Review of nanophotonics approaches using nanostructures and nanofabrication for III-nitrides ultraviolet-photonic devices

Abstract. Group III-nitride semiconductor materials especially AlGaN are key-emerging candidates for the advancement of ultraviolet (UV) photonic devices. Numerous nanophotonics approaches using nanostructures (e.g., nanowires, nanorods, and quantum dots/disks) and nanofabrication (e.g., substrate patterning, photonic crystals, nanogratings, and surface-plasmons) have been demonstrated to address the material growth challenges and to enhance the device efficiencies of photonic devices operating at UV wavelengths. Here, we review the progress of nanophotonics implementations using nanostructured interfaces and nanofabrication approaches for the group III-nitride semiconductors to realize efficient UV-based photonic devices. The existing challenges of nanophotonics applications are presented. This review aims to provide analysis of state-of-the-art nanophotonic approaches in advancing the UV-photonic devices based on group III-nitride semiconductors.

Review of nanophotonics approaches using nanostructures and nanofabrication for III-nitrides ultraviolet-photonic devices

Introduction
The ultraviolet (UV) wavelength regime is typically divided into four wavelength ranges, i.e., UV-A (320 to 400 nm), UV-B (280 to 320 nm), UV-C (200 to 280 nm), and vacuum-UV (10 to 200 nm). Across these ranges, there are huge needs to realize UV-based photonic and electronic devices, e.g., laser diodes (LDs), light-emitting diodes (LEDs), photodetectors (PDs), high-electron-mobility transistors (HEMTs), metal-oxide-semiconductor field effect transistors, and radio-frequency switches. For instance, the market forecast for UV-LEDs is increasing from $90 million in the year 2014 to $800 million by the year 2020 due to the increasing demands in epoxy curing, water purification, food or medical sterilization, and new emerging applications. 1 The new applications include light sources, lithography, gas or chemical or bioagent sensing, free-space nonline-of-sight communication, phototherapy, forgery detection, medical diagnostic, consumer electronics, and horticulture. 2 To realize UV-based photonic and electronic devices, several wide energy bandgap materials have been utilized, e.g., group III-nitride semiconductors, diamond, SiC, β-Ga 2 O 3 , and ZnO. 3,4 The group III-nitride semiconductors have been considered the key UV-material because of their advantages, e.g., wide direct bandgap, material alloying capability for heterostructure and bandgap engineering, high electron mobility, high thermal stability, high breakdown field, *Address all correspondence to: Mohd Sharizal Alias, E-mail: drsharizal@gmail.com; Boon S. Ooi, E-mail: boon.ooi@kaust.edu.sa and high saturation velocity. 5 Furthermore, the material can sustain harsh temperature or radiation, and the material is toxic-free. Due to the alloying capability of the group III-nitride semiconductors, various material systems including AlGaN, InGaN, and InAlGaN have been used as the active region material for UV-based LEDs, LDs, PDs, and HEMTs. [6][7][8][9] Moreover, new boroncontaining III-nitride semiconductors, e.g., BN and BAlN, have emerged recently for UV-photonic and electronic devices. [9][10][11][12][13] Nevertheless, AlGaN has been the mostly used material for UVphotonic devices because it has a large and tunable energy bandgap from 3.4 to 6.1 eV that can cover UV-A, B, and C spectral ranges. Furthermore, AlGaN has the nearest material compatibility with the InGaN material system. This is important to capitalize on the commercial-ready platform and technology maturity stage of InGaN-based photonic and electronic devices. However, to obtain an efficient photonic device (e.g., high external quantum efficiency, EQE), the AlGaN material system encounters several key challenges regarding material growth and device fabrication. Obtaining a high-quality (crack-and defects-free) AlGaN crystal growth is challenging as the crystalline quality degrades with increasing Al composition and the existence of economic lattice-matched substrate is limited. 3,6,7,14,15 Also, the realization of high doping efficiency for the p-doped AlGaN layer is challenging because of the required high activation energy (E A ) in such a layer (e.g., Mg-dopant case). 3,6,15,16 Moreover, in the case of AlGaN-based light emitters, the light-extraction efficiency (LEE) for AlGaN-based devices is low (<10%) due to the anisotropic optical polarization property for Al-rich AlGaN layers and due to the high absorption of the contact layers and encapsulation materials. 14,16 To address these key challenges especially for AlGaN-based photonic devices, numerous nanophotonics schemes related to the material growth and device stucture involving nanostructured interfaces [e.g., nanowires (NWs), nanorods (NRs), and quantum disks (QDisks)] and nanofabrication approaches [e.g., substrate patterning, surface roughening, photonic crystals (PhCs), and surface-plasmon (SP)] have been proposed. 2,14,15,17 Here, we review the progress of nanophotonics implementations using nanostructured interfaces and nanofabrication approaches for the group III-nitride semiconductors (with the main focus on AlGaN material) to realize efficient UV-based photonic devices. We first discuss the AlGaN-related material and device challenges, and their effects on the device efficiency briefly. We then review the progress on using various nanophotonics schemes involving nanostructured interfaces and nanofabrication approaches adopted in the group III-nitride semiconductors for UV-based photonic devices. We will provide an overview of each approach and analysis of the device improvements (e.g., EQE and LEE). Finally, we provide an overall summary and point out the future directions of related research. This review aims to provide a framework of existing nanophotonic approaches and their further development in advancing the UV-photonic devices based on the group III-nitride semiconductors.

Nanophotonics Approaches for Enhancing Device Efficiencies
Focusing on light-emitting devices, e.g., LEDs and LDs, the device performance can be typically characterized by EQE for the LEDs case and by external differential quantum efficiency (EDQE) for the LDs case. The EQE is directly proportional to the internal quantum efficiency (IQE) and LEE. The IQE is a product of the injection efficiency (IE) and radiative recombination efficiency (RRE) of a material or a device. 7 As for the EDQE, it depends on IE, IQE, and mirror losses (which is indirectly associated with the LEE). 18 Thus, in both the LEDs and LDs cases, IE, IQE, and LEE play significant roles in enhancing the device performance. As stated earlier, several challenges limit the efficiency of the UV-photonic devices based on AlGaN material system, e.g., difficulty in crystal growth, p-doping, and LEE. Figure 1 shows the relationship of the above and several other challenges toward the IE, IQE, and LEE of LEDs and LDs.
Most of the challenges shown in Fig. 1 will be discussed further in the following section. Briefly here, the IE limitation is mainly due to the difficulty in obtaining high conductivity of doping (especially for p-type) to realize efficient p − n junction devices, e.g., LEDs or LDs. For example, a high E A (i.e., 310 meV for Al 0.27 Ga 0.73 N) is required to dope a Mg-dopants and will increase further for higher Al composition. 18 The low p-doping efficiency reduces the IE, which results in a high resistance p-contact and poor holes injection across the device. As for IQE, in general, increasing the Al composition to obtain shorter UV wavelengths in the AlGaN material causes a poor crystal quality due to increased density of defects in the form of cracks, dislocations, or impurities. Such defects can generate nonradiative center trap carriers and hence decrease the radiative recombination rate. This will reduce the IQE of the device. Due to the lack of native substrates, AlGaN has been grown on foreign substrates, e.g., sapphire, which is transparent to UV. However, to grow a high-quality AlGaN epitaxial structure on sapphire substrates is challenging due to large lattice mismatch and thermal expansion coefficient mismatch between the AlGaN/sapphire interface. Cracks easily formed with thicker AlGaN epilayer due to the accumulation of strong tensile strain. Compared with the InGaN epilayer on sapphire, AlGaN produced higher threading-dislocation densities (TDDs), i.e., ∼10 10 to 10 11 cm −2 due to the lattice mismatch, together with the large sticking coefficient and low surface mobility of Al adatoms during the crystal growth (in comparison, the TDDs for InGaN is ∼10 8 cm −2 ). 16,19 A recent approach of using bulk AlN substrates to grow AlGaN has been promising to reduce TDDs. However, the cost of AlN substrates is still expensive, and similar to sapphire, it is electrically insulating. To overcome this issue, pand n-contacts of the AlGaN-based devices need to be fabricated on the same side of the wafer (lateral injection). This results in nonuniform current injection across the device (current crowding) and localized heating. Furthermore, in Al-rich AlGaN epilayer, O 2 can be easily incorporated into the lattice, hence increasing the impurities. 20 The spontaneous and piezoelectric polarizations are dominant in group III-nitride semiconductors, and these polarizations are even stronger with higher Al composition (strain from the tensile stress). 19 Both polarizations can reduce the RRE due to the quantum-confined Stark effect (QCSE), i.e., electron-hole recombination decreased from the band bending effect. 19 Also, the LEE of AlGaN light-emitters suffers from a strong anisotropic optical polarization, i.e., the transverse magnetic (TM) polarization becomes dominant with increasing Al composition. Moreover, the p-contact layer (mostly p-GaN is used because of high conductivity and easy to form ohmic contact) and encapsulation materials are absorbing at most of the UV spectral range.
To overcome the above-mentioned challenges to enhance IE, IQE, and LEE for AlGaN-based UV photonic devices, different approaches using nanophotonic schemes have been developed. Figure 2 listed the nanophotonic schemes that can be applied to enhance the efficiency of UV-photonic devices. We categorized these schemes into two main parts, i.e., nanostructured interfaces and nanofabrication approaches. Using a generic UV-LD device structure as an example (as shown in Fig. 3), we list some of the nanophotonic schemes that can be used to improve the device efficiency. For example, the use of nanostructured interfaces (e.g., NWs, NRs, and QDisks) and nanofabrication approaches (e.g., substrate patterning, PhCs, and naogratings) can enhance the IQE and LEE of the device. The detail of the nanophotonic schemes and related device performances will be reviewed next.

Nanophotonics Schemes Using Nanostructured Interfaces
Several of the nanophotonic schemes, e.g., NWs, NRs, NCs, and monolayer are based on a bottomup approach, i.e., material growth. Due to the high crystallinity requirement for the group IIInitride semiconductors including AlGaN, their epitaxial structures have been grown mostly using molecular beam epitaxy (MBE) and metal-organic chemical vapor deposition (MOCVD).

