1.
Introduction
Ultraviolet (UV)-emitting group III–nitride materials hold a promising potential for a variety of multifunctional applications, including solid-state lighting technology[1-4] and water purification and disinfection[5, 6]. With the wide range of wavelength tunability available to UV-emitting group III–nitride materials, the most promising germicidal ultraviolet devices are found in aluminum gallium nitride and its alloys (AlxGa1?xN, where 0 < x < 1), and one of the most crucial applications of AlxGa1?xN-based devices is water sterilization[7–10] , particularly for highly water-stressed countries. Group III–nitride materials are chemically and thermally robust[11-13], exhibit long carrier lifetimes[14], are operationally stable[15], and are the only known materials that have wide and direct bandgaps and are wavelength-tunable within the UV regime of operation (from around 200 to 400 nm). Other material systems do not exhibit these two properties simultaneously, but achieving p-type conduction is challenging for high aluminum-content group III–nitride compounds[16].
The heterogeneous integration of various forms of inorganic materials (which encompasses growing numbers of material types) into one electronic system is based on group III–nitride compound semiconductors[17-28]. Examples include the following: on-chip frequency upconversion[29], nanomechanical optical detection[30, 31], solid-state neutron detection[32–34], piezoelectric resonators and electrical and harmonic generators[35-40], strain-gated transistors (SGTs)[41], multiple-valued logic (MVL) circuits[42], single-photon emission[43-45], water splitting[46-51], solar-blind photodetection[52], pressure[53], gas[54], pH[55], sensors, white light generation from light-emitting diodes (LEDs)[56-58] and from laser diodes (LDs)[59], metal–oxide–semiconductor field-effect transistors (MOSFETs)[60-62] and high-electron-mobility transistors (HEMTs)[63, 64], mechanically demanding applications[65], and thermoelectrics and thermal management[66-70]. All of these show promise in advancing the development of the new "new electronics" industry[71]. While resonant tunneling transport is essential for the operation of ultrafast electronic oscillators and quantum cascade lasers[72-75], this quantum mechanical effect remained elusive in the family of group III–nitride semiconductors until very recently[76, 77]. Encomendero et al. have engineered high-current resonant tunneling transport in group III–nitride heterostructures, which is employed to generate microwave power[76]. Their GaN/AlN resonant tunneling diodes exhibited microwave oscillations with peak current densities up to 220 kA/cm2. A close look at common compound semiconductor materials reveals that zinc oxide (ZnO), a group II–VI semiconductor, has a wide direct bandgap[78], but it suffers from low p-type conductivity. Nevertheless, ZnO is one of the most promising materials for the development of sensor[79] and translucent optoelectronic devices[80]. These predictions are based on the fact that zinc (Zn) is an earth-abundant material; the exciton binding energy of ZnO are unusually high (about 60 meV), and the bandgap energy is around 3.4 eV[81]. From another point of view, it is well known that an internal electrical field can reduce the effective bandgap energy, a phenomenon referred to as the Franz–Keldysh effect[82, 83]. Bridoux et al. have demonstrated that this effect can be intensified when a tensile strain is generated in ZnO thin films[84]. Since the piezoelectric coefficient of ZnO is very high, very small strains can produce electric fields that reduce the effective energy gap on these films, providing a building block to the new field of piezo-phototronics[85-88].
Successful doping of a semiconductive material is a crucial factor in achieving an efficient carrier injection process to realize excellent device performance characteristics[89-96]. AlxGa1?xN-based UV light-emitting devices can be employed in a variety of applications, including water purification (because of the high absorbance by pathogenic DNA at UV wavelengths)[97], medical diagnostics[98], high-efficiency lighting[99], and chemical/biological detection processes[100]. Remarkable progress has been made to improve the quality and performance of devices based on AlxGa1?xN, and their attractive properties. Such properties include the tunability of their bandgap energies within a significant portion of the UV spectral range (namely, UV-C below 280 nm, UV-B between 280 and 315 nm, and UV-A between 315 and 400 nm), high chemical and device operational stability[101] and reliability[102], internal quantum efficiency (IQE), external quantum efficiency (EQE), and wall plug efficiency (WPE) of III–nitride-based devices, especially deep-ultraviolet (DUV) devices[103]. These properties remain relatively low, and the presence of spontaneous and piezoelectric fields limits their potential[104, 105]. The main causes of such low efficiency parameters are the high density of threading dislocations (TDs) extending from the surface of a strained layer system, which causes internal structural cracking and the subsequent increase in nonradiative recombination channels within the device active regions[90, 106-110]. These issues arise mainly from the lattice and thermal mismatches between the grown material and the substrate[111-113]. For devices emitting at short wavelengths below 280 nm, the EQEs do not exceed 20%[104], and hardly reach 20.3%, by designing a DUV LED utilizing a UV transparent magnesium (Mg)-doped Al0.65Ga0.35N top contact layer, a rhodium (Rh) mirror electrode, an AlN template on a patterned sapphire substrate, and encapsulation resin[114]. In addition, AlxGa1?xN-based heterostructures exhibit poor p-type doping behavior[115, 116], and generally suffer from significant light extraction losses, particularly toward the DUV spectral regime[117], while conductive n-type AlxGa1?xN layers can be realized with relative ease[118, 119]. A recent report by Nippert et al. attributed the efficiency droop in Al0.45Ga0.55N quantum wells (QWs) to internal loss mechanism within the QWs, strikingly similar to the well-known case of InxGa1?xN/GaN QWs[120]. They attributed the reduction in IQE at high excitation power densities to the losses associated with direct and indirect[121] Auger processes, and have extracted an Auger recombination coefficient of C = 2.3 × 10–30 cm6/s using time-resolved photoluminescence (TRPL), which is of the same order as determined by some studies investigating InxGa1?xN/GaN QWs for experimental report of
While it has been well established that electron-beam irradiation[91] and thermal processing[128] can activate dopant impurities in InxGa1?xN and remove magnesium-hydrogen complexes in GaN films[129], this raises the following questions: (1) If an electron-beam can ionize impurities in InxGa1?xN, can one develop an in-situ ion activation system for the growth of AlxGa1?xN[130–133]? (2) During the epitaxial growth process, can one irradiate the sample with UV, X-ray, or, perhaps,
In this article, we review the recent progress in the growth and fabrication of UV and DUV group III–nitride optoelectronic devices and materials based on AlxGa1?xN[135-138] and boron aluminum nitride and its alloys (BxAl1?xN, 0 < x <
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Figure1.
(Color online) Graphical abstract reflecting areas explored in this review article.
| Property | AlNa | Sapphireb | Ga2O3c |
| Thermal conductivity (W/(m·K)) | 180–230 | 42 | [100]: 13.6 |
| [010]: 22.8 | |||
| Thermal expansion coefficient (10–6 K–1) | 4.6d | 7.0f | αa = 0.1–2.78 |
| 5.2e | 7.7g | αb = 1.68–5.84 | |
| αc = 1.74–6.27 | |||
| Specific heat capacity (J/(kg·K)) | 720 | 750 | 490 |
| Melting point (°C) | 2200 | 2053 | 1725 |
| Young's modulus (GPa) | 320 | 470 | 230 |
| Vickers hardness (GPa) | 11 | 22.5 | (101): 9.7 |
| ($bar 201 $): 12.5 | |||
| a As reported by MARUWA CO., LTD, Japan. b As reported by Kyocera Corporation, Japan. c As reported by TAMURA Corporation, Japan. Thermal expansion coefficient values are from Ref. [152] and reported in the temperature range of 24.85–926.85 °C. d In the range of 40–400 °C. e In the range of 40–800 °C. f In the range of 40–400 °C, perpendicular to c-axis. g In the range of 40–400 °C, parallel to c-axis. | |||
Table1.
Comparison between thermal and mechanical properties of AlN, sapphire, and Ga2O3 substrates.
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| Property | AlNa | Sapphireb | Ga2O3c |
| Thermal conductivity (W/(m·K)) | 180–230 | 42 | [100]: 13.6 |
| [010]: 22.8 | |||
| Thermal expansion coefficient (10–6 K–1) | 4.6d | 7.0f | αa = 0.1–2.78 |
| 5.2e | 7.7g | αb = 1.68–5.84 | |
| αc = 1.74–6.27 | |||
| Specific heat capacity (J/(kg·K)) | 720 | 750 | 490 |
| Melting point (°C) | 2200 | 2053 | 1725 |
| Young's modulus (GPa) | 320 | 470 | 230 |
| Vickers hardness (GPa) | 11 | 22.5 | (101): 9.7 |
| ($bar 201 $): 12.5 | |||
| a As reported by MARUWA CO., LTD, Japan. b As reported by Kyocera Corporation, Japan. c As reported by TAMURA Corporation, Japan. Thermal expansion coefficient values are from Ref. [152] and reported in the temperature range of 24.85–926.85 °C. d In the range of 40–400 °C. e In the range of 40–800 °C. f In the range of 40–400 °C, perpendicular to c-axis. g In the range of 40–400 °C, parallel to c-axis. | |||
1.1
Aluminum gallium nitride system
While the doping process and its effects on group III–nitride semiconductors are not fully understood yet, we assert that AlN is the hardest to dope when compared to GaN or indium nitride (InN) and their alloys (InxGa1?xN, 0 < x < 1 )[89]. The primary difficulty is efficient doping of grown films with aluminum-rich clusters; the problem is particularly severe for p-type doping, which is essential for the realization of Ohmic contacts in optoelectronic devices. This can be attributed to the empirical fact that magnesium impurities do not incorporate well into AlN and its high-aluminum-content alloys because of the high energy required to activate dopants. In particular, a high density of compensating point defects can be generated in these materials[153, 154]; such defects include nitrogen vacancies (VN)[155], substitutional magnesium (MgGa)–VN complex and other defect complexes[156-158], and magnesium interstitials[159, 160]. Therefore, high-conductivity p-type aluminum-rich AlxGa1?xN films are difficult to obtain because of the low doping efficiency caused by the high activation energies of the acceptor dopant atoms. Using experimental data obtained by different research groups[161-165], the activation energy of silicon (Si) as an n-type dopant was observed to begin increasing dramatically when the incorporated aluminum content increased in AlxGa1?xN beyond 40%, and reaches its maximum in pure AlN at around 280 meV, as shown in Fig. 2. This translates to an undesirable outcome: When the activation energy of silicon is high, only a few percent of silicon impurities are activated in AlxGa1?xN. While increasing the silicon impurity incorporation percentage seems an option, there is a self-compensate effect when the concentration of silicon is too high, as shown in Fig. 3. High silicon concentrations were observed to compromise the crystal integrity of AlxGa1?xN, increasing its dislocation density and internal aluminum and/or VNs. This situation reciprocates for the p-type doping process, where the activation energy of magnesium (
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Figure2.
(Color online) Comparison of donor activation energies of Si-doped AlxGa1?xN obtained experimentally by various research groups[162–165]. Reprinted with permission from Ref. [161]. ?2017, Elsevier.
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Figure3.
(Color online) Variation in AlxGa1?xN resistivity values as a function of the Si/III ratio[166]. The percentages refer to aluminum mole fraction. Reprinted with permission from Ref. [161]. ?2017, Elsevier.
The most common p-type dopant for AlxGa1?xN is magnesium through the introduction of highly-pure bis(cyclopentadienyl)magnesium (MgCp2) into a metalorganic vapor-phase epitaxy (MOVPE) reactor[169]. While group II elements and transition metals could, in theory, constitute p-type dopants for group III–nitride compound semiconductors, the conductivity of magnesium-doped AlxGa1?xN (specifically for high x values) is relatively low at room temperature due to the high
$ p(T) = N_ {{ m A}}{ m{e}}^{-E^{ m {Mg}}_ {{{ m A}}}/kT}, $  | (1) |
Islam et al. demonstrated tunable DUV LEDs using ultrathin GaN quantum dots (QDs) in the device active regions[193], by manipulating the GaN layer thicknesses with monolayer (ML) precision and applying a polarization-induced doping scheme for both n- and p-type carrier injection regions to enhance the electrical injection of carriers into the active regions[194, 178]. The surface-emitted light from GaN ML quantized structures is predominantly transverse electric (TE)-polarized[195], enhancing light extraction process when compared to AlxGa1?xN-based active regions[196]. We note that Reich et al. employed
By growing quasi-one-dimensional (quasi-1D) group III–nitride nanostructures using plasma-assisted molecular beam epitaxy (PA-MBE)[30, 138, 213-216], dislocation- and piezoelectric polarization-free AlxGa1?xN-based light-emitting nanowire structures can be realized. Nanowires have the advantage of allowing for the growth of lattice-mismatched foreign substrates, such as silicon and sapphire[217-221], in addition to achieving nearly defect-free crystals[222, 223], and therefore, opportunity for fabricating high efficiency optoelectronic devices, including LEDs and lasers, because of the highly effective lateral stress relaxation associated with the nanowire large surface-area-to-volume ratios[224-228]. However, because of the higher surface-area-to-volume ratios in lower-dimen sional semiconductor structures and the subsequent formation of deep-level surface trap energy states, the intrinsic properties of d-dimensional semiconductor structures for all d
1.2
Boron aluminum/gallium nitride system
Previous studies on aluminum-based group III–nitride materials have shown desirable structural and electronic properties with tunable direct bandgap that can cover the entire UV spectral band. These studies highlighted the increase in the optoelectronic active area of devices, where UV light emission could be enhanced fourfold by incorporating boron, but they lack any direct empirical demonstration[235]. While Liu et al. have theoretically investigated the spontaneous polarization and piezoelectric constants of wurtzite BxGa1?xN and BxAl1?xN (0 ≤ x ≤ 1) ternary alloys and concluded that they are not piezoelectric for boron concentrations of 75% and 87%, respectively[236], the miscibility of gallium and aluminum in BN are still under study. However, they are projected to have very low miscibility in BN at normal growth temperatures as the boron miscibility in AlN and GaN is evidently very low as well (a maximum of 14.4% boron content in single-phase wurtzite BxAl1?xN films was reported by Li et al.)[237-242]. For both unstrained ternary BxGa1?xN and BxAl1?xN alloys, Teles et al. used first-principles calculations to observe very high critical temperatures of roughly 9000 and 9500 K that cause broad miscibility gaps between BN and binaries GaN and AlN, which is attributed to the large lattice mismatch between these compounds[240, 241].
