3.1.1.

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 ${\mathrm{SiO}}_2$, 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 $\mathrm{GaN}/\mathrm{c}\text{-}{\mathrm{Al}}_2{\mathrm{O}}_3$ wafer that was patterned with a hexagonal array of circular holes using EBL and amorphous ${\mathrm{SiN}}_{\mathrm{x}}$ 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 ${\mathrm{NH}}_3$ 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 ${\mathrm{NH}}_3$ 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 ${\mathrm{NH}}_3$. 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.

Fig. 4

GaN nanostructures evolution with respect to ${\mathrm{NH}}_3$ flow rate show in (a) 150, (b) 25, (c) 10, and (d) 5 sccm. The insets show lower magnification SEM images. Reproduced with permission from Ref. 35.

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 ($\approx 10\mathrm{nm}$) 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.

Fig. 5

(a) Left: arrays of nanoscale opening apertures defined by EBL process on a Ti mask on planar GaN template on the c-plane sapphire substrate. Right: schematic of the SAG of GaN NWs arrays on the nanopatterned substrate. (b) A tilted view SEM image showing GaN NWs arrays selectively grown in the opening apertures. The inset shows a high-magnification SEM image of GaN NWs arrays. Reproduced with permission from Ref. 36.

3.1.2.

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.

3.1.3.

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 $\text{ratio}\gg 1$) 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 $\sim 5$ monolayers (MLs) compared with the sample with $\sim 2$ MLs-thick QDisks. Moreover, the sample with graded top Mg-doped AlGaN layer shows 10 times improvement in the output power.

Fig. 6

(a) Top-view SEM image of the AlGaN NWs-based device structure. The inset shows cross-sectional SEM of these NWs. (b) STEM image of the coalesced AlGaN NWs. (c) A schematic of the device. Reproduced with permission from Ref. 42.

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 the SEM image of NWs grown on a Si substrate with the device schematic shown in Fig. 7(b). Figure 7(c) shows the I–V curve with the inset indicating a low leakage current under reverse bias. The lasing characteristics of AlGaN NWs were investigated by electrical injection under a continuous-wave operation at various temperatures. 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 $\sim 12\mathrm{A}{\mathrm{cm}}^{-2}$.

Fig. 7

(a) SEM image of NWs grown on a Si substrate, (b) schematic of the device, (c) I–V characteristics of the device; the inset shows the I–V curve on a semilog scale, and (d) emission spectra measured at 6 K under different current densities with the inset shows zoom-in of the lasing spectra. Reproduced with permission from Ref. 44.

Using NWs, an enhancement in the LEE for a UV-LED can be achieved. Djavid and Mi⁴⁵ studied the lateral side emission properties of NWs UV-LEDs using a simulation approach. Enhanced LEE of $\sim 72\%$ 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 $2.5\times 2.5$ to $15\times 15{\mu \mathrm{m}}^2$. 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

Fig. 8

(a) LEE versus diameter of NWs, (b) LEE versus spacing of NWs, (c) LEE versus thickness of the p-GaN layer, and (d) simulated LEE for devices with side length variations. Reproduced with permission from Ref. 45.

4.2.1.

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 self-assembled 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.

4.2.2.

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 $\sim 20\%$ 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 $0.9\text{-}\mu \mathrm{m}$-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 $0.9\text{-}\mu \mathrm{m}$-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 cone-shaped 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 $\sim 98\%$ 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 $\sim 46.4$ 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.

Fig. 17

Pyramidal-shaped PSS substrates prepared without lithography. Reproduced with permission from Ref. 97.

4.2.3.

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 dry-etching. 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 flip-chip 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.

Fig. 18

(a) SEM image of fabricated light extraction structures: a triangular lattice of circular AlN cones with subwavelength nanostructures. The insets show an enlarged SEM image and a schematic. (b) Schematic of UV-LED grown on an AlN substrate with nanophotonic light extraction structures. Reproduced with permission from Ref. 91.

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 $\sim 20$ times in the output power over that of a conventional flat-surface LED was achieved.

4.2.4.

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

4.2.5.

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 $7\text{-}\mu \mathrm{m}$ 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.

4.3.1.

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 $\sim 398\mathrm{nm}$ 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).