Nanostructures: Nanowires, Nanorods, and Nanocolumns
The NWs technology has recently pursued as NWs exhibit unique characteristics, e.g., high crystalline quality with low defect density and large surface-to-volume ratio. There are three significant ways to grow such high-quality NWs: the so-called lithography-based selective area growth (SAG), vapor liquid solid (VLS) growth, and the catalyst-free self-assisted growth schemes. Although the researchers have been succeeded to grow GaN NWs using VLS process, unavoidable contamination due to the metal usage in the growth process remained challenging.  To overcome such a problem, many efforts have been explored for NWs growth in the absence of metal particles. 21 In view of this, researchers have used selective-area metal-organic vapor phase epitaxy to grow ordered arrays of III-nitride NWs. 22 On the other hand, the SAG and VLS growth schemes are time-consuming, complicated, lack of control over doping densities, and limited crystalline quality. 23,24 An alternative growth scheme has been utilized to accomplish highdensity self-assembled defect-free group-III nitride NWs at N-rich conditions. 25-30

Selective area growth
Several efforts have been made to grow high-quality GaN NWs, NRs, and NCs on patterned substrates using MOCVD and hydride vapor phase epitaxy. MBE can also be utilized to achieve SAG under enhanced nucleation of III-nitride nanocrystals that were formed through the surface nitridation of an array of group-III metal nanodisks (NDs) on Si. 24 At low nucleation process of MBE, the III-metal sources readily adhere to an SAG mask of SiO 2 , which results in a formation of a large number of self-assembled GaN NCs on counterparts of the patterned area. 31 For instance, Gotschke et al. have shown MBE growth of regular arrays of uniform-size GaN NWs on patterned Si/AlN/Si (111) substrates. The Si top layer has been patterned with e-beam lithography (EBL), resulting in uniform arrays of holes with different diameters and periods. 32 Choi et al. 33 reported the SAG of Ga-polar thin GaN NWs on patterned GaN/sapphire (0001) substrates using MOCVD with a continuous gas supply. GaN NRs formation on Ga-polar GaN by continuous mode MOCVD SAG was demonstrated under a relatively Ga-rich conditions. The growth rates of the semipolar planes, usually slowest growing planes, were controlled by varying the V/III ratio, which later allowed to grow GaN NRs along c-direction with nonpolar sidewalls.
First, the authors used a GaN∕c-Al 2 O 3 wafer that was patterned with a hexagonal array of circular holes using EBL and amorphous SiN x formed by plasma-enhanced chemical vapor deposition (PECVD) as a masking material. 34 Later, GaN nanostructures were grown at a fixed temperature (1125°C) and TMG flow rate (24 sccm), whereas the NH 3 flow rate was varied from 150 to 5 sccm. Figures 4(a)-4(d) show scanning electron microscope (SEM) images for the evolution of the GaN nanostructure morphology as a function of NH 3 flow rate of 150, 25, 10, and 5 sccm, respectively. This revealed that the nanopyramids (NPs) structure defined by {1-101} semipolar planes need to be formed under the higher pressure of NH 3 . As the size of the NPs is observed to be comparable with the original opening diameter, the growth rate of semipolar m-plane {1-101} is extremely low, which determines the formation of nanostructures as NPs.
The SAG of GaN NWs was further utilized as a template to grow AlGaN NWs-based photonic devices. For example, Le et al. 36 have used arrays of nanoscale opening apertures defined by EBL process on a Ti mask on planar GaN template on the c-plane sapphire substrate to grow GaN NWs. Later, this approach has been used as a template to grow defect-free coalesced AlGaN NWs-based UV-LED device structures. Figure 5(a) shows the schematic representation of selective area epitaxy of GaN/AlGaN NWs arrays on a GaN template/c-sapphire substrate by employing a thin (≈10 nm) Ti layer as the growth mask. Hexagonal openings situated as a triangular lattice were created on the Ti mask having a lattice spacing, a and a lateral size, h for each opening. These nanohole arrays were then created on the Ti mask using EBL and reactive ion etching. Later, GaN NWs arrays were selectively grown in the opening apertures by MBE. Figure 5(b) shows the SEM image of GaN NWs, exhibiting a high degree of size uniformity and are vertically aligned along the c-axis with well-defined hexagonal pyramid morphology comprised of six semipolar top facets.

Vapor-liquid-solid growth
The growth of III-V nanostructures has also been achieved using catalyst-assisted routes, e.g., catalytic metal-induced VLS mechanism, where metal droplets enable the material to solidify below the frozen liquid-solid interface. 37,38 The nanostructures can be then formed via the vapor−liquid−solid mechanism with, e.g., Au, Fe, or Ni as growth initiators. 39 Recently, Sivadasan et al. 40 reported on the optical properties of AlGaN NWs grown using the CVD method. They investigated the influence of the distribution of Au catalyst nanoparticles on the size and shape of AlGaN NWs. However, due to the deleterious effects of metal impurities on the structural, optical, and electrical properties of the semiconducting nanostructures grown using the VLS process, 23 researchers seldom used this method for the AlGaN NWs growth.

Catalyst-free self-assembled nanowires on conventional substrates
Another method to grow NWs is self-assembled growth, which does not require any catalyst and prepatterned substrates. Several researcher works have been reported on the growth of such spontaneously formed NWs on Si substrates using the MBE and MOCVD. 28,30,41 As compared with the ELOG, VLS, and SAG growth schemes, employing MBE or MOCVD for NWs growth is less time consuming and more scalable. Under certain conditions, e.g., N-rich and low metal conditions (V/III ratio ≫ 1) the growth of self-assisted dislocation free NWs on sapphire and Si substrates can be realized.
In this context, Janjua et al. 42 reported AlGaN NWs LEDs using the pendeo-epitaxy technique. Highly dense and spontaneously coalesced NWs emitting at 303-nm wavelength were grown on Si substrates. Figure 6(a) shows a top-view SEM image of the AlGaN NWs, whereas the inset shows a cross-sectional SEM of these NWs. The coalescence of the top p-GaN (in the inset) improves the current spreading and injection and reduces carrier leakage without the needs for planarization. Figure 6(b) shows a cross-sectional scanning tunneling electron microscope (STEM) image of an array of AlGaN NWs with device schematic shown in Fig. 6(c). To further optimize the AlGaN NWs LEDs, the self-assembled NWs were grown with varying QDisks thickness to study the carrier separation in the active-region and implemented a graded p-contact-layer to increase the output power. 43 They observed an improvement by 100 times in the direct recombination rate for samples with thicker QDisks thickness of ∼5 monolayers (MLs) compared with the sample with ∼2 MLs-thick QDisks. Moreover, the sample with graded top Mg-doped AlGaN layer shows 10 times improvement in the output power.
In addition to LEDs, NWs-based LD has been recently demonstrated. Li et al. 44 reported a highly stable and electrically pumped lasing emission in the UV-A at low temperature using the defect-free disordered AlGaN core-shell NWs arrays grown on Si substrate. Figure 7(a) shows    Figure 7(d) shows the EL spectra measured at 6 K. Two discrete peaks centered at 332.7 and 334.1 nm were observed. The spectral linewidth is as narrow as 0.2 nm above the lasing threshold of ∼12 A cm −2 .
Using NWs, an enhancement in the LEE for a UV-LED can be achieved. Djavid and Mi 45 studied the lateral side emission properties of NWs UV-LEDs using a simulation approach. Enhanced LEE of ∼72% was accomplished by optimizing the NWs size and density, as well as optimizing the thickness of the p-GaN layer. Figure 8(a) shows the dependence of LEE on the diameter of NWs. Hence, it is evident that there is an optimum diameter at which the light could be efficiently coupled through the NWs for efficient lateral side emission. Figure 8(b) shows the LEE that was further calculated as a function of the spacing between NWs from 163 to 203 nm. The corresponding LEE varied from 9% to 72%. TM polarized light can propagate horizontally through the NWs arrays resulting in significantly enhanced LEE. Figure 8(c) shows the LEE having an overall decreasing trend with increasing p-GaN thickness. Also, the authors have performed a simulation with device sizes in the ranges of 2.5 × 2.5 to 15 × 15 μm 2 . As shown in Fig. 8(d), the LEE decreases from 72% to 54% as device size increases. In addition to the diameter and spacing of NWs, the tapering angle of the grown AlGaN NWs can also enhance the LEE for a UV-LED. 46

Quantum Disks and Quantum Dots
Here, we discuss the research progress on the use of quantum dots (QDs) and QDisks for UV photonics. First, we focus more on the material properties and then we will discuss the advancement of the device based on QDs and QDisks.
In addition to the ability to tune the energy bandgap of the group III-nitride semiconductors by material composition variation, this tuning can also be achieved through quantum confinement using a thin active layer in the form of QDs and QDisks structures. Islam et al. 47 have shown that by changing the thickness of GaN QDisks, a PL emission as short as 219-nm wavelength can be obtained as shown in Fig. 9(a). Recent works reveal that a higher IQE can be achieved in the GaN-based QDs 49,50 and QDisks 47,48,51,52 compared with the AlGaN-based quantum wells structure in the similar wavelength range. Figure 9(b) shows a comparison between GaN and AlGaN heterostructures in terms of IQE. 47 Tanaka et al. 53 grew GaN QDs on AlGaN-layered surfaces with MOCVD and observed a strong PL peak at ∼3.55 eV. In the same year, Seifert et al. 54 demonstrated the growth of III-V and III-nitrides-based QDs by means of the Stranski-Krastanov method using the MOCVD platform. This is followed by the successful growth of GaN QDs using the MBE platform. 55 Next, Ristić et al. 56 reported defect-free and fully relaxed AlGaN NCs on Si (111) substrate grown by MBE. In contrast, GaN QDisks (also grown for a comparison study) are fully strained as indicated by the Raman measurement. 56 A comprehensive study on optical, electrical, and morphological properties of GaN QDisks was later followed by Ristić et al. 48,57 It was found that PL emission red-shifted (due to the QCSE) when the Al composition in the GaN/AlGaN QDisks is increased from 20% to 28%. However, blue-shift in PL emission was observed [as shown in Fig. 9(c)] when the thickness of QDisks is decreased from 4 to 2 nm (using the set of NCs with 28% of Al composition with different QDisks thickness). 48,57 Recently, several UV-photonic devices were demonstrated based on the GaN QDs and QDisks. 42,49,50,51,[58][59][60][61][62][63] One of the challenges to achieving efficient UV emitting devices is to overcome the poor current injection. In this context, Verma et al. 49,50 demonstrated UV-LED based on  GaN QDs as an active layer and utilized tunnel junction (TJ) scheme. Figure 10(a) shows the cross-sectional STEM of the device with the EL characteristic. In Fig. 10 To summarize, the ongoing research on GaN-based QDs and QDisks is paving a promising path to develop highly efficient UV-photonic devices. Table 1 summarizes the findings of research performed so far for both the material optimization and devices development.