Hexagonal boron nitride (h-BN) has attracted a considerable amount of attention because of the unique combination of properties it exhibits[243], including its wide bandgap energy isostructural with graphite[141, 244-248], low relative permittivity[249-251], high thermal conductivity[252, 253], and chemical inertness. It typically exhibits p-type behavior because of acceptor-like vacancy formation[254]; when grown on sapphire substrates, oxygen impurities diffuse from the substrate during high temperature growth, acting as substitutional donors[255]. The presence of these donors causes donor-acceptor pair (DAP) recombination involving carbon deep level acceptors[256, 257]. Despite that and its simple crystal structure, the nature of its bandgap configuration (i.e., direct vs. indirect) remains debatable[258-274]. In 2016, Cassabois, Valvin, and Gil provided evidence that h-BN has 5.955 eV indirect energy bandgap[275]. Through phonon-assisted two-photon absorption measurements, they demonstrated that h-BN exhibits phonon-assisted optical transitions that arise from the observation of a thermal distribution of excitons in the high-energy tail of the different emission lines. More evidence of the indirect nature of h-BN energy bandgap is the observation of different phonon replicas of the free exciton at 5.76 and 5.86 eV, which agree with the calculation results. Although the authors demonstrated the indirect bandgap nature of h-BN, there are many uncertainties that remain to be addressed more deeply and clearly. One of them is the strong emission line at 5.76 eV, at which even stimulated emission has been observed; this seems infeasible if the material has an indirect bandgap. Significantly more investigation need to be conducted to verify the optical transition of h-BN and the mechanism of the phonon-assisted emission if h-BN is an indirect material[275]. In 2017, Laleyan et al. demonstrated that the critical challenges that hindered the development of efficient DUV photonic devices with aluminum-rich AlxGa1?xN can be addressed by magnesium dopant-free AlxGa1?xN/h-BN nanowire heterostructures[276]. They observed enhanced emission from AlN (at around 210 nm [
In 2018, Pierucci et al. demonstrated the controlled growth of h-BN on graphite using MBE via van der Walls epitaxy[279]. Fig. 4 illustrates the obtained film with scanning electron microscopy (SEM) images. They found that this type of heterostructure produces sharp bands, while h-BN and graphite mostly retained their original electronic band structure.
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Figure4.
(Color online) (a) Crystal structure of single layer h-BN. (b) SEM image of h-BN growth on HOPG demonstrating nucleation from HOPG in terrace steps pointed out by the blue arrows (darker contrast areas represent the underlying HOPG substrate while the lighter contrast areas represent regions of h-BN epitaxial growth). (c) High resolution SEM image of an h-BN island displaying areas of single and bi-layer growth and part of exposed HOPG substrate. Reprinted with permission from Ref. [279]. ?2018, AIP Publishing.
2.
Group III–nitride crystal structure and material properties
2.1
Wurtzite crystal structure
Of all groups of III–nitride compound semiconductors, devices based on AlxGa1?xN are ideally suited for UV and DUV device design and fabrication because of the tunability of photoemission and photodetection wavelengths from about 200 to 364 nm by adjusting the AlN mole fraction. Because GaN crystal structures can be assimilated as cubic if the hexagonal planes are slided, the crystal lattice constants of group III–nitride semiconductors can be determined considering that each crystal forms a cubic lattice; hence, the lattice constant can be determined using the following expression:
$ a_{{ m {cubic}}} = sqrt[3]{sqrt{3} a_ {{ m w}}^2 c_ {{ m w}}}, $  | (2) |
where
m w}} $
m w}} $
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Figure5.
(Color online) Plot of bandgap energy versus lattice constant a value of the (Al, In, Ga)N material system[280], α-Ga2O3[281-284], h-BN[275], and diamond, at room temperature.
The wurtzite structure is represented in Fig. 6(a) and is made of two hexagonal close-packed lattices perpendicular to the substrate (c-direction). The lattice constants are a = b
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Figure6.
(Color online) (a) Wurtzite structure. (Reprinted from [https://commons.wikimedia.org/wiki/File:Wurtzite_polyhedra.png]. Image stated to be in the public domain). (b) Wurtzite planes. Reprinted with permission from Ref. [288]. ?2017, IOP Publishing.
Spontaneous polarization in III–V semiconductors has detrimental effects through band bending[289, 290]. It results from the partially ionic bond between nitrogen anions and metal cations, although the chemical bonds in III–V materials are primarily covalent. Polarization vector points toward the [000
| Property | AlN | GaN | InN |
| Lattice constant (?) | a = 3.1127 ± 0.0003 | a = 3.1880 ± 0.0001 | a = 3.53–3.548 |
| c = 4.9816 ± 0.0005 | c = 5.1856 ± 0.0005 | c = 5.69–5.76 | |
| Energy bandgap (eV) | 6.2 | 3.44 | 0.69 |
| Poisson's ratio | 0.287 and 0.216a | 0.37 and 0.33b | 0.14–0.20c |
| Thermal expansion coefficient (10–6 K–1) | $ alpha_parallel $ = 4.2d | $ alpha_parallel $ = 5.59f | $ alpha_parallel $ = 3.1 |
| $ alpha_perp $ = 5.3e | $ alpha_perp $ = 3.17g | – | |
| Thermal conductivity (W/(m·K)) | 319 | 230 | 800 |
| Specific heat (J/(g·K)) | 0.6 | 0.49 | 0.32 |
| Young's modulus (GPa) | C11 = 410 ± 10, C12 = 149 ± 10, C13 = 99 ± 4, C33 = 389 ± 10, C44 = 125 ± 5, C66 = 120 ± 10 | C11 = 390 ± 15, C12 = 145 ± 20 C13 = 106 ± 20, C33 = 398 ± 20 C44 = 105 ± 10, C66 = 123 ± 10 | C11 = 190 ± 7, C12 = 104 ± 3 C13 = 121 ± 7, C33 = 182 ± 6 C44 = 10 ± 1 |
| Index of refraction | 2.1–2.2h | 2.4–2.7i | 2.05–3.06j |
| Melting point (K) | 3100 | 2538l | 1800m |
| a For (0001) and (1210) orientations, respectively. b For (0001) and (1210) orientations, respectively. c In the biaxial strain condition[317]. d In the temperature range of 295–1075 K. e In the temperature range of 295–1075 K. f In the temperature range of 300–700 K. g In the temperature range of 300–700 K. h For $ lambda = $ 300–240 nm. i For $ lambda = $ 600–350 nm. j For $ lambda = $ 1500–600 nm. k Under moderate nitrogen pressure. In vacuum, AlN starts to dissociate at 2200 K. l At atmospheric pressure. Calculated using least-squares fitted result as determined by molecular dynamics simulations. m N2 pressures exceeding 100 kbar are necessary for InN stability[318]. | |||
Table2.
Crystallographic[291–294], mechanical[291, 292, 295–300], thermal[291, 293, 294, 301–314], and optical properties[305, 315, 316] of select wurtzite group III–nitride semiconductors at room temperature.
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| Property | AlN | GaN | InN |
| Lattice constant (?) | a = 3.1127 ± 0.0003 | a = 3.1880 ± 0.0001 | a = 3.53–3.548 |
| c = 4.9816 ± 0.0005 | c = 5.1856 ± 0.0005 | c = 5.69–5.76 | |
| Energy bandgap (eV) | 6.2 | 3.44 | 0.69 |
| Poisson's ratio | 0.287 and 0.216a | 0.37 and 0.33b | 0.14–0.20c |
| Thermal expansion coefficient (10–6 K–1) | $ alpha_parallel $ = 4.2d | $ alpha_parallel $ = 5.59f | $ alpha_parallel $ = 3.1 |
| $ alpha_perp $ = 5.3e | $ alpha_perp $ = 3.17g | – | |
| Thermal conductivity (W/(m·K)) | 319 | 230 | 800 |
| Specific heat (J/(g·K)) | 0.6 | 0.49 | 0.32 |
| Young's modulus (GPa) | C11 = 410 ± 10, C12 = 149 ± 10, C13 = 99 ± 4, C33 = 389 ± 10, C44 = 125 ± 5, C66 = 120 ± 10 | C11 = 390 ± 15, C12 = 145 ± 20 C13 = 106 ± 20, C33 = 398 ± 20 C44 = 105 ± 10, C66 = 123 ± 10 | C11 = 190 ± 7, C12 = 104 ± 3 C13 = 121 ± 7, C33 = 182 ± 6 C44 = 10 ± 1 |
| Index of refraction | 2.1–2.2h | 2.4–2.7i | 2.05–3.06j |
| Melting point (K) | 3100 | 2538l | 1800m |
| a For (0001) and (1210) orientations, respectively. b For (0001) and (1210) orientations, respectively. c In the biaxial strain condition[317]. d In the temperature range of 295–1075 K. e In the temperature range of 295–1075 K. f In the temperature range of 300–700 K. g In the temperature range of 300–700 K. h For $ lambda = $ 300–240 nm. i For $ lambda = $ 600–350 nm. j For $ lambda = $ 1500–600 nm. k Under moderate nitrogen pressure. In vacuum, AlN starts to dissociate at 2200 K. l At atmospheric pressure. Calculated using least-squares fitted result as determined by molecular dynamics simulations. m N2 pressures exceeding 100 kbar are necessary for InN stability[318]. | |||
2.2
Polarization-induced fields
One of the principal causes of low IQE values in group III–nitride-based optoelectronic devices is the significant polarization fields built-in by the non-centro-symmetric nature of the atomic bonds. In a multilayer heterostructure, the spatial gradient of the polarization across the interfaces causes the formation of fixed charges, resulting in band bending and in turn, separation of carriers[319]. Hence, polarization fields are generally undesirable in light-emitting devices, as they hamper the integrity of the electronic and optical properties of the devices, reducing their performances. In the wurtzite crystal structure, spontaneous and piezoelectric polarization fields are caused by the partly ionic bond that induces a small crystal distortion and the presence of strain, respectively. In gallium-polar films the spontaneous polarization (PSP) is parallel to growth direction (positive z-direction), whereas in nitrogen-polar films, PSP is antiparallel to growth direction. Most of the group III–nitride planar crystals are grown gallium-polar, while most of the nanowires are grown nitrogen-polar. However, this depends on the growth conditions, and polarity can be easily interchanged.
Because the piezoelectric polarization field (PPE) is related to the internal crystal strain (
$varepsilon_{xx} = varepsilon_{yy} = frac{a-a_0}{a_0}, quad varepsilon_{zz} = -2frac{C_{13}}{C_{33}}varepsilon_{xx}, $  | (3) |
where a, a0, C13, and C33 are the equilibrium and strained values of the in-plane lattice constant and the elastic constants[320]. The piezoelectric polarization is then expressed as:
$ P_{ m{PE}} = pm 2 varepsilon_{xx}left(varepsilon_{31}-varepsilon_{31}frac{C_{13}}{C_{33}} ight), $  | (4) |
where, depending on the metal or nitrogen polarity, we have the plus or minus sign. For the specific cases of GaN, AlN, and InN in metal-polar condition, the term within the brackets is negative; therefore, PPE is negative (polarization vector points the substrate); while in nitrogen-polar condition it is positive, as summarized in Fig. 7
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Figure7.
(Color online) Polarization field directions and interface polarization charge distribution signs in GaN and AlxGa1?xN structures. Reprinted with permission from Ref. [288]. ?2017, IOP Publishing.
Changes in the normal component of the polarization field, ΔP, cause the formation of fixed charges at the heterointerfaces with density σP = –ΔP and the subsequent carrier separation due to band bending. In the case of MQW structures, the continuity equation of the displacement flux through the different interfaces can be expressed as follows:
$ varepsilon_{ {{ m w}}}F_{ {{ m w}}} + P_{ {{ m w}}} = varepsilon_{ { m b}} F_ { m b} + P_ { m b}, $  | (5) |
where w and b stand for well and barrier, respectively, and ε, F and P are the permittivity of the material, the built-in electric field, and the sum of spontaneous and piezoelectric polarization fields, respectively. Fw and Fb are given by the following:
$ F_ {{ m w}} = frac{(P_ { m b}-P_ {{ m w}})L_ { m b}}{varepsilon_{ m b}L_ {{ m w}} + varepsilon_ {{ m w}}L_ { m b}}, quad F_ { m b} = frac{(P_ {{ m w}}-P_ { m b})L_ {{ m w}}}{varepsilon_{ m b}L_ {{ m w}} + varepsilon_ {{ m w}}L_ { m b}}, $  | (6) |
solved in the limit of an infinitely periodic structure andwhen voltage drop across layers is equal to zero(FwLw+FwLw = 0, L is the layer thickness). Fig. 8 shows the MQW band bending and the misaligned electron–hole wavefunction that imply lower absorption and emission properties. Large field values of few MV/cm are usually present in c-plane III–nitride structures, causing the QCSE to dominate as the gap between conduction band minima and valence band maxima are reduced (red-shifted) by Fw.
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Figure8.
(Color online) Conduction and valence band bending and electron and hole wavefunctions for the case of MQWs. Reprinted with permission from Ref. [288]. ?2017, IOP Publishing.
2.3
Dislocations in AlxGa1?xN
Although InxGa1?xN-based devices demonstrate a remarkable resilience to high dislocation densities[321], dislocations are believed to be a significant factor limiting the efficiency of AlxGa1?xN-based LEDs[322-325]. InxGa1?xN/GaN-based light-emitting devices can be operational even with defect densities of up to 1011 cm–2 because of the effects of carrier localization[95, 321, 326, 327]. In any case, these defects are detrimental for optimal device performance, as they can increase leakage current densities (especially in visible wavelength LEDs)[328, 329], they act as nonradiative recombination centers, they reduce the IQE[330], and they decrease the lifetimes of LDs[331]. There are three types of dislocations in group III–nitrides that can be identified: (1) edge (a-type), mixed ([a+c]-type), and screw (c-type) dislocations[332], and each type of dislocation can be expressed using a Burgers vector (b) as is illustrated in the following[333, 334]:
$ b_{ m{edge} }= frac{1}{3}langle 11bar{2}0 angle, quad b_ { m{mixed}} = frac{1}{3}langle 11bar{2}3 angle, quad b_ { m{screw}} = frac{1}{3}langle 0001 angle. $  | (7) |
The defect density across a structure is directly influenced by epitaxial growth conditions of the materials on designated substrates. Because of its optical and thermal properties as well as its relatively low cost when compared to free-standing GaN substrates, sapphire is commonly employed as a substrate for the growth of group III–nitride light-emitting devices[335-337]. However, large lattice and thermal mismatches hamper their optoelectronic performances. For instance, sapphire and AlN exhibit a 13% lattice mismatch that causes high density of TDs at the interface between a sapphire substrate and an AlN buffer layer[338, 339], propagating toward the device active regions, and also causes lower IQEs[340, 341].
Moreover, due to the higher sticking coefficient and lower surface mobility of aluminum adatoms[342-344], AlxGa1?xN epitaxial growth results are more challenging, compared to the GaN counterpart, having larger adatom mobility. For this reason, the growth of AlxGa1?xN material results in 3D islands as aluminum adatoms are incapable of moving from their point of impact to energy favorable places. As the islands merge with each other, high-density defects such as grain boundaries and dislocations grow (mosaic model)[345], with densities in the range of 1010–1011cm–2 for AlN, much higher compared to GaN (108 cm–2)[346]. Another theory, the nucleation model, that has been supported by transmission electron microscopy (TEM) studies and atomic force microscopy (AFM) studies[347], suggests that the TDs nucleate in the low-temperature nucleation layer extend in subsequent layers. Regarding the techniques used to reduce the TDs in the device active region, we can find superlattices growth[348], insertion layer growth[349], growth of graded layers[350], and temperature- and nitrogen-controlled growth[349, 351].
