the efficiency of green LEDs. One approach is to grow such devices on non-polar or semi-polar planes in order to eliminate/minimize the quantum-confined Stark effect in the InGaN/GaN QWs [133]. However, issues related to the ther-modynamic stability of InGaN in such LEDs are still relevant.

An alternative approach is to use InGaN quantum dots (QDs) as the active region of the device. Since phase separation and

Figure 13. Estimated IQE of various III-nitride alloys and QWs versus the respective alloy composition [115]. Reprinted by permission from Macmillan Publishers Ltd: Nature Materials [115], Copyright (2006).

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atomic ordering are stain driven phenomena and since stain in the QDs is relieved due to the large surface area, one would not expect that the InGaN QDs would have a tendency for phase separation or alloy ordering. However, issues related to quant-um-confined Stark effect are still relevant. Self-assembled InGaN QDs by the Stranski–Krastanov method have being developed by plasma-assisted [134] and ammonia-source molecular beam epitaxy [135], as well as metalorganic vapor-phase epitaxy [136, 137] methods. InGaN/GaN multi-quantum dot LEDs grown by MOVPE [138], ammonia-source MBE [139] and plasma-assisted MBE [140–144] methods have also been reported.

Figure 14(a) shows the electroluminescence spectra of three LEDs, grown by plasma-assisted MBE, and emitting in the blue, green, and red parts of the spectrum [140]. The blue LED, hav-ing an active region consisting of 3 In0.18Ga0.82N/GaN MQWs, has a peak emission at 440 nm with the FWHM of 30 nm under 20 mA injection current. The active regions of the green and the red LEDs are made of 5 periods of InGaN/GaN MQDs with the InGaN composition of 41% and 46%, respectively. The InGaN QD layer height and barrier thickness are 4 nm and 6 nm respectively. As shown in figure 14(a) the peak emission of the green LED is at 560 nm with FWHM of 87 nm while the red LEDs emit at 640 nm with FWHM of 97 nm under 20 mA injection current. Figure 14(b) illustrates the EL peak position as a function of injection current for the blue, green, and red LEDs. For the blue InGaN/GaN MQW LEDs, the emission photon energy is blue shifted by 24 meV when the injection current increases from 9 to 40 mA. For the green InGaN/GaN MQD LED, the EL peak position is blue shifted by 112 meV as the injection current increases from 26 to 70 mA. For the red InGaN/GaN MQD LED, the emission energy is blue shifted by 127 meV as the injection cur rent increases from 25 to 72 mA. Large blue shift in the EL peak position with increasing injec-tion current for an MOVPE grown blue InGaN/GaN MQD LED was also reported [138]. There are two possible reasons that can be attributed to this large EL blue shift with increase in the injection current in the MQD LEDs: (a) the band-filling effect [145]; and (b) the screening of the piezoelectric field by the high injection cur rent [146].

Light-emission efficiency enhancement from InGaN based LEDs can also be obtained by the use of surface plasmon polaritons (SPPs) supported by metallic films or nanostruc-tures [147–154]. This approach can be treated as a two-step process involving plasmon-enhanced spontaneous emission followed by scattering of the emitted SPPs into radiation [155]. First, energy is transferred from the excited electron–hole pairs in the active layer to the available SPP modes via near-field electromagnetic coupling. Given the large field enhance-ments and high densities of modes associated with SPPs, this process can be extremely fast and thus compete favorably with nonradiative recombination mechanisms, thereby lead-ing to enhanced IQE [156–158]. The key requirements in this respect are near-field proximity between the active layer and the metal surface, and close matching between the emission wavelength and the SPP resonance. Second, the emitted SPPs are extracted via scattering into radiation modes from a suit-able geometry providing the required phase matching. As a result of this two-step process, an overall increase in light emission efficiency can be obtained if the SPP extraction effi-ciency is larger than the emission efficiency in the absence of plasmonic nanostructures.

As an example, Henson et al [154] have used diffractive arrays of silver nanocylinders to increase the radiative effi-ciency of InGaN/GaN quantum wells emitting at near-green wavelengths. Large enhancements in luminescence intensity (up to a factor of nearly 5) were measured when the array period exceeds the emission wavelength in the semiconduc-tor material. The experimental results and related numerical simulations indicate that the underlying mechanism is a strong resonant coupling between the light-emitting excitons in the quantum wells and the plasmonic lattice resonances of the arrays. These excitations are particularly well suited to light-emission-efficiency enhancement, compared to localized sur-face plasmon resonances at similar wavelengths, due to their larger scattering efficiency and larger spatial extension across the sample area.

6.1.3. Efficiency droop. Besides the reduction of the LED efficiency in the green and yellow parts of the spectrum,

Figure 14. (a) Room temperature EL spectra of the blue, green and red LEDs under the injection of 20 mA; (b) EL peak position dependence on the injection current for the blue, green and red LEDs [140]. [140] John Wiley & Sons. Copyright © 2007 WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim.

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the efficiency of InGaN-based LEDs is also reduced at high operating currents, a property known as ‘efficient droop’ [112, 159–162]. Figure 15 shows the performance of blue and green-emitting high power InGaN/GaN MQW hetero-structure thin-film flip-chip (TFFC) LED Lamps [112]. The decrease in efficiency is slightly due to increase in forward voltage and heating, but the dominant factor is the decrease in internal quantum efficiency (IQE) as the current density is raised. Many proposals have been advanced to account for this efficiency droop in InGaN based LEDs. Some researchers proposed that the origin is carrier delocalization [163, 164], and others proposed enhanced Auger recombination [165] and electron leakage [166]. However, different methods of sample preparation and device structure as well as different measure-ment conditions and the interpretation of the data with differ-ent models has led to an uncertainty as to the exact origin of this phenomenon.

An excellent review of the various proposed models is presented by Piprek [160]. In this article we present only a brief discussion of the basic model for the IQE of InGaN-based LEDs. The reader is referred to the article by Piprek for a comprehensive review of the various proposed models. The internal quantum efficiency of an LED is given by the ratio of the injected current leading to radiative recombination in the QWs of the device versus the total injected current in the contacts of the device.

IQE =Irad

I=

Irad

Irad + Ilost. (6.1)

In this expression the total current was split into current that generates photons in the QW (Irad) and current that lost to other processes (Ilost). Efficiency droop occurs if Ilost increases much more than Irad as the injection current increases. Carrier losses can occur either inside the QWs or other parts of the device. Non-radiative recombination inside the QWs can occur through states in the middle of the energy gap by the

Shockley–Read–Hall (SRH) recombination (ISRH) or Auger recombination (IAuger) [167]. Carrier recombination outside the QWs are characterized as carrier leakage (Ileak). Thus, the total injection current can be written as:

I = Irad + ISRH + IAuger + Ileak. (6.2)

Carrier recombination inside the QWs is determined by the first three terms of equation (6.2)

IQW = Irad + ISRH + IAuger = qVQW(An + Bn2 + Cn3)

(6.3)Where q is the electronic charge, VQW is the volume of all QWs, and A, B, C are the recombination coefficients for the SRH, radiative, and Auger processes respectively. The leak-age current is difficult to describe by a simple expression. Ozgur et al [168] proposed a phenomenological expression for the leakage current density, Jleak = kJb, where J is the total current density and k and b are adjustable parameters. Piprek [160] has used a similar formula to relate the leakage current, Ileak, to the current IQW injected into the QWs.

Ileak = α (IQW)m

(6.4)

The combination of equations (6.1)–(6.4) leads to the follow-ing expression for the IQE:

IQE =ηinjBn2

An + Bn2 + Cn3 (6.5)

Where the injection efficiency ηinj represents the fraction of carriers that recombine inside the QWs

ηinj =IQW

I=

IQW

IQW + αImQW

. (6.6)

Thus, equation (6.5) takes also into account the carrier leakage.The majority of papers in the literature, which address the

influence of Auger recombination on the efficiency droop are neglecting carrier leakage and thus equation (6.5) becomes:

IQE =Bn2

An + Bn2 + Cn3 . (6.7)

The recombination coefficients for InGaN QWs extracted from measurements using this ABC model are of the order of A ~107 (s−1), B ~ 10−11 (cm3 s−1), and C ~ 10−30 (cm6 s−1) [160]. On the other hand theoretical calculations by Hader et al [169], based on a microscopic many body model and a k.p band structure model, leads to much smaller coefficient for the direct band to band Auger process (C = 3.5 × 10−34 cm6 s−1). Such small value makes the influence of Auger recombi-nation to the efficiency droop to be negligible.

Delaney et al [170] conducted first-principle density-func-tional and many-body-perturbation theory and confirmed the negligible probability for intra-band Auger recombination in InGaN with band gap greater than 2 eV. But they also identi-fied an inter-band Auger process, which exhibits a peak Auger coefficient of C = 2 × 10−30 in bulk InGaN with 2.5 eV band-gap, which could account for the efficient droop. However, the probability of this inter-band Auger recombination decreases rapidly with the changing of the band gap and thus it could not explain the efficient droop in InGaN LEDs observed over

Figure 15. External quantum efficiency and light output versus dc forward current for both blue and green InGaN-GaN QW-based TFFC LED lamps. Chip sizes are 1 × 1 mm [112]. Reproduced with permission from [112]. © Copyright 2007 IEEE.

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a wide range of wavelengths. Thus, the role of Auger recom-bination in accounting for the widely observed ‘efficiency droop’ is still being debated [161].

6.1.4. White LEDs. In current LED technology, white light is generated by combining an InGaN-based blue LED with a yellow phosphor or by combining a near-UV LED with a tricolor phosphor (blue, green and red). Thus, wavelength down-conversion is the process used for most solid-state light-ing (SSL) applications, since the efficiency of direct-emitting green, yellow, and amber LEDs, based on InxGa1−xN and Inx(GayAl1−y)1−xP, remains inferior to the efficiency of phos-phor-converted LEDs (pc-LEDs) for these colors. The reader is referred to more specialized articles addressing this topic [112, 113].

Currently, the majority of white light emission for various applications is realized by conversion using one (yellow in basic) or more phosphor components (green and red for higher quality white light). Whether it is the only phosphor present, or as a component in a phosphor blend, cerium-doped yttrium aluminum garnet [(Y1−xCex)3 Al5O12 or YAG:Ce3+)] is the most popular yellow phosphor used in pc-LEDs. Most often, YAG:Ce3+ is used in powdered form, where the powders are dispersed in a polymer matrix [171]. While highly cost-efficient, a disadvantage of the powder-in-polymer approach is that these phosphor powders are comprised of highly scattering micro-size particles that can lead to significant backscattering losses. More importantly, the powdered phosphors are almost always dispersed in polymers such as silicones or epoxies with low ther-mal conductivity, and these polymers prevent efficient heat dis-sipation, which limits the phosphor performance due to thermal degradation [172]. This problem is particularly troublesome in applications that require high light flux or operation at elevated temperatures (e.g. projection, automotive headlights, etc).

In order to overcome these limitations, scatter-free or low-scattering phosphors with strong blue absorption, high quantum efficiency, and efficient heat dissipation are needed. Recently, polycrystalline ceramic or single crystalline YAG:Ce3+ converters have been developed [173–177], and companies, such as OSRAM and Philips, have commercially available lighting technologies that make use of ceramic converters. While the internal efficiencies of YAG:Ce3+ sin-gle crystals or transparent ceramic converters can be quite high, they have lower light extraction efficiency than porous ceramic converters or conventional powder-in-silicone phos-phors due to the narrow light escape cone, which is defined by the total internal reflection (TIR) critical angle, θc = arcsin(nambient/nconverter) at the planar emitting surface relative to interface normal. For extraction into air, the critical angle is approximately 33° for YAG:Ce3+ (n = 1:82). In such a case, a significant portion of the light is reflected at the converter/air interface, which leads to the trapping of the converted light (the yellow emission) in guided modes where the radiation is either absorbed or emitted out from the plate edges. Only a

very small portion (1/2∫ θc

0 sin (θ) dθ ≈ 8%) can be extracted

from the top surface. Furthermore, from the viewpoint of blue light absorption, there is a decrease of absorption efficiency in such highly translucent phosphor plates compared to those

of higher scattering, which leads to the low generation of con-verted light [178]. For transparent ceramic or single crystal converters to be used in solid state lighting applications, it is essential to improve their light extraction efficiency while maintaining their high internal quantum efficiency.

One approach to improve the light extraction efficiency is to incorporate a photonic crystal lattice (PCL) on the top of the surfaces [179]. As mentioned above, for a smooth sur-face, light emitted from the phosphor comprises two angular regions. The first region is light emitted within the escape cone defined by θc. Outside the escape cone light is emit-ted in a set of sharp guided modes that propagate within the YAG slab. The number of these modes increases with the slab thickness and their sharpness is a function of the material losses and reflectivity of the interfaces. These guided modes are either lost by absorption or collected at the edges of the YAG slab. In the presence of a PCL, the photonic crystal acts as a diffraction grating and diffracts some of these guided modes into the air, increasing the directional light extraction efficiency [180]. Directly forming a PCL in YAG:Ce3+ is dif-ficult with standard reactive ion etching processes, since it is such a chemically and thermally stable material. Therefore, depositing a thin film of another material on the YAG:Ce3+ surface, offers a good alternative. SiO2, TiO2, and Si3N4 are popular photonic bandgap materials used in the formation of PCLs on phosphor plates [181–186] and they can be depos-ited on a large scale with relatively low-cost processes. In particular, the high transparency at visible wavelengths and the relatively large refractive index (~2.4) of TiO2 makes it a favorable material for making PCLs on YAG:Ce3+ ceramic converters. A high refractive index for the PCL material is favorable because the refractive index contrast between the PCL materials, i.e. TiO2 and air, is proportional to the strength of the scattering [179].

A systematic study of the effects of 2D hexagonal-lattice TiO2 photonic crystal layers on the directional light extraction efficiency of low-scattering YAG:Ce3+ ceramic converters has recently been reported by Sun et al [187]. Yellow directional light extraction enhancement by a factor of 4.4 is achieved with a 2D photonic crystal layer having a diameter of 430 nm, a lattice constant of 590 nm, and a height of 350 nm.

6.2. UV LEDs

AlGaN alloys are well suited for UV LEDs, because their energy gap can be tuned by changing the alloy composition to cover all three regions of the UV electromagnetic spectrum UV-A (340–400 nm), UV-B (290–340 nm) and UV-C (200–290 nm). An excellent references on this topic is the recently published book by Kneissl and Rass [16].

Ultraviolet light emitting diodes (UV-LEDs) are crucial for a number of applications such as water/air/food steri-lization, surface disinfection, free-space non-line-of-sight communication, epoxy curing, counterfeit detection, fluores-cence or Raman identification of biological/chemical agents, and various diagnostic and therapeutic medical applications. However, the development of UV LEDs and in particular the deep UV LEDs has not yet reached the same level as the

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visible spectrum LEDs. Despite intense efforts worldwide, the maximum external quantum efficiency (EQE) of fully packaged AlGaN-based deep UV-C LEDs is less than 10% as shown in figure 16 [16, 188–193]. This is to be contrasted with InGaN-based violet-blue LEDs, whose EQE efficiency is more than 50%.

6.2.1. Fundamental materials problems. Based on the data of figure 13, the AlGaN alloys appear to be inferior to those of InGaN alloys for emitter applications. Specifically, according to the data in this figure the IQE of AlGaN bulk films decreases monotonically from 0.5 to 0.05 as the AlN mole fraction varies from 0 to 55%. There are a number of fundamental problems with AlGaN alloys independent of growth method, which are responsible for the relatively low progress in developing effi-cient deep UV LEDs. Compositional inhomogeneities due to phase separation, observed in InGaN alloys, are not expected in AlGaN alloys, since the Al and Ga atoms have identical ionic radius in tetrahedral sides. Furthermore, AlGaN alloys, grown heteroepitaxially on sapphire substrates, generally suffer from high concentration of extended as well as point defects. There are multiple reasons for the high density of extended and point defects in heteroepitaxially grown AlGaN films. Principle amongst which is the high activity of nitrogen produced either by thermal decomposition of ammonia in the MOCVD method or by plasma decomposition of molecular nitrogen in the MBE method [194]. During epitaxial growth this active nitrogen reacts instantly with arriving Al atoms on the substrate and limits their diffusivity. This leads to the nucleation of small AlGaN islands and films with microstruc-tures consisting of small hexagonal columnar domains. Such materials are expected to have high dislocation density, since the plurality of the threading dislocations occur primarily at the boundaries of the hexagonal columnar domains due to their incomplete coalescence [61]. Furthermore, coalescence

of the small islands leads to tensile stress, which promotes the nucleation and propagation of cracks [195]. This is even more important in silicon doped high Al content AlGaN films, since silicon is anti-surfactant and thus, it leads to microstructures with even smaller domains [196].

Another source of the poor internal quantum efficiency (IQE) of deep UV LEDs is the incorporation of oxygen in AlGaN due to the high chemical affinity of aluminum for oxy-gen [197–199]. While oxygen is a shallow donor in GaN and InGaN alloys, it is known to form DX like centers in AlGaN alloys with high AlN mole fraction [200]. Other potential problems are the poor doping efficiency of n- and particularly p-AlGaN with high AlN mole fraction, which is responsible for the poor carrier injection in the active region of the device. This is due to the high ionization energies of Mg acceptors and Si-donors, which are 630 and 280 meV respectively for AlN [201].

Finally, another fundamental problem influencing the EQE in AlGaN-based deep UV LEDs is the difficulty in extract-ing the light from the device due to changes in the valence band structure of AlGaN alloys as the AlN mole fraction in the alloy increases. As shown in figure 17, the band structure of AlGaN alloys changes as a function of AlN mole fraction due to the difference in the crystal field splitting of GaN (+38 meV) and AlN (−219 meV). This reversal in the order of the valence bands in AlN and in AlGaN alloys as the AlN mole fraction increases imposes certain selection rules in the fun-damental optical transitions. Specifically, in GaN this leads to light polarized with the E ⊥ C-axis (TE mode, surface emis-sion), while in AlN the light is polarized with the E||C-axis (TM mode, edge emission) [202, 203].

6.2.2. Growth methods of improving the IQE of AlGaN alloys. Various approaches are currently been pursued to address these problems. A number of researchers reported that

Figure 16. External quantum efficiencies for AlGaN, InAlGaN-, and InGaN quantum well LEDs emitting in the UV spectral range [188]. Reproduced with permission from [188]. Copyright Research Gate.

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the IQE of AlGaN MQWs, grown by the lateral epitaxial over-growth method, increases as the dislocation density decreases from 1010 cm−2 to about 5 × 108 cm−2 in agreement with simulations [204]. Although these reports refer to edge dislo-cations, which as discussed previously are not recombination centers but traps, it is very likely that the density of screw dislocations, which are non-radiative recombination centers [118], decreases by the same amount as the edge dislocations. Monroy et al [205] have employed indium as a surfactant to promote two-dimensional (2D) growth during deposition of AlGaN alloys under slightly N-rich or stoichiometric con-ditions by plasma-assisted MBE. This growth mode should lead to larger islands and thus fewer dislocations as discussed previously.

An alternative approach is to develop methods of intro-ducing compositional inhomogeneities in the AlGaN alloys during their growth in analogy with InGaN alloys. Our group has reported that growth of AlGaN alloys by plasma-assisted MBE under more Ga than required for stoichiometric growth, does not affect the film stoichiometry [206]. In this growth mode the excess Ga accumulates as liquid in the surface of the growing film and thus, the active nitrogen and Al-atoms as well as intentional or unintentional impurities have to dissolve in the liquid Ga and incorporate into the growing film from the liquid phase. Thus, the growth mode has converted from phys-ical vapor deposition to liquid phase epitaxy (LPE) [198, 199]. This growth mode is likely to lead to lateral compositional inhomogeneities due to statistical fluctuations of the thickness of the liquid Ga on the surface of the growing AlGaN film. These compositional inhomogeneities lead to band structure potential fluctuations, which are sufficiently deep that lead to exciton localization even at room temperature. This prevents the carriers from diffusing to point or extended defects and to recombine non-radiatively.