Nanophotonics Schemes Using Nanofabrication Approaches
Various nanophotonic approaches to enhance the material properties (IQE) and device performance (e.g., IE, LEE, and EQE) of UV-photonic devices are discussed in this section. The nanophotonic approaches can be realized using advanced nanofabrication techniques including (1) lithographic patterning, e.g., EBL, nanoimprint lithography (NIL), and laser interference lithography (LIL), followed by wet-or dry-etching, or (2) direct patterning/etching using focused-ion-beam (FIB).

Epitaxial Lateral Overgrowth
Most common obstacles in the growth of group-III nitride semiconductors are the lattice/thermal mismatch between the grown epilayers and foreign substrates (due to the lack of native substrates that results in high TDDs (10 7 to 10 9 cm −2 ). Thus, a key factor to improve the performance of UV-devices is to reduce the TDDs and to solve the cracking issue of AlN and AlGaN films that typically originate from tensile strain beyond the critical layer thickness. Considering the cost and large area limitations of the native substrates (e.g., AlN or GaN), the devices that are grown on the native substrates are difficult to be commercialized. Therefore, to enhance the material quality of epilayers grown on nonnative substrates, different approaches, e.g., epitaxial lateral overgrowth (ELOG), pulsed-ELOG (PLOG), migration-enhanced ELOG (MEELOG), patterned dielectric distributed Bragg reflector (DBR) embedded ELOG, and nanoscale-ELOG schemes, have been suggested and successfully employed. In the ELOG growth scheme, a masked substrate is used in such a way that the material selectively grows on the unmasked region and thereby lateral coalescence takes place at the growth fronts. For instance, Balakrishnan et al. 66 have demonstrated an air bridge-assisted high-temperature ELOG approach to deposit 12-μm-thick, high-quality AlN layers over a patterned 6H-SiC substrate as templates for the UV-LED. The overall dislocation density of the layer was estimated to be in the range of 10 6 cm −2 or less with major portions of the layer formed over the voids are dislocation free. Chen et al. 67 reported on pulsed-ELOG (PLOG) of AlN films on basal plane sapphire substrates. This approach, at temperatures higher than 1150°C, enhanced the adatom migration, thereby significantly increasing the lateral growth rates. This enabled a full coalescence in wing regions as wide as 4 to 10 μm resulting in a reduction of the TDDs from 10 10 to 10 7 cm −2 . Moreover, Adivarahan et al. 68  Figure 11(d) shows comparison between the output power of standard and PLOG grew UV-LEDs; the output powers of PLOG AlN-based LED devices do not saturate for DC pump currents, whereas the standard LED output powers start to saturate. The superior performance of the reported UV-LEDs was ascribed to the reduced thermal impedance and the decrease of overall defect density in the laterally overgrown AlN. Additionally, Jain et al. 69 reported the coalescence of growth fronts using a technique in which the migration-enhanced MOCVD was successfully combined with the ELOG method to grow AlN with less defect density.
Cheng et al. 70 have demonstrated an enhancement of the output power of UV-LED by realizing ELOG on a patterned DBR substrate. The UV-LED samples were grown on c-plane sapphire substrates by a low-pressure MOCVD system. Figure 12  cathodoluminescent (CL) images acquired for a reference sample (without DBR) and with DBR, respectively. The dark spots denote strong nonradiative recombination centers showing low CL intensity result from the TDDs terminated at the surface. The defect density of DBR embedded UV-LED is 36% lower compared with the reference UV-LED. Figure 12(d) shows the EL spectrum, whereas Fig. 12(e) shows the light-current-voltage (L-I-V) characteristic for all samples. The optical power of SiO 2 -and DBR-embedded UV-LEDs have significant 45% and 75% improvement as compared with the reference device. This enhancement was attributed to the improvement in the material quality from ELOG and in the light output coupling from the embedded DBR. Similarly, Huang et al. 71 reported the InGaN/GaN MQW UV-LED with embedded Ta 2 O 5 ∕SiO 2 micromirror array (MMA) in the ELOG of GaN for the LEE enhancement. They achieved a 75.2% LEE enhancement for the patterned SiO 2 LED, and a 102.6% LEE enhancement for the MMA-LED, both were compared with the control sample.
The nanoscale ELOG scheme shows a more TDDs reduction in the overgrown group IIInitride semiconductors as compared with the microscale ELOG. [72][73][74] For example, Conroy et al. 73 reported an inexpensive nanoscale patterning process for ELOG on AlN layers grown by MOCVD. The uniform 1-μm-length NRs structure, formed by dry-etching of monolayer silica spheres, was used as a template for nanoscale ELOG epitaxy. A schematic drawing and an SEM image of the structure are shown in Figs. 13(a) and 13(b), respectively. A massive reduction in TDDs in the overgrown AlN layer is observed when the sample was annealed at 1100°C. This is formed at a much lower temperature compared with that typically required for microscale ELOG. 75 Figure 13(c) shows the FIB section of AlN overgrowth at 1100°C tilted to g ¼ ð0002Þ showing a-type dislocations. Figures 12(e) and 12(f) show the dislocations at higher magnifications, whereas Fig. 13(d) shows the FIB cross sections of typical high-quality three step AlN templates in the g ¼ ð0002Þ without ELOG. Without NRs patterning, the observed density of the edge and mixed dislocations was found to be ∼4.4 × 10 9 cm −2 . In comparison, AlN films produced using the NRs ELOG process had TDDs of <10 8 cm −2 .
Recently, Lee et al. 74 have demonstrated AlGaN-based UV-LED with periodic air-voidsincorporated nanoscale patterns enabled by nanosphere lithography and ELOG on a sapphire substrate. Figure 14(a) shows their fabrication process of AlN/sapphire nanopatterned substrate   function of the injection current measured from 50 LEDs at room temperature. The LOP of the UV-LEDs/NPS exhibited higher LOP by 67% than that of the reference device.
To summarize the ongoing research activities on the ELOG growth scheme, it is an outstanding approach to achieve efficient UV-photonic devices by means of reducing the defects at the growth fronts. These growth fronts were later utilized as templates for the overgrowth of the device structures. Table 2 consolidates the reduction of TDDs in the device structures using the ELOG scheme. Recently, research efforts have been demonstrated to address the problem of large lattice and thermal expansion mismatch by growing GaN on layered transition metal dichalcogenides (TMDs), e.g., MoS 2 and WeS 2 , with the sapphire substrates beneath. [78][79][80][81] The similar route can be used in growing AlGaN for the UV wavelengths in addition to the ELOG scheme.

Substrate Patterning
Various substrate patterning approaches have been suggested to enhance the performance of IIInitrides-based devices grown on sapphire substrates. Among them, the patterned sapphire substrate (PSS), i.e., sapphire substrate surface patterning with nanometer-to-micrometer-sized artificial structures, has been proven to be very promising to enhance the device's performance regarding the IQE and LEE. While the IQE is improved due to the reduction of TDDs by selective-area growth, the LEE is improved due to the photon scattering at the uneven sapphire surface.
A typical patterning process of the sapphire substrate surface involves several steps, as shown in the upper part of Fig. 15. After substrate cleaning, the surface is coated with a resist layer. Then, the desired spatial resolution and shape of the pattern are determined by a designed mask using a relevant lithography method. The patterning process ends with the resist development. During this step, the resist is selectively removed. The second part of the PSS fabrication is to transfer the resist structure, i.e., pattern, to the sapphire substrate. Here, two different ways can be applied as shown in the lower part of Fig. 14; either by an additive technique by depositing a thin film on the substrate surface or using a subtractive technique by etching the substrate surface. Ultimately, the PSS process ends with the resist removal. In the literature, different shapes of PSS have been suggested and fabricated. Figure 16 shows SEM images for some examples of PSS.
Here, we review the use of PSS to enhance the performance of the UV-photonic devices by taking into account the approach used to fabricate the PSS (additive or subtractive techniques). The utilization of PSS for Flip-Chip UV-LED is also discussed. Thereafter, the advantage of using substrate sidewall roughening for UV-LEDs is presented. Ultimately, a few examples of patterning non-sapphire substrates, i.e., Si substrates for UV applications, are discussed.  Table 3 summarizes the reported achievements acquired from the use of PSSs for UV-photonic devices.

PSS fabrication for UV applications: additive techniques
NIL was used to pattern a nanoholes array on a planar AlN/sapphire template. AlGaN/AlN MQW grown on such templates showed a lasing emission at 271.9 nm with a threshold power density lower by 68% than that of the same MQW grown on planar AlN/sapphire templates. 89 This reduction in the threshold power density is attributed to the improvement in the structure quality reflected by the TDDs in the structure grown on nanopatterned AlN/sapphire templates. Nanoscale patterning in an AlN layer, resulting in self-assembled NRs on a sapphire substrate as shown in Fig. 16(f), was prepared by an inductively coupled plasma (ICP) etching using a selfassembled monolayer of silica spheres on the AlN layer as a lithographic mask. 73 Growing an AlN layer above these NRs, a reduction in the edge-type dislocation density by approximately two orders of magnitude is demonstrated. Crack-free UV-LEDs emitting at 280-nm wavelength  have been successfully grown on a 4-in. AlN-NRs/sapphire substrates. 74 The latter was fabricated by utilizing nanosphere lithography [ Fig. 14(a)] and ELOG. The output power of these LEDs is enhanced by 67%. This enhancement is attributed to the formation of embedded periodic air-voids-incorporated nanoscale patterns [ Fig. 14(b)], which cause an improvement in the crystal quality of epitaxial layers. Simultaneously, the LEE of this device is enhanced by breaking the predominant in-plane-guided propagation of the photons.
Recently, AlN NRs were prepared on a sapphire substrate by a catalyst/lithography-free method, which is polarity-selective epitaxy and etching. 92 Core-shell-type AlGaN/AlN MQW emitting at 260-nm wavelength and grown on these AlN NRs show an improvement by 40 times in the PL intensity compared with the same structure grown on planer AlN template.