2.4
Efficiency droop
Efficiency droop is a phenomenon mainly associated with the reduction in InxGa1?xN/GaN-based LED efficiencies as the injection current densities increase. Despite the fact that this phenomenon is very pronounced in blue and green LEDs, it also affects UV devices to a lesser degree[352, 353]. Fig. 9 demonstrates the droop in quantum efficiency with increasing injection current for three LEDs emitting at different wavelength regimes, namely UV, blue, and green LEDs. The decrease in quantum efficiency is caused by the increase in bias voltages and subsequent device heating[354-358], which in turn reduces quantum efficiencies. Such efficiency droops were attributed to carrier delocalization[359, 360], electron leakage attributed to polarization mismatches[357, 361, 362], drift-leakage mechanism[363] tunneling leakage currents[364], poor carrier injection[365, 366], low hole mobility in heavily p-doped aluminum-rich AlxG1?xaN[367], and Auger recombination[122, 123, 368, 369]. However, there is still no conclusive evidence to pinpoint the origin of this problem[370]. For the special case of UV LEDs, the droop is supposedly less prominent, which can be attributed to the significantly reduced Auger losses in wide-bandgap semiconductors because of the substantial reduction in the coupling between electron and hole bands with increasing bandgap. However, the Coulomb interaction that assists in the Auger transitions is significantly stronger in these materials because of the smaller dielectric constants and the subsequent larger exciton binding energies, rendering the importance and strength of this argument inconspicuous[371-373].
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Figure9.
(Color online) Efficiency versus injection current curves of GaN-based UV, blue, and green LEDs, demonstrating a decrease in quantum efficiency with increasing injection current. Green LEDs were shown to have the most prominent efficiency droop. Reprinted with permission from Ref. [352]. ?2013, John Wiley & Sons.
The IQE of a light-emitting device takes into account how efficiently holes and electrons can be extracted from their respective injection layers and the percentage of carriers that recombine radiatively to emit photons, and is expressed as[104]
$ {{ m{IQE}}} = frac{ m{rate of photon generation}}{ m{rate of carrier injection into active region}}. $  | (8) |
The IQE can also be expressed as the ratio of the injected current that leads to radiative recombination in the device active region to the total injected current (Itot) as follows:
$ { m{IQE}} = frac{I_ {{ m r}}}{I_ { m{tot}}} = frac{I_ {{ m r}}}{I_ {{ m r}}+I_ { m{nr}}}, $  | (9) |
where Inr is the current that is dissipated in nonradiative processes and transitions. When Inr dominates (i.e., Inr > Ir), efficiency droop takes place. Inr includes the carrier losses attributed to SRH recombination, Auger recombination, and carrier leakage outside the QWs. Hence, Itot can be expressed as
$ I_ { m{tot}} = I_ {{ m r}} + I_ { m{SRH}} + I_ { m{Auger}} + I_ { m{leakage}}. $  | (10) |
Inside the QWs, the carrier recombination current can be expressed as
$ I_ { m{QW}} = I_ {{ m r}} + I_ { m{SRH}} + I_ { m{Auger}} = qV_{ { m{QW}}}(AN + BN^2 + CN^3), $  | (11) |
where q is the elementary charge value and VQW is the total volume of the QW region. A, B, C are the three recombination coefficients of SRH, radiative, and third-order processes (such as Auger losses), respectively, and N is the excess density of carriers involved in the respective recombination processes. The leakage current can be represented by[366]
$ I_ { m{leakage}} = alpha(I_ { m{QW}})^2. $  | (12) |
By combining Eq. (9) with Eq. (12), we can express the IQE as follows
$ { m{IQE}} = frac{eta_ { m{inj}}+BN^2}{AN + BN^2 + CN^3}, $  | (13) |
where
m{inj}} $
2.5
Thermodynamic photoinduced disorder
In the case of group III–nitride semiconductors, it is broadly assumed that the photocarrier relaxation dynamics consist of a fast initial decay process in the subpicosecond range, followed by a slower decay described by the ultrafast carrier thermalization dynamics, the carrier trapping by surface states, and the slower carrier cooling effects, respectively[377]. This process allows for radiant refrigeration, suggesting that they are solid-state radiative heat pumps[378] as illustrated by David et al. for the case of group III–nitride LEDs[379]. Hence, a more in-depth understanding of the opto-electrothermal properties of group III–nitride semiconductors through an entropic point of view is necessary to further understand their applicability and operational and thermal stability[380] in multifunctional applications[381]. In 2017, we employed temperature-dependent photoluminescence to examine the photoinduced entropy of an ensemble of thick Al0.18Ga0.82N nanowires[382] and In0.32Ga0.68N/GaN p–i–n double-heterostructure nanowire photodiodes[383], shown in Fig. 10, by correlating the energy exchange during the photoexcitation and photoemission processes of the light–solid reaction and the generation of photoinduced entropy of the nanowires using temperature-dependent (6 to 290 K) photoluminescence. We defined the photoinduced entropy as a thermodynamic quantity that represents the unavailability of a system's energy for conversion into work due to luminescence refrigeration[384]. We also studied the ultrafast dynamics of photocarriers using wavelength-integrated time-resolved photoluminescence down to the subnanosecond regime. We investigated the thermodynamic behavior alongside the photocarrier dynamics in GaN-based nanowires because an in-depth understanding of the factors that govern energy transfer processes, and the transport of free carriers through them, is crucial to realizing creative designs of reliable and efficient high-power electronics and LEDs. Strong exciton localization in metal-rich clusters, carrier trapping by surface defect states, and thermodynamic entropy effects were examined and related to the photocarrier dynamics.
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Figure10.
(Color online) (a) Schematic and layer structure of the In0.32Ga0.68N/GaN p–i–n nanowires and (b) the evolutions in the total carrier recombination lifetime and the amount of entropy generation with temperature. (Reprinted with permission from Ref. [383]. ?2017, AIP Publishing). (c) Schematic and layer structure of the Al0.18Ga0.82N nanowires and (d) the integrated and low-energy-peak-related evolution of the amount of entropy generation with temperature. Reprinted with permission from Ref. [382]. ?2017, AIP Publishing.
In In0.32Ga0.68N/GaN nanowires, we observed a rising trend in the amount of generated photoinduced entropy of the system above 250 K, while a fluctuating trend in the generated entropy of the system below 250 K was observed. The fluctuations in the generated entropy stabilized between 200 and 250 K. It was supposed that the amount of generated photoinduced entropy of the In0.32Ga0.68N active region increase as more nonradiative channels became activated, and more shallowly localized carriers settle into deeply localized states (allowed for through the strong localization in indium (In)-rich clusters)[385, 386]; thereby, additional degrees of uncertainty related to the energy of states involved in thermionic transitions were attained. For the Al0.18Ga0.82N nanowires, we observed a fluctuating trend in the generated entropy of the system below 200 K, with a fluctuation frequency that was significantly lower than what we had previously observed in In0.32Ga0.68N. In contrast to the sharp increase in generated entropy at temperatures close to room temperature in In0.32Ga0.68N, an insignificant increase was observed in Al0.18Ga0.82N, indicating lower degrees of disorder-induced uncertainty in the wider bandgap semiconductor. We conjectured that the improved atomic ordering in AlxGa1?xN[387-390] induced lower degrees of disorder-related uncertainty related to the energy of states involved in thermionic transitions; in keeping with this conjecture, we observed lower fluctuation frequency below 200 K.
3.
Innovative material growth mechanisms and device fabrication techniques
Advancements in growth techniques made it possible to achieve highly improved optoelectronic device performance through high-quality single-crystal material growth[391-393]. Since GaN exhibits a smaller degree of lattice mismatch with the lattice of AlN compared with that of sapphire (a = 4.7576 ? and c = 12.9834 ? at 25 °C)[293], the epitaxial growth of GaN films on AlN/sapphire templates can be conveniently carried out[394, 395]. This assisted in the reduction of dislocations on overgrown active regions. In 1983, Yoshida, Misawa, and Gonda achieved significant improvements in the electrical and luminescent properties of reactive MBE-grown GaN/AlN heterostructures on sapphire substrates[396], while in 1989, Akasaki et al. achieved high-crystal-quality and crack-free overgrown GaN epitaxial layers on AlN acting as a low-temperature buffer layer using MOVPE in 1986[397], while in 1989, Akasaki et al. prepared GaN and AlxGa1?xN (0≤ x ≤0.4) films by the preceding deposition of the AlN buffer layer, where they considerably reduced the mosaicity of and microscopic fluctuations in the crystallite orientation[339]. Another milestone step for growing high quality group III–nitrides was accomplishing p-type doping in GaN. In 1989, Amano et al. realized magnesium-doped GaN by low-energy electron-beam irradiation (LEEBI) treatment, and the properties of the GaN p–n junction LED were reported for the first time[91]. The first blue LEDs consisting of p-GaN/n-InxGa1?xN/n-GaN double-heterostructures were fabricated in 1993 by Nakamura, Senoh, and Mukai[398]. Later, in 1996, Nakamura et al. demonstrated the first violet laser emitting at 417 nm based on In0.20Ga0.80N/GaN/AlxGa1?xN heterostructures[94]. This breakthrough demonstrated the use of group III–nitride materials beyond the blue and green LEDs, though achieving p-type doping for AlxGa1?xN-based UV devices remains difficult as magnesium dopant demands more activation energy in group III–nitride semiconductors. For the first time, Asif Khan et al. reported the fabrication and characterization of a high-frequency GaN/Al0.13Ga0.87N-based heterojunction FET in 1994[399].
3.1
Nanostructured layer growth
3.1.1
Molecular beam epitaxy
In 2016, while addressing challenges encountered during growth of aluminum-rich AlxGa1?xN nanowires for DUV optoelectronic devices, Zhao et al. demonstrated that such nanowires with significantly enhanced compositional uniformity can be realized through a new growth prototype other than the conventional nitrogen-rich growth conditions. They argued that they can a achieve precise control on the optical bandgap energy of the ternary AlxGa1?xN nanowires by employing a GaN nanowire template and varying the substrate temperature while improving the aluminum/gallium compositional uniformity. They demonstrated AlxGa1?xN nanowire LEDs, with emission wavelengths spanning from 236 to 280 nm[400]. Their key results are summarized in Fig. 11. The improved growth paradigm revolves around the idea that low nitrogen flow rates would enhance the surface migration of aluminum adatoms, causing a more uniform aluminum/gallium compositional distribution. This resulted in an accurate control on the emission wavelengths of ternary AlxGa1?xN nanowires, which can be accomplished by altering the substrate temperature, instead of the conventional epitaxial way of varying aluminum and gallium beam equivalent pressures (BEPs)[99, 401-403].
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Figure11.
(Color online) (a) Schematic illustration of AlxGa1?xN nanowires grown on GaN nanowire templates (with a low nitrogen flow rate of 0.4 sccm) on a Si substrate. (b) Schematic of direct growth of AlxGa1?xN nanowires on a Si substrate, forming an AlxGa1?xN quasi-film. (c) Elevation-view SEM image of the AlxGa1?xN/GaN nanowires. (d) Elevation-view SEM image of highly coalesced AlxGa1?xN nanowires grown directly on a Si substrate. (e) Low-magnification scanning transmission electron microscope (STEM) image of a single AlxGa1?xN/GaN nanowire and (f) the color-coded (green-bordered) electron energy loss spectroscopy (EELS) maps depicting the elemental distribution of Al and Ga. (g) High-resolution image of the AlxGa1?xN segment in (e), showing the relative uniformity of aluminum distribution, with the thin bright band being the p-GaN contact layer. (h) Low-magnification STEM image of a single AlxGa1?xN/GaN nanowire grown with a high nitrogen flow rate of 1.0 sccm (the inset shows a high-resolution image taken from the AlxGa1?xN segment) manifesting the remarkable atomic-scale compositional modulation. (i) PL spectra of ensembles of nanowires grown under substrate temperatures between 895 and 960 °C with a nitrogen flow rate of 0.4 sccm. (j) Evolution of the peak PL wavelength with AlN mole fraction. (k) PL spectrum (blue curve) of a nanowire sample grown with similar conditions as the sample emitting at 232 nm in (i) but with an increased nitrogen flow rate of 1.3 sccm. The latter is shown in (k), red curve. Reprinted with permission from Ref. [400]. ?2016, AIP Publishing.
Vuong et al. studied the optoelectronic properties of h-BN grown by PA-MBE. By combining AFM, spectroscopic ellipsometry, and photoluminescence spectroscopy in the DUV regime, they compared the quality of h-BN grown on sapphire and highly oriented pyrolytic graphite substrates. They were able to demonstrate DUV emission in h-BN, with an emission spectra peak at 235 nm, indicating the high optical quality of the MBE-grown h-BN. The h-BN epitaxial layers grown on highly oriented pyrolytic graphite have shown superior performance in the DUV regime down to 210 nm, compared to the same films grown on sapphire[404].
Liu et al. have demonstrated large-area AlN nanowall 214 nm LEDs grown on a sapphire substrate. Through temperature-dependent and power-dependent photoluminescence measurements and rate equation analysis, a relatively high internal quantum efficiency of about 60% was determined for the AlN nanowall structures at room temperature. A consistent blueshift in the emission wavelengths was observed with decreasing nanowall widths because of the reduced distribution in tensile strain. The LEDs exhibited excellent current–voltage I–V characteristics, including a turn-on voltage of 7 V and current densities of greater than 170 A/cm2 at 12 V[405].
3.1.2
Metalorganic vapor phase epitaxy
While AlxGa1?xN-based structures show significant importance in applications requiring DUV light sources, the absorption of UV light by p-type layers results in these devices suffering from low light extraction efficiencies (LEEs)[114, 136, 406-410]. To solve this issue, various methods have been proposed by a large body of research[114, 411, 412]. In this context, Jo, Maeda, and Hirayama utilized a transparent p-Al0.70Ga0.30N layer that was synthesized primarily for enhancing the UV LEE, paving a promising path toward zero absorption property of output UV light[413]. The transparent p-Al0.70Ga0.30N layer served as a contact layer and led to a higher LEE by using reflective metal electrodes. Conventionally, blue and UV LED structures incorporate p-type GaN contact layers, but these layers are not suitable for DUV LEDs because GaN layers absorb most of the output UV emission. Higher aluminum mole fraction AlxGa1?xN epitaxial layers with highly transparent and highly conductive aluminum metal layers are crucial to obtain excellent DUV devices, but the fabrication process of realizing high quality p-type AlxGa1?xN layers is difficult due to issues associated with the large activation energy of acceptors, resulting in low hole concentrations and low formation energies of compensation defects. Nevertheless, two 260 nm LEDs were fabricated on sapphire substrates by MOVPE, with one employing a nickel (Ni)/gold (Au) as metal electrode, while the other employed a Ni/aluminum metal electrode[413]. Both LEDs utilize p-type Al0.70Ga0.30N layers with a magnesium concentration of 5 × 1019 cm–3 and have been thermally annealed at 950 °C in N2 ambience. Regardless of the low hole concentration < 1016 cm–3, by using highly transparent (> 95% transmittance) p-type Al0.70Ga0.30N contact layer, the LED demonstrated an efficiency of 2%, similar to that of the LED employing p-type GaN contact layers. Furthermore, an enhancement of 3.4% in EQE was obtained in the LED employing a Ni/aluminum metal electrode compared to that employing a Ni/Au electrode. The aforementioned improved LED performance was attributed to hopping conduction by acceptor sites[184, 414, 415]. To date, numerous efforts have been made to improve the conductivity property of the p-type AlxGa1?xN grown layers by increasing the hole concentration in magnesium-doped AlxGa1?xN and decreasing its acceptor activation energy. In spite of several approaches developed to overcome this obstacle, including magnesium-delta doping[180, 416], co-doping[417], and polarization-induced doping[418], these structures involve wurtzite crystallographic orientations and suffer from high polarization effects, which result in even larger acceptor activation energies. During the AlxGa1?xN growth process, applying nitrogen-rich conditions with large V/III ratios is favorable to achieve better magnesium incorporation and smoother morphology structures.