Also, in the proposed LPE growth mode the incorporation of impurities in the AlGaN film requires that that their solubil-ity in Ga at the growth temperature to be relatively high. We have reported previously that the concentration of impurities such as O, H and C is in the 1018–1019 cm−3 range when the GaN films are grown under nitrogen rich conditions, while they are two to three orders of magnitude less when the films are grown under Ga-rich conditions [198, 199]. The reduction in oxygen incorporation can be accounted for by the formation of volatile gallium oxides, which would then desorb. Similarly,

hydrogen may form volatile gallium hydrides. Regarding the reduction in carbon impurities one has to assume that the solu-bility of carbon in Ga must be very low, since gallium carbides compounds are not known to exist.

Bhattacharyya et al [207] investigated the effect of the excess liquid-Ga on the IQE of AlGaN MQWs by growing a series of identical in thickness and composition Al0.70Ga0.30N/AlN MQWs on the c-plane of sapphire substrates and stud-ying their photoluminescence efficiency as a function of temper ature. Specifically, during the growth of the wells the flux of Ga was varied from that corresponding to stoichio-metric conditions (III/V ~ 1) to (III/V 1). The thickness of the barriers and wells were 6 nm and 1.5 nm respectively. The nor malized luminescence spectra of these samples are pre-sented in figure 18(a). These data show that all MQW samples exhibit only a sharp near band-edge emission. It is important to stress that although the thicknesses and compositions of the wells and barriers were the same for all samples, the emis-sion wavelengths varies from 225 nm to 250 nm as the Ga-flux increased from III/V close to one to much greater than one. This red shift of the PL spectra for these identical MQWs indicates that the excess Ga during the growth of the wells introduces band structure potential fluctuations whose depth increases with the amount of excess Ga. Consistent with this interpretation is also the increase of the spectral width from 11 nm to 16 nm. These potential fluctuations are much deeper than the statistical ones due to alloy disorder and thus, they can cause carrier localization and efficient radiative recombi-nation even at room temperature.

The IQE of these MQWs was determined by measuring the photoluminescence spectra as a function of temperature, using the method described by Chichibu et al [115], and the results are shown in figure 18(b). Thus, the IQE for these MQWs varies from 5% for peak emission at 225 nm, to 50% for peak emission at 250 nm. This increase in the IQE for identical Al0.7Ga0.3N/AlN MQWs is attributed to the localiza-tion of the excitons due to band structure potential fluctua-tions introduced during the growth of the wells under Ga-rich conditions.

Zhang et al [208] also investigated the growth and prop-erties of deep UV emitting AlGaN/AlN MQWs, grown on the c-plane of 6H-SiC substrates. Growth of AlGaN alloys and QWs on SiC has the advantage that the lattice mismatch between SiC and AlN is only ~1%. These authors reported

Figure 17. The band structure near the Γ point of AlxGa1−xN alloys for (a) x = 0, (b) x = 0.25, and (c) x = 1 [202]. Reprinted with permission from [202]. Copyright (2004), AIP Publishing LLC.

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that such Al0.7Ga0.3N /AlN MQWs emitting at 245 nm have an IQE of 68% [208].

Based on this discussion, it appears that the IQE of AlGaN based deep UV LEDs can be resolved either by using methods to reduce the density of dislocations [204, 205] or by growing the active region of the LED under conditions, which intro-duce band structure potential fluctuations [196, 207–212]. Thus, the light extraction efficiency is currently what limits the EQE efficiency of the deep UV LEDs.

7. Lasers

The demonstration of blue–violet diode lasers by Nakamura and coworkers at Nichia in 1996 [40] represents a major achievement in the field of III-nitride semiconductor opto-electronics. Prior to this work, the only semiconductor lasers emitting at blue wavelengths were based on II–VI materials and suffered from severe reliability issues, so that their poten-tial for commercialization was limited. On the other hand, in a short time span after their initial demonstration, InGaN MQW lasers emitting in the 400-nm wavelength range have reached sufficient levels of performance and reliability to become attractive for practical applications [40, 213–219]. In fact, an operating lifetime of over 10 000 h was reported as early as in 1997 [214]. The main driving market of these devices is high-density optical data storage, as well as applications in the areas of full-color displays, projectors, and printers. As a clear indication of their commercial impact, the develop-ment of III-nitride diode lasers has enabled the creation of the Blu-ray Disc format providing unprecedented capacity for the entertainment industry [217].

In most high-performance III-nitride lasers, the gain medium consists of InGaN/GaN MQWs with relatively low InN mole fraction, although GaN/AlGaN or AlInGaN MQW laser struc-tures have also been developed. As already pointed out in this review, a unique property of III-nitride semiconductors is the

ability to tune their emission wavelength from the near-infrared to the UV range by varying the alloy composition. As a result, extensive research efforts are currently devoted to extending the spectral reach of lasers based on these mat erials. On the long-wavelength side, green-emitting diode lasers have been dem-onstrated using higher-In-content InGaN/GaN MQW active materials [220–226], with a maximum emission wavelength above 530 nm. Similar to the case of green LEDs, the main challenge here is related to the decrease in structural quality and the increase in piezoelectric polarization with increasing InN mole fraction in InGaN/GaN QWs. To address the latter issue, semi-polar GaN substrates are often used for the growth of these devices [221, 222, 224, 226]. In addition, the refractive-index contrast between the core and cladding layers of a typical InGaN/GaN laser waveguide decreases with increasing wave-length, which causes a reduction in the optical confinement fac-tor and therefore an increase in the laser threshold.

At the same time, the development of nitride UV lasers is also being widely investigated [227–236], motivated by applications in bio-chemical sensing, material processing, water treatment, and non-line-of-sight optical communica-tions. The shortest room-temperature emission wavelength reported to date under electrical injection is 336 nm, which was obtained with an AlGaN/AlGaN MQW gain medium [234]. Laser emission at even shorter wavelengths is possi-ble using AlGaN alloys of larger bandgap energy; however at present it remains a considerable challenge mostly due to the well-known difficulty in doping AlGaN alloys of high Al content, as required for the purpose of electrical injection. In fact, room-temperature lasing at deep-UV wavelengths down to 214 nm has already been demonstrated with high-quality AlN layers using optical pumping [231]. Optical gain at wavelengths below 250 nm has also been measured with the variable-stripe-length method in AlGaN/AlN QWs featuring strong band-structure potential fluctuations [210, 211, 237], grown by plasma-assisted MBE under Ga-rich conditions as described previously in this article.

Figure 18. (a) Room-temperature luminescence spectra from Al0.7Ga0.3N/AlN MQWs, with identical well and barrier widths as described in the text; (b) room temperature IQE of these MQWs [207]. Reprinted with permission from [207]. Copyright (2009), AIP Publishing LLC.

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Due to their advanced level of development and wide range of applications, a comprehensive description of III-nitride diode lasers is beyond the scope of this review. In fact, entire books on the subject have already been published [3, 7, 238] and can be consulted for further details. The following discus-sion is limited to a brief review of the more distinctive aspects of the physics of these devices, as compared to lasers based on more ‘traditional’ III–V semiconductors (i.e. As- or P-based materials). Perhaps the most important question in the context of III-nitride diode lasers is related to the very nature of their optical gain. In particular, it has been argued that stimulated emission in InGaN QWs arises from localized excitons cre-ated by compositional and layer-thickness fluctuations and/or by phase separation of InN and GaN [119]. These structural phenomena produce potential energy minima providing 3D car-rier confinement (similar to the case of quantum dots), where a population inversion of bi-excitons can be established. The carriers confined in such energy minima are also prevented from nonradiative recombination at dislocations, which would explain the surprisingly high quantum efficiency of InGaN QWs despite their large defect densities [239]. In any case, this picture remains a subject of some controversy and the opposing view of optical gain due to interband recombination (as in more traditional III–V semiconductors) has also been argued [240].

Regarding the active material of typical III-nitride lasers, QWs are commonly used because of their reduced volume and 2D density of states, which result in lower threshold cur-rent density and larger differential gain compared to analogous bulk gain media [241]. It should be noted, however, that several characteristic features of III-nitride semiconductors partly limit the advantages that are commonly associated with QWs in the context of diode lasers. First, due to the large effective masses of these materials, significant quantization effects only occur in very narrow QWs, which are extremely sensitive to the quality of the hetero-interfaces. Narrow QWs are also needed to mini-mize the quantum confined Stark effect, which would otherwise strongly decrease the oscillator strength and therefore increase the laser threshold. Second, because the ‘heavy-’ and ‘light-hole’ (usually denoted A and B) valence bands have comparable effective masses along the c direction, no significant splitting between the A and B subbands occur in (c-plane) QWs. As a result, the use of these QWs does not produce a reduced density of valence-band states near the emission wavelength, as is com-monly observed in the case of As-based materials and which is very desirable for the purpose of decreasing the laser threshold. Finally, the application of strain in (c-plane) nitride QWs does not appreciably decrease the hole effective masses and therefore the valence-band density of states, because of their small spin-orbit splitting [242]. Therefore, unlike the case of diode lasers based on zincblende materials, the performance of III-nitride QW lasers does not substantially benefit from the presence of strain in the active layer. Incidentally, some of these considera-tions (in addition to the absence of polarization-induced intrin-sic electric fields) also motivate the ongoing efforts to develop nitride QW lasers grown along nonpolar directions [243, 244].

Regarding the device structure, a schematic cross-sectional view and band lineup of a typical in-plane InGaN/GaN MQW laser are shown in figure 19 (where band-bending due to the

intrinsic electric fields is not included for simplicity). Here the fundamental optical mode is confined along the growth direc-tion by the refractive-index discontinuity of the n-GaN/n-AlGaN and p-GaN/p-AlGaN interfaces immediately below and above the MQW region, respectively. Lateral confinement, on the other hand, is provided by the ridge structure etched through the upper cladding layer. A p-doped GaN/AlGaN superlattice is often used in place of AlGaN in this layer, for the purpose of increasing the fraction of ionized acceptors and therefore the hole density. As in typical LEDs, an undoped AlGaN electron-blocking layer is also grown over the QWs to avoid electron overflow outside of the active region at high bias currents. Finally, the top p-GaN and bottom n-GaN serve as the contact layers. If the growth sub-strate itself is conducting (e.g. n-doped free-standing GaN), the n-metal contact can be deposited on its bottom surface.

Initial nitride lasers were mostly grown on sapphire, possi-bly using epitaxial lateral overgrowth (ELO) in order to reduce the dislocation density, whereas more recently SiC as well as free-standing GaN substrates have also been employed. Besides its large lattice and thermal mismatch with GaN, (c-plane) sapphire has the disadvantage of not sharing any common cleavage plane with the overlaying epitaxial films. As a result, high-quality laser facets in devices grown on sap-phire cannot be obtained simply by cleaving, which provides an additional fabrication challenge. Furthermore, due to the relatively large refractive-index difference between sapphire and GaN, a complex pattern of higher-order transverse modes can be supported by the laser epitaxial film, leading to unde-sirable multi-lobe far-field patterns. Finally, it should be noted that the use of ELO techniques and/or substrates providing a better lattice match is beneficial from the point of view of laser reliability, as a direct correlation between reduced dislocation density and increased device lifetime has been observed [245].

p-AlGaN

substrate

n-GaN

n-AlGaN

p-contact

n-contact

InGaN/GaN MQWsn-GaN

p-GaN

p-GaN

Ec

Ev

Figure 19. Schematic cross-sectional view (top) and band lineup (bottom) of a typical InGaN/GaN MQW diode laser (not drawn to scale). In the band diagram, band bending due to the intrinsic electric fields is not included for simplicity.

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8. UV Photodetectors

8.1. Introduction

Historically, the development of high quantum efficiency semiconductor UV detectors has been hampered by the extremely strong absorption and radiation induced aging effects in most semiconductor materials. The AlGaN material system is well suited for UV photodetectors because its direct bandgap can be tuned from 360 nm to 200 nm, by changing the alloy composition and thus, enabling true visible-blind or solar-blind detectors. As semiconductor devices, such detec-tors are lightweight, compact, have low power requirements and low internal losses. In addition nitride semiconductors exhibit good electronic transport properties; they are physi-cally robust, chemically inert, have high corrosion resistance and are non-toxic. These properties make them also attrac-tive for use in hostile environments and at high temperatures. AlGaN alloys have been used to fabricate either photocon-ductive or photovoltaic UV detectors [246–255]. Excellent reviews on UV photodetectors were presented by Razeghi and Rogalski [248], Campbell et al [250] and Monroy et al [252].

8.2. Photoconductive detectors

8.2.1. GaN ‘visible blind’ UV photoconductive detectors. In this section we discuss ‘visible blind’ photoconductive detec-tors based on GaN thin films with interdigitated Ohmic contacts placed on the surface of the material to maximize light transmis-sion while minimizing transit time. When photons with energy greater than the band-gap are incident on it, they are absorbed and generate electron–hole pairs within the material. An elec-tric field, applied to a photoconductor by an external voltage source, causes the carriers to move towards the electrodes cre-ating a current that is proportional to the incident photon flux.

In the following we discuss the properties of a number of such GaN photoconductive detectors produced by plasma-assisted MBE [247, 251, 254]. A number of n-type auto-doped GaN films were grown by this method and their resistivity was varied from 10 to 107 (Ω · cm) by varying the ratio of the Ga to active nitrogen flux produced by the plasma discharge. The relatively high resistivity of these films is due to reduc-tion of oxygen incorporation during growth under excess Ga. These films were fabricated into photoconductive detectors by depositing lithographically interdigitated electrodes with elec-trode spacing ranging from 5 µm to 20 µm. The Ohmic metal contacts were formed by depositing Ti (20 Å)/Al (2000 Å)/Au(5000 Å) by e-beam evaporation and annealed to 850 C to form Ohmic contacts.

8.2.1.1. Spectral response. The spectral response of a typical GaN photodetector produced by this method is shown in fig-ure 20. The photoresponse shows a sharp increase at 365 nm, corresponding to the bandgap of GaN, and visible light rejec-tion of more than three orders of magnitude. The spectral response of GaN is quite remarkable because it remains con-stant for wavelengths down to 200 nm, which was the limit of our measurement apparatus. This needs to be contrasted with conventional semiconductors, like Si and GaAs, where

the photoresponse drops dramatically at short wavelengths, a result attributed to surface recombination in these materials. The constant photoresponse at short wavelengths in GaN is evidence that surface recombination is less important in GaN, consistent with our earlier conclusion that the surface states in GaN have moved towards the band edges, where they act as traps rather than recombination centers.

8.2.1.2. Photoconductive gain. The photocurrent in photo-conductive detectors is given by the expression:

Ip = qη(

Popt

hν

)(µτVd2

)= qη

(Popt

hν

)G (8.1)

Where Ip is the measured photocurrent, q is the electronic charge, η is the quantum efficiency, Popt is the optical power, hν is the photon energy, µ is the electron mobility, τ is the carrier lifetime, V is the applied bias voltage, d is the inter-electrode spacing, and G is the photoconductive gain. The photoconductive detectors have gain because the minority carriers are trapped while the majority carriers go around the circuit many times before recombination. According to equa-tion (8.1) the gain of a photoconductive detector of known dimensions and at a known applied voltage is a measure of the mobility-lifetime (µτ) product of the material, which is the real figure of merit of the material.

Figure 21 shows the mobility-lifetime (µτ) products meas-ured for a number of n-GaN detectors as a function of film resis-tivity. Thus the µτ product varies monotonically from 10−2 cm2 V−1 to 10−7 cm2 V−1 as the film resistivity changes from 10 to 107 (Ω · cm). Based on Hall Effect measurements performed on these materials, this large variation in µτ products is attributed mainly to variation in electron lifetime. Films with the higher gain correspond to longer lifetimes and films with relatively smaller gains correspond to shorter lifetimes. The behavior described in figure 21 is typical of photoconductive detectors, for which the gain-bandwidth product remains constant.

8.2.1.3. Responsivity. Responsivity of a detector is the ratio of the photocurrent flowing through the detector divided by

Figure 20. Spectral response of GaN photoconductive detector fabricated on an n-GaN film grown by plasma-assisted MBE [254]. Reproduced with permission from [254]. © (2007) COPYRIGHT SPIE–The International Society for Optical Engineering.

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the incident optical power. Thus, according to equation (8.1) the responsivity is given by the expression:

R = qηGhν

. (8.2)

The responsivity of the various GaN detectors was deter-mined by measuring the photocurrent using a lamp with emis-sion at 254 nm. The photon flux intensity of the lamp was determined by measuring the photocurrent from a calibrated UV enhanced silicon photodiode (Hamamatsu S5226-8BQ), illuminated under identical conditions as the photodetectors. The UV responsivity of a number of GaN detectors measured at 254 nm as a function of detector dark current is shown in figure 22. These responsivities are significantly larger than the responsivities for commercial Si (0.1 A W−1 at 200 nm), GaP (0.03 A W−1 at 254 nm) and GaAsP (0.035 A W−1 at 254 nm) UV photovoltaic detectors [256]. Of course the larger respon-sivity is related to the fact that photoconductive detectors have gain. Based on these results, the investigated detectors may be broadly divided into two categories. The first type has low

dark resistivity and high responsivity and the second type has high resistivity and relatively lower responsivity.

8.2.1.4. Dynamic range. Dynamic range is a detector param-eter that indicates the range over which the detector response remains linear with respect to the incident light intensity. The effect of decreasing the light intensity is to move the quasi-Fermi levels of electrons and holes closer to each other towards the middle of the gap and thus, probe the interaction of the trap levels and recombination centers within the gap of the semiconductor [257]. Thus, the measurement of photocon-ductivity as a function of light intensity is a powerful tool to investigate the distribution of traps and recombination centers in the semiconductor bandgap. If we denote Ip = f γ , where Ip is the photocurrent and f is the generation rate, then the exponent γ is a measure of the linearity of the photoresponse.

The dynamic range of a number of n-GaN photodetectors was investigated by varying the intensity of the incident laser beam (He–Cd, 10 mW at 325 nm) over several orders of mag-nitude using calibrated neutral density filters. In semi-insulat-ing GaN films the response was found to be practically linear (γ = 0.93) over five orders of magnitude as shown in figure 23.

In general, the exponent was found to vary from γ = 0.5 for conducting films to γ = 1 for semi-insulating films as shown in figure 24. The dependence of the exponent γ on the position of the dark Fermi level for the investigated films is shown in figure 25. The variation of γ between 0.5 and 1.0 can be accounted for by considering an exponential distribution of traps from the band edges using a model first proposed by Rose [257]. According to this model this type of light intensity dependence can be accounted for by considering a single type of recombination center lying below the dark Fermi level and an exponential distribution of electron traps extending from the conduction band edge. The exponential electron trap dis-tribution is expressed as:

Nt = N0exp(−Ec − Et

kT1

). (8.3)

Figure 22. UV responsivity as a function of detector dark current for several n-GaN photodetectors [254]. Reproduced with permission from [254]. © (2007) COPYRIGHT SPIE–The International Society for Optical Engineering.

Figure 23. Photoresponse as a function of excitation intensity for a GaN photoconductive detector fabricated from semi-insulating GaN films [254]. Reproduced with permission from [254]. © (2007) COPYRIGHT SPIE–The International Society for Optical Engineering.

Figure 21. Mobility-lifetime products as a function of film dark resistivity for n-GaN photodetectors fabricated on films produced by plasma-assisted MBE [254]. Reproduced with permission from [254]. © (2007) COPYRIGHT SPIE–The International Society for Optical Engineering.

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The characteristic temperature T1 can be adjusted to make the exponential density of states to vary more or less rapidly with energy.