PSS fabrication for UV applications: subtractive techniques
Stripe PSSs have been suggested to improve the light-output efficiency of near-UV-LEDs emitting at around 409-nm wavelength. 93 These substrates were fabricated by a standard photolithography method and etched by the ICP etching. The LEDs grown on such a stripe PSS showed an enhancement in the LOP ∼20% over the same LEDs grown on the nonpattern sapphire substrate. In addition to the reduction in TDDs, this enhancement is attributed to the geometric effect of the stripe PSS adding the escape probability of the guided light. Furthermore, Pan et al. studied the effects of both the stripe depth and the stripe orientation on the performance of GaN-dased LEDs emitting at 400-nm wavelength. 85 They employed a set of PSSs with periodic stripe patterns prepared by photolithography and wet chemical etching. The output power of LEDs is enhanced by 87% when they were grown on a 0.9-μm-deep PSS with stripes along <11-20> direction. Further an improvement of 40% in the output power was achieved when the LEDs were grown on a 0.9-μm-deep PSS with stripes along the <1-100> direction.
Hemispherical-and volcano-shaped PSSs, shown in Figs. 16(c) and 16(d), respectively, were also suggested to enhance the LEE of GaN-based LEDs. 83 The LEE of LEDs emitting at 380 nm and grown on the optimized volcano-shaped PSS is higher by about 2.8 times than that of the LED grown on the planar sapphire substrate and is enhanced by about 1.6 times in comparison with that of the LED grown on the hemispherical PSS. Zhou et al. 94 studied the effects of the of PSS size on the performance of UV-LEDs emitting at 375 nm. In this study, they used coneshaped PSS with different sizes that were fabricated by a combination of the thermal-reflow photoresist technique and an ICP etching process. They found that with increasing both the pattern size and PSS fill factor, the TDDs in UV-LEDs decrease. This is due to a reduced GaN island density that causes fewer grain boundaries and also due to an increased lateral overgrowth area of GaN. Consequently, an improvement in the IQE and LOP of the LED is observed; the output power of the UV-LED grown on a large PSS with a fill-factor (FF) of 0.71 was 131.8% higher than that of the same LED grown on a small PSS with a FF of 0.4. In 2013, Dong et al. 86 fabricated nano-PSSs by nanosphere lithography and have demonstrated the advantage of using these PSSs to enhance the performance of AlGaN-based deep UV-LEDs. In particular, the output power of the UV-LED emitting at 282-nm wavelength of about 3 mW with 3.45% EQE is enhanced by ∼98% compared with that of the UV-LED grown on the nonpatterned substrate. LIL combined with dry-etching and MOCVD has been used to fabricate highly periodic AlN lateral polar structures on a sapphire substrate with submicron domain widths. 95 Such structures can be used for quasiphase matching to exploit frequency doubling via second harmonic generation in UV-C.
Using a double-sided PSS for GaN-based LEDs, an improvement in the device performance is numerically predicted for different pattern shapes. 96 Experimentally, nanoscale patterns on a double-sided sapphire substrate were fabricated using NIL and ICP etching. The output power of the LED grown on the double-sided PSS was 52% greater than the control sample. Furthermore, the use of a double-sided patterned substrate results in an increase in the heat dissipation efficiency and hence a decrease in the LED thermal resistance. Pyramidal-shaped PSS substrates were prepared by a wet-etching without lithography and dry-etching processes. 97 Figure 17 shows the pattern morphology of three substrates used for this study. LEDs grown on these substrates showed an enhancement in the output power by ∼46.4 to 51.5% compared with that of the same LEDs grown on flat sapphire substrates. Furthermore, LEDs are grown on the n-PSS(III) substrate ( Fig. 17) with 73% pattern coverage has the highest output power.

PSS for flip-chip UV-LEDs
Flip-chip designs have been used for AlGaN-based UV-LEDs because conventional LED contact layer material is absorbing in the UV spectral range. To enhance the LEE of flip-chip UV-LEDs, an approach to integrate monolithic arrays of the microlens to sapphire substrates was suggested. 87 In particular, UV-LEDs were bonded on high thermal conductive AlN ceramic submounts to improve the thermal dissipation, whereas the emitted UV light was extracted through sapphire substrates with microlenses. The latter were fabricated by photoresist reflow and dryetching. An enhancement by 55% in the output power at 20 mA is reported for a UV-LED emitting at 280-nm wavelength with microlenses compared with the same structure fabricated without microlenses. 87 This enhancement is attributed to the reduced internal reflections of the light caused by the microlens surface profile.
Groove-shaped PSSs were fabricated using photolithography and dry-etching. A high-quality lateral epitaxy overgrowth-AlN templates grown on these PSSs have enabled to realize flipchip AlGaN-based UV-LEDs that are emitting at 266-and 278-nm wavelengths with an output power of 5.3 mW with 1.9% EQE and 8.4 mW with 3.4% EQE, respectively. 88 Flip-chip UV-LEDs on AlN substrates show a relatively low light extraction and low output power. This is due to the total internal reflection (Fresnel reflection) at the interface between the AlN surface and the ambient medium and to the internal optical absorption by the p-GaN contact layer and the bulk AlN substrate. To overcome the problem of the Fresnel reflection, the AlN bottom-side surface was constructed to have a configuration composed of a hybrid structure of PhCs and subwavelength nanostructures. This hybrid structure is shown in Fig. 18(a) and was realized using EBL and subsequent ICP etching. 91 Growing a UV-LED emitting at 265 nm on the resulting AlN substrate consequently led to an enhancement in the LED output power by 196% (the output power of 90 mW at 350 mA). Figure 18(b) shows a schematic drawing of the UV-LED used for this study.
Despite the advantage of this relatively high output power achieved in the aforementioned study, 91 the employed fabrication method, i.e., EBL, is time-consuming, and hence this limits the patterned device size. To overcome this issue, a large-area AlN nanophotonic light-extraction structure fabricated by NIL is used to realize UV-LEDs emitting at 265 nm. 90 Furthermore, an improved design of a multiple-strip p-electrode geometry for these LEDs results in a uniform current density distribution. Consequently, an output power higher than 150 mW has been recorded with low-efficiency droop. Also, an increase by ∼20 times in the output power over that of a conventional flat-surface LED was achieved.

Substrate sidewall roughening
Substrate sidewall roughening is an approach suggested to enhance the LEE of AlGaN-based UV-LEDs. 98 This approach has revealed that a roughening region for effective light extraction enhancement exists on the substrate's sidewall. Roughening outside the effective roughening region will degenerate the sidewall light extraction due to inward photon scattering and absorption. Picosecond laser dicing was used to form roughening layers on the substrate's sidewalls without introducing thermal damage. A UV-LEDs emitting at 275-nm wavelength with three roughening layers shows 13.2% higher output power than the same LEDs with two roughening layers. 98

Patterning non-sapphire substrate for UV applications
A UV-LED emitting at 359-nm wavelength was fabricated on a stripe-patterned Si substrate. 99 To achieve this, a fully coalesced ELOG AlN layer with a 7-μm thickness was used as a template for the UV-LED growth. This LED was fabricated in back-emission configuration, and the Si substrate was removed after the flip-chip bonding. An enhancement in the output light extraction by 50% is observed for the back-emission LED compared with the top emission one. Likewise, a high crystal quality AlN template was used to grow an AlGaN-based UV-LED on a microcircle- patterned Si substrate. 100 The latter was achieved using standard lithography and ICP etching. The fabricated UV-LED emits at 325 nm with an EQE of about 0.03%.
Cantilever epitaxy method was used to grow AlN layers on 200-nm-nanopatterned Si substrate. 101 This substrate was patterned into nanostripes with nominally 150-nm wide ridges and 50-nm wide trenches Si, with an etching depth of 100 nm using EBL and dry-etching. The resultant AlN/Si structures were used as templates to grow UV-LEDs emitting at 344-nm wavelength with a peak pulse power of about 0.7 mW.

Surface Roughening
Surface roughening has been proven to be an effective approach to enhance the LEE of UV-LEDs. Next, we grouped the surface roughening into surface texturing and surface patterning.

Surface texturing
Surface texturing architectures help to improve the surface emission area and increase the probability of photons to escape out of the surface by the angular randomization. 102 Lee et al. 102 fabricated a UV-LED device emitting at 365 nm and obtained a total enhancement by a factor of 7.8 compared with the conventional UV-LED. This is attributed to the surface structuring, which exhibits random scattering effects. Figure 19(a) shows the schematic of the devices showing the roughening surface with root-mean-square (rms) roughness of ∼398 nm by immersing the sample in 45 wt. % KOH at 110°C for 180 s. 102 Figure 19(b) shows the L-I-V characteristics of three samples emphasizing the enhancement in the overall performance toward surface-roughened LED. Furthermore, Chiang et al. 103 achieved an enhancement in the output power by 525% and in the EQE by 6.5 times for a UVLED emitting at 365 nm, compared with a conventional LED, by a combination of the noninsulation current blocking layer, wafer bonding, laser lift-off, and surface treatment processes. Figures 19(d) and 19(e) show a schematic drawing of the fabrication steps and SEM images for the roughened surface, respectively. A high improvement in the L-I-V characterization was realized as shown in Fig. 19(e).
Inoue et al. 90 reported a high output power of 150 mW of a UV-LED operating at 265 nm. The UV-LED structure was fabricated on AlN substrate to minimize the lattice and thermal expansion mismatches between the substrate and the epitaxial layers. However, the high refractive index (RI) of the AlN substrate results in a small light escape cone angle. Imprinted polymer patterns by NIL were performed to create a large-area AlN nanophotonic structure, which was capable of reaching an improved EQE by ∼4 times compared with a flat LED, as shown in Figs. 19(f)-19(h). Chiang et al. 104 have suggested an abraded surface structuring method for a thin-GaN (NTG) UV-LED as shown in Fig. 19(i). As a result, the LED emitting at 365 nm shows an output power enhancement of ∼18.3% to 27.4% and EQEs of 20% to 27.2% at an injected current of 700 and 1500 mA, respectively. The enhancements in the output power and the EQE were attributed to the elimination of the electrode effect due to a micrometer-range abraded surface modification after an etching process using KOH solution. Additionally, a 2.5-times light extraction gain has been reported by Zhou et al. 105 by utilizing a roughening process using photoelectrochemical etching. As a result, an enhancement output power of ∼4.6 and ∼2.5 times was achieved for 280-and 325-nm emitting LEDs, respectively.
Khizar et al. 87 have reported an enhancement by 55% in the output power of the UV-LED emitting at 280 nm by integrating monolithic microlens arrays. This is due to curvature surface profile of the microlenses, which helps in increasing light reflections, in which more light can be extracted and as a result, LEE is enhanced. In a subsequent report, Kim et al. 106 demonstrated an enhanced LEE by 23% for UV-LED emitting at by utilizing SiO 2 nanosphere arrays. The nanosphere arrays act like nanolenses with a hexagonal-like pattern when immersed into a BCB resin and undergo a heat treatment. The hemispherical pattern leads to the random direction of the light rays with the LEE increased by multiscattering effects. Recently, Liu at al. 107 characterized a UV-LED that operates at 214 nm via nanowalls structures and calculated an IQE of 60%. Alternative nanostructures features have been detailed by others, e.g., large-scale nanostructure nanosphere, 108 silica glass with tailored nanostructures, 109 and truncated cones. 110