Metal-source flow-rate modulation epitaxy (FME) provides unique opportunities to realize high hole concentrations in magnesium-doped AlxGa1?xN. Luo et al. were able to demonstrate a hole concentration of about 2.3 × 1017 cm–3 at room temperature, nearly ten times higher than conventional growth procedures, with resistivities as low as 12.7 Ω·cm by implementing metal-source FME[419]. Owing to its smooth step-flow growth, improved surface morphology and crystalline quality were observed as a result of employing FME capability in enhancing the adatoms migration as seen in Figs. 12(a) and 12(b). The working principle of the FME technique is based on the interrupted flow rate of metal while a nitrogen source is continuously suppling, ensuring high nitrogen-rich conditions and keeping maximum V/III flow ratio. A slow growth rate of about 4 nm/min was observed by employing FME interrupted metal flow rate, and hence the Al(Ga) adatoms had sufficient time to migrate on terraces toward the steps. As a result, a highly efficient magnesium doping, with acceptor activation energies ranging between 20 and 22 meV, and a smooth surface morphology, with AFM root mean square (RMS) roughness of about 1 nm, compared to 5.4 nm by conventional growth method, were obtained. This can enhance quantum efficiency of DUV LEDs. The FME principle was also utilized to grow AlxGa1?xN/AlN distributed Bragg reflectors (DBR) with a reflectivity of 97% and a stopband of 6–9 nm at DUV wavelengths ranging between 220 nm and 250 nm[420]. The AlN molar fraction was regulated to obtain the required bandgap wavelength associated with the peak reflective wavelength of the AlxGa1?xN/AlN DBRs. The grown DBRs of optical quarter wavelength of the desired peak wavelength were grown using MOVPE at a temperature of 1130 °C and a pressure of 100 mbar, while assisted through the FME technique by continuously supplying trimethylaluminum (TMAl) and NH3, while frequently interrupting the supply of trimethylgallium (TMGa). The AlN molar fractions in two samples were chosen to be 0.48 and 0.61 by controlling the growth thickness of the respective AlxGa1?xN layer to be 400 nm and by varying the molar flowrate ratio TMAl/III between 0.35 and 0.44, respectively. For the FME-AlxGa1?xN, the AlN molar ration was controlled by tuning TMGa and TMAl period thicknesses and TMAl single period thickness. As a result, the atomic steps of DBR structure was fully strained to the AlN layer and became a defect-free surface and, hence, improved DBR reflectivity was attained. In this regard, we also reported the significant role of yttrium oxide (Y2O3)-doped hafnium oxide (HfO2) on silicon dioxide (YDH/SiO2) DBR having a high reflectivity of 99.9% at a wavelength of 240 nm with a stopband of 50 nm[421]. The YDH/SiO2 DBRs were deposited by RF magnetron sputtering (RF power of 75 W and argon gas flow of 25 sccm) on UV-grade sapphire substrates. The YDH/SiO2 DBR thicknesses were kept at 30 nm with RMS roughness of about 1.2 nm and 1.07 nm for YDH and SiO2, respectively. Because of its higher bandgap energy resulting from doping HfO2 by Y2O3, YDH films are semi-transparent in the UV-C band, rendering them prime candidates for use in DUV devices to enhance optical performance.
onerror="this.onerror=null;this.src='http://www.jos.ac.cn/fileBDTXB/journal/article/jos/2019/12/PIC/18100007-12.jpg'" class="figure_img" id="Figure12"/>
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Figure12.
(Color online) AFM images of the magnesium-doped AlxGa1?xN layers showing enhanced RMS roughnesses of (a) 5.4 nm for x = 0.35 by continuous growth method and (b) 1 nm for x = 0.43 by metal-source FME. Reprinted with permission from Ref. [419]. ?2018, AIP Publishing. XRD characteristics of B-doped AlN/AlN/sapphire, AlN/sapphire, and B-doped AlN/sapphire samples. (c, d) Symmetric 0002 AlN reflections. (e, f) Asymmetric
Instead of employing p-type AlxGa1?xN layers, several efforts have been made to substitute these low conductivity layers with BN[422]. One of the advantages of epitaxially grown BN is its small magnesium acceptor activation energy of approximately 31 meV[423] compared to about 150–300 meV in mixed-phase BN[424-426], and approximately 170–510 meV in AlxGa1?xN[162, 167, 169, 427]. Imura et al. conducted a detailed investigation on boron incorporation on the growth of high-quality AlN layers (Figs. 12(c)–12(g))[154]. They have confirmed a homogeneous boron concentration of 2 × 1021 cm–3 (≈ 2% of the atomic composition) in the MOVPE-grown AlN epitaxial layers. A conventional AlN/sapphire sample exhibited different emissions, namely near-band-edge emission at 208.4 nm and likely defect-related emissions at 315 nm and 400–450 nm originating from complexes of aluminum vacancies ((V
m {Al}}} $
m{Al}}} $
Hakamata et al. investigated the combination of sputtering and MOVPE to study surface quality and optical properties of MOVPE-grown AlN and AlxGa1?xN epitaxial layers on sputtered and annealed AlN/sapphire templates[128]. Sputtering is known for its capability to deliver high quality and large size AlN layers through inexpensive means by adopting annealing at high temperatures. The cause of the high crystallinity in sputtered films is post annealing, which enhances the coalescence of AlN columnar structure and reduces possible dislocations in sputtered layers[432-435]. In their study, an AlN thin film was sputtered on c-plane sapphire substrates and then annealed (SP-AN) at 1700 °C. Then, the MOVPE growth processes of AlN and AlxGa1?xN epitaxial layer on SP-AN AlN/sapphire were detailed. Homogeneous epitaxial layers (homoepilayers) of AlN were grown in two stages. Stage I occurred when growth temperature reached and kept at 1250 °C for three minutes in H2 and NH3 ambience. Stage II occurred when a 1-
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Figure13.
(Color online) (a–c) Surface roughness enhancement trend (from 0.54 to 0.13 nm). In (a) an AFM scan of a sputtered AlN/sapphire template after high-temperature annealing is shown, while in (b), the AlN/sapphire template in (a) scanned after the temperature had been increased to 1250 °C and maintained for three minutes in an NH
Sun et al. have conducted AlN growth experiments by MOVPE to investigate the effects of TMAl pretreatment on the surface roughness and crystalline quality of grown AlN layers on c-plane sapphire substrates[439]. In their study, AFM results revealed that precise pretreatment time is the most important factor to realize TD-free smooth surfaces on AlN films. The pretreatment time of five seconds showed the smoothest surfaces, whereas samples without pretreatment or those that underwent a pretreatment time of greater than ten minutes suffered from TDDs, pits, and rougher surfaces. Mixed nitrogen- and aluminum-polarity exhibited in the AlN films was believed to be responsible for low surface roughness on the untreated sample because of the presence of AlxOyN interlayer between AlN nucleation layer and the sapphire substrate[440, 441]. For 0–5 s of TMAl pretreatment, RMS roughness values of less than 1 nm indicated excellent morphologies because the nitrogen-polarity areas vanished by inhibiting the oxygen diffusion from sapphire substrate[442]. By increasing the pretreatment time beyond 40 s, RMS roughness values increased as nanocolumns were observed at the surface as a result of the formation of 3D islands at the nucleation layer and the presence of 3D growth mode[443]. Finally, the MOVPE growth of single-crystalline AlN-on-sapphire-based ring resonators with high optical Q factors were demonstrated by Bruch et al.[36]. A 1.1
3.2
Device fabrication
In this section, we discuss various methodologies to enhance the performance of DUV devices. Most challenging to realizing growth of group III–nitride semiconductors are the high TDDs resulting from the lattice mismatch and/or thermal expansion mismatch between foreign substrates such as sapphire or silicon and the grown AlN epitaxial layers. On the other hand, the growth process of AlN on native substrates such as AlN or GaN is limited due to the high cost and unavailability of inexpensive commercial substrates. Various approaches have been discussed to address the sapphire substrates' TDDs[137]. Patterned sapphire substrates (PSS) are one of these promising approaches to suppress TDs where nanometer-to-micrometer-sized patterns are transferred to the sapphire substrate by additive or subtractive techniques. The main advantage of the PSS is its ability to minimize TDDs through selective area growth and by improving the photon scattering by uneven sapphire surface. Recently, crack-free DUV LED epitaxial layers were obtained on a four inch c-plane sapphire substrate by epitaxial lateral overgrowth (ELO) on periodic air-voids-incorporated nanoscale patterns as shown in Figs. 14(a)–14(c)[445]. Since the ELO technique demonstrates unique properties in hindering TDD effects and increasing the output power of UV LEDs, it has been widely utilized in micro-patterned substrates. However, it is considered an expensive growth process due to first, higher temperature (above 1300 °C) requirements, which are necessary for the migration of aluminum adatoms, and second, long epitaxial time necessary to realize coalesces thick of 10
onerror="this.onerror=null;this.src='http://www.jos.ac.cn/fileBDTXB/journal/article/jos/2019/12/PIC/18100007-14.jpg'" class="figure_img" id="Figure14"/>
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Figure14.
(Color online) (a) Schematic diagram of silica nanosphere lithography, (b) plan-view, and (c) bird's-eye-view SEM images of the fabricated NPS. (d) Plot of the average LOP as a function of the injection current measured from 50 LEDs at room temperature. The relative EQE, estimated by dividing the photocurrent by the injection current, is also shown in the inset. The OP of the DUV LEDs on NPS shows much higher LOP by 67% than that of the reference DUV LED at the injection current of 20 mA. Reprinted with permission from Ref. [445]. ?2017, AIP Publishing.
Another method to reduce the fabrication cost of the ELO process of AlN is decreasing the coalescence thickness via nanosphere lithography (NSL). In this regard, Dong et al. successfully utilized nanopatterned sapphire substrates (NPSS) that coalesce after only 3 μm of growth, compared to 10 μm for microstructure patterning methods, and hence, reduce fabrication time[323]. They were able to enhance the performance of 282-nm Al0.40Ga0.60N/Al0.50Ga0.50N-based DUV LED by NSL technique, where TDs in the vicinity of nanovoids are suppressed as they bend in the direction of voids sidewalls as depict in Figs. 15(a)–15(c). The EQE was achieved around 3.45% with an enhancement of approximately 98%, compared with conventional at sapphire substrate, by adopting NSL and wet etching mechanisms. In another report, Zhou et al. demonstrated the effect of different sizes of cone-shaped-PSS on performance of UV LED emitting at 375 nm[446]. They noticed that TDDs are in a decreased trend associated with increasing pattern sizes and with the PSS filling factor. The measured light output power (LOP) of the UV LED grown on a large PSS with a fill factor of 0.71 was 131.8% higher than that of the UV LED grown on a small PSS with a fill factor of 0.4. This behavior can be attributed to the reduction of GaN island density and to an increased lateral overgrowth area of GaN, which leads to higher optical performance of the device. The growth process of GaN epitaxial layers on larger PSS including AlN nucleation layer (15 nm of thickness) increases the transition time from 3D island to 2D coalescence, which resulted in a better IQE performance. In a subsequent report, a light-extraction mechanism based on an Rh mirror electrode and a transparent p-type Al0.65Ga0.35N:Mg contact layer was studied by Takano et al.[114]. They achieved a 20% enhancement in EQE at 275 nm of wavelength. In their design, a conventional AlxGa1–xN-based LED is modified into a UV LED structure by utilizing the following features: (1) crack-free AlN on the PSS, (2) Al0.65Ga0.35N:Mg p- type contact layer, (3) mirror electrode, and (4) silicon encapsulation resin. The encapsulation resin helped to reduce optical loss and reached 89% transmission. Implementing Rh mirrors into the LED structure results in an improvement in output power from 3.9 to 18.3 mW at 20 mA, and in the EQE from 4.3% to 20.3%, in comparison with conventional LEDs. Increased multireflection of light in the chip and absorption reduction effectively improved the device performance (Figs. 15(d) and 15(e)).
onerror="this.onerror=null;this.src='http://www.jos.ac.cn/fileBDTXB/journal/article/jos/2019/12/PIC/18100007-15.jpg'" class="figure_img" id="Figure15"/>
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Figure15.
(Color online) (a) Schematic of the fabrication process flow to create nanopatterns on a sapphire substrate. SEM images of the (b) patterned PR and (c) wet-etched NPSS. The inset in (c) shows the line profile of the patterns of NPSS by AFM measurement. (Reprinted with permission from Ref. [323]. ?2013, AIP Publishing). (d) Schematics of conventional (left) and novel UV LED structures (right). In the novel UV LED structure, a transparent p-type Al0.65Ga0.35N:Mg contact layer, a Rh mirror electrode, a PSS, and encapsulation resin are introduced. (e) Output power–current and (f) EQE–current characteristics. Blue and red dots show the characteristics of the conventional and novel UV LED structures, respectively. Inset in (e) shows the EL spectra of the UV LEDs at a DC of 20 mA. (Reprinted with permission from Ref. [114]. ?2017, The Japan Society of Applied Physics).
Fig. 16 shows a moth-eye micro-arrays which was fabricated on the back of sapphire platform to significantly increase the LEE by weakening the total internal reflection (TIR)[447]. Consequently, a high degree of optical polarization (DOP) of 81.8% is obtained for DUV LED utilizing moth-eye micro-arrays, compared to a standard DUV LED of DOP of 64.7%. The high performance of LEE (Fig. 16(f)) is due to the impact of moth-eye micro-arrays on improving the TE/transverse magnetic (TM) modes.
onerror="this.onerror=null;this.src='http://www.jos.ac.cn/fileBDTXB/journal/article/jos/2019/12/PIC/18100007-16.jpg'" class="figure_img" id="Figure16"/>
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Figure16.
(Color online) (a) TE and (b) TM mode light propagation characteristics in moth-eye DUV LEDs reported by Wang et al.. The purple line represents TE/TM mode light emission intensity profile inside the MQWs, whereas the shadow areas correspond to the sapphire surface and sidewall radiation areas, as well as TIR radiation areas. (c) I–V characteristics; (d) normalized spectra; (e) LOP; (f) LEE enhancement factor. The inset depicts the peak wavelengths at different current injection levels. (g) SEM and (h) AFM image of moth-eye microstructure. (i) The structure size along the dash line in the AFM image. Reprinted with permission from Ref. [447]. ?2018, American Chemical Society.
Similar to sapphire substrate integration in the DUV LED fabrication process, growing AlN epitaxial layers on silicon substrate is of critical importance due to the availability, suitability, and low-cost of silicon substrates. Nevertheless, the lattice mismatches between silicon (111) and AlN (0001) are approximately 19% in the
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Figure17.