As the light intensity increases, more and more of the Nt states are converted from trapping states to recombina-tion centers. The conversion occurs as the quasi-Fermi level sweeps through the Nt states towards the conduction band. As the density of recombination states increases, the electron life-time decreases. The density of the new recombination cent-ers can be calculated by integrating equation (8.3) between the original Fermi level EF and the quasi-Fermi level EFn. Knowing the density of recombination centers the electron lifetime can be calculated. Thus, the steady state value of the excess carriers, if the generation rate is f, is given by [257]:

∆n = f τn ∼ f T1/(T+T1). (8.4)

Thus, according to this model if T1 is equal to or larger than T, the exponent γ = T1/ (T + T1) lies between 0.5 and 1.0.

In summary for low resistivity films, the dark Fermi level lies close to the conduction band. Upon illumination, the quasi-Fermi level sweeps through only a small range of energies towards the conduction band. Thus the number of traps that convert to recombination centers is small. In addition, states closer to the band edge are less efficient as recombination cent-ers. Therefore, the electron lifetimes are long, leading to higher values of µτ products. For high resistivity films, the dark Fermi level lies deeper in the bandgap. Upon illumination the quasi Fermi level moves closer to the conduction band, sweeping through a large band of energy. Thus, the concentration of states which are converted from traps to recombination centers is larger. Furthermore, these states being closer to the middle of the gap are more efficient recombination centers, yielding shorter lifetimes for electrons and thus, smaller µτ products.

Similar behavior was observed by Bube and Redfield for amorphous hydrogenated silicon [258]. The exponential dis-tribution of traps in amorphous silicon was attributed to dis-order and dangling bonds. In the GaN films the exponential distribution of traps has its origin to potential fluctuations pos-sibly due to inhomogeneous distribution of impurities, point defects and extended defects like stacking faults.

8.2.2. AlGaN ‘solar blind’ UV photoconductive detec-tors. AlGaN-based UV detectors are making strides in detecting UV radiation from 360 nm to x-rays, as well as alpha particles [247]. Such detectors were found to have improved sensitivity, high spectral selectivity and low noise. Specifi-cally, AlGaN-based UV detectors are finding use in areas such as the detection of UV flames for combustion control, surveil-lance of rockets and intercontinental ballistic missiles, secure space-to-space communication, detection of UV scintillation for medical imaging, monitoring of pollutants like nitrous oxide (NO2) and sulfur dioxide (SO2) in the ionosphere, space-based instrumentation for UV astronomy and in UV photolithography for semiconductor processing.

As discussed previously in section 6, AlGaN alloys tend to undergo partial long range atomic ordering. In such alloys Iliopoulos [130] has shown that the energy gap is reduced as the degree of ordering increases. The observed reduction of the band gap with ordering was confirmed by Dudiy and Zunger using first-principal calculation combined with large-scale atomistic empirical-pseudopotential simulations [259].

A number of partially ordered undoped AlGaN films were fabricated into photoconductive detectors using the method described previously. The resistivity of these AlGaN films was found to increase by 7 orders of magnitude as the AlN mole increases from 0 to 50%. The (µτ) product of these detectors was calculated from the measurement of the photocurrent and the results are shown in figure 26 [131]. Plotted in the same figure are the (µτ) product for n-GaN detectors taken from figure 21. Thus, the (µτ) product for the AlxGa1−xN alloys vary from 10−2 to 10−5 cm2 V−1 as the composition of the films is changed from x = 0 to x = 0.5.

From the data of figure 26 it is apparent that that in both GaN and AlGaN the gain decreases monotonically as the resistivity of the films increases. In both cases, the variation in the µτ product with resistivity is attributed mainly to changes

Figure 25. Variation of the exponent γ as a function of the position of the dark Fermi level [254]. Reproduced with permission from [254]. © (2007) COPYRIGHT SPIE–The International Society for Optical Engineering.

Figure 24. Photoresponse as a function of excitation intensity for a number of GaN photoconductive detectors fabricated from n-GaN films of various resistivities [254]. Reproduced with permission from [254]. © (2007) COPYRIGHT SPIE–The International Society for Optical Engineering.

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in the lifetime of the photo-generated carriers. Specifically, the change in the lifetime with changes in the dark resistivity is attributed to the observation that the films have an exponen-tial distribution of defects states and that the relative density of occupied states and recombination centers changes as the quasi-Fermi levels change upon illumination.

It is important to note that the AlGaN alloys have higher (µτ) products than GaN even though they are expected to have lower mobility than GaN due to alloy scattering and shorter lifetime than GaN due to their more defective nature. Misra et al [131] accounted for this anomaly by postulating that the ordered and random AlGaN domains form a Type-II hetero-structure as indicated in figure 27. In such band alignment the photogenerated electron–hole pairs are physically separated in the conduction the valence bands of the ordered and ran-dom domains respectively leading to an increase in recombi-nation lifetime by orders of magnitude.

Consistent with these experimental results Dudiy and Zunger predicted by atomistic empirical pseudopotential sim-ulations that the band alignment between random and ordered domains in AlxGa1−xN alloys changes from Type I to type II at x = 0.4 leading to an increase in recombination lifetime of electrons and holes [132]. Recombination occurs only when both electrons and holes occupy the same domain by thermal emission. Similar results of ordering on carrier lifetimes have been reported for Ga0.47In0.53As alloys [260].

In conclusion, photoconductive detectors based on GaN and partially ordered AlGaN alloys were fabricated end evaluated. In both cases the mobility-lifetime product, which is an intrin-sic figure of merit of the photoconductive gain, were found to decrease monotonically with the film resistivity. In both cases the decrease of lifetime as the material becomes more insulating is attributed to exponential band tails due to potential fluctua-tion arising from defects or the growth mode. In the case of par-tially ordered AlGaN alloys, the mobility-lifetime product for films with AlN mole fraction close to 50% is about two orders of magnitude higher than that of GaN films with comparable resistivity. This result was accounted for by the longer lifetime of the photo-generated carriers in the partially ordered alloys.

8.3. Photovoltaic UV detectors

Such detectors based on AlGaN have been fabricated in the form of Schottky barriers, p-i-n diodes, photo-transistors, avalanche photodiodes [248, 250, 252]. For example, ‘visible-blind’ GaN p-i-n photodiode arrays with peak responsivity of 0.11 A W−1 at 360 nm corresponding to 48% internal quantum efficiency have been reported, with visible rejection of 3–4 orders of magnitude in the 400–800 nm spectral range [261]. Furthermore, by alloying GaN with AlN one can design these devices to be completely ‘solar-blind’. To design such devices one has to take into account that there is a reduction in the solar radiation in the 240 nm–280 nm spectral region due to atmospheric absorption. Thus, for an AlGaN p-i-n diode, the composition of the n-layer should be chosen to absorb any UV light at wavelengths shorter than 240 nm. Also, the i-layer is chosen to have an absorption edge at 280 nm [262].

9. UV electroabsorption modulators

9.1. Introduction

Electroabsorption modulators are semiconductor electro-optic devices in which a change in absorption coefficient Δα (or refractive index Δn) is induced by an externally applied electric field [263]. In the past several years, near-infrared electroab-sorption modulators based on zincblende III–V compounds (arsenide, phosphide and antimonide semiconductors) have been developed with impressive performance metrics, for a variety of applications in fiber-optic communications such as data transmission, photonic switching and optical intercon-nects [263]. In general, particularly strong modulation can be obtained when the absorption edge is dominated by excitonic effects, due to the sharp nature of the resulting absorption fea-tures. However, in mostly covalent smaller-bandgap bulk semi-conductors such as GaAs the exciton binding energy (~4 meV) is substantially less than the thermal energy kBT at room temper-ature. As a result, the excitonic nature of the absorption edge at room temperature is not obvious due to thermal broadening, and the associated benefits for electroabsorption modulation are

Figure 27. Schematic representation of a potential band structure in a photoconductive detector based on partially ordered AlGaN alloys [131]. Reprinted with permission from [131]. Copyright (1999), AIP Publishing LLC.

Figure 26. Variation of mobility-lifetime product of AlGaN and GaN photodetectors with film resistivity [131]. Reprinted with permission from [131]. Copyright (1999), AIP Publishing LLC.

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lost. On the other hand, the exciton binding energy becomes substantially larger in quantum-well (QW) structures, lead-ing to well resolved excitonic absorption peaks even at room temper ature. High-performance optical modulators have there-fore been developed over the years based on the quantum con-fined Stark effect (QCSE) in QWs, where large changes in the excitonic resonance are obtained through the application of an electric field along the growth direction [263–265].

On the other hand, visible/ultraviolet (UV) electroabsorp-tion modulators based on wide-bandgap III-nitride semicon-ductors have so far received limited attention [266–272], despite a number of promising potential applications. Of par-ticular interest is their development for non-line-of-sight free-space optical communications based on atmospheric light scattering [273], where the use of short-wavelength radiation is advantageous due to its large atmospheric scattering cross-section. External optical modulators in these systems would allow for higher transmission rates, without the deleterious transient heating effects that are typically associated with direct current modulation of semiconductor light sources. Nitride electroabsorption modulators incorporated within a laser cavity could also be used for the generation of short pulses of visible/UV radiation via Q-switching [266].

A basic property of nitride semiconductors that is par-ticularly important in this context is provided by their large exciton binding energies (about 25 meV in GaN and even higher in ternary AlGaN alloys). This is a direct consequence of the heavy electron and hole effective masses of these mat-erials, which in turn are directly related to their large bandgap energies. As a result, even in bulk samples at room temperature the optical absorption edge is dominated by excitonic effects [274], so that strong electroabsorption of near-bandgap radia-tion can be expected. The first UV optical modulator based on a 0.4 µm-thick GaN film grown by metal-organic chemical vapor deposition (MOCVD) has been reported by Oberhofer et al [267]. However, this device was found to require a pro-hibitively large applied voltage (>80 V) to produce any appre-ciable change in transmission, with a maximum modulation depth under normal-incidence operation of 18% at 305 V bias. Significantly stronger electroabsorption modulation at similar wavelengths (near 350 nm) has been demonstrated based on the QCSE in GaN/AlGaN QWs [268]. In particular, a com-parable normal-incidence modulation depth of about 20% was obtained in this device with a modest reverse bias of only 10 V. Similar devices operating at blue/violet and deep-UV

wavelengths have also been demonstrated based on InGaN/GaN and GaN/AlGaN QWs, respectively [269, 270].

9.2. Electroabsorption modulators based on bulk GaN films

The device described in this section was grown by RF plasma-assisted MBE on (0 0 0 1) sapphire [271]. A schematic cross-sectional view of the device is shown in figure 28(a). A relatively thick (0.5 µm) AlN film was initially grown on the substrate, so that all subsequent epitaxial layers are under compressive strain, which reduces their probability of developing cracks. A transparent contact layer consisting of Si-doped n-Al0.16Ga0.84N was then deposited, followed by a nominally intrinsic Al0.3Ga0.7N film, whose function is to electrically isolate the GaN active region from the bottom contact layer. The GaN active region is also undoped and has a thickness of 0.4 µm.

Prior to the device fabrication, the material optical absorp-tion spectrum was determined via transmission measurements at room temperature and the results are shown in figure 29. These data clearly show that even without cryogenic cool-ing the absorption edge of this bulk sample is dominated by excitonic effects, leading to a sharp peak at a photon energy of about 3.47 eV. The abrupt increase in absorption at about 3.7 eV is due to the 1 µm-thick Al0.16Ga0.84N film.

Figure 28. (a) Schematic cross-sectional view of a bulk GaN electroabsorption modulator; (b) top-view micrograph of a fabricated device [271]. Reprinted with permission from [271]. Copyright (2011), AIP Publishing LLC.

Figure 29. Absorption spectrum of the electroabsorption modulator shown in figure 1(a) [271]. Reprinted with permission from [271]. Copyright (2011), AIP Publishing LLC.

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Electroabsorption modulators based on the device geometry of figure 28(a) were fabricated using standard photolithographic techniques. 800 µm × 800 µm mesa structures were first formed by inductively-coupled plasma etching in a chlorine-based chemistry. A Ti/Al/Ni/Au multilayer film was then deposited by electron-beam evaporation on the exposed n-Al0.16Ga0.84N around each mesa, and annealed at 850 °C to form an Ohmic contact. Finally, a thin semi-transparent Ni/Au film was depos-ited over the top surface of each mesa, with a thick Au pad added in one corner for wire bonding. The resulting junction with the underlying GaN film is of the Schottky type, so that the external field across the active layer can be controlled through the appli-cation of a reverse bias on the Ni/Au contact. A micrograph of a fabricated modulator is shown in figure 28(b).

The measured transmission spectra through each device were normalized to similarly measured transmission spectra through a sapphire substrate. The resulting traces for a repre-sentative GaN modulator under various reverse bias volt ages from 0 to 14 V are shown in figure 30(a). A zoom-in of the transmission spectra near the excitonic resonance is shown in the inset of figure 30(a). As the applied voltage is increased,

the excitonic absorption resonance is broadened and quenched, leading to an increase in transmission near the exciton peak and to a decrease in transmission at sufficiently detuned photon energies. From these traces we obtain a maximum modulation depth M of about 30% at a photon energy of about 3.45 eV, where M is defined as the ratio [T(V) – T(0)]/T(0) and T(V) is the device transmission as a function of bias voltage V.

The measured transmission spectra can also be used to calculate the corresponding changes in absorption coefficient Δα(V) = α(V) − α(0) versus photon energy, for different val-ues of the applied voltage. Specifically, since T(V) is propor-tional to e−α(V)d, where d is the thickness of the absorbing layer (0.4 µm in this case), we can write

∆α (V) = −1dln

[T(V)

T(0)

]. (9.1)

Several spectra of Δα obtained with this procedure are shown in figure 30(b). The maximum change in absorption coeffi-cient here is found to increase with applied voltage up to a peak (absolute) value of about 7 × 103 cm−1 at V = 12 V. As the reverse bias is further increased, a non-negligible amount of leakage current begins to flow across the active layer. This leads to a thermal modulation of the band edges via resistive heating, which tends to compensate the field-induced changes in the absorption edge. As a result, no further increase in Δα is obtained at higher voltages. The origin of this leakage cur-rent is likely due to structural defects in the epitaxial material and/or surface states on the mesa sidewalls. In any case, the results presented in figure 30 clearly demonstrate the potential of nitride bulk films for electroabsorption modulation, consis-tent with their uniquely large exciton binding energies.

9.3. Electroabsorption modulators based on GaN/AlGaN quantum wells

Electroabsorption modulators based on various multiple-QW structures were also developed, using the same growth, fabri-cation, and characterization procedures described in the previ-ous section. In the following we present results obtained with the device structure shown schematically in figure 31. Here the bottom contact layer consists of a Si-doped n-Al0.15Ga0.85N film, and the active region comprises 50 pairs of nominally undoped 40 Å GaN wells and 50 Å Al0.2Ga0.8N barriers. The overall thickness of this multiple-QW layer (0.45 µm) is therefore similar to that of the GaN film in the device structure of figure 28. Once again a semi-transparent Ni/Au Schottky contact, deposited over the topmost barrier, is employed to modulate the voltage across the active region.

As already discussed, QW electroabsorption modulators rely on the QCSE, whereby an external electric field along the growth direction is used to change the exciton oscillator strength and simultaneously shift the absorption edge. An important difference between the present GaN/AlGaN devices and similar modulators based on zincblende III–V compounds is related to the characteristic intrinsic electric fields of nitride QWs grown along the (0 0 0 1) direction. These fields are due to the presence of spontaneous and piezoelectric polarizations of different magnitudes in the well and barrier layers (leading to the creation of surface charge on the heterointerfaces), and

Figure 30. (a) Normalized transmission spectra through a bulk GaN modulator for different values of the applied reverse bias voltage from 0 to 14 V in steps of 2 V. The inset shows a zoom-in of these traces near the excitonic resonance. (b) Differential absorption spectra derived from the traces of (a) [271]. Reprinted with permission from [271]. Copyright (2011), AIP Publishing LLC.

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can be as large as several MV cm−1. In the device geometry of figure 31, the intrinsic electric fields in the GaN well layers point towards the substrate, and are therefore decreased by the application of a reverse bias on the Schottky contact. As a result, the spatial overlap between the electron and hole wave-functions in each QW increases with increasing reverse volt-age, leading to an increase in the excitonic oscillator strength and absorption; at the same time, a blue shift of the absorption edge is also expected. This is the exact opposite of what hap-pens in modulators based on typical zincblende QWs, where no intrinsic electric fields exist so that under zero bias the wells and barriers have a rectangular profile.

In figure 32(a) we plot the normalized transmission spectra of a device based on the structure of figure 31 for different values of the applied reverse bias. The oscillations observed in the high-transmission portion of these traces are due to thin-film interference effects. The exciton absorption peak is centered at a photon energy of about 3.51 eV in this sample. The resulting transmission dip is barely noticeable under zero or low applied bias, due to the weak overlap of the electron and hole wavefunctions caused by the intrinsic electric fields, and then becomes more and more pronounced as the voltage is increased, as expected. Near the excitonic resonance, the device modulation depth M is therefore negative. Its absolute value increases with increasing applied voltage up to a bias of about 15 V, after which the device performance becomes once again limited by electrical leakage, as confirmed by the electrical characteristics shown in the inset of figure 32(a). When the applied voltage is 12 V, the magnitude of M has a maximum value of about 34%, which is slightly larger than the value obtained with the same bias in the bulk devices of figure 28.

In figure 32(b) we plot the differential absorption spectra, Δα, versus photon energy, obtained from the traces of fig-ure 32(a) using equation (9.1). For the thickness d of the absorb-ing layer here we use the sum of the well-layer thicknesses only (0.2 µm), since the electron and hole wavefunctions in

these QWs are strongly localized inside the wells (particularly because of their triangular potential-energy profiles). A maxi-mum absorption change of over 2 × 104 cm−1 is observed in this figure. This value compares favorably with that of typical infrared electroabsorption modulators for similar bias voltages [264], which highlights the promise of the present devices for practical applications. In this respect, optical modulators based on nitride semiconductors therefore benefit from the very large near-band-edge absorption coefficients that are typical of these materials [274].

At the same time, the photon energy of maximum absorp-tion change in figure 32(b) is found to undergo a small blueshift of about 10 meV as the applied voltage is increased from 0 to 15 V, reflecting the occurrence of a comparable shift in the excitonic resonance. These results are similar to those reported in [268], where a theoretical study was then presented indicating that a larger blueshift is in prin-ciple expected. The likely source of this discrepancy is the presence of significant population of background carriers in the QWs, which in turn can be attributed to unintentionally

Figure 32. (a) Normalized transmission spectra through a GaN/AlGaN multiple-QW modulator for different values of the applied reverse bias voltage (0, 2, 4, 6, 8, 10, 12, 14, and 15 V). The inset shows the current-voltage characteristics of the same device. (b) Differential absorption spectra derived from the traces of (a) [271]. Reprinted with permission from [271]. Copyright (2011), AIP Publishing LLC.

Figure 31. Schematic cross-sectional view of a GaN/AlGaN multiple-QW electroabsorption modulator [271]. Reprinted with permission from [271]. Copyright (2011), AIP Publishing LLC.

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incorporated oxygen donor impurities in the AlGaN barriers [197]. At low bias, these carriers occupy the QW bound states leading to substantial screening of the internal electric fields in the well layers. As the reverse bias is increased, the QWs become increasingly depopulated due to thermionic emis-sion and tunneling through the barriers downstream (which become more and more tilted in the presence of the exter-nal field). Correspondingly, the internal fields in the wells layers are no longer screened causing a redshift in the band edge, which partly counteracts the blueshift produced by the applied field [268].