Fig. 19

(a) Schematic of the conventional LED (top) and the roughened LED (bottom) with SEM image (inset). Reproduced with permission from Ref. 102. (b) L–I–V characteristics of the samples. Reproduced with permission from Ref. 102. (c) Schematic of the fabrication steps for the UV-LED. Reproduced with permission from Ref. 103. (d) SEM images under different etching conditions. Reproduced with permission from Ref. 103. (e) L–I–V characteristics of UV-LEDs. Reproduced with permission from Ref. 103. (f) Schematic of the fabrication steps for the nanophotonic light-extraction structure. Reproduced with permission from Ref. 90. (g) Output power and enhancement factor. Reproduced with permission from Ref. 90. (h) EQE as a function of the current. Reproduced with permission from Ref. 90. (i) Schematic and SEM images of devices with the micro abraded surface (insets). Reproduced with permission from Ref. 104.

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 $\sim 4$ 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 $\sim 18.3\%$ 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 $\sim 4.6$ and $\sim 2.5$ 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 ${\mathrm{SiO}}_2$ 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

4.3.2.

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 –${\mathrm{CF}}_3$ 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.

Fig. 20

(a) Schematics of conventional (left) and UV-LEDs (right). Reproduced with permission from Ref. 111. (b) Output power (top) and EQE (bottom) characteristics. Reproduced with permission from Ref. 111. (c) Encapsulated UV-LED in hemispherical resin. Reproduced with permission from Ref. 111. (d) Schematic of fabrication steps for UV-LED with micro-ring active mesa. Reproduced with permission from Ref. 112. (e) EL spectra of TE and TM polarized light with the angle-dependent output power profile (inset) and L–I–V curve of the devices. Reproduced with permission from Ref. 112. (f) Electrical field distribution of the devices (left) and the calculated output power of the devices as a function of the etched area to mesa ratio (right). Reproduced with permission from Ref. 112.

A 70% improvement in output power at peak emission of 276 nm has been reported by Fayisa et al. by employing an inclined sidewall at the p-GaN with ${\mathrm{MgF}}_2/\mathrm{Al}$ omnidirectional reflectors.¹¹² Figure 20(d) shows the schematic fabrication process of AlGaN-based UV-LEDs with a $5\times 5$ array of the microring active mesa. The microring array consists of a truncated cone-shaped active mesa, which is covered by the inclined ${\mathrm{MgF}}_2$ 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 sidewall-heading TM polarized light via the ${\mathrm{MgF}}_2/\mathrm{Al}$ 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 (${\lambda }_o$), initial threshold current ($I_o$), and initial output power ($P_o$).

Table 4

Recent progress of using surface roughening approach for UV-LEDs improvements.

4.4.1.

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),¹²⁹130.^–131 edge-emitting laser (EEL),¹³⁹ microdisk LD,¹³⁷ resonant-cavity LED (RCLED),¹³² resonant-cavity PD (RCPD),¹³³ and LEDs.^70,134135.^–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 UV-based 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.

Table 5

UV-dielectric DBRs and nitride DBRs employed in various photonic devices operating in the UV regime.

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 ($>30\mathrm{nm}$) 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., ${\mathrm{ZrO}}_2/{\mathrm{SiO}}_2$,¹⁴⁰ ${\mathrm{Ta}}_2{\mathrm{O}}_5/{\mathrm{SiO}}_2$,¹²⁹ ${\mathrm{TiO}}_2/{\mathrm{SiO}}_2$,¹³⁰ and ${\mathrm{Si}}_3{\mathrm{N}}_4/{\mathrm{SiO}}_2$¹⁴¹ for UV-A; ${\mathrm{Al}}_2{\mathrm{O}}_3/{\mathrm{SiO}}_2$¹⁴² for UV-B; ${\mathrm{HfO}}_2/{\mathrm{SiO}}_2$¹⁴³ and $\mathrm{YDH}/{\mathrm{SiO}}_2$¹⁴⁴ for UV-C, respectively.

Table 6

Examples of UV-dielectric DBRs demonstrated in UV wavelengths.