(Color online) AFM images showing the effect of several growth rates on two AlN layers, where the growth rates are (a) 250 nm/h and (b) 50 nm/h. Sides-sectional STEM images depicting two 420 nm-thick AlN/Si(110) samples (c) with and (d) without V-shape pits by controlling the growth rate. Reprinted with permission from Ref. [448]. ?2017, Royal Society of Chemistry.
Micro-circle-patterned silicon substrate (mPSIS) is another configuration technique to enhance the quality of grown AlN template on silicon substrate[449]. Different mPSIS sizes were fabricated to grow thick AlN template with the help of NH3 pulsed-flow and ELO in the MOVPE growth process. During the lithography process, the ICP bias powers were modified to obtain different mPSIS sizes. MPSIS pattern sizes of 1.7 and 2.0
4.
Critical review of the current status of nitride-based optoelectronic devices
4.1
Light-emitting diodes
AlxGa1?xN-based DUV LEDs have been dominating research in the last few years[450-462]. Most notably, using the growth paradigm described earlier, Zhao et al. were able to fabricate defect-emission-free AlxGa1?xN nanowire LEDs operating in the UV-C regime, and attributed the enhanced current conduction characteristics to the improved acceptor magnesium dopant incorporation in the AlxGa1?xN nanowires and the subsequent acceptor impurity band conduction[184, 400, 463]. Fig. 18(a) shows the device schematic, while the I–V characteristics of an AlxGa1?xN nanowire LED with emission wavelength around 240 nm are depicted in Fig. 18(b). Fig. 18(c) depicts the room-temperature EL spectra measured from several devices. The output power measurements (L–I–V) of a UV LED device are shown in Fig. 18(d).
onerror="this.onerror=null;this.src='http://www.jos.ac.cn/fileBDTXB/journal/article/jos/2019/12/PIC/18100007-18.jpg'" class="figure_img" id="Figure18"/>
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Figure18.
(Color online) (a) Schematic illustration of the p–i–n AlxGa1?xN nanowire-based LEDs grown on titanium (Ti)/Au on Si by Zhao et al. (b) Linear and semilogarithmic scale I–V characteristics curve of an AlxGa1?xN nanowire UV LED emitting around 240 nm. (c) EL spectra measured from several nanowire LEDs with different AlN mole fractions under an injection current of 20 mA. (d) L–I–V and relative EQE curves of a UV device emitting around 240 nm. Reprinted with permission from Ref. [400]. ?2016, AIP Publishing.
4.2
Photodetectors
Taking advantage of high absorption coefficients (>105 cm–2) and wide bandgap energy, group III–nitride semiconductor alloys have attracted remarkable attention for UV photon detection owing to their prospective applications in solar radiation monitoring[464], missile plume detection[465-467], biological treatment[468], and non-line-of-sight (NLOS) communication[469]. Compared to other wide bandgap materials (e.g., ZnO, SiC, and Ga2O3), the large bandgap tunability of group III–nitrides also allows the absorption or cut-off wavelength to be engineered by simply changing the alloy composition, an easy tailoring for demanding applications that require high signal-to-noise ratios (SNR) and spectrum selectivity. Apart from that, high thermal and chemical stabilities of group III–nitride material system stands out in harsh environments, compared to conventional silicon-based detectors, which have limited operating temperature range and require extra protective packaging[470-473]. In addition, silicon-based photodetectors typically have low photoresponsivity in the UV regime of operation due to low penetration depth of UV-wavelength photons in silicon[474]. Photomultiplier tubes (PMTs) were also suggested as alternative devices in UV photon detection[475]; however, the associated bulky volume and high operating voltage hinders further applications of PMTs[476]. Following the development of a more mature epitaxial growth for group III–nitride compounds, AlxGa1?xN-based photodetectors (PDs), which can cover the wavelength range from 200 to 360 nm depending on the aluminum composition[477], show promise for achieving highly reliable UV photodetectors and in alleviating the aforementioned issues, and are the subject of active research in UV and DUV photon detection, opening up a plethora of potential applications in the future.
The early work on GaN-based PDs started in in 1992 by Khan et al., where a photoconductive UV detector with high spectral responsivity (R) of 2000 A/W (200 to 365 nm) was demonstrated[478]. This is the first reported UV sensor based on as-deposited single-crystal insulating GaN films. Since then, various types of AlxGa1?xN-based PDs, particularly for photon detection in the solar-blind region (i.e., 230 to 280 nm), such as Schottky-type PDs[479-481], metal–semiconductor–metal (MSM) PDs[482-486], HEMTs[487], and p–i–n heterojunction PDs[488-491], have been demonstrated. Among these, MSM-based PDs attract noteworthy attention, owing to their ease of fabrication, low stray capacitances, and high switching speed for monolithic integration on photonic circuits[480, 492, 493]. Another notable advantage of MSM-based PDs is that p-type doping of AlxGa1?xN is not necessary, putting it in dominance when compared to p–i–n structured PDs. In addition, depending on the device design, MSM-based PDs can be illuminated through either front or bottom illumination. For example, Wang et al. compared performance differences in terms of responsivity and EQE of Al0.40Ga0.60N-based MSM UV PD, in both illumination modes[494]. For a photodetector, the EQE can be determined using
$ {{ m {EQE}}} = frac{I_{ { m PC}}}{P_{ { m {IL}}}} frac{hc}{qlambda} = {R} frac{hc}{qlambda}, $  | (14) |
where IPC is the generated photocurrent, PIL is the illuminated optical power, h is Planck's constant, c is speed of light, and
ho_ {
m {screw}} $
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Figure19.
(Color online) (a) Transmission spectrum of Ni/Au interdigitated electrodes used in an MSM-based PD. Inset shows the typical structure of an AlxGa1?xN-based MSM PD. (b) Corresponding responsivity spectrum of a bottom-illuminated Al0.40Ga0.60N-based MSM PD. Inset shows the comparison of measured responsivity for front and bottom-illumination mode. (Reprinted with permission from Ref. [494]. ? 2013, AIP Publishing). Comparison between n-AlxGa1?xN nanoflower, nanowire, and nanorod PDs in terms of (c) photocurrent and (d) photoresponsivity spectra. (Reprinted with permission from Ref. [508]. ?2018, Royal Society of Chemistry).
Despite the advantages of bottom-illuminated configuration PDs, the thermal and lattice mismatches between conventional sapphire substrates and group III–nitride materials still persist. While paving the way toward the solar-blind regime of operation, high aluminum composition of greater than 40% are required to provide the aforementioned applications in the solar-blind region. High TDDs and thermal structural cracking developed during the epitaxial growth, especially at increased aluminum contents, hinder the device performance of AlxGa1?xN-based PDs, which requires low charge generation rates and stray capacitances. In order to alleviate these issues, various methods have been reported, including selective area growth[509], pulse atomic epitaxial technique (PALE)[510], and crack-free AlN template layer use. Cicek et al. demonstrated a solar-blind Al0.40Ga0.60N-based p–i–n PD with a high EQE of up to 89% (at 5 V reverse bias voltage) and a responsivity of around 176 mA/W (at zero bias voltage) by using various optimized growth and structural design methods[511]. Improvements in structural characteristics were attributed to the use of a crack-free AlN template layer, a highly conductive silicon-indium co-doped[512] Al0.50Ga0.50N, and an improved magnesium-doped Al0.38Ga0.62N layer grown using MOVPE. In a recent work, Han et al. proposed the use of metal–insulator–semiconductor (MIS) structures in solar-blind Al0.45Ga0.55N-based PDs, as a competitive alternative to the conventional p–i–n structure[513]. The latter is impeded by growth of high-aluminum-content p-doped AlxGa1?xN layers. The demonstrated bottom-illuminated MIS-based structure exhibited a high EQE of 70.6% (at 3 V reverse bias voltage) at 270 nm with a relatively fast rise time of 1.4 ns. The significance of optimized AlN template in growing high quality AlxGa1?xN PD structures has been substantially highlighted. Chen et al. reported a solar-blind AlxGa1?xN-based p–i–n PD grown and fabricated on a high quality and low residual compressively stressed AlN template[514]. They attributed the improved performance of the device (EQE at 68.8% and response speed of 6.5 ns) to the insertion of a mesothermal AlN (MT-AlN) interlayer to assist in strain relief along the growth direction. Similar work on optimized growth conditions for producing high quality and crack-free AlxGa1?xN layers have also been reported. High specific detectivity (
m Hz}} $
Other than planar structures, AlxGa1?xN-based PDs constructed using lower-dimensional structures, such as nanowires and nanorods, have also received significant attention[518]. In 2017, Kang et al. reported on a two-step growth process for single-crystalline silicon-doped Al0.45Ga0.55N nanorods as a solar-blind PD[519]. The vertically-aligned and compositionally uniform nanorod structure displayed high responsivity of up to about 115 mA/W and high sensitivity of approximately 64% in the UV-C region[519]. Later, the same group demonstrated improved photoresponsivity of up to 0.72 A/W, fabricated using n-type Al0.45Ga0.55N nanowires with flower-like morphology, also called as nanoflowers[508]. Figs. 19(c) and 19(d) show improved photocurrent by close to one order of magnitude and high responsivity of up 0.72 A/W in Al0.45Ga0.55N nanoflower-structured PD, compared to that of nanowire-structured PD (0.13 A/W) and nanorods (0.09 A/W). The improved performance of nanostructured PDs was largely attributed to improved photon absorption due to high aspect ratios and compact density distribution[138]. These works paved the way for high performance low-dimensional AlxGa1?xN-based PDs in various applications[520]. Furthermore, group III–nitride-based avalanche photodetectors (APD), which show promise in providing higher photocurrent gains combined with high speeds and low operating voltages, are also of contemporary interest in the community[521-529]. Moreover, the use of Geiger-mode operation in group III–nitride-based APDs has also received significant attention for its potential employment in high sensitivity single-photon counting for various applications, such as quantum computing and astrophysics[530, 531]. Compared to UV-enhanced silicon-based APDs, AlxGa1?xN-based APDs promise high wavelength tunability and selectivity. Low leakage currents and high avalanche gains (5 × 105) in an Al0.05Ga0.95N APDs grown on free-standing GaN substrates were demonstrated by Kim et al. in 2015,[532] highlighting a significant improvement in APD performance, compared to other APD structures grown on sapphire templates. Later in 2016, Wu et al. reported a true solar-blind Al0.40Ga0.60N-based APD structure with zero-bias EQE of up to 52.7% with an avalanche gain of higher than 2 × 104 at a relatively high voltage bias of
onerror="this.onerror=null;this.src='http://www.jos.ac.cn/fileBDTXB/journal/article/jos/2019/12/PIC/18100007-20.jpg'" class="figure_img" id="Figure20"/>
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Figure20.
(Color online) (a) Device structure of Al0.40Ga0.60N-based p–i–n–i–n separate absorption and multiplication (SAM) APD. (b) I–V curves and multiplication gain achieved using Al0.40Ga0.60N-based APD. (Reprinted with permission from Ref. [522]. ?2014, IEEE). (c) Responsivity spectrum of an h-BN-based MSM PD with peak responsivity of 0.1 mA/W at 212 nm and a clear cut-off wavelength at around 225 nm. Inset shows the schematic illustrations of the device structure, where h-BN layers were transferred onto SiO
However, to tailor PDs for the DUV wavelength regime of operation, high quality and crack-free high-aluminum-content layers are pivotal. Up until now, no solution has been found to address high defect densities associated with the incorporation of AlxGa1?xN layers with high aluminum contents. Walker et al. was one of the earliest to demonstrate AlxGa1?xN-based PDs operating in DUV wavelengths with a responsivity of 0.11 A/W at 232 nm (5 V reverse voltage) and an IQE of 90%[536]. In their device structure, high-aluminum-content AlxGa1?xN layers of up to x = 70% were successfully grown using low-pressure chemical vapor deposition (LPCVD) on sapphire substrates. Subsequently, in 2007, G?kkavas et al. demonstrated Al0.75Ga0.25N-based with relatively high responsivity of up to 0.53 A/W at 222 nm[537]. The high responsivity obtained in this work can be attributed to the high-quality Al0.75Ga0.25N layer, demonstrating a low dark current of less than 100 pA at very high bias voltages of up to 350 V. Alternative nitride-based materials, such as AlN and BN, have also been widely discussed[477, 538-542]. AlN-based photodetectors received significant attention in the early 2000s because of their high bandgap energies (around 6.1 eV) and high SNRs with the suppression of visible background for DUV detection. The demonstration of AlN-based MSM PD for operation below 230 nm was reported by Li et al.[477]. A sharp cut-off wavelength was observed at 207 nm with peak responsivity of 0.4 A/W (at 100 V voltage bias) at 200 nm, with a rejection ratio of four orders of magnitude. The insertion of superlattice structures after the AlN nucleation layer have also been introduced by Nikishin et al. to reduce the inversion domains (IDs) in AlN-based PDs[543]. The fabricated large-area MSM PD exhibited a low dark current of 50 fA up until 30 V voltage bias with the insertion of AlN/GaN short-period superlattices (SPSLs), with a peak responsivity of 0.08 A/W at 202 nm. Following these discoveries, large area (about 3.1 mm2) DUV PDs based on AlN films were demonstrated by Barkad et al. with titanium nitride (TiN) as a Schottky contact[544]. The cut-off wavelength was reported to be around 203 nm, the shortest ever cut-off wavelength reported in AlN-based PDs, with a rejection ratio of up to three orders of magnitude. Other than the above, the first surface acoustic wave (SAW)-based DUV sensor based on AlN films for photodetection below 200 nm has also been reported[545]. In recent years, BN with similar bandgap energy (around 6 eV) as AlN have also received increasing attention for DUV applications. In 2018, Soltani et al. demonstrated DUV PD with cut-off wavelength below 200 nm using cubic BN (c-BN)[546]. The MSM-based PD demonstrated a very short cut-off wavelength at 193 nm because of the wide bandgap of c-BN (about 6.4 eV), while the peak responsivity appears at 32 mA/W at 180 nm with 35 V voltage bias. Although the achieved responsivity is still considered relatively low, the cut-off wavelength was the shortest at that time, and a high rejection ratio of up to four orders of magnitude was also recorded. Later in 2012, Li et al. explored h-BN epitaxial layers as UV photodetectors with sharp cut-off wavelengths at around 230 nm with high absorption coefficients (approximately 7 × 105 cm–1) than that of AlN (around 2 × 105 cm–1)[547]. In addition to their sufficient performance as thermal interface materials[548], optically transparent (down to 210 nm) BN nanosheets (BNNSs) have also been examined for UV detection below 300 nm[549]. This study showed that the BNNSs-based DUV PDs were blind to photons with energies of less than 4 eV, with a PD output power of around 2
As discussed earlier, further improvements in the performance of group III–nitride-based PDs were largely inhibited by the low crystal quality and low conductivity characteristics originating from the growth of PD structure materials. To alleviate these issues, hybridization of GaN-based PDs with 2D materials, such as graphene, have been proposed. Phototransistors based on graphene/GaN structures with relatively high responsivity of 0.36 A/W at 325 nm were reported by Tian et al. and shown in Fig. 21(a). Their reported device characteristic values were significantly higher than those realized without the integration of graphene layers (up to 700-fold improvement); the high achieved gain can be attributed to high carrier mobility effects exhibited in graphene, which allow for multiple carrier transport to generate high photocurrent densities[552]. Without the graphene layer, it can be expected that generated carriers would be trapped in the lower-mobility GaN layer, limiting the achievable photocurrents and gains in PD devices. Similar work was later published by Tian et al. They demonstrated a significantly enhanced responsivity of up to 5.83 A/W with illumination at 325 nm and an EQE of up to 2200%, which they accredited to multiple carrier circulation and long carrier lifetimes in the graphene layer using a similar hybrid phototransistor structure[553]. However, it was also emphasized that responsivity considerably decreased with increasing power of incident light, as shown in Fig. 21(b), because of dampening of the net built-in electric fields by opposite electric fields formed by accumulated photogenerated carriers. Nevertheless, both reports provide insight into improving the performance of UV phototransistors without stringent requirements on the crystal quality of light-absorbing layers. Notably, a similar strategy was employed in DUV PD devices to improve photoconductive gains, with graphene layers achieving transparencies greater than 80% within the DUV regime[554]. The incorporation of BNNSs and 1D cellulose nanofibers has also been demonstrated as flexible solar-blind PDs by Lin et al. Although their demonstrated paper-based device exhibited a relatively low responsivity of 0.05 mA/W at 185 nm illumination, the device displayed near-zero responsivity between UV-B (300 nm) up to NIR (1000 nm) wavelengths, manifesting its true solar-blind detection capabilities[52]. Furthermore, taking advantage of the high thermal conductivity of h-BN (about 146 W/mK), the paper-based device was capable of functioning at high temperatures of up to 200 °C, which represents a significant step forward in reaching the goal of flexible and wearable electronic devices. Table 3 provides a summary of key performance characteristics of state-of-the-art group III–nitride-based and hybrid UV and DUV PDs.