The key conclusion of the present study is that, when similar active-layer thicknesses and device geometries are employed, electroabsorption modulators based on bulk nitride films can provide similar performance levels to those of GaN/AlGaN QW devices. This should be compared to the case of arsenide, phosphide, or antimonide semiconductors, where the electroabsorption effect in QWs is an order-of-magnitude stronger than in the bulk [264]. The key distinctive property of nitride compounds in this respect is provided by their very large exciton binding energies, which are comparable to the thermal energy kBT at room temperature. As a result, strong excitonic effects are observed in these materials even at room temperature, which is desirable for efficient electroabsorption.

It should be noted that the presence of quantum confine-ment in QWs is still in principle advantageous, because it can further strengthen the exciton binding. However, the perfor-mance of the nitride QW modulators presented in this work is also likely limited by a number of other related factors. First, the bias-dependent screening of the internal polariza-tion fields described in the previous section effectively limits the net Stark shift and related modulation of the absorption edge induced by the applied bias. Second, excitonic absorp-tion features in QWs can experience additional broadening due to thickness fluctuations and interface roughness scatter-ing. While this is a general concern, it may be particularly important in nitride QWs, where relatively small layer widths are typically used to avoid excessive separation of the electron and hole wavefunctions due to the internal electric fields. As a general rule, the narrower a QW the more sensitive its bound states are to interface fluctuations. Finally, the internal electric fields also fundamentally limit the exciton binding strength, because either the electron–hole wavefunction overlap is sig-nificantly reduced (near zero applied voltage), or the excitons can be effectively dissociated via thermionic emission and tunneling (under reverse bias), or both (under forward bias).

In summary the results presented in here clearly indicate that high-performance UV optical modulators can be devel-oped using nitride films and QWs. As already mentioned, our measured absorption changes near the exciton resonance compare favorably with those of typical infrared modulators. Correspondingly, if the nitride active layers of these devices were incorporated in planar dielectric waveguides, order-of-magnitude changes in transmission could be achieved with reasonably low applied voltages of order 1 V. The same mat-erials are also suitable to integration in asymmetric Fabry–Perot cavities, which would allow further increasing the modulation depth and reducing the drive voltage [275].

10. Intersubband transitions in AlGaN quantum wells

10.1. Basic device physics

In recent years, significant research efforts have also been devoted to the study of the basic physics and device appli-cations of intersubband (ISB) transitions in MBE-grown GaN/AlGaN QWs. These transitions involve quantized elec-tronic states derived from the same energy band (typically the conduction band), and feature many desirable properties for optoelectronic device applications [276]. Most notably, the ISB transition wavelength can be tuned by design over a broad range independent of the bandgap energy of the underlying materials. ISB transitions also generally feature very large oscillator strengths, and therefore large absorp-tion or gain coefficients and giant optical nonlinearities. Furthermore, the nonradiative ISB relaxation lifetimes are typically ultrafast (sub-picosecond) due to LO-phonon emis-sion processes, which is particularly desirable for all-optical switching applications. As a result of these favorable prop-erties, ISB transitions have enabled the realization of several high-performance devices such as quantum-well infrared pho-todetectors (QWIPs) and quantum cascade (QC) lasers oper-ating at mid- and far-infrared wavelengths [276]. At present, ISB devices are mostly based on GaAs/AlGaAs or GaInAs/AlInAs QWs, whose material and device technologies are par-ticularly mature. At the same time, GaN/AlGaN QWs provide a promising new materials platform that may allow extend-ing the spectral reach and functionality of these devices. Their potential advantages are mainly related to two basic material properties, namely their large conduction-band offsets ΔEc and large LO-phonon energies hνLO.

The former parameter determines the energy range over which a quantum structure can support confined states derived from the conduction band, and therefore it provides a funda-mental limit to the short-wavelength end of the accessible spec-tral region for ISB transitions. In As-based heterostructures ΔEc is typically less than 0.5 eV, and as a result ISB devices based on these materials are only suitable in practice to opera-tion at wavelengths longer than about 3 µm. In GaN/AlGaN QWs ΔEc can be as large as 1.75 eV, and therefore these het-erostructures can accommodate ISB transitions at record short wavelengths throughout the mid-infrared and well into the near-infrared spectral region, including the low-loss transmis-sion window of optical fibers. As a result, GaN/AlGaN QWs could be used to bridge the existing gap at short-wave infra-red wavelengths between high-performance interband and ISB devices. Furthermore, they allow taking advantage of the unique properties of ISB transitions for device applications in fiber optics, such as ultrafast all-optical switching.

At the same time, the large LO-phonon energies of nitride semiconductors are particularly attractive for the development of THz light sources [277–280]. Existing QC lasers based on GaAs/AlGaAs QWs have rapidly become the leading laser technology for the THz spectral range [281]. However, their operation fundamentally requires cryogenic cooling, with maximum reported temperatures below 200 K [282]. Furthermore, they can only provide incomplete coverage of

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the THz spectral range. Both of these limitations are intrin-sic, as they are related to the presence in (Al)GaAs of opti-cal phonon modes with THz frequencies (e.g. νLO ≈ 8.5 THz in GaAs). The excitation of these vibrational modes provides both a prohibitively large absorption mechanism for light at nearby frequencies (reststrahlen absorption) and an ultra-fast nonradiative decay mechanism between laser subbands separated by THz-range energies. As a result, GaAs-based devices are completely precluded from operation at frequen-cies within and immediately around the ~8–9 THz GaAs reststrahlen band. GaAs/AlGaAs QC lasers emitting at lower frequencies can be realized, but are limited to operation at cry-ogenic temper atures due to the process of thermally activated LO-phonon emission.

This nonradiative decay mechanism is illustrated sche-matically in figure 33, where the traces labeled |u⟩ and |l⟩ denote the upper and lower laser subbands of a generic THz QC structure emitting at a frequency ν below the LO-phonon frequency νLO. At cryogenic temperatures, a large population inversion can be established in this structure, since the major-ity of electrons in |u⟩ occupy states near the bottom of the sub-band, from which they cannot decay nonradiatively into |l⟩ via LO-phonon emission because of the requirement of energy conservation. As the temperature is increased, however, more and more electrons in |u⟩ gain enough thermal energy that scattering into |l⟩ via LO-phonon emission becomes allowed. Since these scattering processes are ultrafast, the end result is a dramatic reduction in the device population inversion and optical gain. It follows from this argument that the smaller the LO-phonon energy of the QW material, the more effectively the population inversion is degraded with increasing temper-ature, and vice versa.

As a result, significantly improved THz spectral cover-age and high-temperature performance can in principle be obtained with QC lasers based on GaN, where hνLO ≈ 91 meV versus 36 meV in GaAs. The latter expectation was quantified in a rigorous Monte Carlo study [279, 280], where two otherwise identical QC structures consisting of GaN/AlGaN and GaAs/AlGaAs QWs were designed, and their population inversion at the optimal bias point was computed as a function of temperature. From the simulation results,

the maximum operating temperatures of the GaAs and GaN devices were estimated to be approximately 200 K (close to the exper imental state-of-the-art [282]) and well over room temperature, respectively. These results suggest that nitride ISB devices have the potential for a revolutionary impact in the field of THz photonics, which has so far been hampered by the lack of practical sources that can operate without cryo-genic cooling. On the other hand, the growth and experimental demonstration of GaN/AlGaN THz QC structures remain a considerable challenge, particularly with regard to obtaining efficient tunneling-based electrical injection into the upper laser subbands.

10.2. Infrared device applications

Initial efforts in the study of ISB transitions in nitride semi-conductors have focused on measuring the infrared transmis-sion spectra of various GaN/Al(Ga)N n-doped multiple-QW samples [283–288]. Pronounced ISB absorption peaks were reported in these studies, centered at near-infrared wavelengths that could not be accessed using As-based materials. Room-temperature optically pumped ISB light emission has also been demonstrated with similar samples, as reported in [289] and [290]. The latter work is summarized in figures 34(a) and (b), where we show a cross-sectional transmission electron microscopy (TEM) image of the multiple-QW sample and a measured emission spectrum, respectively. The GaN/AlN QWs in this structure were designed to support at least three bound subbands, as shown in the calculated conduction-band diagram plotted in the inset of figure 34(b). To demonstrate light emission, a pulsed tunable optical parametric oscilla-tor (OPO) was used to resonantly pump electrons from the ground-state subband |1⟩ to the second-excited subband |3⟩ of each QW, followed by radiative relaxation into the first-excited subband |2⟩. It should be noted that the |1⟩ → |3⟩ ISB trans-itions are only allowed because of the strong built-in electric fields of these QWs, which partially break the symmetry of the bound-state envelope functions. As shown in figure 34(b), the measured luminescence spectra are peaked near 2.05 µm, which is the shortest ISB emission wavelength reported with any QW materials system. These results therefore confirm the promise of nitride semiconductors to extend the spectral reach of ISB devices (including ultimately light sources) to record short wavelengths.

A particularly important device application of near-infrared ISB transitions is all-optical switching at fiber-optic commu-nication wavelengths. This functionality is widely regarded as a key ingredient of future ultra-broadband communication networks, which would enable the processing of information directly in the optical domain. While various technologies have been investigated for this purpose, ISB transitions are ideal candidates due to their intrinsically ultrafast relaxation life-times and large optical nonlinearities. In its simplest form, an ISB all-optical switch can be implemented using QWs embed-ded within an optical waveguide and designed to provide ISB absorption at the data-transmission wavelength (1.55 µm in most broadband fiber-optic communication networks). Signal pulses carrying information at this wavelength are therefore

Momentum

hνLO

PHOTONEMISSION

LO PHONONEMISSION

THERMALEXCITATION

Ene

rgy

hν

|u⟩|l⟩

Figure 33. Schematic illustration of the competition between photon emission and thermally activated LO-phonon emission by electrons in the upper laser subband of a THz QC laser.

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strongly attenuated as they propagate through the waveguide, unless a strong control pulse at comparable wavelength is also simultaneously injected, in which case the ISB absorption is temporarily saturated and the signal transmittance increases. These devices have been developed based on GaN/AlN QWs and used to demonstrate all-optical switching with ultrafast response times [291, 292]. As an illustration, in figure 35 we show the measured signal transmittance through such a device as a function of signal-control delay time, for different val-ues of the input control-pulse energy [293]. The waveguide used in this work is illustrated schematically in the inset of the figure: its active layer consists of 30 GaN/AlN QWs with ISB absorption spectrum peaked near 1.55 µm, sandwiched between an AlN template layer and a GaN cap layer which

is then etched to form a ridge. As clearly shown in the figure, the control pulses can produce a large instantaneous increase in signal transmittance, which only lasts for a time interval on the order of a few hundred femtoseconds. Therefore, these devices can perform all-optical switching at ultrafast bit rates of several hundred Gbit/sec.

Other device functionalities that have been reported based on near- and mid-infrared ISB transitions in GaN/Al(Ga)N QWs include photodetection [294–298], electro-optic modu-lation [299, 300], and electroluminescence [301, 302]. III-nitride ISB light detection has been demonstrated with QWIPs [294–296] as well as QC photodetectors [297, 298]. The lat-ter devices are photovoltaic detectors where light is absorbed via ISB transitions between two QW bound states, followed by sequential tunneling transport from one active-layer repeat unit to the next through a suitably designed injector. Their demonstration with nitride QWs is therefore especially impor-tant as the first example of a device functionality based on QC structures with this materials system. High-speed opera-tion at near-infrared wavelengths near 1.55 µm has also been reported with these devices [298], as well as with GaN/AlN QWIPs [296], again consistent with the ultrafast nature of ISB transitions. Another more recent milestone is the demon-stration of mid-infrared electroluminescence in GaN/AlGaN QC structures, at wavelengths near 5 µm in excellent agree-ment with simulations [302]. This work also highlighted the dominant role played by interface roughness scattering in III-nitride QWs, and the importance to include the resulting effective interface grading in the QC design.

10.3. Terahertz device applications

THz technologies have important applications in security screening, biomedical sensing, industrial quality control, and spectroscopic imaging. However, their widespread emer-gence has so far been limited by the lack of suitable devices, most notably efficient solid-state sources. As discussed in

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Figure 34. (a) Cross-sectional TEM image of a GaN/AlN multiple-QW sample used to demonstrate optically pumped ISB light emission [290]. (b) Measured emission spectrum. The sample conduction-band diagram is shown in the inset. The solid line is a modified-Gaussian fit of the experimental data. Reprinted with permission from [290]. Copyright (2009), AIP Publishing LLC.

Figure 35. Signal transmittance through a GaN/AlN QW waveguide as a function of signal-control delay time, for different values of the input control-pulse energy. Inset: schematic cross-sectional view of the waveguide, and calculated field-intensity profile of the fundamental guided mode [293]. Reproduced with permission from [293]. © (2010) COPYRIGHT SPIE–The International Society for Optical Engineering.

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section 10.1, III-nitride ISB devices can operate across the entire THz spectrum, including frequencies that are funda-mentally inaccessible to their arsenide counterparts due to reststrahlen absorption. Furthermore, GaN THz QC lasers could potentially reach room-temperature operation, unlike existing GaAs devices which are intrinsically limited to cryo-genic temper atures. Motivated by these considerations, sub-stantial research efforts have been devoted in recent years to the study of THz ISB transitions in GaN/AlGaN QWs via absorption spectr oscopy measurements [303–306]. The initial demonstrations of THz ISB photodetection [307, 308] and electroluminescence [309] have also been reported.

From a design perspective, an important challenge in this context is provided by the strong spontaneous and piezo-electric fields that exist in QWs grown along the c direction. As described in section 4, these internal fields produce a trap-ezoidal conduction- and valence-band lineup which tends to blue-shift the ISB transition energies, and therefore can signif-icantly complicate the development of long-wavelength ISB devices. In addition, in wide QWs designed for THz-range ISB transitions, the first-excited states necessarily fall within the triangular portion of the band lineup, well below the top of the barriers. As a result, carrier extraction from these subbands into the overlaying continuum of unbound states (as needed to produce the photocurrent in QWIPs) becomes extremely diffi-cult. For the same reason, interwell tunneling transport among the lower-energy states of these QWs (which forms the basis of QC laser operation) is equally problematic.

To address this design challenge, initial work has focused on the development of step-QW structures [303, 305], where two or more layers of different (Al)GaN compositions are used in each well to create a more rectangular potential energy profile. In [307], a QWIP operating near 13 THz was dem-onstrated based on a variation of this approach, where two different AlGaN layers are also used in each barrier to allow for a bound-to-quasi-bound QWIP design (i.e. with the first-excited subbands nearly resonant to the top of the barriers). A photocurrent spectrum measured with this device near 20 K

is shown in figure 36(a), together with the conduction-band diagram of the active QWs under bias. The peak responsiv-ity was estimated to be 7 mA W−1, which is reasonable for THz QWIPs [310, 311]. However, a potential drawback of these step-QW structures is the presence of additional hetero-interfaces in each QW period, which makes the bound states especially vulnerable to interface roughness scattering. As a result, extremely fast ISB decay lifetimes and proportionally large spectral broadening are produced.

An alternative approach to address the same design chal-lenge is the use of QWs grown along semi-polar (or nonpo-lar) crystallographic directions, where the undesirable internal electric fields are significantly reduced (or completely elimi-nated). This approach is now being widely investigated, as allowed by the development and commercialization of high-quality semi-polar and nonpolar free-standing GaN substrates. Nonpolar m-plane or a-plane III-nitride QWs have already been used for the measurement of THz ISB absorption [304, 306], as well as the investigation of mid-infrared ISB trans-itions [312–315]. Figure 36(b) shows the low-temperature responsivity spectrum (peaked at about 10 THz) and conduc-tion-band diagram of a GaN/Al0.06Ga0.94N QWIP grown on a semi-polar (2 0 2 1) GaN substrate [308]. In these QWs the internal electric fields are an order of magnitude smaller than in identical heterostructures grown along the c direction, so that a fairly rectangular QW potential-energy profile is obtained under zero-bias conditions. Compared to the step-QW device of figure 36(a), a much narrower linewidth is obtained with this semi-polar QWIP (16.1 meV versus 38.3 meV full width at half maximum), indicative of reduced scattering from inter-face roughness and other defects and/or improved uniformity of the active QWs. It should also be noted that the spectra of both QWIPs overlap the reststrahlen band of GaAs (~33 to 37 meV), clearly illustrating the ability of GaN devices to cover this frequency range that has so far been inaccessible to ISB optoelectronics.

The other key ingredient for the further development of III-nitride THz ISB devices (especially QC lasers) is the ability to

Figure 36. (a) Photocurrent spectrum of a double-step III-nitride QWIP measured at 20 K under an applied voltage bias of +0.8 V (solid line), and Gaussian fit (dashed line) [307]. The sharp cutoff near 70 meV is due to the reststrahlen band of GaN. (b) Photocurrent spectrum of a semi-polar III-nitride QWIP measured at 10 K under an applied voltage of 1.2 V [308]. The conduction-band diagrams of both devices are shown in the insets. (a) Reprinted with permission from [307]. Copyright (2012), AIP Publishing LLC. (b) Reprinted with permission from [308]. Copyright (2016), AIP Publishing LLC.

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obtain efficient tunneling transport. In the past several years, this transport mode has also been extensively studied, par-ticularly in the context of GaN/Al(Ga)N double-barrier struc-tures [316–320]. In the initial reports, the measured electrical characteristics were found to suffer from strong hysteresis and to rapidly degrade upon repeated voltage scans—a behav-ior that was generally attributed to filling of traps related to threading dislocations. More recently, however, stable and repeatable negative differential resistance due to resonant tunneling has been demonstrated through concerted efforts aimed at minimizing the formation of strain-induced defects [318–320]. Sequential tunneling transport in simple GaN/AlGaN QC structures has also been investigated through the measurement of highly nonlinear electrical characteristics, with the experimental turn-on voltages well reproduced by simulation results [321]. In the same report, a weak electrolu-minescence peak near 10 THz was presented, consistent with a picture of photon-assisted tunneling (i.e. interwell ISB light emission) in the active QWs. Significantly stronger emission at 1.4 THz with a linewidth as small as 0.134 THz was measured in a different QC structure in [309] and explicitly attributed to ISB transitions (albeit not the intended radiative transitions in the active layer design). We believe that the overall progress reviewed in this section, while still rather early-stage, is quite encouraging for the prospects of the future demonstration of III-nitride THz QC lasers.

11. Recent developments

11.1. GaN substrates

As discussed in the previous sections the majority of the opto-electronic devices based on III-nitride semiconductors were grown heteroepitaxially on foreign substrates such as (0 0 0 1) sapphire and SiC. In the past several years significant prog-ress has been made to form bulk GaN substrates. An excel-lent review on the topic was presented by Amano [322]. One method includes the growth of thick GaN films by the HVPE method on foreign substrates (sapphire, GaAs, LiGaO3, SiC and Si) and then separating them from the substrate to produce a freestanding wafer. A number of techniques have been devel-oped for the separation of GaN from the substrate. The laser liftoff technique, has been used to separate GaN from sapphire [323]. The void-assisted separation method facilitated by the deposition on the substrate a TiN interlayer is an alternative approach [324]. Our group developed the method of self-sep-aration during cooling from the growth temperature due to the different thermal expansion coefficient of GaN and the sap-phire substrate [325]. Such quasi-GaN substrates became com-mercially available by a number of vendors in the past several years with dislocation density of about 5 × 106 cm−2.

An alternative approach to forming free standing GaN substrates is the high pressure solution growth method [326, 327]. In this method single crystals of GaN are produced by spontaneous nucleation with very low defect density (dis-location density ~102 cm−2). The produced wafers by this method were generally small (~1 cm2). By increasing the pressure further Inoue et al [328] reported the growth of free

standing GaN substrates with a size about 2.5 cm. The major problem with this method is that it requires very high pres-sures (~20 000 atm.)