Surface patterning
Recently, a light-extraction mechanism based on an Rh mirror electrode and a transparent p-type AlGaN:Mg contact layer has been investigated by Takano et al. 111 They successfully achieved a 20% enhancement in EQE at an emission wavelength of 275 nm. In this work [ Fig. 20(a)], a conventional AlGaN LED is modified into a UV-LED structure by utilizing the following features: (1) crack-free AlN on the patterned sapphire substrate, (2) AlGaN:Mg p-type contact layer, (3) mirror electrode, and (4) Si encapsulation resin. The encapsulation resin helped to reduce the 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 the conventional LEDs. Increased multireflection of the light in the chip and absorption reduction helps to improve the device performance. Figure 20(b) shows the device characteristics (output power and EQE). The encapsulated LED in hemispherical resin on a Si submount is shown in Fig. 20(c). Encapsulation using other UV-transparent material, e.g., AlN-doped fluoropolymer and BVE polymer with -CF 3 ends, has been demonstrated recently. 113,114 Additionally, several research groups demonstrated pixel-LEDs (by meshing the p-contact metals) to enhance the output power and EQE characteristics. 115,116 However, the pixels were fabricated at the microscale and introduced complexity into the LED fabrication.
A 70% improvement in output power at peak emission of 276 nm has been reported by  cone-shaped active mesa, which is covered by the inclined MgF 2 sidewall of Ti/Al/Ni/Au ohmic contacts. The p-GaN-layer etching was performed by ICP etching to reflect photons down to the substrate without being absorbed. Decreasing the light absorption and directing strong sidewallheading TM polarized light via the MgF 2 ∕Al reflectors leads to significant enhancement in the LEE and output power as shown in Figs. 20(e) and 20(f), respectively. Table 4 summarizes the recent applications of surface roughening to enhance the performance of UV-LEDs. In this table, the device characteristics are described regarding initial emission wavelength (λ o ), initial threshold current (I o ), and initial output power (P o ).

Photonic Crystals
PhCs are inhomogeneous dielectric structures with periodic modulation of the RI, which allows the control and manipulation of the light propagation. PhCs can form a photonic bandgap where the light cannot propagate through the structure within a range of frequencies. Concerning the geometry of the structure, PhCs can be divided into one-, two-, and three-dimensional (1-D, 2-D, and 3-D) PhCs. In 1-D PhCs, the periodic modulation of RI occurs in one direction only, whereas RI stays uniform in the other two directions. 1-D PhCs have been widely used for several photonic devices, e.g., LDs, LEDs, and PDs in the form of a distributed Bragg reflector (DBR) as a highly reflective mirror. The most important aspect of the 2-D/3-D PhCs design is the lattice pattern choice. Most widely studied two-dimensional (2-D) PhCs either are with a hexagonal (triangular) or square lattices. In general, 1-D PhCs can be formed using thin-films deposition at various growth temperatures to suit the material or process requirements. [117][118][119][120][121][122][123] As for the 2-D/3-D PhCs, it can be fabricated using EBL, photolithography, NIL, or direct laser writing, followed by various etching techniques to form the air holes or pillars of a layer. [124][125][126][127][128]

1-D Photonic crystals: distributed Bragg reflector
A DBR is formed by multiple stacking of two alternate high and low RI layers with the thickness of each layer equivalent to the quarter-wave of the designated wavelength. DBRs have been used as high reflective mirrors in many UV-based photonic devices (summarized in Table 5), e.g., vertical-cavity surface-emitting laser (VCSEL), 129-131 edge-emitting laser (EEL), 139 microdisk LD, 137 resonant-cavity LED (RCLED), 132 resonant-cavity PD (RCPD), 133 and LEDs. 70,[134][135][136]138 These DBRs were formed using oxides named as UV-dielectric DBR and group III-nitride semiconductors, named as UV-nitride DBR. Most of the aforementioned UVbased photonic devices require DBR(s) with high reflectivity, R > 99%. To achieve such a high R, the contrast in RI must be high to have a smaller period and broad stopband, the chosen materials exhibit a high transparency at the device operating wavelength, the material interfaces must be abrupt with smooth surface to reduce the optical loss, and the chosen materials are compatible with lattice-matching to subsequent material deposition/growth. Compared with UV-nitride DBRs, UV-dielectric DBRs offer more extensive material choices (no lattice-matching limitation) and flexibility in the material growth or deposition technique, e.g., sputtering, PECVD, and electron beam evaporation. Importantly, highly reflective DBR (R > 99%) and broad stopband (>30 nm) can be easily obtained within a small number of repetitions (typically 10 to 15 periods) due to the high RI contrast among oxide materials. Table 6 lists some examples of oxides-based DBRs that have been demonstrated in the UV spectral range, e.g., ZrO 2 ∕SiO 2 , 140 Ta 2 O 5 ∕SiO 2 , 129 TiO 2 ∕SiO 2 , 130   is required (>40 periods) and the DBR stopband is not so broad due to insufficient RI contrast. Replacing Al 2 O 3 by HfO 2 as the high RI layer for the DBR is interesting because of high RI contrast. However, HfO 2 is absorbing for wavelengths below 250 nm. Recently, a highly reflective UV-C dielectric DBR has been demonstrated using a YDH∕SiO 2 material system. 144 Figure 21(a) shows that YDH thin films are less absorbing compared with HfO 2 thin films across the UV wavelengths. In Fig. 21(b), the YDH thin films exhibit wider energy bandgap of ∼6.21 eV compared with the HfO 2 thin films as measured using a phase-modulated spectroscopic ellipsometry. 145 This increase in energy bandgap indicates the transparency of the YDH material in the UV-C wavelengths especially below 250 nm. Figure 21(c) shows the reflectivity spectrum of 15 periods YDH∕SiO 2 DBR and the transmission spectrum of single-layer YDH thin films as a function of the wavelength. The reflectivity peak of R ∼ 99.9% of the DBR is obtained before the absorption from YDH becomes dominant. The achieved stopband of ∼50 nm is quite extensive within the UV-C regime. The inset of Fig. 21(c) shows the crosssectional SEM micrograph of the DBR structure indicating a uniform period thickness distribution with minor roughness introduced. Despite the advantages of the UV-dielectric DBRs, e.g., extensive material choices, fabrication flexibility, highly reflective, high RI contrast, and the wide stopband, the major drawback is that they are not electrically conductive and the thermal conductivity is mostly low. A complex current injection device scheme, e.g., intracavity 132,134 or flip-chip, 138 is required to pump a LEDs with dielectric DBRs electrically. Moreover, the poor thermal conductivity of dielectric DBRs deters the dissipation of heat generated during lasing operation. To overcome limitations, epitaxially grown UV-nitride DBRs have been developed using lattice-mismatched material [(Al)GaN/Al(Ga)N-based, e.g., GaN/AlN, 146 GaN/AlGaN, [147][148][149][150] AlGaN/AlGaN, 151 AlGaN/ AlN, [152][153][154] and boron-containing nitride, i.e., BAlN/AlN 155 ], lattice-matched material (AlInN/AlGaN), [156][157][158][159] and nitride-porous (AlGaN/air) material. 135 The advantage of the UVnitride DBRs is that they can be monolithically grown and doped (either por n-doped for electrical pumping), and they are furthermore thermally conductive. However, obtaining crack-free DBRs is challenging due to the limitation in crystal growth (substrate compatibility, i.e., latticeand thermal expansion-mismatch, and accumulation of tensile stress) and inadequate RI contrast, i.e., higher periods are required which is challenging for crystal growth in addition to the narrow DBR stopband. It is becoming more critical when the wavelength is shorter (at UV-C), where AlGaN/AlGaN DBR needs to be used to avoid absorption issue, but at the same time facing more difficulty in growing high-quality crystal (penalty of higher Al composition) and reduction in RI contrast. Recently, several strain-management techniques (as summarized in Table 7 and shown in Fig. 22) have been demonstrated for growing lattice-mismatched UV-nitride DBRs to suppress the cracks formation and to reduce the mismatch-induced tensile stress.  In Fig. 22(a), a crack-free 20 periods of GaN/AlGaN DBR is realized by inserting three sets of Ga/AlGaN superlattices (SLs) at several positions during the DBR growth. 146 The reflectivity at 399 nm wavelength is ∼97%. A similar technique was demonstrated by Moudakir et al., 149 where a multilayer (can be considered as SLs) of AlN/GaN/AlN was inserted at the midpoint of a GaN/AlGaN DBR to suppress the crack generation. For the crack suppression using the buffering technique, Someya and Arakawa 148 [ Fig. 22(b)] found that a highly reflective and crack-free GaN/AlGaN DBR can be grown as long as the thickness of the GaN buffer layer is <400 nm. Using a thick AlN buffer layer (∼1 μm), a highly reflective and crack-free DBR based on AlGaN/AlN with 29.5 periods was obtained at 356-nm operating wavelength [see Fig. 22(c)]. 152 The use of a thick AlN buffer layer keeps the AlGaN layers in compressive stress, whereas the AlN layers were strain-free. A double-layer buffer (using AlN/AlGaN) was also proposed to grow a crack-free and highly reflective GaN/AlGaN DBR on a SiC substrate. 147 Additionally, an interlayer of GaN or AlN at any point of the UV-nitride DBR structure can be utilized to suppress the crack existence in UV-nitride DBR. Using a thin interlayer, the existence of tensile strain can be compensated. 150 Figure 22(d) shows that a crack-free AlGaN/AlGaN DBR can be realized after a GaN interlayer is grown after the initial growth of the AlN template on a sapphire substrate. 151 A highly reflective DBR is obtained using 45 periods. The crack suppression is still well managed for a thicker GaN/AlGaN DBR (with 60 periods) using few AlN interlayers at several midpoints during the DBR growth. 150 In Fig. 22(e), Abid et al. 155 reported a boron-containing nitride DBR using wider energy bandgap material in the form of BAlN/AlN DBR. This material system has higher RI contrast, which resulted in a wider stopband (16 to 20 nm) and exhibits less optical absorption than the AlGaN-based DBRs. However, obtaining a high-quality BAlN layer is challenging, which explains the low reflectivity characteristic of the grown DBRs. Recently, several UV-nitride DBRs were demonstrated for UV-C spectral range using advanced techniques, e.g., trilayer period, 160 balanced strains, 153 and AlN low-temperature (LT) renucleation. 154 In the technique of trilayer period [ Fig. 22(f)], a period of the DBR consists of three layers, i.e., AlGaN, AlInN, and AlInGaN was used to accommodate the lattice mismatch instead of simply using bilayer, which is typical for a DBR. As for the balanced strains technique [ Fig. 22(g)], a crack-free AlGaN/AlN DBR is obtained by counterbalancing the tensile and compressive stress build-up (using a certain Al composition). As for the AlN low-temperature (LT) renucleation [ Fig. 22(h)], an Al layer (grown at LT) was inserted at the midpoint of an AlGaN/AlN DBR structure. All of these advanced UV-nitride DBRs were highly reflective at sub-280-nm wavelength. However, the stopbands were considered narrow ranging from 9 to 16 nm. To overcome the lattice-mismatch and strain limitations in the UV-nitride DBRs, nearly lattice-matched DBRs using AlInN/AlGaN material system have been demonstrated (summarized in Table 8) as an alternative approach. Figures 23(a) and 23(b) exhibit a UV-A and UV-B latticematched nitride DBRs with a peak reflectivity at 361-and 246-nm wavelength, respectively. A wider DBR stopband (18 to 20 nm) was obtained for AlInN/AlGaN DBR compared with GaN/ AlGaN DBR due to higher RI contrast. Although crack-free DBRs can be realized using a lattice-matched material system, the AlInN/AlGaN-based DBRs faced complexity during the growth due to the needs of numerous temperature ramping processes (thus resulting in long growth time) and inhomogeneity of In atoms distribution at such growth temperature. 155 Furthermore, all of the AlInN/AlGaN-based UV-DBRs can only be grown on AlGaN or AlN templates due to the large built-up strain in the case of GaN templates.
Recently, a technique using nitride-porous, i.e., a combination of epitaxial growth and chemically wet-etching resulting in AlGaN/air, was reported for UV-nitride DBRs. In Fig. 24(a), the cross-sectional SEM image shows the alternating AlGaN epilayer (high RI layer) and porous structure (low RI layer) formed after the etching process. The reflectivity of the AlGaN/porous DBR was measure to be as high as 90% with a peak wavelength of 379 nm as shown in Fig. 24(b). A wide DBR stopband of ∼37.2 nm is achieved because of the higher RI contrast of nitride and air. The high RI contrast also resulted in a highly reflective DBR by only using 12 periods of AlGaN/air. Nevertheless, there are still limitations for nitride-porous DBRs, e.g., maintaining mechanical support for the porous layer, process scalability, time-consuming etching, and realizing current injection.