| Device structure | Material system | Bias (V) | Peak responsivity (A/W) | EQE | Rise/Fall time (s) | Year | Ref. |
| MSM | (Al0.28Ga0.72)2O3 | 2.5 | 1.17 (230 nm) | 631% | – | 2019 | [551] |
| MSM | $ beta $-Ga2O3/TiN | 15 | 277.08 (250 nm) | 1.37 × 105% | – | 2019 | [555] |
| MSM | $ beta $-Ga2O3 | 10 | 8.41 (254 nm) | 4.11 × 103%* | 18 s/1.6 s | 2019 | [556] |
| MSM | $ beta $-Ga2O3 | 20 | 150 (254 nm) | 7.40 × 104% | 1.8 s/0.3 s | 2018 | [557] |
| MSM | Al0.50Ga0.50N | 20 | 0.094 (270 nm) | 43.2%* | <20 ms/ <20 ms | 2018 | [503] |
| MSM (flexible) | h-BN | 10 | 0.00005 (185 nm) | 0.03%* | 0.267 s/0.393 s | 2018 | [52] |
| MSM | h-BN | 20 | 0.0001 (212 nm) | 0.06%* | 0.32 s/0.63 s | 2018 | [550] |
| MSM | Al0.50Ga0.50N | 10 | 5 (264 nm) | 2.35 × 103%* | 0.52 s/44 s | 2017 | [558] |
| MSM | Al0.50Ga0.50N | 5 | 106 (250 nm) | 4.96 × 108%* | 0.2 s/1000 s | 2017 | [507] |
| MSM (SiO2-passivated) | Al0.25Ga0.75N/GaN | 10 | 0.27 (365 nm) | 91.78%* | – | 2015 | [504] |
| MSM | Al0.25Ga0.75N | 5 | <0.01 (310 nm) | <4% * | – | 2013 | [502] |
| MSM | Al0.42Ga0.60N | 20 | 0.045 (270 nm) | 21.5% | – | 2013 | [494] |
| 0.085 (270 nm) | 39% | – | |||||
| MSM | TiN/AlN | 100 | 0.002 (200 nm) | 1.24%* | – | 2010 | [544] |
| MSM | AlN | 10 | 0.08 (202 nm) | 49.14%* | – | 2009 | [543] |
| MSM | c-BN | 35 | 0.032 (180 nm) | 22.06%* | – | 2008 | [546] |
| MSM | Al0.75Ga0.25N | 50 | 0.53 (222 nm) | > 250% | – | 2007 | [537] |
| p–i–n | Al0.46Ga0.54N | 0 | 0.15 (271 nm) | 68.8% | 6.5 ns/0.95 $ mu $s | 2018 | [514] |
| p–i–n | Al0.34Ga0.66N | 0.5 | 0.1 (250 nm) | 50% | 0.4 $ mu $s/0.4 $ mu $s | 2017 | [517] |
| p–i–n | AlxGa1–xN | 0 | 0.21 (360 nm) | 70% | 1.7 ns/4.5 ns | 2014 | [515] |
| p–i–n | Al0.40Ga0.60N | 5 | 0.192 (275 nm) | 89% | – | 2013 | [511] |
| p–i–n | Al0.40Ga0.60N | 5 | 0.129 (279 nm) | 57% | – | 2013 | [509] |
| p–i–i | Al0.55Ga0.45N | 5 | 0.075 (254 nm) | 37% | – | 2012 | [510] |
| p–i–n | Al0.70Ga0.30N | 5 | 0.11 (232 nm) | 58.83%* | – | 2000 | [536] |
| p–i–n–i–n | Al0.40Ga0.60N | 15 | 0.15 (280 nm) | 50% | – | 2014 | [522] |
| p–i–n APD | Al0.40Ga0.60N | 0 | 0.114 (278 nm) | 52.7% | – | 2016 | [533] |
| p–i–n APD | Al0.05Ga0.95N | 0 | 0.043 (354 nm) | 16% | – | 2015 | [532] |
| MIS | Al0.45Ga0.55N | 3 | 0.154 (270 nm) | 70.6% | 1.4 ns/55 $ mu $s | 2018 | [513] |
| MIM | (Al0.28Ga0.72)2O3 | 2.5 | 0.4 (230 nm) | 218% | – | 2019 | [551] |
| p–n | PtSe2/GaN | 0 | 0.193 (265 nm) | 90.36%* | 0.172 $ mu $s/284 $ mu $s | 2019 | [559] |
| p–n | MoS2/GaN | 0 | 0.187 (265 nm) | 87.55%* | 46 $ mu $s/114 $ mu $s | 2018 | [560] |
| Heterojunction | graphene/Ga2O3 nanowires | 5 | 0.185 (258 nm) | 88.96%* | 9 ms/8 ms | 2018 | [561] |
| Heterojunction | graphene/$ beta $-Ga2O3 | 6 | 12.8 (254 nm) | 6.68 × 103% | 1.5 ms/2 ms | 2018 | [562] |
| Heterojunction | graphene/h-BN/GaN | 1 | 1915 (245 nm) | 9.69 × 105%* | 8 s/6 s | 2016 | [563] |
| Schottky | graphene/$ beta $-Ga2O3/graphene | 10 | 9.66 (254 nm) | 4.72 × 103%* | 0.96 s/0.81 s | 2017 | [564] |
| Schottky | graphene/Al0.25Ga0.75N/GaN | 2 | 0.56 (300 nm) | 231.60%* | – | 2016 | [565] |
| Phototransistor | graphene/GaN | 10 | 0.361 (325 nm) | 138% | 3.2 ms/1.2 ms | 2018 | [552] |
| Phototransistor | graphene/GaN | 10 | 5.83 (325 nm) | > 2.20 × 103% | 2.7 ms/4.6 ms | 2018 | [553] |
| Nanorods | Al0.45Ga0.55N | 3 | 0.115 (250–276 nm) | 53.84%* | – | 2017 | [519] |
| Nanowires (Schottky) | GaN | 1 | 25 (357 nm) | 8.69 × 103%* | – | 2013 | [518] |
| Nanoflowers (MSM) | Al0.45Ga0.55N | 2 | 0.72 (265 nm) | 340% | – | 2018 | [508] |
Table3.
Performance comparison of select group III–nitride-based and hybrid UV and DUV PDs at room temperature (* Indicates an EQE value we calculated using Eq. (14)).
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| Device structure | Material system | Bias (V) | Peak responsivity (A/W) | EQE | Rise/Fall time (s) | Year | Ref. |
| MSM | (Al0.28Ga0.72)2O3 | 2.5 | 1.17 (230 nm) | 631% | – | 2019 | [551] |
| MSM | $ beta $-Ga2O3/TiN | 15 | 277.08 (250 nm) | 1.37 × 105% | – | 2019 | [555] |
| MSM | $ beta $-Ga2O3 | 10 | 8.41 (254 nm) | 4.11 × 103%* | 18 s/1.6 s | 2019 | [556] |
| MSM | $ beta $-Ga2O3 | 20 | 150 (254 nm) | 7.40 × 104% | 1.8 s/0.3 s | 2018 | [557] |
| MSM | Al0.50Ga0.50N | 20 | 0.094 (270 nm) | 43.2%* | <20 ms/ <20 ms | 2018 | [503] |
| MSM (flexible) | h-BN | 10 | 0.00005 (185 nm) | 0.03%* | 0.267 s/0.393 s | 2018 | [52] |
| MSM | h-BN | 20 | 0.0001 (212 nm) | 0.06%* | 0.32 s/0.63 s | 2018 | [550] |
| MSM | Al0.50Ga0.50N | 10 | 5 (264 nm) | 2.35 × 103%* | 0.52 s/44 s | 2017 | [558] |
| MSM | Al0.50Ga0.50N | 5 | 106 (250 nm) | 4.96 × 108%* | 0.2 s/1000 s | 2017 | [507] |
| MSM (SiO2-passivated) | Al0.25Ga0.75N/GaN | 10 | 0.27 (365 nm) | 91.78%* | – | 2015 | [504] |
| MSM | Al0.25Ga0.75N | 5 | <0.01 (310 nm) | <4% * | – | 2013 | [502] |
| MSM | Al0.42Ga0.60N | 20 | 0.045 (270 nm) | 21.5% | – | 2013 | [494] |
| 0.085 (270 nm) | 39% | – | |||||
| MSM | TiN/AlN | 100 | 0.002 (200 nm) | 1.24%* | – | 2010 | [544] |
| MSM | AlN | 10 | 0.08 (202 nm) | 49.14%* | – | 2009 | [543] |
| MSM | c-BN | 35 | 0.032 (180 nm) | 22.06%* | – | 2008 | [546] |
| MSM | Al0.75Ga0.25N | 50 | 0.53 (222 nm) | > 250% | – | 2007 | [537] |
| p–i–n | Al0.46Ga0.54N | 0 | 0.15 (271 nm) | 68.8% | 6.5 ns/0.95 $ mu $s | 2018 | [514] |
| p–i–n | Al0.34Ga0.66N | 0.5 | 0.1 (250 nm) | 50% | 0.4 $ mu $s/0.4 $ mu $s | 2017 | [517] |
| p–i–n | AlxGa1–xN | 0 | 0.21 (360 nm) | 70% | 1.7 ns/4.5 ns | 2014 | [515] |
| p–i–n | Al0.40Ga0.60N | 5 | 0.192 (275 nm) | 89% | – | 2013 | [511] |
| p–i–n | Al0.40Ga0.60N | 5 | 0.129 (279 nm) | 57% | – | 2013 | [509] |
| p–i–i | Al0.55Ga0.45N | 5 | 0.075 (254 nm) | 37% | – | 2012 | [510] |
| p–i–n | Al0.70Ga0.30N | 5 | 0.11 (232 nm) | 58.83%* | – | 2000 | [536] |
| p–i–n–i–n | Al0.40Ga0.60N | 15 | 0.15 (280 nm) | 50% | – | 2014 | [522] |
| p–i–n APD | Al0.40Ga0.60N | 0 | 0.114 (278 nm) | 52.7% | – | 2016 | [533] |
| p–i–n APD | Al0.05Ga0.95N | 0 | 0.043 (354 nm) | 16% | – | 2015 | [532] |
| MIS | Al0.45Ga0.55N | 3 | 0.154 (270 nm) | 70.6% | 1.4 ns/55 $ mu $s | 2018 | [513] |
| MIM | (Al0.28Ga0.72)2O3 | 2.5 | 0.4 (230 nm) | 218% | – | 2019 | [551] |
| p–n | PtSe2/GaN | 0 | 0.193 (265 nm) | 90.36%* | 0.172 $ mu $s/284 $ mu $s | 2019 | [559] |
| p–n | MoS2/GaN | 0 | 0.187 (265 nm) | 87.55%* | 46 $ mu $s/114 $ mu $s | 2018 | [560] |
| Heterojunction | graphene/Ga2O3 nanowires | 5 | 0.185 (258 nm) | 88.96%* | 9 ms/8 ms | 2018 | [561] |
| Heterojunction | graphene/$ beta $-Ga2O3 | 6 | 12.8 (254 nm) | 6.68 × 103% | 1.5 ms/2 ms | 2018 | [562] |
| Heterojunction | graphene/h-BN/GaN | 1 | 1915 (245 nm) | 9.69 × 105%* | 8 s/6 s | 2016 | [563] |
| Schottky | graphene/$ beta $-Ga2O3/graphene | 10 | 9.66 (254 nm) | 4.72 × 103%* | 0.96 s/0.81 s | 2017 | [564] |
| Schottky | graphene/Al0.25Ga0.75N/GaN | 2 | 0.56 (300 nm) | 231.60%* | – | 2016 | [565] |
| Phototransistor | graphene/GaN | 10 | 0.361 (325 nm) | 138% | 3.2 ms/1.2 ms | 2018 | [552] |
| Phototransistor | graphene/GaN | 10 | 5.83 (325 nm) | > 2.20 × 103% | 2.7 ms/4.6 ms | 2018 | [553] |
| Nanorods | Al0.45Ga0.55N | 3 | 0.115 (250–276 nm) | 53.84%* | – | 2017 | [519] |
| Nanowires (Schottky) | GaN | 1 | 25 (357 nm) | 8.69 × 103%* | – | 2013 | [518] |
| Nanoflowers (MSM) | Al0.45Ga0.55N | 2 | 0.72 (265 nm) | 340% | – | 2018 | [508] |
onerror="this.onerror=null;this.src='http://www.jos.ac.cn/fileBDTXB/journal/article/jos/2019/12/PIC/18100007-21.jpg'" class="figure_img" id="Figure21"/>
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Figure21.
(Color online) (a) Schematic of the device structure and bandgap energy diagram of hybrid graphene/GaN UV PD. (Reprinted with permission from Ref. [552]. ?2018, The Optical Society). (b) Photocurrent and responsivity curves of a hybrid graphene/GaN UV-A PD at 10 V reverse bias (with and without graphene). (Reprinted with permission from Ref. [553]. ?2018, AIP Publishing).