${\mathrm{SiO}}_2$ is commonly used as the low RI layer in UV-dielectric DBRs as it offers the lowest RI among the dielectrics and has wider energy bandgap. However, most of the materials used for high RI layer in the UV-A DBRs are not transparent at the UV-B and UV-C spectral ranges. Although ${\mathrm{Al}}_2{\mathrm{O}}_3$ is transparent for extensive range till vacuum-UV,¹⁴² a high number of period is required ($>40$ periods) and the DBR stopband is not so broad due to insufficient RI contrast. Replacing ${\mathrm{Al}}_2{\mathrm{O}}_3$ by ${\mathrm{HfO}}_2$ as the high RI layer for the DBR is interesting because of high RI contrast. However, ${\mathrm{HfO}}_2$ is absorbing for wavelengths below 250 nm. Recently, a highly reflective UV-C dielectric DBR has been demonstrated using a $\mathrm{YDH}/{\mathrm{SiO}}_2$ material system.¹⁴⁴ Figure 21(a) shows that YDH thin films are less absorbing compared with ${\mathrm{HfO}}_2$ thin films across the UV wavelengths. In Fig. 21(b), the YDH thin films exhibit wider energy bandgap of $\sim 6.21\mathrm{eV}$ compared with the ${\mathrm{HfO}}_2$ 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 $\mathrm{YDH}/{\mathrm{SiO}}_2$ DBR and the transmission spectrum of single-layer YDH thin films as a function of the wavelength. The reflectivity peak of $R\sim 99.9\%$ of the DBR is obtained before the absorption from YDH becomes dominant. The achieved stopband of $\sim 50\mathrm{nm}$ is quite extensive within the UV-C regime. The inset of Fig. 21(c) shows the cross-sectional SEM micrograph of the DBR structure indicating a uniform period thickness distribution with minor roughness introduced.

Fig. 21

The $\mathrm{YDH}/{\mathrm{SiO}}_2$ DBR characteristics showing (a) absorbance spectra comparison for YDH and ${\mathrm{HfO}}_2$ single-layer films in the UV regime, (b) Tauc plot indicates the energy bandgap for YDH and ${\mathrm{HfO}}_2$ single-layer films, and (c) reflectivity spectrum for 15 periods DBR and the transmission spectrum for a single-layer YDH films with the inset shows the cross-sectional SEM of the DBR structure. Reproduced with permission from Ref. 144.

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,¹⁴⁷148.149.^–150 AlGaN/AlGaN,¹⁵¹ AlGaN/AlN,¹⁵²153.^–154 and boron-containing nitride, i.e., BAlN/AlN¹⁵⁵], lattice-matched material (AlInN/AlGaN),¹⁵⁶157.158.^–159 and nitride-porous (AlGaN/air) material.¹³⁵ The advantage of the UV-nitride DBRs is that they can be monolithically grown and doped (either $p$- or $n$-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., lattice- and 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.

Table 7

Strain management techniques demonstrated for lattice-mismatched UV-nitride DBRs.

Fig. 22

Various strain-management techniques demonstrated for lattice-mismatched nitride-DBRs resulted in high reflectivity mirror at UV wavelengths and crack-free DBR structure as indicated by the cross-sectional SEM micrographs. (a) GaN/AlN SLs. Reproduced with permission from Ref. 146. (b) Thin GaN single buffer. Reproduced with permission from Ref. 148. (c) AlN single buffer. Reproduced with permission from Ref. 152. (d) GaN interlayer. Reproduced with permission from Ref. 151. (e) Wider energy bandgap. Reproduced with permission from Ref. 155. (f) Trilayer period. Reproduced with permission from Ref. 160. (g) Balanced strains. Reproduced with permission from Ref. 153. (h) AlN-LT nucleation. Reproduced with permission from Ref. 154.

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 $\sim 97\%$. 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 Arakawa¹⁴⁸ [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 $<400\mathrm{nm}$. Using a thick AlN buffer layer ($\sim 1\mu \mathrm{m}$), 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 lattice-matched 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.

Table 8

Examples of UV-nitride DBRs utilizing lattice-matched material combination.

Fig. 23

Examples of lattice-matched UV-nitride DBRs using AlInN/AlGaN material system showing (a) peak reflectivity at 361-nm wavelength, reproduced with permission from Ref. 157, and (b) peak reflectivity at 246-nm wavelength, reproduced with permission from Ref. 159.

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 $\sim 37.2\mathrm{nm}$ 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.