2-D Photonic crystals: in-plane pattern lattices
Kashima et al. 161 reported an AlGaN-based UVC-LED emitting at 283-nm wavelength with an EQE as high as 10% using PhCs etched into the top p-contact as shown schematically in Fig. 25(a). The cross-sectional STEM for the PhC structure is shown in Fig. 25(b).  Furthermore, the advantage of using PhCs to improve the device performance has been reported for AlInGaN/AlInGaN-based LEDs emitting at 333-162 and 340-nm 163 wavelengths. An enhancement in the output power by a factor of 2.5 is found for the LED (emitting at 333 nm) with a PhC compared with the control LED without a PhC. For the UV-LED emitting  at 340 nm, by integrating the PhCs into the device structure, the forward voltage decreased by ∼1.6 V and the power enhanced by 95% at 20 mA. 163 Other PhC works base their research on improving the PhC optical properties, e.g., in the case of exploring the effect of PhCs on time-resolved 333 nm UV-LED EL spectra 164 and microresonators. 165 Specifically, an in-plane PhC waveguide with a line defect is made [ Fig. 26(a)]. Additionally, central holes are displaced by a maximum of 12 nm laterally, and the lattice periodicity is 170 nm. The experimental PL data are shown in Fig. 26(b) together with the mode profiles obtained by 3D-FDTD simulation, manifesting one sharp high fundamental mode resonance in near UV and lower one in UV-A regime. Such studies and the PhCs concept can be successfully applied, for example, in AlN microdisks with embedded QDs exhibiting a tunable resonance spectrum consisting of numerous modes, highlighting the high-quality Q factors [ Fig. 26(c)]. PhCs also have been successfully implemented in devices for the use as PDs. 168 Figure 26(g) shows a PD/LED emission with hexagonal beam pattern due to the hexagonal lattice pattern PhC on top of the device. It is capable of sensing the absolute photovoltage (PV), and it is revealed that PV strongly depends on the device orientation with respect to the detected beam's direction, i.e., on the incidence angle θ and azimuth angle φ, as seen in Fig. 26(h). There we can see the sixfold periodic pattern as φ rotates a full circle, as is suspected due to the hexagonal type symmetry of the PhC that couples the light into PD layers. Last, the responsivity curves versus wavelength at different angles are shown in Fig. 26(i), showcasing the capability of PV difference sense. Thus, one of the possible applications of this PD includes orientation/direction-aware and localization (as in positioning) solutions.  166,167 Although this may not be considered a true general 3-D case, it is a hybrid structure consisting of a PhC partially etched through an AlN/AlGaN DBR. It establishes complex light-structure interaction that depends at a minimum on three orthogonal spatial degrees of freedom, thus 3-D. As can be seen from the results [ Fig. 26(e)], the PL signal is greatly enhanced. The PL enhancement is 5 times compared with DBR without the PhC. The peak wavelength of this structure is 374 nm-inside the stopband of DBR. The UV-DBR helps to eliminate resonant modes outside the DBR high reflection region, which in turn decreases energy waste. The simulated luminescence spectrum [ Fig. 26(f)] confirms the experimental resonant peak with high accuracy. Table 9 summarizes the reported UV devices utilizing 2-D or 3-D PhCs. The table compares typical device structure, working wavelength, and a general enhancement factor.

Nanogratings
Nanogratings (NGs), when properly implemented on a surface of a light-emitting device, is a way to improve the LEE. Alias et al. 172 demonstrated NGs for UV-LED as NWs gratings around the device lateral configuration as shown in Fig. 27(a). The NWs gratings were patterned using gas-assisted FIB (GAFIB) etching with I 2 halide gaseous precursor. The GAFIB etching can enhance etching, reduce surface damage, and decrease the redeposition effect. 173,174 Figure 27(b) shows a substantial increase in the LOP detected from the top. Figure 27(c) shows the calculated intensity pattern of the devices with and without DBR NGs.

Subwavelength gratings
Subwavelength gratings (SWGs) are not only applied to enhance the light output, but also to control the polarizing properties. Although SWGs support only zeroth-order diffraction, they still exhibit polarization dependence. When an SWG is designed as a polarizer, one often targets to achieve high polarization selectivity: ideally, one polarization component is completely cutoff, and the other one is diffracted with 100% efficiency. The Ge SWG polarizer that is shown in Fig. 28(a) is fabricated with extinction ratio (ER) of 15 to 17.4 dB in UV-A range with a low aspect ratio (AR). 175 Works by Weber et al. show feasibility of efficient wire grid polarizers (WGPs) in UV-B and UV-C regimes. 176,177,179,180 They have reported different material grating structures, e.g., Ir, 177 Si, 180 and Al 176,179,180 gratings down to 250 nm and W 179 grating down to 193-nm wavelength, as seen in Fig. 28(b) with given examples in SEM images of Ir structures on polymer gratings [ Fig. 28(c)]. Reference 178 shows insights into WGP design using material complex RI analysis and optical simulations presents TiO 2 as a better alternative [ Fig. 28(d)] to W and Ir polarizers for UV-C range. The TiO 2 -based WGP showed a superior ER of 384 (25.8 dB) and transmittance of 10% at 193 nm as shown in Fig. 28(e). Another work investigates Cr 2 O 3 SWGs as WGPs further in depth for UV-C with a grating period equal to ∼90 nm and measured ER of 138 (21.4 dB). 181 The performance was compared with Al 2 O 3 , Ta 2 O 5 , TiO 2 , and Al gratings at 193 nm. The Cr 2 O 3 grating showed the highest absorption thus suppressing the TE mode the most while having practically the same transmission as the other materials. Reference 182 by numerical simulation considered AlGaN UV-A LED enhancement and fabricated a structure with a simple Si SWG on top of the LED that resulted in ER of ∼16 and increase in TM polarized EL intensity by ∼20%. The same author also reported in Ref. 183 a reverse ER, i.e., s/p ratio of ∼4 at 360-nm wavelength, but by etching the top contact to form a grating structure.