In recent years significant progress has also been made in the development of free standing GaN substrates by the ammonothermal method. This method, which was first pro-posed by Dwilinski et al [329, 330], is based on solvother-mal synthesis like the method used to synthesize quartz. Specifically supercritical NH3 is used as a solvent for Ga and the solubility of GaN is increased by using a mineralizer to deposit GaN. The reaction takes place at about 400–500 °C and 2000–3000 atm. GaN substrates up to 2-in in diameter produced by this method are commercially available with dis-location density <105 cm−2. As important the curvature radius of the substrate is 1000 m or larger, while that of HVPE grown GaN substrates is 2–30 m [322]. Non-polar m-plane and semi-polar (2 0 −2 1) GaN substrates, made by this method, are also commercially available.

11.2. LEDs on GaN substrates

As discussed in section 6 all commercially available LEDs until now are grown heteroepitaxially on foreign substrate such as sapphire and silicon carbide. The high concentration of defects due to lattice mismatch between substrate and epi-taxial film is in general undesirable since it limits the inter-nal quantum efficiency (IQE) and reliability of the device. Another key limitation to the efficiency of GaN-based LEDs is the ‘efficiency droop’. For the best commercial LEDs, the IQE peaks at a current density of a few A cm−2, but is reduced by about 25% at high current density of 100 A cm−2.

The recent availability of free standing GaN substrates presented the opportunity of developing the second genera-tion LEDs grown homoepitaxially on such substrates. In a series of papers researchers at Soraa Corp [331–335]. have reported GaN based LED on GaN substrates with a record power conversion efficiency of 84% at 85 °C. Furthermore, the efficiency remains high at high current density, owing to low current-induced droop and low series resistance. These types of devices have now become commercially available.

11.3. LEDs on non-polar and semipolar GaN substrates

The commercially available InGaN/GaN based LEDs are cur-rently grown along the polar c-axis direction of the wurtzite crystal, and thus, their IQE is reduced due to internal fields asso-ciated with the spontaneous and piezoelectric polarization [127]. This is one of the sources of the observed reduction in IQE of the green LEDs, since they require higher InN mole fraction in the wells. An alternative approach is to develop such devices on nonpolar or semipolar GaN substrates, which recently became commercially available. An excellent review addressing both material and device issues of non-polar and semipolar Group-III nitride has been published by Speck and Chichibu [336].

InGaN-based light-emitting diodes (LEDs) grown on semi-polar orientations of GaN are an attractive alternative to con-ventional c-plane devices for long-wavelength emission due to reduced internal electric fields. The polarized emission

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from such devices may also enable high-efficiency color dis-plays and optically tuned white light [337]. However, semi-polar LEDs have yet to displace conventional c-plane LEDs [338]. This is largely due to the limited size and high cost of bulk semipolar substrates, as well as materials challenges such as the formation of misfit dislocations and basal-plane stacking faults [339]. On the other hand, as discussed in sec-tion 7, high power green lasers have produced on semipolar GaN substrates [221, 222, 226].

11.4. LEDs on silicon substrates

Significant research effort over the past 20 years has also been dedicated in the development of GaN-based LEDs on silicon substrates [340–350]. The fundamental issue of developing GaN based LEDs on Si is cracking due to the large thermal mismatch between GaN and Si. Various approaches have been used to reduce the thermal stress. This include the insertion of low-temperature AlN interlayers, introducing multiple AlGaN/GaN interlayers, and growing on pre-patterned substrates. Using such approaches growth of crack-free GaN based LEDs on pre-patterned Si (1 1 1) has been demonstrated with reasonable performance [345, 346]. In more recent times UV lasers based on III-nitride nanowires on Si (1 1 1) have been demonstrated [351, 352].

Acknowledgments

The authors would like to thank our collaborators and gradu-ate students, whose work was referenced in this review article. This work was supported in part by the NSF Division of Elec-trical, Communications, and Cyber Systems standard Grant No. 1408364, overseen by Dr Mahmoud Fallahi. Certain images in this publication have been obtained by the author(s) from the Wikipedia/Wikimedia website, where they were made available under a Creative Commons licence or stated to be in the public domain. Please see individual figure cap-tions in this publication for details. To the extent that the law allows, IOP Publishing disclaim any liability that any person may suffer as a result of accessing, using or forwarding the image(s). Any reuse rights should be checked and permis-sion should be sought if necessary from Wikipedia/Wikimedia and/or the copyright owner (as appropriate) before using or forwarding the image(s).

References

[1] Edgar J H (ed) 1995 Properties of Group III Nitrides (EMIS Data Review vol 11) (London: INSPEC)

[2] Pearton S J (ed) 1997 GaN and Related Materials (New York: Gordon and Breach)

[3] Nakamura S and Fasol G 1997 The Blue Laser Diode (Berlin: Springer)

[4] Pankove J I and Moustakas T D (ed) 1998 Gallium nitride (GaN) I Semiconductors and Semimetals vol 50 (San Diego: Academic)

[5] Pankove J I and Moustakas T D (ed) 1999 Gallium nitride (GaN) II Semiconductors and Semimetals vol 57 (San Diego: Academic)

[6] Gil B (ed) 1998 Group III Nitride Semiconductor Compounds: Physics and Applications (Oxford: Clarendon)

[7] Nakamura S, Pearton S and Fasol G 2000 The Blue Laser Diode 2nd edn (Berlin: Springer)

[8] Doppalapudi D and Moustakas T D 2002 Handbook of Thin Film Materials (Semiconductor and Superconductor Thin Films vol 4) ed H S Nalwa (New York: Academic) ch 2

[9] Razeghi M and Henini M (ed) 2004 Optoelectronic Devices: III-Nitrides (Amsterdam: Elsevier)

[10] Morkoc H 2008 Handbook of Nitride Semiconductors and Devices, Vol. 1,: Materials Properties, Physics and Growth (New York: Wiley)

[11] Morkoc H 2009 Handbook of Nitride Semiconductors and Devices, Vol. 2 Electronic and Optical Processes in Nitrides (New York: Wiley)

[12] Morkoc H 2009 Handbook of Nitride Semiconductors and Devices, Vol. 3, GaN-Based Optical and Electronic Devices (New York: Wiley)

[13] Morkoc H 2009 Nitride Semiconductors and Devices 2nd edn (Berlin: Springer)

[14] Veal T D, McConville C F and Schaff W J (ed) 2009 Indium Nitride and Related Alloys (Boca Raton, FL: CRC Press)

[15] Piprek J (ed) 2007 Nitride Semiconductor Devices: Principle and Simulations (New York: Wiley)

[16] Kneissl M and Rass J (ed) 2016 III-Nitride Ultraviolet Emitters: Technology and Applications (Springer Series in Materials Science vol 227) (Berlin: Springer)

[17] Juza R and Hahn H 1938 Über die Kristallstrukturen von Cu3N, GaN und InN Metallamide und Metallnitride Anorg. Allgem. Chem. 234 282

[18] Grimmeiss H and H-Koelmans Z 1959 Z. Naturf. 14a 264 [19] Maruska H P and Tietjen J J 1969 The preparation and

properties of vapor-deposited Single Crystalline GaN Appl. Phys. Lett. 15 367

[20] Seifert W, Franzheld R, Butter E, Sabotta H and Reide V 1983 On the origin of free carriers in high-conducting n-GaN Cryst. Res. Technol. 18 383–90

[21] Pankove J I, Miller E A and Berkeyheiser J E 1972 Electroluminescence in GaN J. Lumin. 5 84

[22] Maruska H P, Stevenson D A and Pankove J I 1973 Violet luminescence in Mg-doped GaN Appl. Phys. Lett. 22 303

[23] Moustakas T D, Molnar R J, Lei T, Menon G and Eddy C R 1992 A comparative study of GaN films grown on different faces of sapphire by ECR-assisted MBE Mater. Res. Soc. Symp. Proc. 242 427–32

[24] Moustakas T D, Lei T and Molnar R J 1993 Growth of GaN by ECR-assisted MBE Physica B 185 36–49

[25] Amano H, Sawaki N, Akasaki I and Toyoda Y 1986 Metaorganic vapor phase epitaxial growth of a high quality GaN film using an AIN buffer layer Appl. Phys. Lett. 48 353

[26] Menon G 1990 Growth of intrinsic monocrystalline gallium nitride thin films by electron cyclotron resonance microwave plasma-assisted molecular beam epitaxy MS Thesis Boston University

[27] Menon G, Molnar R, Berkowitz S J, Moustakas T D and Scanlon J 1991 Growth of GaN Films on the a-plane of sapphire by ECR-assisted MBE Bull. Am. Phys. Soc. 36 544

[28] Lei T, Fanciulli M, Molnar R J, Moustakas T D, Graham R J and Scanlon J 1991 Epitaxial growth of zincblende and wurtzitic gallium nitride thin films on (0 0 1) silicon Appl. Phys. Lett. 59 944

[29] Nakamura S 1991 GaN growth using GaN buffer layer Japan. J. Appl. Phys. 30 L1705

[30] Amano H, Akasaki I, Kozawa T, Hiramatsu K, Sawaki N, Ikeda K and Ishii Y 1988 Electron beam effects on blue luminescence of zinc-doped GaN J. Lumin. 40–1 121

[31] Amano H, Kito M, Hiramatsu K and Akasaki I 1989 P-type conduction in Mg-doped GaN treated with low-energy electron-beam irradiation (LEEBI) Japan. J. Appl. Phys. 28 L2112

Rep. Prog. Phys. 80 (2017) 106501

Review

33

[32] Van Vechten J A, Zook J D, Hornig R D and Goldenberg B 1992 Defeating compensation in wide gap semiconductors by growing in H that is removed by low-temperature de-ionizing radiation Japan. J. Appl. Phys. 31 3662

[33] Pankove J I, Zanzucchi P J, Magee C W and Lucovsky G 1985 Hydrogen localization near boron in silicon Appl. Phys. Lett. 46 421

[34] Moustakas T D 1988 The role of hydrogen in trace impurity doping of amorphous silicon 1st Int. Symp. on Physics and Applications of Amorphous Semiconductors ed F Demichelis (Singapore: World Scientific) pp 399–411

[35] Akasaki I, Amano H, Kitoh M and Hiramatsu K 1991 Photoluminescence of Mg-doped p-type GaN and electroluminescence of GaN p-n junction LED J. Lumin. 48–9 666

[36] Nakamura S, Mukai T, Senoh M and Iwasa N 1992 Hole compensation mechanism of p-type GaN films Japan. J. Appl. Phys. 31 1258

[37] Moustakas T D and Molnar R J 1993 Growth and doping of GaN films by ECR-assisted MBE Mater. Res. Soc. Symp. Proc. 281 753

[38] Brandt M S, Johnson N M, Molnar R J, Singh R and Moustakas T D 1994 Hydrogenation of p-type gallium nitride Appl. Phys. Lett. 64 2264

[39] Nakamura S, Mukai T and Senoh M 1994 Candela-class high-brightness InGaN/AlGaN double-heterostructure blue-light-emitting diodes Appl. Phys. Lett. 64 1687

[40] Nakamura S, Senoh M, Nagahama S, Iwasa N, Yamada T, Matsushita T, Kiyyoku H and Sugimoto Y 1996 InGaN-based multi-quantum-well-structure laser diodes Japan. J. Appl. Phys. 35 L74

[41] Trampert A, Brandt O and Ploog K H 1998 Crystal structure of group-III nitrides Gallium Nitride (GaN) I: Semiconductors and Semimetals vol 50 ed J I Pankove and T D Moustakas (San Diego: Academic) ch 7

[42] File:Wurtzite polyhedra.png—Wikipedia, the free encyclopedia [43] Xia Q, Xia H and Ruoff A L 1993 Pressure-induced rocksalt

phase of aluminum nitride: a metastable structure at ambient condition J. Appl. Phys. 73 8198

[44] Perlin P, Jauberthie-Carillon C, Itie J P, San Miguel A, Grzegory I and Polian A 1992 Raman scattering and x-ray absorption spectroscopy in gallium nitride under high pressure Phys. Rev. B 45 83

[45] Ueno M, Yoshida M, Onedera A, Shimommura O and Takemura K 1994 Stability of the wurtzite-type structure under high pressure:GaN and InN Phys. Rev. B 49 14

[46] Lei T, Moustakas T D, Graham R J, He Y and Berkowitz S J 1992 Epitaxial growth and characterization of zinc-blende gallium nitride on (0 0 1) silicon J. Appl. Phys. 71 4933

[47] Brandt O 1998 Cubic group III nitrides Group III Nitride Semiconductor Compounds ed B Gil (Oxford: Clarendon) ch 11

[48] Yeh C-Y, Lu Z W, Froyen S and Zunger A 1992 Zinc-blende-wurtzite polytypism in semiconductors Phys. Rev. B 46 10086

[49] Ng H M and Cho A Y 2002 Investigation of Si doping and impurity incorporation dependence on the polarity of GaN by molecular beam epitaxy J. Vac. Sci. Technol. B 20 1217

[50] Hellman E S 1998 The polarity of GaN: a critical review MRS Internet J. Nitride Semicond. Res. 3 1–11

[51] Nye J F 1985 Physical Properties of Crystals (Oxford: Oxford University Press)

[52] Lei T, Ludwig K F and Moustakas T D 1993 Heteroepitaxy, polymorphism and faulting in GaN thin films on silicon and sapphire substrates J. Appl. Phys. 74 4430

[53] Sun C J, Kung P, Saxler A, Ohsato H and Razeghi M 1994 A crystallographic model of (00.1) aluminum nitride epitaxial thin film growth on (00.1) sapphire substrate J. Appl. Phys. 75 3964

[54] Kung P, Sun C J, Suxler A, Ohsato H and Razeghi M 1994 Crystallography of epitaxial growth of wurtzite-type thin films on sapphire substrates J. Appl. Phys. 75 4515

[55] Grandjean N, Massies J, Vennegues P, Laugt M and Leroux M 1997 Epitaxial relationship between GaN and Al2O3 (0001) substrates Appl. Phys. Lett. 70 643

[56] Basu S N, Lei T and Moustakas T D 1994 Microstructures of GaN films deposited on (001) and (111) Si using ECR-MBE J. Mater. Res. 9 2370

[57] Araki T, Saito Y, Yamaguchi T, Kurouchi K, Nanishi Y and Naoi Y 2004 Radio frequency molecular beam epitaxial growth of InN epitaxial films on (0001) sapphire and their properties J. Vac. Sci. Technol. B22 2139

[58] Bhuiyan A G, Hashimoto A and Yamamoto A 2003 Indium nitride (InN): a review on growth, characterization and properties J. Appl. Phys. 94 2779–808

[59] Venables J 2000 Introduction to Surface and Thin Film Processes (Cambridge: Cambridge University Press)

[60] Landau L D and Lifshitz I M 1980 Statistical Physics 3rd edn (Oxford: Pergamon)

[61] Moustakas T D 2013 The role of extended defects on the performance of optoelectronic devices in nitride semiconductors Phys. Status Solidi A 210 169–74

[62] Matthews J W and Blakeslee A E 1974 Defects in epitaxial multilayers: I. Misfit dislocations J. Cryst. Growth 32 265

[63] Akasaki I and Amano H 1998 ‘Lasers’ in gallium nitride (GaN) I Semiconductors and Semimetals vol 50, ed J I Pankove and T D Moustakas (San Diego: Academic) ch 15

[64] Okumura H, Kimoto T and Suda J 2012 Over 700 nm critical thickness of AlN grown on 6H-SiC (0001) by molecular beam epitaxy Appl. Phys. Express 5 105502

[65] Keller S, Keller B P, Wu Y F, Heying B, Kapolnek D, Speck J S, Mishra U K and Denbars S P 1996 Influence of sapphire nitridation on properties of gallium nitride grown by metallorganic chemical vapor deposition Appl. Phys. Lett. 68 1525

[66] Grandjeam N, Massies J and Leroux M 1996 Nitridation of sapphire. Effect on the optical properties of GaN epitaxial overlayers Appl. Phys. Lett. 69 2071

[67] Vennegues P and Beaumont B 1999 Transmission electron microscopy study of the nitridation of the (0001) sapphire surface Appl. Phys. Lett. 75 4115

[68] Georgakilas A, Mikroulis S, Cimalla V, Zervos M, Kostopoulos A, Androulidaki M, Komninou Ph, Kehagias Th and Karakostas Th 2001 Effects of the sapphire nitridation on the polarity and structural properties of GaN layers grown by plasma-assisted MBE Phys. Status. Solidi a 188 567

[69] Mikroulis S, Georgakilas A, Kostopoulos A, Cimalla V, Dimakis E and Komninou Ph 2002 Control of the polarity of molecular beam epitaxy grown GaN thin films by the surface nitridation of Al2O3 (0001) substrates Appl. Phys. Lett. 80 2886

[70] Wang Y, Ozcan A S, Ozaydin G, Ludwig K F Jr, Bhattacharyya A, Moustakas T D, Zhou H, Headrick R L and Siddons D P 2006 Real time synchrotron x-ray studies of low and high temperature nitridation of c-plane sapphire Phys. Rev. B 74 235304

[71] Doppalapudi D, Iliopoulos E, Basu S N and Moustakas T D 1999 Epitaxial growth of gallium nitride thin films on a-plane sapphire by molecular beam epitaxy J. Appl. Phys. 85 3582–9

[72] DenBaars S P and Keller S 1998 Metallorganic chemical vapor deposition (MOCVD) of group III nitrides Gallium Nitride (GaN) I ed J I Pankove and T D Moustakas (Semiconductors and Semimetals vol 50) (San Diego: Academic) ch 2

[73] Romano L T, Krusor B S, Singh R J and Moustakas T D 1997 Structure of GaN films grown by molecular beam epitaxy on (0001) sapphire J. Electron. Mater. 26 285

Rep. Prog. Phys. 80 (2017) 106501

Review

34

[74] Munkholm A, Thompson C, Foster C M, Eastman J A, Auciello O, Stephenson G B, Fini P, DenBaars S P and Speck J S 1998 Determination of the cubic to hexagonal fraction in GaN nucleation layers using grazing incidence x-ray scattering Appl. Phys. Lett. 72 2972

[75] Wu H, Kapolnek D, Tarsa E J, Heyung B, Keller S, Keller B P, Mishra V K, DenBaars S P and Speck J S 1996 Nucleation layer evolution in metal-organic chemical vapor deposition grown GaN Appl. Phys. Lett. 68 1371

[76] Keller S, Kapolnek D, Keller B P, Wu Y, Heying B, Speck J S, Mishra V K and DenBaars S P 1996 Effect of the trimethylgallium flow during nucleation layer growth on the properties of GaN grown on sapphire Japan. J. Appl. Phys. 35 L285

[77] Huang D, Visconti P, Jones K M, Reshcikov M A, Yun F, Baski A A, King T and Morkoc H 2001 Dependence of GaN polarity on the parameters of the buffer layer grown by molecular beam epitaxy Appl. Phys. Lett. 78 4145

[78] Okada N, Murata T, Tadatomo K, Chang H C and Watanabe K 2009 Growth of GaN layer and characterization of light emitting diodes using random-cone patterned sapphire substrate Japan. J. Appl. Phys. 48 122103

[79] Li Y, You S, Zhu M, Zhao L, Hou W, Detchprohm T, Taniguchi Y, Tamura N, Tanaka S and Wetzel C 2011 Defect-reduced green GaInN/GaN light emitting diodes on nanopatterned sapphire Appl. Phys. Lett. 98 151102

[80] Song B, Sun H, Kao C-K and Moustakas T D 2016 Growth mode of nitride semiconductors on nano-patterned sapphire substrates by molecular beam epitaxy Phys. Status Solidi c 13 195

[81] Bykhovski A, Gelmot B and Shur M 1993 Strain and charge distribution in GaN-AlN-GaN semiconductor-insulator-semiconductor structure for arbitrary growth orientation Appl. Phys. Lett. 63 2243