4.3
Lasers
Poor current conduction has also prevented the achievement of electrically pumped QW lasers operating in the UV-B (280–315 nm) and UV-C (200–280 nm) bands[566-570]. Because they are still being extensively investigated in research laboratories around the globe, the development of DUV LDs is still in its infancy stage, as these researchers almost exclusively attempt to exploit the AlxGa1?xN material system[571, 572]. To realize an LD, energy transfer into the gain medium of the laser from an external source to produce excited states in the atoms (referred to as the laser pumping process) is necessary. The laser pumping method can be classified into two categories:
1. Electrical pumping: LDs are naturally pumped using electric currents. Researchers have demonstrated electrically pumped lasers at relatively long emission wavelengths, such as 330 nm[573] and 342 nm[574], and also at 275 nm[575]. The former two operate within the UV-A region, while the latter emits a UV-C light suitable for germicidal applications. The challenges for realizing electrically injected DUV LDs using AlxGa1?xN MQW structures include the high densities of dislocations in the lattice structure, large spontaneous and piezoelectric polarization fields, and poor p-type doping characteristics. These issues hinder improvement in the DUV LD EQE and the output power of electrically injected AlxGa1?xN MQW-based DUV LDs. Some researchers have focused their attention on nanowire-based LDs because of the potentially high crystalline quality that they can achieve, the high LEEs, and the low strain-induced polarization exhibited in nanowire structures[576]. Random lasers based on nanowire arrays have been demonstrated at peak lasing wavelengths of 239 nm[577] and 262.1 nm[578]. However, their optical output powers are still considerably low, while peak emitting wavelengths are mostly unstable because the laser cavities are defined through random multiple scattering.
2. Optical pumping: Pumping LDs using an optical pumping source is an artificial way to operate a laser, although dye lasers are optically pumped[579-583]. DUV LDs emitting at the sub 250 nm regime have been demonstrated[584-586], but producing light using this method would require another optical power source with higher optical power outputs and shorter excitation wavelengths, rendering this mechanism impractical in real-life applications.
Given these hindrances, conventional optically pumped DUV LDs and electrically pumped random DUV LDs are impossible to practically exploit in vital applications, such in the study of geology and climate in outer space. An alternative way to drive an LD active material is by using an electron beam, through which electron–hole pairs are generated by the energy transfer of high-energy electrons emitted from an electron gun. One advantage of this approach is that the incorporation of p- and n-type doped layers is not necessary, avoiding the undesired optical absorption commonly associated with these layers. In 2015, Klein et al. demonstrated high-power green and blue electron-beam-pumped LDs, with a peak optical powers of 5.9 W (
5.
Substrate technology and heterogeneous integration of new materials
To realize efficient optoelectronic devices for applications necessitating DUV light emission, high-crystalline-quality AlN templates are the key enablers of success in growing upper AlxGa1?xN epitaxial layers. Various growth techniques were developed to either achieve efficient UV and DUV light emission or merely improve the crystalline quality of AlxGa1?xN layers[589, 590].
5.1
High quality AlN templates
Kataoka, Funato, and Kawakami developed 3D Al0.50Ga0.50N QWs on MOVPE-grown AlN templates for polychromatic DUV LEDs[591]. They attempted a maskless regrowth of AlN on a trench-patterned AlN template procedure to create the AlN multi-facets of (0001), {
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Figure22.
(Color online) (a) Schematic illustration of an Al0.50Ga0.50N-based polychromatic UV LED structure on patterned AlN template. Reprinted with permission from Ref. [591]. ?2017, The Japan Society of Applied Physics. (b) Schematic illustration of a flip-chip mounted AlxGa1?xN-based UV LED with uneven MQWs. Reprinted with permission from Ref. [594]. ?2017, The Japan Society of Applied Physics. (c) Temperature-programmed desorption mass spectroscopy analysis for S- (left) and A-type (right) resin samples. For S-type sample, no decomposition observed. However, A-type sample showed CO, CO2, and H2O peaks after DUV irradiation. Inset shows the schematics of DUV LED with underfilling and encapsulation. Reprinted with permission from Ref. [410]. ?2015, The Japan Society of Applied Physics. (d) PL versus pump power of metacavity-MQWs. Narrow lasing peaks were shown from a broad emission spectrum. Inset shows PL peak intensity versus pump power indicating the lasing threshold. Reprinted with permission from Ref. [630]. ?2018, John Wiley & Sons.
Chang et al. demonstrated the growth of high-quality AlN templates on NPSSs via graphene-assisted quasi-van der Waals epitaxy[614-617]. This process allowed for rapid film coalescence, reducing growth times. Because of the presence of graphene, AlN films tend to be 2D laterally expanded on the NPSS, resulting in reduction of TDD (from 5.5 × 109 cm–2 to 1.5 × 109 cm–2 on the surface for layer thickness of 1.5
5.2
Hybrid integration on Ga2O3
Gallium oxide and its ternary alloys ((InxAlyGaz)2O3, x + y + z = 1 ) are wide-bandgap oxides of post-transitional metals with ionic bonds. With superior breakdown fields of around 8 MV/cm compared with SiC (3 MV/cm) and GaN (3.8 MV/cm), their intrinsic band gap spans from 4.9 to 5.3 eV; there are five known polytypes of Ga2O3, α, β, γ, δ, and ε[618]. Among these, β-Ga2O3 is the most stable structure with a superior melting point of up to 1820 °C; it became available as a wafer substrate by using conventional melt growth techniques. Despite its large bandgap ranging from 4.7 to 5.0 eV, its electrical properties can be tuned from an insulator to an n-type semiconductor with carrier concentration of up to 1019 cm–3 for DUV device applications[619, 620]. In contrast,
High n-type conductivity and transparency characteristics reveal that Ga2O3 is an ideal candidate for the realization of DUV devices as a transparent conducting oxide (TCO) substrate. Even though Ga2O3 does not have a hexagonal structure, it is relatively easy to induce a hexagonal atomic arrangement for optoelectronic device applications. In 2015, epitaxial growth of planar GaN layer on (100) a-plane
As part of a GaN epitaxial layer growth study on
5.3
Polymeric materials as key enablers for DUV light-emitting device mass production
Polymeric materials, such as optically isotropic and stable amorphous fluorine resin[627, 628] and graphene oxide (GO)-based fluoropolymer composites[629], may constitute a suitable substitute to other materials for the mass-production and packaging of AlxGa1–xN-based UV and DUV LEDs because of their decomposition characteristics below 330 nm[410, 457, 627, 628]. DUV LED dies are generally flip-chip bonded on submounts with gold bumps, and polymer-based materials can be used between the submount and dies as underfilling materials, as well as for the rest of entire surfaces (as an encapsulation). Yamada et al. revealed that instead of conventional resins, C-F-based polymers, so called S-type amorphous fluorine resins (terminated with a trifluoromethyl), are ideal candidates for the underfilling and encapsulating materials for the DUV LEDs[410]. Fig. 22(c) shows the temperature-programmed desorption mass spectroscopy analysis for S- and A-type resin with an inset indicating the schematics of flip-chip LEDs with underfilling and encapsulation. From this experiment, S-type resin tests showed no peaks after 261 nm light irradiation. However, A-type resin showed CO, CO2, and H2O peaks after DUV irradiation, which means decomposition of materials. With the optically isotropic characteristics of S-type resin for encapsulation, it showed no degradation over 3000 h of operation time.
5.4
Nanolasers realized via nanoscale plasmonic effects
Increasing demand of miniature-sized LD are of importance for the optical communications and data storage industries. One of the possible approaches for the realization of nanoscale lasers is using surface plasmon polaritons, which contain strong localized fields at the dielectric/metal interfaces. Recently, Shen et al. demonstrated DUV lasers with hyperbolic metamaterial (HMM) including multiple stacks of dielectric (20 nm of MgF2)–metal (20 nm of aluminum) interfaces[630, 631]. Those specially designed hyperbolic metacavity arrays helped to achieve a high-quality photonic density of state (PDOS), and eventually all excited plasmon oscillations coupled in one lasing mode. Focused ion beam (FIB)-fabricated 200 nm2 hyperbolic metacavity arrays were formed on MOVPE-grown high aluminum-content AlxGa1–xN MQW active regions, and the top of the structures were capped with annealed aluminum films as reflectors. Using a 266 nm pulsed laser as an exciting source, plasmonic nanolaser was demonstrated at lasing peak of 289 nm with narrow linewidth of 0.8 nm as depicted in Fig. 22(d). The lasing threshold is shown in the inset of Fig. 22(d).
5.5
Integration of 2D MX2/group III–nitride interfaces for the optoelectronic applications
In addition to UV light-emitting devices, researchers have also focused their efforts on advancing the integration of new materials for the realization of DUV photodetection. As a new approach for this application, the interesting characteristics exhibited at the interfaces of 2D TMDs and 3D group III–nitride semiconductors have recently emerged as the central theme of various studies[632-636]. The fabrication of 2D platinum diselenide (PtSe2)/GaN heterojunctions for self-powered DUV photodetectors has also been reported[559]. Group III–nitride semiconductor materials and 2D TMDs can be bonded through the weak out-of-plane van der Waals interactions in the absence of surface dangling bonds on their surfaces[637], providing an easy path for the integration of MX2/group III–nitride interfaces for the optoelectronic applications, where M = a transition metal (Mo, W, Nb, Ta, Ti, Re) and X = S, Se, Te[638, 639]. Specifically, the direct bandgap[640-642] 2D layered molybdenum disulfide (MoS2) is the most promising candidate for this purpose and widely explored for PD applications[560, 638, 643-649]. With a relatively small lattice mismatch of 0.8% between MoS2 and GaN layer, free-dangling bond surface of MoS2 provides excellent heterojunctions[637]. Zhuo et al. investigated p–n MoS2/GaN heterojunctions as self-powered PDs and developed a detection device for 265 nm DUV light operating at zero bias voltage. The naturally formed band bending at the MoS2/GaN interface results in the alignment of the Fermi levels of MoS2 and GaN[650], inducing a built-in electric field that enables the device operation at zero bias voltage. Moreover, the strong interaction of the MoS2 layer with incident light manifested excellent photoresponsivity to DUV light, with demonstrated high responsivity of 187 mA/W, high specific detectivity of 2.34 × 1013 Jones, and high ON/OFF current ratio of over 105[560].
6.
Electrical carrier injection and band alignment
In this section, the band offsets of heterojunctions and their effects on charge carrier injection associated with the electrical transport properties are reviewed. We also review the doping and injection efficiencies of charge carriers. We emphasize the role of band-offset parameters as important parameters to understand the carrier injection efficiency and confinement effect of devices employing heterointerfaces. We further discuss the state of current injection efficiency in AlxGa1–xN-based UV LEDs.
6.1
Band alignment
The state and band parameters at semiconductor heterointerfaces are essential to the effective design of active regions in contemporary electronic and optoelectronic devices[331, 635, 651]. Band discontinuities present at heterojunctions, formed by adjoining two different semiconductor materials, have unequal bandgap energies that act as potential barriers, which play an important role in controlling the charge carrier transport properties[652]. The feasibility of charge carrier injection across such heterointerfaces relies on the interface quality and the band-offset parameters (i.e., CBOs and valence band offsets (VBOs))[653, 654]. Based on these parameters, heterostructures are classified based on three alignment types: straddling gap (referred to as type I), staggered gap (type II), and broken gap (type III) junctions[655]. In general, type I band alignments with high refractive index contrasts are the most favorable for electron and optical mode confinement, thereby facilitating the carrier recombination process in the active regions of AlxGa1–xN-based optoelectronic devices. The other two alignment types are necessary to realize other electronic devices, such as hot electron transistors, provided they exhibit reasonably low CBOs, which allows for improved tunneling probabilities and lower turn-on voltages[656]. Researchers have extensively employed X-ray photoelectron spectroscopy (XPS) to determine band offset parameters and interface alignment types formed by group III–nitrides using their bandgap energy values. Several studies can also be found on the band offset parameters of dissimilar heterojunctions formed between compound III–nitrides and other families of semiconductor materials, such as oxides and 2D materials[635, 657]. For instance, King et al. determined the band alignment type for InN/AlN and InN/GaN heterojunctions to be type I heterointerfaces using high-resolution XPS (HRXPS) measurements and the bandgap properties of the constituent materials[658, 659]. The corresponding band alignment schematic is shown in Fig. 23. Table 4 summarizes a thorough literature survey of band offset parameters for compound III–nitrides-based and hybrid heterostructures. It is noteworthy that irrespective of the polarity and crystal structure, all group III–nitride semiconductors exhibit type I heterojunctions, which satisfies the fundamental requirement to design optoelectronic devices. However, although the nature of type I heterojunctions at wide bandgap compound III–nitrides interfaces facilitates carrier and optical mode confinement, there are other bottleneck issues, such as line defects, structural cracking, and low p-type doping efficiencies, hindering the overall device efficiencies.
onerror="this.onerror=null;this.src='http://www.jos.ac.cn/fileBDTXB/journal/article/jos/2019/12/PIC/18100007-23.jpg'" class="figure_img" id="Figure23"/>
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Figure23.
Schematic representations of the band alignment at (a) InN/AlN and (b) InN/GaN interfaces acquired using HRXPS studies. Reprinted with permission from Refs. [659, 658], respectively. ?2007, AIP Publishing and ?2008, American Physical Society.
| Heterojunction structure | VBO (eV) | CBO (eV) | Heterojunction type | Measurement method | Ref. |
| Wurtzite GaN/AlN | 0.7 ± 0.24 | – | I | XPS | [660] |
| Cubic GaN/AlN | 0.5 ± 0.1 | 1.4 ± 0.1 | I | Theory (ab initio) | [661] |
| Cubic AlN/GaN (zincblende (001) SLs) | 1.02 | 0.68 | I | Theory (Quasiparticle) | [204] |
| Wurtzite InN/AlN | 1.52 ± 0.17 | 4.0 ± 0.2 | I | XPS | [659] |
| 1.81 ± 0.2 | – | I | XPS | [660] | |
| a-plane GaN/AlN | 1.33 ± 0.16 | – | I | XPS | [662] |
| a-plane AlN/GaN | 0.73 ± 0.16 | – | I | XPS | [662] |
| Semi-polar AlN/GaN | 0.7 ± 0.2 | 2.1 ± 0.2 | I | XPS | [663] |
| InN/GaN | 0.58 ± 0.08 | 2.22 ± 0.1 | I | HRXPS | [658] |
| 0.85 | 1.82 | I | Internal photoemission | [664] | |
| 1.04, 0.54 | – | I | XPS | [665] | |
| 1.05 ± 0.25 | – | I | XPS | [660] | |
| 0.5 ± 0.1 | 2.2 ± 0.1 | I | XPS and PL | [666] | |
| 1.07 | 1.68 ± 0.1 | I | Photocurrent spectroscopy | [667] | |
| 0.72 ± 0.28 | – | – | XPS | [668] | |
| InAl0.83N/GaN | 0.2 ± 0.2 | 1 | I | XPS | [669] |
| In0.17Al0.83N/GaN | 0.15 | – | I | XPS | [670] |
| In0.25Al0.75N/GaN | 0.1 ± 0.2 | 0.4 | I | XPS | [399] |
| In0.3Al0.7N/GaN | 0.0 ± 0.2 | 0.2 | I | XPS | [399] |
| AlxGa1–xN/GaN | 0.3x | – | I | Theory (tight binding) | [671] |
| B0.14Al0.86N/GaN | 0 | 2.1 ± 0.2 | I | HRXPS | [672] |
| B0.14Al0.86N/Al0.70Ga0.30N | 0.4 ± 0.05 | 0.1 ± 0.05 | II | HRXPS | [673] |
| InN/($ bar{2}01 $) $ beta $-Ga2O3 | –0.55 ± 0.11 | –3.35 ± 0.11 | I | HRXPS | [674] |
| ITO/($ bar{2}01 $) $ beta $-Ga2O3 | –0.78 ± 0.3 | –0.32 ± 0.13 | I | HRXPS | [675] |
| ITO/(010) $ beta $-(Al0.14Ga0.86)2O3 | –1.18 ± 0.2 | 0.32 ± 0.05 | I | HRXPS | [676] |
| Al2O3/(010) $ beta $-(Al0.14Ga0.86)2O3 | 0.23 ± 0.04 | 1.67 ± 0.3 | I | HRXPS | [677] |
| Wurtzite InN/h-BN | –0.3 ± 0.09 | 4.99 ± 0.09 | II | XPS | [678] |
| GaN/single-layer MoS2 | 1.86 ± 0.08 | 0.56 ± 0.1 | II | HRXPS | [636] |
| In0.15Al0.85N/MoS2 | 2.08 ± 0.15 | 0.6 ± 0.15 | I | HRXPS | [635] |
Table4.