Other devices and applications
Recent work demonstrates homogenous NG fabrication by exciting SP polaritons (SPPs) with 266-nm UV fs pulses thus allowing a formation of gratings with a wavelength-dependent period smaller than half of the excitation wavelength, showcasing another possibility of fabricating  SWGs. 184 This team managed to achieve gratings with around 50-to 60-nm period lengths. Another work 185  The nanograting structures in this section are summarized in Table 10. The table compares typical device structure, working wavelength, and a general enhancement factor.

Microdisks and Nanopillars
Another approach that has been suggested for the development of III-nitride-based UV emitting devices is the use of microdisks and nanopillars. Similar to the previously described NGs and PhCs, microdisks and nanopillars provide a resonant cavity for the light to be confined and amplified. Moreover, both the microdisks and nanopillars act as a micro-and nanoscale Emitting at 365 nm (electrical pumping) light source, thus becoming a promising approach for applications in UV micro-and nanophotonics. In 2003, Jiang and Lin 193 reviewed III-nitride microemitters, then in 2014, Néel et al. 165 briefly reviewed AlN microdisks, and in 2015, Zhang et al. 194 reviewed III-nitride microdisk lasers over the UV and visible spectrum. Here, we provide a historical context of a broad set of microdisks and nanopillars structures focusing exclusively on III-nitride UV emitters fabricated by a top-down approach.

Microdisks
A semiconductor microdisk laser is a circular resonator cavity formed by top-down fabrication of an active material, i.e., semiconductor material or structure, where both physical and chemical etching are used to pattern the structures and release them into a mushroom-like configuration. The microdisk made from the active material is sustained on a pillar made from the substrate Here, we present and discuss the experimental demonstrations and historical progress toward UV microlasers based on microdisk structures. The first top-down attempt to realize microdisks of GaN/AlGaN MQW was the work of Mair et al. 195 Employing pulsed optical pumping under LT condition (10 K), they were able to compare the PL (continuous and time-dependent) of the GaN/AlGaN material before and after microdisk fabrication [ Fig. 30(a)]. This study showed that the recombination lifetime and the quantum efficiency (QE) of the GaN/AlGaN intrinsic transitions, were improved by a factor of 10 after the formation of the microdisks. Additionally, the contribution of the impurity transition of misfit dislocations was drastically reduced. These findings confirmed the influence of microdisks structure on the light emitting properties of GaN-based materials and set from that moment a path for the further investigation of such structures. Using the same material system, the same workgroup was able to observe, for the first time, the formation of UV-A radial modes at room temperature and whispering-gallery (WG) modes at 10 K. 201 The first UV-A microdisk pulsed-laser was achieved by Chang et al. 202 demonstrating the first lasing WG features using a GaN microdisk as well as a decrease in the stimulated emission threshold by 10 times as compared with a GaN squared-shaped structure.
In addition to efforts mentioned above, Kneissl et al. 196 209 also demonstrated a mature computational approach to study the microdisk UV lasers.
One of the most recent report on microdisk lasers by Zhu et al. 198 illustrates a concept of spontaneous formation of a sidewall grating [Figs. 30(g) and 30(h)], which enables the single-mode lasing at 379 nm of a GaN microdisk on Si substrate. Ultimately, it is important to highlight the work of Sellés et al., 200,211 holding the record for the shortest wavelength III-nitrides microdisk laser of 275 nm [ Fig. 30(k)]. This was achieved by integrating ultra-thin GaN/AlN MQW (0.7/5 nm) on Si. As both GaN and AlN are binary compounds, the quality of the material is not compromised as compared with alloyed materials, e.g., ternary compounds, e.g., AlGaN and InGaN.

Nanopillars and nanowires
Similarly to microdisks, nanopillars and NWs are structures with the potential to function as a lasing element in the UV spectral range. As seen in Table 12, the lasing of such structures has been achieved in a variety of pillar/wire dimensions and configurations.

Table 11
Historical progress in III-nitrides-based microdisk for UV photonics.    [215][216][217]219 These NWs were then isolated and used in several configurations to demonstrate UV-A lasing. First, a single NW (∅ ¼ 0.14 μm, L ¼ 4.7 μm) was investigated under optical pumping at room temperature. After the pumping threshold, a single-mode lasing at 371-nm wavelength was emitted from the longitudinal facets produced by the Fabry-Perot (FP) oscillations along the wire. Moreover, by setting two discrete GaN NWs one besides the other, a coupled cavity was created with single-mode lasing at λ c ¼ 370 nm [Figs. 31(a), 31(b), and 31(e)]. The lasing emission from such a coupled cavity is associated with the Vernier effect, i.e., suppression of resonant modes due to destructive interference. 220 Due to the nature of this coupling effect, the lasing threshold was increased to 874 kW∕cm 2 ; this is about three times higher than the lasing threshold of a single NW. Similarly, a single mode lasing in a discrete GaN NW (∅ ¼ 0.35 μm, L ¼ 5.3 μm) was achieved by a method of mode-selection based on the UV ohmic loss of the GaN-NW/Au interface. By modeling and calculating the propagation loss of different resonant modes within the NW cavity in contact with the metal, a specific mode (the one with less exposure to the NW/metal interface) is found to experience a loss significantly lower than any other optical mode and to be low enough to be overcome by the modal gain. It was quantified that using this method, the lasing threshold increases only by 13%, reporting a threshold value of 276 kW∕cm 2 . Ultimately, a GaN NW (∅ ¼ 0.2 μm, L ¼ 5.0 μm) of the same nature was characterized by optical pumping while lying on a SiN NGs. It was observed that when the NW cavity is perpendicular to the grating pitch, the substrate grating works as a distributed feedback grating (DFG) enabling the single-mode lasing of the NW at a threshold power density of ∼300 kW∕cm 2 . It is important to highlight that the diameter and length of the GaN NWs are critical parameters toward single mode emission.
The reports by Li et al. 214   due to the high UV reflectivity of Al film. In Ref. 214, a WG-based single mode lasing at 373-nm wavelength is achieved at an optical pumping power density of 420 kW∕cm 2 . The WG nature of the lasing mode is verified through the peripheral far-field emission measurements around the set of nanopillars, relying on the fact that the single-mode WG lasing must be directional. Additionally, FDTD simulations were performed to understand the origin of the resonant peak wavelength. The emission spectra showed two resonant wavelengths for the nanopillars (λ c1 ¼ 365 nm, λ c2 ¼ 373 nm). The emission of the peak at 365 nm belongs to the firstorder WG, which is exposed to the walls of the nanopillar being prone to losses [ Fig. 31

Other unique micro-and nanostructures
Additionally, to what has been discussed earlier, we have found a set of unique structures that can significantly impact the development of UV photonics. Zeng et al. 221 demonstrated the first InGaN/GaN-based microring resonator in UV-A regime. Due to the shape of these structures, radial modes are not allowed, and only WG resonant modes can be produced. This fact represents an advantage of the ring-structures when developing WG resonators. Following the microring, Wang et al. 222 demonstrated a GaN nanoring [ Fig. 32(a)] enhanced by a coating of thin Al (50 nm). In this report, experimental and computational characterizations agree that the device presented hybrid WG-SP lasing modes (λ c ∼ 365 nm) created from the combination of the ringshaped semiconductor and the evanescent fields on the semiconductor/metal interface. Although this effect is not desired for single-mode lasing devices, the nature of the plasmonic resonances opens up the possibility of using this kind of nanorings for sensing at the nanoscale. 227 The same research group fabricated a GaN nanostripe, coated with 20-nm SiO 2 and 60 nm Al [ Fig. 32(b)]. 223 Here, they observed lasing at λ c ¼ 370 nm under a relatively low optical power density threshold of 0.042 kW∕cm 2 , at room temperature. Due to the waveguide-like shape of the nanostripe, and because of the highly reflective internal walls (Al), lateral confinement exists, and an FP cavity is formed in the transversal plane. An innovative approach is presented by Wang et al. demonstrating high Q (>6000) WGmode lasing at room temperature in a GaN micropyramid (λ c ¼ 367.2 nm, lasing threshold = 400 to 500 kW∕cm 2 ). 225 These GaN pyramids are created by KOH wet etching of GaN, previously transferred onto a Ni-Ag-Pt-Au/Cu substrate. This nonconventional method allows the formation of high-quality GaN micropyramids standing on a metallic reflector [ Fig. 32(d)]. Because of this configuration, the microcavity is capable of sustain "quasi-WG" lasing modes fed by the pyramid faces and the bottom mirror. Another unique microstructure was reported by Zhang et al. 224 including a very comprehensive study of the material properties, the optical cavity, the power injection, as well as the carrier dynamics. Mushroom-like structures (hemispherical microcavities) [ Fig. 32(c)] were achieved in an InGaN/GaN-on-Si system showing WG lasing modes with the strongest lasing peak at λ c ¼ 381 nm. One benefit of such hemispherical structure is the elimination of parallel facets (like those in microdisks), ideally removing any possibility of observing FP oscillations, creating pure WG/ radial modes. The last structure discussed is the nanospiral created by Liao et al. [Fig. 32(e)]. 226 This structure is made of GaN coated with 30-nm SiN and 50-nm Al thin film. Circularly polarized UV-A lasing at λ c ¼ 363.5 nm was demonstrated at a low lasing threshold of 0.017 and 0.016 kW∕cm 2 , respectively, for left and right circular polarizations. Similar to the previous structures coated with Al, the metal significantly enhances the transversal cavity and light confinement [ Fig. 32(f)]; however, in this spiral, the role of the longitudinal traveling modes along the waveguide-like structure is critical to creating polarization rotation. This polarization rotation could be generated by the minor changes of curvature along the GaN-Al interface, impinging a small phase change to the lateral confined modes. Moreover, these observations open opportunities for further research and development of nanosources of circularly polarized UV light. To the present day, this nanoscale circularly polarized UV laser, represent by itself a breakthrough in UV nanophotonics.