[82] Bernardini F, Fiorentini V and Vanderbilt D 1997 Spontaneous polarization and piezoelectric constants of III–V nitrides Phys. Rev. B 56 R10024

[83] Ambacher O et al 2000 Two dimensional electron gases induced by spontaneous and piezoelectric polarization in undoped and doped AlGaN/GaN heterostructures J. Appl. Phys. 87 334

[84] Chuang S L 2009 Physics of Photonic Devices (Hoboken: Wiley)

[85] Vurgaftman I and Meyer J R 2003 Band parameters for nitrogen-containing semiconductors J. Appl. Phys. 94 3675

[86] Shockley W 1939 On the surface states associated with periodic potential Phys. Rev. 56 317

[87] Dhesi S S, Stagarescu C B, Smith K E, Doppalapudi D, Singh R and Moustakas T D 1997 Surface and bulk electronic structure of thin film wurtzite GaN Phys. Rev. B 56 10271

[88] Ryan P, Chao Y-C, Downes J, McGuinness C, Smith K E, Sampath A V and Moustakas T D 2000 Surface electronic structure of p-type GaN (0001) Surf. Sci. 467 L827

[89] Foresi J S and Moustakas T D 1993 Metal contacts to gallium nitride Appl. Phys. Lett. 62 2859

[90] Spitzer W G and Mead C A 1963 Barrier height studies on metal-semiconductor systems J. Appl. Phys. 34 3061

[91] Mead C A and Spitzer W G 1963 Photoemission from Au and Cu into CdS Appl. Phys. Lett. 2 74

[92] Mead C A and Spitzer W G 1963 Fermi level position at semiconductor surfaces Phys. Rev. Lett. 10 471

[93] Mead C A 1965 Surface states on semiconductor crystal: barriers on the Cd (Se:S) system Appl. Phys. Lett. 6 103

[94] Kurtin S, McGill T C and Mead C A 1969 Fundamental transitions in the electronic nature of solids Phys. Rev. Lett. 22 1433

[95] Pauling L 1967 The nature of the Chemical Bond 3rd edn (New York: Cornell University Press)

[96] Pankove J I and Schade H E P 1974 Photoemission from GaN Appl. Phys. Lett. 25 53

[97] Boer K W 1990 Survey of Semiconductor Physics (New York: Van Nostrand)

[98] Greco G, Iucolano F and Roccaforte F 2016 Ohmic contacts to gallium nitride materials Appl. Surf. Sci. 383 324

[99] Sampath A, Ng H M, Korakakis D and Moustakas T D 1998 Metal contacts to n-AlxGa1−xN Mater. Res. Soc. Symp. Proc. 482 1095

[100] Song J-O, Kim S-H, Kwak J S and Seong T-Y 2002 Formation of vanadium-based ohmic contacts to n-GaN Appl. Phys. Lett. 83 1154

[101] Schweitz K O, Wang P K, Mohney S E and Gotthold D 2002 V/Al/Pt/Au ohmic contacts to n-AlGaN/GaN heterostructures Appl. Phys. Lett. 80 1954

[102] Wang J H, Mohney S E, Wang S H, Chowdhury U and Dupuis R D 2004 Vanadium-based ohmic contacts to n-type Al0.6Ga0.4N J. Electron. Mater. 33 418

[103] France R, Xu T, Chen P, Chandrasekaran R and Moustakas T D 2007 Vanadium-based ohmic contacts to n-AlGaN in the entire alloy composition Appl. Phys. Lett. 90 062115

[104] Mori K, Takeda K, Kusafuka T, Iwaya M, Takeuchi T, Kamiyama S, Akasaki I and Amano H 2016 Low-ohmic contact resistance V-based electrode for n-type AlGaN with high AlN molar fraction Japan. J. Appl. Phys. 55 05FL03

[105] Schweitz K O and Mohney S E 2001 Phase equilibria in transition metal Al-Ga-N system and thermal stability of contacts to AlGaN J. Electron. Mater. 30 175

[106] Fomenko V S 1966 Handbook of Thermionic Properties: Electronic Work Functions and Richardson Constants of Elements and Compounds (New York: Plenum) p 103

[107] Pookpanratana S, France R, Bar M, Weihardt L, Fuchs O, Blum M, Yang W, Denlinger J D, Moustakas T D and Heske C 2008 Intermixing and chemical structure at the interface between n-GaN and V-based contacts Appl. Phys. Lett. 93 172106

[108] Pookpanratana S et al 2010 Chemical structure of vanadium-based contact formation on n-AlN J. Appl. Phys. 108 24906

[109] Pookpanratana S et al 2012 Microstructure of vanadium-based contacts on n-GaN J. Phys. D: Appl. Phys. 45 105401

[110] Nakamura S, Senoh M and Mukai T 1993 High-power lnGaN/GaN double-heterostructure violet light emitting diodes Appl. Phys. Lett. 62 2390

[111] Nakamura S, Senoh M, Iwasa N and Nagahama S-I 1995 High-brightness InGaN blue, green and yellow light-emitting diodes with quantum well structure Appl. Phys. Lett. 67 1868

[112] Krames M R, Shchekin O B, Mueller-Mach R, Mueller G O, Zhou L, Harbers G and Craford M G 2007 Status and future of high-power light-emitting diodes for solid-state lighting J. Displ. Technol. 3 160–75

[113] Harbers G, Bierhuizen S and Krames M R 2007 Performance of high power light emitting diodes in display illumination applications J. Displ. Technol. 3 98–109

[114] Joyce W B, Bachrach R Z, Dixon R W and Sealer D A 1974 Geometrical properties of random particles and the extraction of photons from electroluminescence diodes J. Appl. Phys. 45 2229

[115] Chichibu S F et al 2006 Origin of defect-insensitivity emission probability in In-containing (Al,In,Ga)N alloy semiconductors Nat. Mater. 5 810

[116] Lester S D, Ponce F A, Craford M G and Steigerwald D A 1995 High dislocation densities in high efficiency GaN-based light-emitting diodes Appl. Phys. Lett. 66 1249

[117] Ponce F A and Bour D P 1997 Nitride-based semiconductors for blue and green light-emitting devices Nature 386 351

Rep. Prog. Phys. 80 (2017) 106501

Review

35

[118] Abell J and Moustakas T D 2008 The role of dislocations as non-radiative recombination centers in InGaN MQWs Appl. Phys. Lett. 92 091901

[119] Chichibu S, Azuhata T, Sota T and Nakamura S 1996 Spontaneous emission of localized excitons in InGaN single and multiquantum well structures Appl. Phys. Lett. 69 4188

[120] Doppalapudi D, Basu S N, Ludwig K F and Moustakas T D 1998 Phase separation and ordering in InGaN alloys grown by molecular beam epitaxy J. Appl. Phys. 84 1389

[121] Doppalapudi D, Basu S N and Moustakas T D 1999 Domain Structure in chemically ordered InxGa1−xN alloys grown by molecular beam epitaxy J. Appl. Phys. 85 883

[122] Nakamura S, Senoh M, Iwasa N and Nagahama S 1995 High-brightness InGaN blue, green and yellow light-emitting diodes with quantum well structures Japan. J. Appl. Phys. 34 L797

[123] Jiang Y et al 2015 Realization of high-luminous efficiency InGaN light-emitting diodes in the ‘green gap’ range Sci. Rep. 5 10883

[124] Auf der Maur M, Pecchia A, Penazzi G, Rodrigues W and Di Carlo A 2016 Efficient drop in green InGaN /GaN light emitting diodes: the role of random alloy fluctuations Phys. Rev. Lett. 116 027401

[125] Lobanova A, Kolesnikova A, Romanov A, Karpov S Y, Rudinsky M and Yakovlev E 2013 Mechanism of stress relaxation in (0001) InGaN/GaN via formation of V-shaped dislocation half-loops Appl. Phys. Lett. 103 152106

[126] Hammersley S, Kappers M J, Massabuau F C-P, Sahonta S-L, Dawson P, Oliver R A and Humphreys C J 2015 Effects of quantum well growth temperature on the recombination efficiency of InGaN/GaN multiple quantum wells that emit in the green and blue spectral regions Appl. Phys. Lett. 107 132106

[127] Della Sala F, Di Carlo A, Lugli P, Bernardini F, Fiorentini V, Scholz R and Jancu J-M 1999 Free-carrier screening of polarization fields in wurtzite GaN/InGaN laser structures Appl. Phys. Lett. 74 2002

[128] Takeuchi T, Sota S, Katsuragawa M, Komori M, Takeuchi H, Amano H and Akasaki I 1997 Quantum-confined stark effect due to piezoelectric fields in GaInN strained quantum wells Japan. J. Appl. Phys. II 36 L382

[129] Bai J, Wang T and Sakai S 2000 Influence of the quantum well thickness on the radiative recombination of InGaN/GaN quantum well structures J. Appl. Phys. 88 4729

[130] Iliopoulos E 2002 Growth kinetics and Investigations of spontaneous formation of superlattices in AlGaN alloys PhD Dissertation Boston University

[131] Misra M, Korakakis D, Ng H M and Moustakas T D 1999 Photoconductive detectors based on partially ordered AlxGa1−xN alloys grown by molecular beam epitaxy Appl. Phys. Lett. 74 2203

[132] Dudiy S V and Zunger A 2004 Type I to type II transition at the interface between random and ordered domains of AlxGa1−xN alloys Appl. Phys. Lett. 84 1874

[133] Pynn C D, Kowsz S J, Oh S H, Gardner H, Farrell R M, Nakamura S, Speck J S and DenBaars S P 2016 Green semipolar III-nitride light emitting diodes grown by limited area epitaxy Appl. Phys. Lett. 109 041107

[134] Adelmann C, Simon J, Feuillet G, Pelekanos N T and Daudin B 2000 Self-assembled InGaN quantum dots grown by molecular beam epitaxy Appl. Phys. Lett. 76 1570

[135] Damilano B, Grandjean N, Vezian S and Massies J 2001 InGaN heterostructures grown by molecular beam epitaxy: from growth mechanism to optical properties J. Cryst. Growth 227/228 466

[136] Tachibana K, Someya T and Arakawa Y 1999 Nanometer-scale InGaN self-assembled quantum dots grown by metalorganic chemical vapor deposition Appl. Phys. Lett. 74 383

[137] Kim H J et al 2004 Growth of In-rich InGaN/GaN quantum dots by metalorganic chemical vapor deposition J. Cryst. Growth 269 95–9

[138] Ji L-W, Su Y K, Chang S J, Chang C S, Wu L W, Lai W C, Du X L and Chen H 2004 InGaN/GaN multiple-quantum dot light emitting diodes J. Cryst. Growth 263 114

[139] Damilano B, Grandjean N, Massies J, Siozade L and Leymarie J 2000 InGaN/GaN quantum wells grown by molecular beam epitaxy emitting from blue to red at 300 K Appl. Phys. Lett. 77 1268

[140] Xu T, Nikiforov A, France R, Thomidis C, Williams A and Moustakas T D 2007 Blue-green-red LEDs based on InGaN quantum dots (QDs) grown by molecular beam epitaxy Phys. Status Solidi a 204 2098

[141] Xu T, Nikiforov A Yu, France R, Thomidis C, Williams A, Moustakas T D, Zhou L and Smith D J 2007 Blue-green-red LEDs based on InGaN quantum dots by plasma-assisted MBE using GaN QDs for dislocation filtering Material Research Society Symp. Proc. vol 955 0955-I05-05

[142] Moustakas T D, Tao Xu, Thomidis C, Yu Nikiforov A, Zhou L and Smith D J 2008 Growth of III-Nitride quantum dots and their applications to blue-green LEDs Phys. Status Solidi a 205 2560

[143] Thomidis C, Nikiforov A Yu, Xu T and Moustakas T D 2008 InGaN-based LEDs grown by plasma-assisted MBE on (0001) sapphire with GaN QDs in the nucleation layer Phys. Status Solidi c 5 2309

[144] Frost T, Banerjee A, Sun K, Chuang S L and Bhattacharya P 2013 InGaN/GaN quantum dot red (λ = 630 nm) laser IEEE J. Quantum Electron. 49 923

[145] Ding Y J, Reynolds D C, Lee S J, Khurgin J B, Rabinovich W S and Katzer D S 1997 Evidence for strong spatially localized band-filling effects at interface islands Appl. Phys. Lett. 71 2581

[146] Perlin P, Kisielowski C, Iota V, Weinstein B A, Mattos L, Shaprio N A, Kruger J, Weber E R and Yang J 1998 InGaN/GaN quantum wells studied by high pressure, variable temperature, and excitation power spectroscopy Appl. Phys. Lett. 73 2778

[147] Okamoto K, Niki I, Shvartser A, Narukawa Y, Mukai T and Scherer A 2004 Surface-plasmon-enhanced light emitters based on InGaN quantum wells Nat. Mater. 3 601–5

[148] Lu Y C, Chen C Y, Yeh D M, Huang C F, Tang T Y, Huang J J and Yang C C 2007 Temperature dependence of the surface plasmon coupling with an InGaN/GaN quantum well Appl. Phys. Lett. 90 193103

[149] Kwon M K, Kim J Y, Kim B H, Park I K, Cho C Y, Byeon C C and Park S J 2008 Surface-plasmon enhanced light-emitting diodes Adv. Mater. 20 1253–7

[150] Yeh D M, Huang C F, Chen C Y, Lu Y C and Yang C C 2008 Localized surface plasmon-induced emission enhancement of a green light-emitting diode Nanotechnology 19 345201

[151] Oh T S, Jeong H, Lee Y S, Kim J D, Seo T H, Kim H, Park A H, Lee K J and Suh E K 2009 Coupling of InGaN/GaN multiquantum-wells photoluminescence to surface plasmons in platinum nanocluster Appl. Phys. Lett. 95 111112

[152] Henson J, Heckel J C, Dimakis E, Abell J, Bhattacharyya A, Chumanov G, Moustakas T D and Paiella R 2009 Plasmon-enhanced light emission from InGaN quantum wells via coupling to chemically synthesized silver nanoparticles Appl. Phys. Lett. 95 151109

[153] Henson J, Dimakis E, DiMaria J, Li R, Minissale S, Negro L D, Moustakas T D and Paiella R 2010 Enhanced near-green light emission from InGaN quantum wells by use of tunable plasmonic resonances in silver nanoparticle arrays Opt. Express 18 21322

[154] Henson J, DiMaria J, Dimakis E, Moustakas T D and Paiella R 2012 Plasmon-enhanced light emission based on lattice resonances of silver nanocylinder arrays Opt. Lett. 37 79

Rep. Prog. Phys. 80 (2017) 106501

Review

36

[155] Sun G, Khurgin J B and Soref R A 2009 Practical enhancement of photoluminescence by metal nanoparticles Appl. Phys. Lett. 94 101103

[156] Barnes W L 1998 Fluorescence near interfaces: the role of photonic mode density J. Mod. Opt. 45 661–99

[157] Gontijo I, Boroditsky M, Yablonovitch E, Keller S, Mishra U K and DenBaars S P 1999 Coupling of InGaN quantum-well photoluminescence to silver surface plasmons Phys. Rev. B 60 11564–7

[158] Neogi A, Lee C-W, Everitt H O, Kuroda T, Tackeuchi A and Yablonovitch E 2002 Enhancement of spontaneous recombination rate in a quantum well by resonant surface plasmon coupling Phys. Rev. B 66 153305

[159] Laubsch A et al 2009 On the origin of IQE droop in InGaN LEDs Phys. Status Solidi c 6 S913

[160] Piprek J 2010 Efficiency droop in nitride-based light-emitting diode Phys. Status Solidi a 207 2217

[161] Verzellesi G, Saguatti D, Meneghini M, Bertazzi F, Goano M, Meneghesso G and Zanoni E 2013 Efficiency droop in InGaN/GaN blue light-emitting diodes: physical mechanisms and remedies J. Appl. Phys. 114 071101

[162] Nippert F, Karpov S Yu, Callsen G, Galler B, Kure T, Nenstiel C, Wagner M R, Straßburg M, Lugauer H-J and Hoffmann A 2016 Temperature-dependent recombination coefficients in InGaN light-emitting diodes: hole localization, Auger processes, and the green gap Appl. Phys. Lett. 109 161103

[163] Mukai T, Yamada M and Nakamura S 1999 Characteristics of InGaN-based UV/blue/green/amber/red light-emitting diodes Japan. J. Appl. Phys. 38 3976

[164] Yang Y, Cao X A and Yan C 2008 Investigation of the nonthermal mechanisms of efficient rolloff in InGaN light-emitting diodes IEEE Trans. Electron Devices 55 1771

[165] Shen Y C, Mueller G O, Watanabe S, Gardner N F, Munkholm A and Krames M R 2007 Auger recombination in InGaN measured by photoluminescence Appl. Phys. Lett. 91 141101

[166] Kim M H, Schubert M F, Dai Q, Kim J K, Schubert E F, Piprek J and Park Y 2007 Origin of efficiency droop in GaN-based light-emitting diodes Appl. Phys. Lett. 91 183507

[167] Piprek J 2003 Semiconductor Optoelectronic Devices: Introduction to Physics and Simulation (San Diego: Academic)

[168] Ozgur U, Liu H, Li X, Ni X and Morkoc H 2010 GaN-based light-emitting diodes: efficiency at high injection levels Proc. IEEE 98 1180–96

[169] Hader J, Moloney J V, Pasenow B, Koch S W, Sabathil M, Linder N and Lutgen S 2008 On the importance of radiative and Auger losses in GaN-based quantum wells Appl. Phys. Lett. 92 261103

[170] Delaney K T, Rinke P and Van de Walle C G 2009 Auger recombination rates in nitrides from first principles Appl. Phys. Lett. 94 191109

[171] Schlotter P, Schmidt R and Schneider J 1997 Luminescence conversion of blue light emitting diodes Appl. Phys. A 64 417

[172] Tamura T, Setomoto T and Taguchi T 2000 Illumination characteristics of lighting array using 10 candela-class white LEDs under AC 100 V operation J. Lumin. 87–9 1180

[173] Murota R, Kobayashi T and Mita Y 2002 Solid state light source fabricated with YAG:Ce single crystal Japan. J. Appl. Phys. 41 887

[174] Bechtel H, Schmidt P, Busselt W and Schreinemacher B S 2008 Lumiramic: a new phosphor technology for high performance solid state light sources Proc. SPIE 7058 70580E

[175] Raukas M et al 2011 Luminescence ceramics for LED conversion Proc. SPIE 7954 795415

[176] Nishiura S, Tanabe S, Fujioka K and Fujimoto Y 2011 Properties of transparent Ce:YAG ceramic phosphors for white LED Opt. Mater. 33 688

[177] Raukas M, Kelso J, Zheng Y, Bergenek K, Eisert D, Linkov A and Jermann F 2013 Ceramic Phosphors for light conversion in LEDs ECS J. Solid State Sci. Technol. 2 R3168

[178] Park H K, Oh J R and Do Y R 2011 2D SiNx photonic crystal coated Y3Al5O12:Ce3+ ceramic plate phosphor for high-power white light-emitting diodes Opt. Express 19 25593

[179] Joannopoulos J, Johnson S, Win J and Meade R 2008 Photonic Crystals—Molding the Flow of Light (Princeton, NJ: Princeton University Press)

[180] Wierer J J Jr, David A and Megens M M 2009 III-Nitride photonic-crystal light-emitting diodes with high extraction efficiency Nat. Photon. 3 163

[181] Park B K, Park H K, Oh J H, Oh J R and Do Y R 2012 Selecting morphology of Y3Al5O12:Ce3+ phosphors for minimizing scattering loss in the pc-LED package J. Electrochem. Soc. 159 J96

[182] Oh J R, Lee Y K, Park H K and Do Y R 2009 Effect of symmetry, shape, and structural parameters of 2D SiNx photonic crystal on the extracted light from Y2O3:Eu3+ film J. Appl. Phys. 105 043103