Summary of group III–nitride, III–oxide, and hybrid material heterointerface properties.
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| Heterojunction structure | VBO (eV) | CBO (eV) | Heterojunction type | Measurement method | Ref. |
| Wurtzite GaN/AlN | 0.7 ± 0.24 | – | I | XPS | [660] |
| Cubic GaN/AlN | 0.5 ± 0.1 | 1.4 ± 0.1 | I | Theory (ab initio) | [661] |
| Cubic AlN/GaN (zincblende (001) SLs) | 1.02 | 0.68 | I | Theory (Quasiparticle) | [204] |
| Wurtzite InN/AlN | 1.52 ± 0.17 | 4.0 ± 0.2 | I | XPS | [659] |
| 1.81 ± 0.2 | – | I | XPS | [660] | |
| a-plane GaN/AlN | 1.33 ± 0.16 | – | I | XPS | [662] |
| a-plane AlN/GaN | 0.73 ± 0.16 | – | I | XPS | [662] |
| Semi-polar AlN/GaN | 0.7 ± 0.2 | 2.1 ± 0.2 | I | XPS | [663] |
| InN/GaN | 0.58 ± 0.08 | 2.22 ± 0.1 | I | HRXPS | [658] |
| 0.85 | 1.82 | I | Internal photoemission | [664] | |
| 1.04, 0.54 | – | I | XPS | [665] | |
| 1.05 ± 0.25 | – | I | XPS | [660] | |
| 0.5 ± 0.1 | 2.2 ± 0.1 | I | XPS and PL | [666] | |
| 1.07 | 1.68 ± 0.1 | I | Photocurrent spectroscopy | [667] | |
| 0.72 ± 0.28 | – | – | XPS | [668] | |
| InAl0.83N/GaN | 0.2 ± 0.2 | 1 | I | XPS | [669] |
| In0.17Al0.83N/GaN | 0.15 | – | I | XPS | [670] |
| In0.25Al0.75N/GaN | 0.1 ± 0.2 | 0.4 | I | XPS | [399] |
| In0.3Al0.7N/GaN | 0.0 ± 0.2 | 0.2 | I | XPS | [399] |
| AlxGa1–xN/GaN | 0.3x | – | I | Theory (tight binding) | [671] |
| B0.14Al0.86N/GaN | 0 | 2.1 ± 0.2 | I | HRXPS | [672] |
| B0.14Al0.86N/Al0.70Ga0.30N | 0.4 ± 0.05 | 0.1 ± 0.05 | II | HRXPS | [673] |
| InN/($ bar{2}01 $) $ beta $-Ga2O3 | –0.55 ± 0.11 | –3.35 ± 0.11 | I | HRXPS | [674] |
| ITO/($ bar{2}01 $) $ beta $-Ga2O3 | –0.78 ± 0.3 | –0.32 ± 0.13 | I | HRXPS | [675] |
| ITO/(010) $ beta $-(Al0.14Ga0.86)2O3 | –1.18 ± 0.2 | 0.32 ± 0.05 | I | HRXPS | [676] |
| Al2O3/(010) $ beta $-(Al0.14Ga0.86)2O3 | 0.23 ± 0.04 | 1.67 ± 0.3 | I | HRXPS | [677] |
| Wurtzite InN/h-BN | –0.3 ± 0.09 | 4.99 ± 0.09 | II | XPS | [678] |
| GaN/single-layer MoS2 | 1.86 ± 0.08 | 0.56 ± 0.1 | II | HRXPS | [636] |
| In0.15Al0.85N/MoS2 | 2.08 ± 0.15 | 0.6 ± 0.15 | I | HRXPS | [635] |
6.2
Doping and injection efficiency of charge carriers
From the ABC model, the IQE is directly proportional to
m {inj}} $
m {inj}} $
With experimental evidence, Zhang et al. computationally demonstrated that silicon step-doping of quantum barriers would lower polarization-induced fields and reduce energy barriers for hole transport[694]. They compared three devices differing solely in terms of doping in barrier layers. Device I is a standard LED with unintentionally doped barriers, while devices II and III consist of silicon-doped layers and step-doped barriers, respectively, as shown in Fig. 24(a). Figs. 24(b) and 24(c) respectively show the experimentally measured and numerically simulated optical output power and EQE as functions of injection currents for devices I, II, and III. Considering the entire current range, device III outperformed I and II. The silicon step-doping suppressed QCSE and enhanced hole injection, consequently improving LED performance.
onerror="this.onerror=null;this.src='http://www.jos.ac.cn/fileBDTXB/journal/article/jos/2019/12/PIC/18100007-24.jpg'" class="figure_img" id="Figure24"/>
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Figure24.
(Color online) (a) Schematic representation of the LEDs. Device I is a standard LED with unintentionally doped barriers, device II is designed with 12 nm-thick barriers, each fully doped with Si, and device III features step-doped barriers (6 nm undoped and followed by 6 nm doped). (b) and (c) respectively show the experimentally measured and numerically simulated optical output power and EQE as a function of current for devices I, II, and III. Reprinted with permission from Ref. [694]. ?2013, IEEE.
6.3
State of current injection efficiency and AlxGa1?xN-based UV LEDs
With the advancement of technology, there is a colossal demand for high power and highly efficient UV optoelectronic devices because of their diverse application (water/air/food sterilization, surface disinfection, free-space non-line-of-sight communication, epoxy curing, counterfeit detection, fluorescence or Raman identification of biological/chemical agents, and various diagnostic and therapeutic medical functions)[695, 402]. With such a diversity in potential applications, as exciting as it may sound, UV optoelectronic devices are far from perfected. Thus, there is a plenty to improve in materials system (choice of material, synthesis/growth), device structure (fabrication), and packaging.
EQE of a light-emitting device (Fig. 25(a)) is the product of electrical injection efficiency (EE), IQE, and LEE, and typical values are shown in Fig. 25(b)[696]. EQE for the UV LEDs and LDs has been improved over time by tweaking materials or structures; however, it is still noticeably lacking[402]. One of the major bottlenecks for abysmal EQE is poor current injection. In this section, we highlight the techniques utilized to improve current injection of UV optoelectronic devices.
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Figure25.
(Color online) (a) Schematic diagram of an AlxGa1?xN-based UV LED. (Reprinted with permission from Ref. [652]. ?2017, MDPI). (b) Illustration of EQE of UV LEDs under different circumstances. (Reprinted with permission from Ref. [696]. ?2014, IOP Publishing).
Current injection is affected by numerous factors: metal contacts, interfaces in between contact material and electron/hole supplier layers, concentration of electron/hole in electron/hole supplier layers, electrical conductivity, and electric field (piezoelectric field) developed across the active region. Both p- and n-type doping became challenging in DUV AlxGa1?xN-based devices due to high aluminum concentrations[697]. High ionization energy of magnesium and lower mobility worsen the hole injection, which creates imbalance in charge carrier injection[698, 91]. Thus, research problems, achievements, and prospects on the enhancement of hole injection are worthy of discussion. The key factors that require attention for the enhancement of hole injection are outlined in Fig. 26.
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Figure26.
(Color online) An outline that underscores essential factors for enhancement in hole injection current.
When a device is electrically biased, the very first hurdle that a hole encounters is the interface of the p-type metal contact and the hole supplier layer[699]. In 1999, Ho et al. demonstrated a low specific contact resistance to p-type GaN where a Ni/Au thin film was deposited on p-type GaN and annealed at 500 °C in air to form NiO and obtain specific contact resistances as low as 4 × 10–6
In the majority of UV LED devices, a hole supplier layer consists of p-type GaN and AlxGa1?xN layers. As a hole reaches the hole supply layer, it encounters a high discontinuity in the valance band at the heterojunction, which makes the hole transport challenging. Graded heterojunctions[707], AlxGa1?xN/ AlyGa1?yN (x > y) superlattice structures[708-710], polarized induced electric fields[178, 711] (as shown in Fig. 27), and magnesium-delta doping[416, 180] are some of the major techniques employed to enhance the hole transport within the hole supplier layer, lowering the activation energy of magnesium dopants to as low as 17 meV. In addition, drift velocity can be significantly increased utilizing the strong field produced at the interface (for aluminum-rich p-AlxGa1?xN and p-GaN) of heterojunction in the hole supply layer[712].
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Figure27.
(Color online) (a) Measured temperature-dependent resistivity for different AlxGa1?xN samples highlighting the polarization boost in p-type conductivity. Sample a is a magnesium-doped GaN sample, while samples b and c are doped with the same magnesium concentration but linearly graded from x = 0 to x = 0.16 and x = 0 to x = 0.3, respectively. (b) relative output power intensity with increasing drive current for the graded AlxGa1?xN p–n junction and the control bulk-doped p–n junction. Schematic energy-band diagrams of (c) a conventional LED device and (d) a polarization-doped device. Reprinted with permission from Ref. [178]. ?2010, American Association for the Advancement of Science.
Electron leakage under high bias currents is inevitable. An EBL is not only thought to be capable of minimizing electron leakage, but also hindering hole transport because of the discontinuity of the valance band[652, 707, 713]. Mehnke et al. implemented AlN electron blocking heterostructures (EBHs) in AlxGa1?xN-based MQW LEDs with peak emission below 250 nm[714]. Owing to the EBHs, the emission from the QW was enhanced, whereas long wavelength parasitic luminescence was significantly reduced. Consequently, adequate hole injection was maintained by optimizing thickness of the EBH. Using such device structure, UV-C LED with peak emission at 234 nm, optical power of 14.5
Holes in the active regions of visible LEDs tend to accumulate close to the interface between the MQW region and the p-type EBL; however, this is not the case in DUV LEDs[716, 717]. Furthermore, hole distribution in the active region can be modified through manipulation of aluminum composition in AlxGa1?xN QWs and QBs. It has been observed that band offset is reduced, which is beneficial for hole transport across the active region[717]. Enhancement of hole concentration in each QW appears more easily accomplished than uniform hole distribution across an entire active region. The thicknesses of the QWs and QBs also affect hole concentration. Fewer than a couple nanometers of QW thickness and less than 8 nm for QB are the thicknesses recommended for the high hole concentrations and excellent spatial overlap of the electron–hole wavefunctions[718-721].
7.
Outlook and future challenges
While sapphire is intrinsically nonconductive, Ga2O3 substrates can be highly conductive and UV transparent, though the carrier mobilities are relatively low compared to GaN[722, 723]. Background impurities do not make sapphire conductive and only increase absorption in the UV region, while AlN templates are currently highly resistive, and AlN substrates suffer from point crystal defects within the UV regime[724-728]. Therefore, from the prospective of device fabrication, optoelectronic devices fabricated on AlN/sapphire require an etch down process that is cumbersome to implement, resulting in laterally-oriented devices exhibiting poor electron injection. AlN substrates might be the ideal choice in development of next-generation AlxGa1?xN-based DUV optoelectronic devices because AlxGa1?xN TDDs can be substantially reduced to values between 105 cm–2 and 105 cm–2, but cost-wise, they are not currently available for practical applications[99, 324, 391, 729-739]. Tokuyama Corporation in Japan reported EQEs of 2.2% and 2.4% for DUV LEDs grown on AlN substrates at 20 mA and 250 mA injection currents, respectively[740, 741], whereas Crystal IS, Inc. in the United States reported up to 4.9% and 5.9% EQEs at 300 mA and 50 mA, respectively, demonstrating low thermal roll-off of up to a 300 mA injection current[742]. Vertically-oriented devices with backside n-type contacts, on the other hand, demonstrate better current injection characteristics. Nanowire structures exhibit unique characteristics, such as lateral strain relaxation and the absence of lattice mismatches with the substrate and structural cracking. Hence, we anticipate that future efficient DUV optoelectronic devices will take advantage of the unparalleled features unleashed by combining nanostructured devices and Ga2O3 substrates. 2D materials and hybrid systems, such as phosphorene[743, 744], phosphorene-like structures[745, 746], 2D TMD heterojunctions[747-752], and graphene/h-BN heterostructures[753, 754], have also been used to improve the use of contact interfacial layers as tunnel layers[755-758], and are occasionally utilized for light coupling in visible and IR regimes because of their narrow bandgaps and high absorption of the signal[759]. We summarized key performance characteristics of state-of-the-art group III–nitride-based and hybrid UV and DUV PDs, as demonstrated in Table 3. Compared to recently emerging wide bandgap energy materials, most notably Ga2O3 and its alloys[760-781], AlxGa1?xN-based PDs remain the mature technology for eventual integration into various applications. Although Ga2O3-based PDs have been demonstrated with high responsivity (1.8 A/W at 236 nm),[782] the spectral response is relatively broad and is not easily tunable, compared to that of AlxGa1?xN-based material. Apart from that, slow response speeds of up to few seconds were also observed in Ga2O3-based PDs due to surface trap states and slow hole transport[782, 783]. These issues would eventually hinder their applicability in systems requiring high selectivity or fast response speed, particularly for wavelength-selective communication channels[784, 785]. BN-based PDs, on the other hand, may still require extensive fundamental study and device optimization before realizing a reliable device with high responsivity and response speed comparable to what AlxGa1?xN-based PDs can realize. Nevertheless, we envision that future development of AlxGa1?xN-based PDs or APDs will heavily rely and focus on 1) epitaxial growth optimization of high-aluminum-content layers for DUV detection, 2) maximizing transient response for communication applications, 3) increasing photocurrent gain, and 4) lowering operating requirements through device optimization, in order to achieve integration on UV or DUV photonic integrated circuits and applications.
Acknowledgements
We acknowledge the financial support from the King Abdulaziz City for Science and Technology (KACST) under grant no. KACST TIC R2-FP-008. This work was partially supported by the King Abdullah University of Science and Technology (KAUST) baseline funding no. BAS/1/1614-01-01, and MBE equipment funding no. C/M-20000-12-001-77 and KCR/1/4055-01-01.