Surface Plasmons
SPs are the collective oscillations of free electrons at a metal-dielectric or metal-semiconductor interface, where the real part of the dielectric function changes its sign. 228 Plasmonics attributes optical technology of coupling SPs to electromagnetic waves. This merged photonic-electronic phenomenon excites propagating SPPs on planar surfaces or localized SPs (LSPs) on nanoparticles. Both of the SPPs and LPSs manifest themselves as evanescent electromagnetic waves with a penetration depth of only a few tens of nanometers in metal, at optical frequencies. The dispersion properties of SPPs and LSPs are unique as they significantly enhance the incident electromagnetic field near the metallic structures. 228,229 SPs can boost the EQE of UV light emitters through two ways, enhancing its LEE and IQE. [230][231][232][233] In the first way, generated photons captured within the device are intentionally directed toward a metallic layer to excite SPPs on its surface. 230 The condition of this process is that the plasmonic resonance wavelength of the metal matches the wavelength of the generated light. If the metal surface is rough enough to scatter SPPs as light out of the device, the LEE of the UV emitter can soar because of extracting the SP-induced magnified light out of the emitter. The second way dramatically increases the emitter IQE through the exciton-SP coupling. 231,232 This way is contingent on placing a QW near a metallic plasmonic surface such that electronhole pairs are located within the near field of the metallic surface. In this scenario, some electronhole pairs recombine and directly transfer their energy to the electrons oscillations on the metal surface to produce SPs instead of photons. This new recombination path, along with the typical electron-hole recombination which generates photons, decreases the spontaneous emission lifetime, and thus enhances the emitter IQE. Here, we discuss the SPs approaches for UV-photonic devices, which include approaches of thin films, a patterned layer, metamaterial, and nanostructures.

Thin plasmonic film
Gao et al. followed the first way mentioned above to enhance the device performance using AlGaN-based UV-LED [ Fig. 33(a)]. 230 In this structure, a 5-nm-thick Al layer is deposited above the MQW structure. When electron-hole pairs recombine in AlGaN alloys, they generate photons with a dominant polarization parallel to the crystal axis (Ekc). Consequently, as shown in Fig. 33(a), only weak light signal of polarization E ⊥ c can be extracted through the escaping cone. The remaining light signals (entire Ekc and part of E ⊥ c polarization) are captured within the LED and propagate parallel to the Al layer. The Ekc polarization one is a TM wave for the Al layer and can excite SPPs on its surface at 294-nm wavelength. Figures 33(b) and 33(c) show the atomic force microscopy (AFM) images for the top surfaces of the UV-LED and Al layer, respectively. Figure 33(d) compares the PL of the UV-LED without Al, after depositing the Al, and after oxidizing the Al. SP-enhancement of 217% is acquired for the PL at 294-nm wavelength.
Lin et al. 231 presented an example for SP-enhanced IQE of UV emitter. A schematic of the AlGaN/GaN QW structure is shown in the inset of Fig. 33(a). The QW is fabricated using MBE and coated with Ag, Au, or Al to excite SPPs on their surfaces. The top AlGaN layer is thin enough to raise the spontaneous emission decay rate of the carriers through direct transfer of the electron-hole energies to the electrons oscillations on the metals surfaces. Before metal deposition onto the device, the PL intensity and lifetime of the carriers were measured at 10 K [ Fig. 34(a)]. Interestingly, this research group discovers the ability to tune the plasmonic resonance wavelength of Ag from the visible to UV region by tuning the Al content of the AlGaN barrier [ Fig. 34(b)]. When the Al content within the barrier layer is ∼23%, the Ag plasmonic resonance matches the emission wavelength of the GaN QW. In Fig. 34(c), the SPenhancement of the PL emission from the GaN QW is compared for three different metals. A threefold increase in the UV-light emission is obtained when using Ag as a plasmonic layer.

Patterned plasmonic layer
Cho et al. 232 reported an SP-enhanced IQE of AlGaN-based UV-LED using a patterned plasmonic layer. The LED structure is schematically shown in Fig. 35(a). To provide an exciton-SP   coupling, the top most p-AlGaN and p-GaN layers are patterned with vertical holes via photolithography and dry-etching. Afterward, Al pillars are deposited inside the holes to be in the very close vicinity to the MQW, as shown in Fig. 35(a). Figure 35(b) shows the top-view SEM image after Al deposition. Then, Al and Ti/Au layer were deposited on top of the device to work as reflectors, and thus light is extracted from the backside, as shown in Fig. 35(a). The output power of the SP-enhanced LED is 45% higher than that of the control LED [ Fig. 35(c)]. This enhancement is attributed to the increase in the spontaneous emission decay rate by the exciton-SP coupling.

Hyperbolic metamaterial
An alternative approach is introduced by Shen et al. 233 to use resonant SPs modes excited within hyperbolic metamaterial (HMM) to enhance simultaneously the IQE and LEE of AlGaN-based UV-LED. In comparison with the typical SPs structures, HMM is made of multiple metal-dielectric interfaces. Excitons from MQW close to the HMM can excite its resonant SPs modes, which increase their spontaneous emission decay rate and the IQE. A significant PL enhancement by 520% is achieved for HMM-based LED in comparison with the control LED.

Plasmonic nanostructures
Son et al. investigated the influence of plasmonic nanostructures on the performance of AlGaN/ GaN LED. 234 The plasmonic nanostructure used in these LEDs consists of Al nanorings, which are assembled directly above the MQW as shown in Fig. 36(a) to allow direct exciton-SP coupling. The fabrication steps of the Al nanorings are summarized in Fig. 36(b). The diameter of the nanorings can be tuned through the diameter of the silica nanospheres until their plasmonic resonance wavelength matches the emission wavelength of the LED. Figure 36(c) shows the AFM of the LED, where the Al nanorings were fabricated using 300-nm diameter silica nanospheres. The PL spectra [ Fig. 36(d)] show various enhancements due to the IQE improvement from the exciton-SP coupling. Recently, Huang et al. 235 used LSPs excitation from Al nanoparticles to improve the LEE of a UV-LED. Using the oblique angle deposition method, Al nanoparticles were deposited at the surface of the LED. At 60-deg deposition angle, maximum PL enhancement (329%) is recorded at 282-nm emission wavelength. This enhancement occurs because when the photons excite LSPs on the surfaces of the Al nanoparticles, these LSPs recombine and generate photons with stronger intensities. For the non-LED device, Shetty et al. 236 investigated the ability of plasmonic nanostructures to improve the performance of UV-PD. Incident UV light on the PD excites LSPs on the surface of Al nanostructures. The magnified fields near the metallic structures are then scattered to increase the PD photocurrent.
In Table 13, we summarize the different trials of enhancing UV-photonic devices using SPs. The table provides a comparison regarding the fabrication technique, device type, structure, characteristics, and the enhancement percentage.

Conclusions and Outlook
The continuously increasing development in UV-photonic devices has mainly resulted from the intensive research on nanophotonics and the advanced nanofabrication approaches applied to realize nanostructures. In this review, we have given a comprehensive description of nanophotonics implementations using nanostructured interfaces and nanofabrication approaches for the group III-nitride semiconductors to realize efficient UV photonic devices.
The NWs-based technology has recently been pursued due to the unique properties of NWs, e.g., high crystalline quality with low defect density and large surface-to-volume ratio. We have reviewed the most significant ways to grow high-quality NWs for UV applications. Lithographybased SAG, VLS growth, and the catalyst-free self-assisted growth schemes have been presented. The key element to improve the performance of UV-photonic devices is to reduce the TDDs and to solve the cracking issue of AlN and AlGaN films, which stems from the tensile strain beyond the critical layer thickness. Therefore, to enhance the material quality of epilayers grown on nonnative substrates, different ELOG schemes have been suggested and successfully employed to enhance the performance of UV devices. We show that research on GaN-based QDs and QDisks is very promising to develop highly efficient UV-photonic devices. Such devices are already outperforming their counterparts regarding material quality, IQE, LEE, responsivity, sensitivity, and the tunability of the wavelength. However, the optical output power of such devices still needs to be improved. The performance of UV-photonic devices can be also enhanced by utilizing surface roughening technologies, e.g., surface texturing and surface patterning. These technologies have shown a high impact on the LEE and device characteristics due to their abilities in increasing the extraction of photons, controlling the lateral waveguiding emission, and in improving the transparent p-contact layer. Further advances in lithography techniques will facilitate the roughening fabrication process and ensure high-resolution structure/surface features. We have also reviewed different shapes of PSSs, e.g., stripe-shaped, volcano-shaped, and groove-shaped, fabricated by the combination of a lithography method and either an additive technique (depositing a thin film on the substrate surface) or a subtractive technique (etching the substrate surface). We have shown that the utilization of PSSs is very effective to enhance the performance of the lightemitting UV devices in terms of decreasing the TDDs and increasing the LEE.
LEE, IQE, and EQE of the UV-photonic devices can also be improved using PhCs structures. In the case of DBRs which is a 1-D PhCs, we have reviewed the recent development of dielectric, epitaxial nitrides, and nitride/porous DBRs. Achieving a truly crack-free, broad stopband, and electrically conducting DBRs for the UV spectral range remains the main challenge. Several achievements on the usage of 2-D and 3-D PhCs on UV-LEDs have been shown. Furthermore, several types of NGs were reviewed, e.g., SWGs in the form of WGPs and 1-D/2-D grating patterns, and a DBR grating. Continuous advancement in the field of direct laser writing, for example, will ease and speed up true 3-D structure prototyping and testing, allowing for greater precision and smaller nanofeatures to be fabricated.
For the case of micro-and nanostructures, such as microdisk, nanopillars, and NWs, we have reviewed the reported structures and devices based on III-nitrides operating in the UV spectral range. We expect that the continuously improving quality of III-nitride materials, as well as cross-conceptual approaches, will further enhance the performance of UV lasing devices for future applications. We have discussed different plasmonic-based approaches used to improve the performance of UV devices. We have shown that Al-based plasmonic structures have been primarily used to enhance the EQE of UV emitters significantly. One major hurdle of using SP-enhanced UV devices is Al oxidation, which inhibits their feasibility. Therefore, it is necessary to discover other metallic materials that have plasmonic resonance within the UV range and do not oxidize in air. Further, we recommend trying to incorporate plasmonic structures, such as Al nanoparticles, within the active region, which would significantly increase IQE of UV emitters.