[183] Mao A, Schaper C D and Karlicek R F Jr 2013 Nanopatterning using a simple bi-layer lift-off process for the fabrication of a photonic crystal nanostructure Nanotechnology 24 085302

[184] Mao A and Karlicek R F Jr 2013 Surface patterning of nonscattering phosphors for light extraction Opt. Lett. 38 2796

[185] Park H K, Yoon S W, Choi D Y and Do Y R 2013 Fabrication of wafer-scale TiO2 nanobowl arrays via a scooping transfer of polystyrene nanospheres and atomic layer deposition for their application in photonic crystals J. Mater. Chem. C 1 1732

[186] Yoon S W, Park H K, Oh J H and Do Y R 2014 Full extraction of 2-D photonic crystal assisted Y3Al5O12:Ce ceramic plate phosphor for highly efficient white LEDs IEEE Photonics J. 6 8400110

[187] Sun H, Piquette A, Raukas M and Moustakas T D 2016 Enhancement of yellow light extraction efficiency of Y3Al5O12:Ce3+ ceramic converters using a TiO2 photonic crystal IEEE Photonics J. 8 4500310

[188] Kneissl M 2015 UV LED efficiency www.researchgate.net/publication/280131929 (Retrieved: 6 October 2015) last update 19 July 2015

[189] Khan A, Balakrishnan K and Katona T 2008 Ultraviolet light-emitting diodes based on group III nitrides Nat. Photonics 2 77

[190] Pernot C et al 2010 Improved efficiency of 255–280 nm AlGaN-based light-emitting diodes Appl. Phys. Express 3 061004

[191] Hirayama H, Tsukada Y, Maeda T and Kamata N 2010 Marked enhancement in the efficiency of deep-ultraviolet AlGaN light-emitting diodes by using a multiquantum-barrier electron blocking layer Appl. Phys. Express 3 031002

[192] Grandusky R, Gibb S R, Mendrick M C, Moe M, Wraback C and Schowalter L J 2011 High output power from 260 nm pseudomorphic ultraviolet light-emitting diodes with improved thermal performance Appl. Phys. Express 4 082101

[193] Shatalov M et al 2012 AlGaN deep-ultraviolet light-emitting diodes with external quantum efficiency above 10% Appl. Phys. Express 5 082101

[194] Wright A N and Winkler K A 1968 Active Nitrogen (New York: Academic)

[195] Nix W D and Clemens B M 1999 Crystallite coalescence: a mechanism for intrinsic tensile stresses in thin films J. Mater. Res. 14 3467

Rep. Prog. Phys. 80 (2017) 106501

Review

37

[196] Moustakas T D, Liao Y, Kao C-K, Thomidis C, Bhattacharyya A, Bhattarai D and Moldawer A 2012 Deep UV-LEDs with high IQE based on AlGaN alloys with strong band structure potential fluctuations Proc. SPIE 8278 82780L

[197] Elsass C R, Mates T, Heying B, Poblenz C, Fini P, Petroff P M, DenBaars S P and Speck J S 2000 Effects of growth conditions on the incorporation of oxygen in AlGaN layers grown by plasma-assisted molecular beam epitaxy Appl. Phys. Lett. 77 3167

[198] Moustakas T D and Bhattacharyya A 2011 Experimental evidence that the plasma-assisted MBE growth of nitride alloys is a liquid phase epitaxy process ECS Trans. 35 63–71

[199] Moustakas T D and Bhattacharyya A 2012 The role of liquid phase epitaxy during growth of AlGaN by MBE Phys. Status Solidi c 9 580–3

[200] McCluskey M D, Johnson N M, Van de Walle C G, Bour D B, Kneissl M and Walukiewicz W 1998 Metastability of oxygen donors in AlGaN Phys. Rev. Lett. 80 4008

[201] Taniyasu Y, Kasu M and Makimoto T 2006 An aluminum nitride light-emitting diode with wavelength of 210 nm Nature 441 325

[202] Nam K B, Li J, Nakarmi M L, Lin J Y and Jiang H X 2004 Unique optical properties of AlGaN alloys and related ultraviolet emitters Appl. Phys. Lett. 84 5264

[203] Shakya J, Knabe K, Kim K H, Li J, Lin J Y and Jiang H X 2005 Polarization of III-nitride blue and ultraviolet light-emitting diodes Appl. Phys. Lett. 86 091107

[204] Karpov S and Makarov Y N 2002 Dislocation effect on light emission efficiency in gallium nitride Appl. Phys. Lett. 81 4721

[205] Monroy E, Daudin B, Bellet-Amalric E, Gogneau N, Jalabert D, Brault J, Barjon J and Dang L S 2003 Surfactant effect of In for AlGaN growth by plasma-assisted molecular beam epitaxy J. Appl. Phys. 93 1550

[206] Iliopoulos E and Moustakas T D 2002 Growth kinetics of AlGaN films by plasma assisted molecular beam epitaxy Appl. Phys. Lett. 81 295

[207] Bhattacharyya A, Moustakas T D, Zhou L, Smith D J and Hug W 2009 Deep ultraviolet emitting AlGaN quantum wells with high internal quantum efficiency Appl. Phys. Lett. 94 181907

[208] Zhang W et al 2012 MBE growth of AlGaN quantum wells on 6H-SiC substrates with high internal quantum efficiency J. Vac. Sci. Technol. B 30 02B119

[209] Liao Y, Thomidis C, Kao C-K and Moustakas T D 2011 AlGaN based deep ultraviolet light emitting diodes with high internal quantum efficiency grown by molecular beam epitaxy Appl. Phys. Lett. 98 081110

[210] Pecora E F, Zhang W, Zhou L, Smith D J, Yin J, Paiella R, Dal Negro L and Moustakas T D 2012 Sub-250 nm room-temperature optical gain from AlGaN/AlN multiple quantum wells with strong band-structure potential fluctuations Appl. Phys. Lett. 100 061111

[211] Pecora E F, Zhang W, Nikiforov A Yu, Yin J, Paiella R, Dal Negro L and Moustakas T D 2013 Sub-250 nm light emission and optical gain in AlGaN materials J. Appl. Phys. 113 013106

[212] Moustakas T D 2016 Ultraviolet optoelectronic devices based on AlGaN alloys grown by MBE MRS Commun. 6 247–69

[213] Nakamura S, Senoh M, Nagahama S, Iwasa N, Yamada T, Matsushita T, Sugimoto Y and Kiyoku H 1996 Room-temperature continuous-wave operation of InGaN multi-quantum-well structure laser diodes Appl. Phys. Lett. 69 4056

[214] Nakamura S et al 1997 InGaN/GaN/AlGaN-based laser diodes with modulation-doped strained-layer superlattices Japan. J. Appl. Phys. 36 L1568

[215] Bour D P, Kneissl M, Romano L T, McCluskey M, Van de Walle C G, Krusor B S, Donaldson R, Walker J, Dunnrowicz C and Johnson N M 1998 Characteristics of InGaN-AlGaN multiple-quantum-well laser diodes IEEE J. Sel. Top. Quantum Electron. 4 498

[216] Chichibu S F et al 1999 Optical properties of InGaN quantum wells Mater. Sci. Eng. B 59 298

[217] Uchida S, Takeya M, Ikeda S, Mizuno T, Fujimoto T, Matsumoto O, Goto S, Tojyo T and Ikeda M 2003 Recent progress in high-power blue-violet lasers IEEE J. Sel. Top. Quantum Electron. 9 1252

[218] Nagahama S, Sugimoto Y, Kozaki T and Mukai T 2005 Recent progress of AlInGaN laser diodes Proc. SPIE 5738 57

[219] Lee S N et al 2006 High-power GaN-based blue-violet laser diodes with AlGaN∕GaN multiquantum barriers Appl. Phys. Lett. 88 111101

[220] Miyoshi T, Masui S, Okada T, Yanamoto T, Kozaki T, Nagahama S and Mukai T 2009 510–515 nm InGaN-based green laser diodes on c-plane GaN substrate Appl. Phys. Express 2 062201

[221] Enya Y et al 2009 531 nm green lasing of InGaN based laser diodes on semi-polar 20–21 free-standing GaN substrates Appl. Phys. Express 2 082101

[222] Tyagi A et al 2010 AlGaN-cladding free green semipolar GaN based laser diode with a lasing wavelength of 506.4 nm Appl. Phys. Express 3 011002

[223] Lutgen S, Avramescu A, Lermer T, Queren D, Muller J, Bruederl G and Strauss U 2010 True green InGaN laser diodes Phys. Status Solidi a 207 1318

[224] Raring J W et al 2010 High-efficiency blue and true-green-emitting laser diodes based on non-c-plane oriented GaN substrates Appl. Phys. Express 3 112101

[225] Sizov D, Bhat R and Zah C-E 2012 Gallium indium nitride-based green lasers J. Lightwave Technol. 30 679

[226] Takagi S et al 2012 High-power (over 100 mW) green laser diodes on semipolar 20–21 GaN substrates operating at wavelengths beyond 530 nm Appl. Phys. Express 5 082102

[227] Kneissl M, Treat D W, Teepe M, Miyashita N and Johnson N M 2003 Ultraviolet AlGaN multiple-quantum-well laser diodes Appl. Phys. Lett. 82 4441

[228] Masui S, Matsuyama Y, Yanamoto T, Kozaki T, Nagahama S and Mukai T 2003 365 nm ultraviolet laser diodes composed of quaternary AlInGaN alloy Japan. J. Appl. Phys. 42 L1318

[229] Takano T, Narita Y, Horiuchi A and Kawanishi H 2004 Room-temperature deep-ultraviolet lasing at 241.5 nm of AlGaN multiple-quantum-well laser Appl. Phys. Lett. 84 3567

[230] Kamiyama S, Iida K, Kawashima T, Kasugai H, Mishima S, Honshio A, Miyake Y, Iwaya M, Amano H and Akasaki I 2005 UV laser diode with 350.9-nm-lasing wavelength grown by hetero-epitaxial-lateral overgrowth technology IEEE J. Select. Quantum Electron. 11 1069

[231] Shatalov M, Gaevski M, Adivarahan V and Khan A 2006 Room-temperature stimulated emission from AlN at 214 nm Japan. J. Appl. Phys. 45 L1286

[232] Kneissl M, Yang Z, Teepe M, Knollenberg C, Schmidt O, Kiesel P and Johnson N M 2007 Ultraviolet semiconductor laser diodes on bulk AlN J. Appl. Phys. 101 123103

[233] Yoshida H, Yamashita Y, Kuwabara M and Kan H 2008 A 342-nm ultraviolet AlGaN multiple-quantum-well laser diode Nat. Photon. 2 551

[234] Yoshida H, Yamashita Y, Kuwabara M and Kan H 2008 Demonstration of an ultraviolet 336 nm AlGaN multiple-quantum-well laser diode Appl. Phys. Lett. 93 241106

[235] Jmerik V N, Mizerov A M, Sitnikova A A, Kop’ev P S, Ivanov S V, Lutsenko E V, Tarasuk N P, Rzheutskii N V and Yablonskii G P 2010 Low-threshold 303 nm lasing

Rep. Prog. Phys. 80 (2017) 106501

Review

38

in AlGaN-based multiple-quantum well structures with an asymmetric waveguide grown by plasma-assisted molecular beam epitaxy on c-sapphire Appl. Phys. Lett. 96 141112

[236] Wunderer T, Chua C L, Yang Z, Northrup J E, Johnson N M, Garrett G A, Shen H and Wraback M 2011 Pseudomorphically grown ultraviolet C photopumped lasers on bulk AlN substrates Appl. Phys. Express 4 092101

[237] Pecora E F, Zhang W, Yin J, Paiella R, Dal Negro L and Moustakas T D 2012 Polarization properties of deep-ultraviolet optical gain in Al-rich AlGaN structures Appl. Phys. Express 5 032103

[238] Nakamura S and Chichibu S F (ed) 2000 Introduction to Nitride Semiconductor Blue Lasers and Light Emitting Diodes (London: Taylor and Francis)

[239] Nakamura S 1998 The roles of structural imperfections in InGaN-based blue light-emitting diodes and laser diodes Science 281 956

[240] Hangleiter A 2002 Optical gain in wide-bandgap group-III nitrides M Osiński, ed Gallium-Nitride-Based Technologies (Bellingham, WA: SPIE Optical Engineering Press)

[241] Coldren L A and Corzine S W 1995 Diode Lasers and Photonic Integrated Circuits (New York: Wiley)

[242] Jeon J B, Lee B C, Sirenko Y M, Kim K W and Littlejohn M A 1997 Strain effects on optical gain in wurtzite GaN J. Appl. Phys. 82 386

[243] Okamoto K, Ohta H, Chichibu S F, Ichihara J and Takasu H 2007 Continuous-wave operation of m-plane InGaN multiple quantum well laser diodes Japan. J. Appl. Phys. 46 L187

[244] Schmidt M C, Kim K C, Farrell R M, Feezell D F, Cohen D A, Saito M, Fujito K, Speck J S, DenBaars S P and Nakamura S 2007 Demonstration of nonpolar m-plane InGaN/GaN laser diodes Japan. J. Appl. Phys. 46 L190

[245] Mukai T, Nagahama S, Yanamoto T and Sano M 2002 Wide-bandgap group-III nitride lasers Gallium-Nitride-Based Technologies ed M Osiński (Bellingham, WA: SPIE Optical Engineering Press)

[246] Khan M A, Kuznia J N, Olson D T, Van Hove J M and Blasingame M 1992 High-responsivity photoconductive ultraviolet sensors based on insulating single-crystal GaN epilayers Appl. Phys. Lett. 60 2917–9

[247] Misra M, Moustakas T D, Vaudo R P, Singh R and Shah K S 1995 Photoconducting ultraviolet detectors based on GaN films grown by electron cyclotron resonance molecular beam epitaxy. X-ray and ultraviolet sensors and applications Proc. SPIE 2519 78–86

[248] Razeghi M and Rogalski A 1996 Semiconductor ultraviolet detectors J. Appl. Phys. 79 7433–73

[249] Carrano J C, Li T, Grudowski P A, Eiting C J, Dupuis R D and Campbell J C 1998 Comprehensive characterization of metal-semiconductor-metal ultraviolet photodetectors fabricated on single crystal GaN J. Appl. Phys. 83 6148–60

[250] Campbell J C, Collins C J, Wong M M, Chowdhury U and Dupuis R D 2002 Back-illuminated solar-blind AlxGa1−xN p-i-n photodiodes Gallium-Nitride-Based Technologies ed M Osiński (Bellingham, WA: SPIE Optical Engineering Press)

[251] Misra M and Moustakas T D 2000 Photoconductivity recombination kinetics in GaN films Mater. Res. Soc. Proc. 662 T5.4.1

[252] Monroy E, Calle F, Pau J L, Muñoz E, Omnès F, Beaumont B and Gibart P 2001 AlGaN-based UV photodetectors J. Cryst. Growth 230 537–43

[253] Kuryatkov V, Chandolu A, Borisov B, Kipshidze G, Zhu K, Nikishin S and Temkin H 2003 Solar-blind ultraviolet photodetectors based on superlattices of AlN/AlGa(In)N Appl. Phys. Lett. 82 1323–5

[254] Moustakas T D and Misra M 2007 Origin of the high photoconductivity gain in AlGaN films Proc. SPIE 6766 67660C

[255] Misra M, Bhattacharyya A and Moustakas T D 2008 Nitride-based UV detectors improve photodetection Laser Focus World (November 2008) pp 64–6 (www.laserfocusworld.com)

[256] Hamamatsu Photodiode Catalog 1996 [257] Rose A 1963 Concepts in Photoconductivity and Allied

Problems (New York: Interscience Publishers) ch 3 [258] Bube R H and Redfield D 1989 Variation of

photoconductivity with doping and optical degradation in hydrogenated amorphous silicon J. Appl. Phys. 66 3074

[259] Dudiy S V and Zunger A 2003 Optical consequences of long-range order in wurtzite AlxGa1−xN alloys Phys. Rev. B 68 041302

[260] Ahrenkiel R K, Ahrenkial S P, Arent D J and Olson J M 1997 Carrier transport in ordered and disordered In0.53Ga0.47As Appl. Phys. Lett. 70 756

[261] Van Hove J M, Hickman R, Klaassen J J and Chow P P, and Ruden P P 1997 Ultraviolet-sensitive, visible-blind GaN photodiodes fabricated by molecular beam epitaxy Appl. Phys. Lett. 70 2282

[262] McClintock R, Yasan A, Mayes K, Shiell D, Darvish S R, Kung P and Razeghi M 2004 High quantum efficiency AlGaN solar-blind p-i-n photodiodes Appl. Phys. Lett. 84 1248

[263] Chemla D S, Miller D A B and Smith P W 1987 Nonlinear optical properties of multiple quantum well structures for optical signal processing Semiconductors and Semimetals vol 34, ed R Dingle (New York: Academic) ch 5

[264] Wood T H 1988 Multiple quantum well (MQW) waveguide modulators J. Lightwave Technol. 6 743

[265] Cunningham J E 1999 Recent developments and applications in electroabsorption semiconductor modulators Mater. Sci. Eng. Res. 25 155

[266] Kneissl M, Paoli T L, Kiesel P, Treat D W, Teepe M, Miyashita N and Johnson N M 2002 Two-section InGaN multiple quantum well laser diode with integrated electroabsorption modulator Appl. Phys. Lett. 80 3283

[267] Oberhofer A E, Muth J F, Johnson M A L, Chen Z Y, Fleet E F and Cooper G D 2003 Planar gallium nitride ultraviolet optical modulator Appl. Phys. Lett. 83 2748

[268] Friel I, Thomidis C and Moustakas T D 2005 Ultraviolet electroabsorption modulator based on AlGaN/GaN multiple quantum wells J. Appl. Phys. 97 123515

[269] Sari E, Nizamoglu S, Ozel T and Demir H V 2007 Blue quantum electroabsorption modulators based on reversed quantum confined Stark effect with blueshift Appl. Phys. Lett. 90 011101

[270] Ozel T, Sari E, Nizamoglu S and Demir H V 2007 Violet to deep ultraviolet InGaN/GaN and GaN/AlGaN quantum structures for UV electroabsorption modulators J. Appl. Phys. 102 113101

[271] Kao C-K, Bhattacharyya A, Thomidis Ch, Paiella R and Moustakas T D 2011 Electroabsorption modulators based on bulk GaN films and GaN/AlGaN multiple quantum wells J. Appl. Phys. 109 083102

[272] Kao C-K, Bhattacharyya A, Thomidis Ch, Moldawer A, Paiella R and Moustakas T D 2012 A comparative study of UV electroabsorption modulators based on bulk III-nitride films and multiple quantum wells Phys. Status Solidi c 9 770–3

[273] Shaw G A, Siegel A M, Model J and Greisokh D 2005 Recent progress in short-range ultraviolet communication Proc. SPIE 5796 214

[274] Muth J F, Lee J H, Shmagin I K, Kolbas R M, Casey H C, Keller B P, Mishra U K and DenBaars S P 1997 Absorption coefficient, energy gap, exciton binding energy, and recombination lifetime of GaN obtained from transmission measurements Appl. Phys. Lett. 71 2572

Rep. Prog. Phys. 80 (2017) 106501

Review

39

[275] Yan R-H, Simes R J and Coldren L A 1991 Surface-normal electroabsorption reflection modulators using asymmetric Fabry–Perot structures IEEE J. Quantum Electron. 27 1922

[276] Paiella R (ed) 2006 Intersubband Transitions in Quantum Structures (New York: McGraw-Hill)

[277] Jovanovic V D, Indjin D, Ikonic Z and Harrison P 2004 Simulation and design of GaN/AlGaN far-infrared (λ ~ 34 µm) quantum-cascade laser Appl. Phys. Lett. 84 2995

[278] Sun G, Soref R A and Khurgin J B 2005 Active region design of a terahertz GaN/Al0.15Ga0.85N quantum cascade laser Superlattices Microstruct. 37 107

[279] Bellotti E, Driscoll K, Moustakas T D and Paiella R 2008 Monte Carlo study of GaN versus GaAs terahertz quantum cascade structures Appl. Phys. Lett. 92 101112

[280] Bellotti E, Driscoll K, Moustakas T D and Paiella R 2009 Monte Carlo simulation of terahertz quantum cascade laser structures based on wide-bandgap semiconductors J. Appl. Phys. 105 113103

[281] Williams B S 2007 Terahertz quantum-cascade lasers Nat. Photon. 1 517

[282] Fathololoumi S, Dupont E, Chan C W I, Wasilewski Z R, Laframboise S R, Ban D, Mátyás A, Jirauschek C, Hu Q and Liu H C 2012 Terahertz quantum cascade lasers operating up to ~ 200 K with optimized oscillator strength and improved injection tunneling Opt. Express 20 3866

[283] Gmachl C, Ng H M, Chu S-N G and Cho A Y 2000 Intersubband absorption at λ~1.55 µm in well- and modulation-doped GaN/AlGaN multiple quantum wells with superlattice barriers Appl. Phys. Lett. 77 3722

[284] Kishino K, Kikuchi A, Kanazawa H and Tachibana T 2002 Intersubband transition in (GaN)m/(AlN)n superlattices in the wavelength range from 1.08 to 1.61 µm Appl. Phys. Lett. 81 1234

[285] Zhou Q, Chen J, Pattada B, Manasreh M O, Xiu F, Puntigan S, He L and Ramaiah K S and Morkoç H 2003 Infrared optical absorbance of intersubband transitions in GaN/AlGaN multiple quantum well structures J. Appl. Phys. 93 10140

[286] Friel I, Driscoll K, Kulenica E, Dutta M, Paiella R and Moustakas T D 2005 Investigation of the design parameters of AlN/GaN multiple quantum wells grown by molecular beam epitaxy for intersubband absorption J. Cryst. Growth 278 387

[287] Tchernycheva M, Nevou L, Doyennette L, Julien F H, Warde E, Guillot F, Monroy E, Bellet-Amalric E, Remmele T and Albrecht M 2006 Systematic experimental and theoretical investigation of intersubband absorption in GaN∕AlN quantum wells Phys. Rev. B 73 125347

[288] Driscoll K, Bhattacharyya A, Moustakas T D, Paiella R, Zhou L and Smith D J 2007 Intersubband absorption in AlN ∕ GaN ∕ AlGaN coupled quantum wells Appl. Phys. Lett. 91 141104

[289] Nevou L, Tchernycheva M, Julien F H, Guillot F and Monroy E 2007 Short wavelength (λ = 2.13µm) intersubband luminescence from GaN∕AlN quantum wells at room temperature Appl. Phys. Lett. 90 121106

[290] Driscoll K, Liao Y, Bhattacharyya A, Zhou L, Smith D J, Moustakas T D and Paiella R 2009 Optically pumped intersubband emission of short-wave infrared radiation with GaN/AlN quantum wells Appl. Phys. Lett. 94 081120

[291] Iizuka N, Kaneko K and Suzuki N 2006 All-optical switch utilizing intersubband transition in GaN quantum wells IEEE J. Quantum Electron. 42 765

[292] Li Y, Bhattacharyya A, Thomidis C, Moustakas T D and Paiella R 2007 Ultrafast all-optical switching with low saturation energy via intersubband transitions in GaN/AlN quantum-well waveguides Opt. Express 15 17922

[293] Paiella R, Driscoll K, Li Y, Liao Y, Bhattacharyya A, Thomidis C, Zhou L, Smith D J, Bellotti E and Moustakas T D 2010 Intersubband device applications of nitride quantum structures. Quantum sensing and nanophotonic devices VII Proc. SPIE 7608 76080N

[294] Hofstetter D, Schad S-S, Wu H, Schaff W J and Eastman L F 2003 GaN/AlN-based quantum-well infrared photodetector for 1.55 µm Appl. Phys. Lett. 83 572

[295] Baumann E, Giorgetta F R, Hofstetter D, Lu H, Chen X, Schaff W J, Eastman L F, Golka S, Schrenk W and Strasser G 2005 Intersubband photoconductivity at 1.6µm using a strain-compensated AlN∕GaN superlattice Appl. Phys. Lett. 87 191102

[296] Giorgetta F R, Baumann E, Guillot F, Monroy E and Hofstetter D 2007 High frequency (f = 2.37 GHz) room temperature operation of 1.55 µm AlN/GaN-based intersubband detector Electron. Lett. 43 185

[297] Vardi A, Bahir G, Guillot F, Bougerol C, Monroy E, Schacham S E, Tchernycheva M and Julien F H 2008 Near infrared quantum cascade detector in GaN∕AlGaN∕AlN heterostructures Appl. Phys. Lett. 92 011112

[298] Vardi A et al 2008 High-speed operation of GaN/AlGaN quantum cascade detectors at λ ≈ 1.55 µm Appl. Phys. Lett. 93 193509

[299] Baumann E, Giorgetta F R, Hofstetter D, Leconte S, Guillot F, Bellet-Amalric E and Monroy E 2006 Electrically adjustable intersubband absorption of a GaN∕AlN superlattice grown on a transistorlike structure Appl. Phys. Lett. 89 101121

[300] Nevou L et al 2007 Short-wavelength intersubband electroabsorption modulation based on electron tunneling between GaN∕AlN coupled quantum wells Appl. Phys. Lett. 90 223511

[301] Hofstetter D, Bour D P and Kirste L 2014 Mid-infrared electro-luminescence and absorption from AlGaN/GaN-based multi-quantum well inter-subband structures Appl. Phys. Lett. 104 241107

[302] Song A Y, Bhat R, Allerman A A, Wang J, Huang T-Y, Zah C-E and Gmachl C F 2015 Quantum cascade emission in the III-nitride material system designed with effective interface grading Appl. Phys. Lett. 107 132104

[303] Machhadani H, Kotsar Y, Sakr S, Tchernycheva M, Colombelli R, Mangeney J, Bellet-Amalric E, Sarigiannidou E, Monroy E and Julien F H 2010 Terahertz intersubband absorption in GaN/AlGaN step quantum wells Appl. Phys. Lett. 97 191101

[304] Edmunds C, Shao J, Shirazi-Hd M, Manfra M J and Malis O 2014 Terahertz intersubband absorption in non-polar m-plane AlGaN/GaN quantum wells Appl. Phys. Lett. 105 021109

[305] Beeler M, Bougerol C, Bellet-Amalric E and Monroy E 2014 Pseudo-square AlGaN/GaN quantum wells for terahertz absorption Appl. Phys. Lett. 105 131106

[306] Lim C B, Ajay A, Bougerol C, Haas B, Schoermann J, Beeler M, Laehnemann J, Eickhoff M and Monroy E 2015 Nonpolar m-plane GaN/AlGaN heterostructures with intersubband transitions in the 5–10 THz band Nanotechnology 26 435201

[307] Sudradjat F F, Zhang W, Woodward J, Durmaz H, Moustakas T D and Paiella R 2012 Far-infrared intersubband photodetectors based on double-step III-nitride quantum wells Appl. Phys. Lett. 100 241113

[308] Durmaz H, Nothern D, Brummer G C, Moustakas T D and Paiella R 2016 Terahertz intersubband photodetectors based on semi-polar GaN/AlGaN heterostructures Appl. Phys. Lett. 108 201102

[309] Terashima W and Hirayama H 2011 Spontaneous emission from GaN/AlGaN terahertz quantum cascade laser grown on GaN substrate Phys. Status Solidi c 8 2302

Rep. Prog. Phys. 80 (2017) 106501

Review

40

[310] Graf M, Scalari G, Hofstetter D, Faist J, Beere H, Linfield E, Ritchie D and Davies G 2004 Terahertz range quantum well infrared photodetector Appl. Phys. Lett. 84 475

[311] Luo H, Liu H C, Song C Y and Wasilewski Z R 2005 Background-limited terahertz quantum-well photodetector Appl. Phys. Lett. 86 231103

[312] Gmachl C and Ng H M 2003 Intersubband absorption at λ ~ 2.1 µm in A-plane GaN/AlN multiple quantum wells Electron. Lett. 39 567

[313] Pesach A, Gross E, Huang C-Y, Lin Y-D, Vardi A, Schacham S E, Nakamura S and Bahir G 2013 Non-polar m-plane intersubband based InGaN/(Al)GaN quantum well infrared photodetectors Appl. Phys. Lett. 103 022110

[314] Kotani T, Arita M and Arakawa Y 2014 Observation of mid-infrared intersubband absorption in non-polar m-plane AlGaN/GaN multiple quantum wells Appl. Phys. Lett. 105 261108

[315] Lim C B, Beeler M, Ajay A, Laehnemann J, Bellet-Amalric E, Bougerol C and Monroy E 2015 Intersubband transitions in nonpolar GaN/Al(Ga)N heterostructures in the short- and mid-wavelength infrared regions J. Appl. Phys. 118 014309

[316] Kikuchi A, Bannai R, Kishino K, Lee C-M and Chyi J-I 2002 AlN/GaN double-barrier resonant tunneling diodes grown by rf-plasma-assisted molecular-beam epitaxy Appl. Phys. Lett. 81 1729–31

[317] Golka S, Pflügl C, Schrenk W, Strasser G, Skierbiszewski C, Siekacz M, Grzegory I and Porowski S 2006 Negative differential resistance in dislocation-free GaN∕AlGaN double-barrier diodes grown on bulk GaN Appl. Phys. Lett. 88 172106

[318] Bayram C, Vashaei Z and Razeghi M 2010 Reliability in room-temperature negative differential resistance characteristics of low-aluminum content AlGaN/GaN double-barrier resonant tunneling diodes Appl. Phys. Lett. 97 181109

[319] Boucherit M, Soltani A, Monroy E, Rousseau M, Deresmes D, Berthe M, Durand C and De Jaeger J-C 2011 Investigation of the negative differential resistance reproducibility in AlN/GaN double-barrier resonant tunnelling diodes Appl. Phys. Lett. 99 182109

[320] Li D, Tang L, Edmunds C, Shao J, Gardner G, Manfra M J and Malis O 2012 Repeatable low-temperature negative-differential resistance from Al0.18Ga0.82N/GaN resonant tunneling diodes grown by molecular-beam epitaxy on free-standing GaN substrates Appl. Phys. Lett. 100 252105

[321] Sudradjat F, Zhang W, Driscoll K, Liao Y, Bhattacharyya A, Thomidis C, Zhou L, Smith D J, Moustakas T D and Paiella R 2010 Sequential tunneling transport characteristics of GaN/AlGaN coupled-quantum-well structures J. Appl. Phys. 108 103704

[322] Amano H 2013 Progress and prospect of the growth of wide-band-gap group III nitrides: development of the growth method for single-crystal Bulk GaN Japan. J. Appl. Phys. 52 050001

[323] Kelly M K, Ambcher O, Dimitrov R, Handschuh R and Stutzman M 1997 Optical process for liftoff of group III-nitride films Phys. Status Solidi a 159 R3

[324] Oshima Y, Eri T, Shibata M, Sunakawa H, Kobayashi K, Ichihashi T and Usui A 2003 Preparation of freestanding GaN wafers by hydride vapor phase epitaxy with void-assisted separation Japan. J. Appl. Phys. 42 L1

[325] Williams A and Moustakas D T 2007 Formation of large-area freestanding gallium nitride substrates by natural stress-induced separation of GaN and sapphire J. Cryst. Growth 300 37

[326] Karpinski J, Porowski S and Miotkowska S 1982 High pressure vapor growth of GaN J. Cryst. Growth 56 77

[327] Porowski S 1996 High pressure growth of GaN-new prospects for blue lasers J. Cryst. Growth 166 583

[328] Inoue T, Seki Y, Oda O, Kurai S, Yamada Y and Taguchi T 2001 Pressure-controlled solution growth of bulk GaN crystals under high pressure Phys. Status Solidi b 223 15

[329] Dwilinski R, Wysmolek A, Baranowski J, Kaminska M, Doradzinski R and Jacobs H 1995 Acta Phys. Pol. A 88 833

[330] Dwilinski R, Doradzinski R, Garczynski J, Sierzputowski L P, Puchalski A, Kanbara Y, Yagi K, Minakuchi H and Hayashi H 2008 Excellent crystallinity of truly bulk ammonothermal GaN J. Cryst. Growth 310 3911–6

[331] Cich M J, Aldaz R I, Chakraborty A, David A, Grundmann M J, Tyagi A, Zhang M, Steranka F M and Krames M R 2012 Bulk GaN based violet light-emitting diodes with high efficiency at very high current density Appl. Phys. Lett. 101 223509

[332] David A, Hurni C A, Aldaz R I, Cich M J, Ellis B, Huang K, Steranka F M and Krames M R 2014 High light extraction efficiency in bulk-GaN based volumetric violet light-emitting diodes Appl. Phys. Lett. 105 231111

[333] Hurni C A et al 2015 Bulk GaN flip-chip violet light-emitting diodes with optimized efficiency for high-power operation Appl. Phys. Lett. 106 031101

[334] David A, Hurni C A, Young N G and Craven M D 2016 Electrical properties of III-nitride LEDs: recombination-based injection model and theoretical limits to electrical efficiency and electroluminescent cooling Appl. Phys. Lett. 109 083501

[335] David A, Hurni C A, Young N G and Craven M D 2016 Carrier dynamics and Coulomb-enhanced capture in III-nitride quantum heterostructures Appl. Phys. Lett. 109 033504

[336] Speck J S and Chichibu S F (ed) 2009 Nonpolar and semipolar group-III nitride-based materials MRS Bull. 34 304–52

[337] Zhao Y et al 2011 High optical polarization ratio from semipolar (20-2-1) blue-green InGaN/GaN light emitting diodes Appl. Phys. Lett. 99 051109

[338] Saito S, Hashimoto R, Hwang J and Nunoue S 2013 InGaN light-emitting diodes on c-face sapphire substrates in green gap spectral range Appl. Phys. Express 6 111004

[339] Farrell R M, Young E C, Wu F, DenBaars S P and Speck J S 2012 Materials and growth issues for high-performance nonpolar and semipolar light-emitting devices Semicond. Sci. Technol. 27 024001

[340] Guha S and Bojarczuk N A 1998 Ultraviolet and violet GaN light emitting diodes on silicon Appl. Phys. Lett. 72 415

[341] Nakada Y and Aksenov I and Okumura H 1998 GaN heteroepitaxial growth on silicon nitride buffer layers formed on Si (111) surfaces by plasma-assisted molecular beam epitaxy Appl. Phys. Lett. 73 827

[342] Nikishin S A, Faleev N N, Antipov V G, Francoeur S, Grave de Peralta L, Seryogin G A, Temkin H, Prokofyeva T I, Holtz M and Chu S N G 1999 High quality GaN grown on Si (111) by gas source molecular beam epitaxy with ammonia Appl. Phys. Lett. 75 2073

[343] Feltin E, Beaumont B, Laugt M, de Mierry P, Vennegues P, Lahreche H, Leroux M and Gibart P 2001 Stress control in GaN grown on silicon (111) by metalorganic vapor phase epitaxy Appl. Phys. Lett. 79 3230

[344] Krost A and Dadgar A 2002 GaN-based optoelectronics on silicon substrates Mater. Sci. Eng. B 93 77

[345] Zhang B, Liang H, Wang Y, Feng Z, Ng K W and Lau K M 2007 High-performance III-nitride blue LEDs grown and fabricated on patterned Si substrates J. Cryst .Growth 298 725

[346] Hsu Y P, Chang S J, Chen W S, Sheu J K, Chu J Y and Kuo C T 2007 Crack-free high-brightness InGaN/GaN LEDs on Si (111) with initial AlGaN buffer and two LT-Al interlayers J. Electrochem. Soc. 154 H191

Rep. Prog. Phys. 80 (2017) 106501

Review

41

[347] Le Louarn A, Vezian S, Semond F and Massies J 2009 AlN buffer layer growth for GaN epitaxy on (111) Si: Al or N first? J. Cryst. Growth 311 3278–84

[348] Dadgar A, Blasing J, Diez A and Krost A 2011 Crack-free, highly conducting GaN layers on Si substrates by Ge doping Appl. Phys .Express 4 011001

[349] Bessolov V N, Konenkova E V, Kukushkin S A, Osipov A V and Rodin S N 2014 Semipolar gallium nitride on silicon: technology and properties Rev. Adv. Mater. Sci. 38 75–93

[350] Hoke W E, Kennedy T D, Torabi A, Lyman P S, Alan Howsare C and Schultz B D 2014 Highly uniform AlGaN/GaN HEMT films grown on 200-mm silicon substrates by plasma molecular beam epitaxy J. Vac. Sci. Technol. B 32 030605

[351] Li K H, Liu X, Wang Q, Zhao S and Mi Z 2015 Ultralow-threshold electrically injected AlGaN nanowire ultraviolet lasers on Si operating at low temperature Nat. Nanotechnol. 10 140

[352] Zhao S, Djavid M and Mi Z 2015 Surface emitting, high efficiency near-vacuum ultraviolet light source with aluminum nitride nanowires monolithically grown on silicon Nano Lett. 15 7006

[353] Leszcynski M, Suski T, Perlin P, Tolsseyre H, Grzegory I, Bockowski M, Jun J, Porowski S and Major J 1995 Lattice constants, thermal expansion and compressibility of gallium nitride J. Phys. D: Appl. Phys. 28 A149

[354] Morkoc H 1998 Gallium Nitride (GaN) I: Semiconductors and Semimetals vol 50, ed J I Pankove and T D Moustakas (San Diego: Academic) ch 8

[355] Azuhata T, Sota T and Suzuki K 1996 Elastic constants of III - V compound semiconductors: modification of Keyes’ relation J. Phys: Condens. Mater 8 3111

[356] Van Vechten J A 1973 Quantum dielectric theory of electronegativity in covalent systems. III. Pressure-temperature phase diagrams, heats of mixing, and distribution coefficients Phys. Rev. B 7 1479

Theodore D. Moustakas is the inaugural Distinguished Professor of Photonics and Optoelectronics at Boston University. He received his PhD from Columbia University in 1974. He held research positions at Harvard University and Exxon Corporate Research Laboratory prior to joining Boston University in 1987 as a Professor of Electrical and Computer Engineering. He is currently an Emeritus Professor of the Electrical and Computer Engineering, of the Division of Materials Science and Engineering and Physics Department at Boston University. Dr. Moustakas is a Fellow of the American Physical Society, the Electrochemical Society, the Institute of Electrical and Electronic Engineers-IEEE, and a Charter Fellow of the National Academy of Inventors. He holds an honorary doctoral degree from the Aristotle University; he received the MBE Innovator Award in 2010; in 2011 he received the Distinguished Scholar Award from the BU College of Engineering and in 2013 he received the Boston University Innovator of the Year Award.

Roberto Paiella is a Professor of Electrical and Computer Engineering at Boston University. He received his B.S. and M.S. degrees in Electrical Engineering from Columbia University, and his Ph.D. degree in Applied Physics from Caltech in 1998. Prior to joining Boston University, he worked at Bell Laboratories, Lucent Technologies, where he developed high-speed mid-infrared quantum cascade lasers for ultrafast pulse generation and optical wireless communications. His current research is focused on the development of novel optoelectronic devices based on semiconductor quantum-confined systems and photonic nanostructures, including work with III-nitride quantum wells, group-IV semiconductor nanomembranes, and plasmonic metasurfaces. He has authored or co-authored more than 80 articles in peer-reviewed journals, has delivered over 20 invited talks, and is the editor of a book entitled Intersubband Transitions in Quantum Structures.

Rep. Prog. Phys. 80 (2017) 106501