Fig. 24

Demonstration of nitride/porous UV-DBR showing (a) alternating AlGaN epilayer and the porous structure and (b) reflectance spectra of the DBR. Reproduced with permission from Ref. 135.

4.4.2.

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-nm¹⁶³ 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 $\sim 1.6\mathrm{V}$ and the power enhanced by 95% at 20 mA.¹⁶³

Fig. 25

(a) AlGaN UV-LED structure with and without a PhC. (b) Cross-sectional STEM image of the PhC. Reproduced with permission from Ref. 161.

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 spectra¹⁶⁴ 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 $\theta$ and azimuth angle $\phi$, as seen in Fig. 26(h). There we can see the sixfold periodic pattern as $\phi$ 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.

Fig. 26

(a) Top-view SEM image of the AlN PhCs waveguide cavity. (b) PL spectrum of an AlN width modulated waveguide cavity and the magnetic field amplitude profiles of the fundamental even and odd cavity modes. (c) PL of a $5\text{-}\mu \mathrm{m}$-diameter microdisc with AlN QDs. The insets show SEM image of a $2\text{-}\mu \mathrm{m}$ microdisc on a Si post; and a zoomed part of PL at one resonance. Reproduced with permission from Ref. 165. (d) Structure of the square PhCs lattice etched into DBR (left) and the top view SEM image (right). (e) Measured reflectivity spectrum of the AlN/AlGaN DBR structure and PL spectra of bulk GaN and patterned UV-DBR structure. Reproduced with permission from Ref. 166. (f) The corresponding band diagram and luminescence spectra (simulation). Reproduced with permission from Ref. 167. (g) Photo of LED with GaN nanostructure pattern emission (left) and SEM cross section (right). (h) Photovoltage scan at the wavelength of 388-nm versus the incidence angle $\theta$ and azimuth angle $\phi$ values. (i) Photovoltage versus wavelength at different $(\theta ,\phi )$ values. Reproduced with permission from Ref. 168.

4.4.3.

3-D Photonic crystals: hybrid structures

Figure 26(d) shows an example of 3-D PhC on the GaN-based sample for UV regime.^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.

Table 9

Summary of UV-photonic devices with 2-D/3-D PhCs approaches.

4.5.1.

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 cut-off, 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

Fig. 28

(a) Schematic of the SWG with top-view SEM image of the fabricated Ge-SWG and the measured transmission and ER of the Ge-SWG. Reproduced with permission from Ref. 175. (b) Measured values of the ER of WGPs. Reproduced with permission from Ref. 176. (c) SEM images of the fabrication process for the WGPs with the overcoated polymer gratings by ALD (top) and by sputter deposition (bottom). Reproduced with permission from Ref. 177. (d) Comparison of application wavelength ranges of several WGPs from various metals and oxides. (e) Measured transmittances and ER of the fabricated titanium dioxide WGP. Reproduced with permission from Ref. 178.

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 W¹⁷⁹ 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 ${\mathrm{TiO}}_2$ as a better alternative [Fig. 28(d)] to W and Ir polarizers for UV-C range. The ${\mathrm{TiO}}_2$-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 ${\mathrm{Cr}}_2{\mathrm{O}}_3$ SWGs as WGPs further in depth for UV-C with a grating period equal to $\sim 90\mathrm{nm}$ and measured ER of 138 (21.4 dB).¹⁸¹ The performance was compared with ${\mathrm{Al}}_2{\mathrm{O}}_3$, ${\mathrm{Ta}}_2{\mathrm{O}}_5$, ${\mathrm{TiO}}_2$, and Al gratings at 193 nm. The ${\mathrm{Cr}}_2{\mathrm{O}}_3$ 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 $\sim 16$ and increase in TM polarized EL intensity by $\sim 20\%$. The same author also reported in Ref. 183 a reverse ER, i.e., s/p ratio of $\sim 4$ at 360-nm wavelength, but by etching the top contact to form a grating structure.

4.5.2.

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 work¹⁸⁵ uses SPPs phenomenon to reduce PL threshold by 75% for lasing in GaN by introducing a defect as a cavity in patterned GaN NG structure with a coated 50-nm Al metal layer [Figs. 29(g)–29(i)]. Although a simpler approach to improving LEE of UV-LEDs is contact patterning in a regular structure like a grating demonstrated in Refs. 189 and 190 [Figs. 29(e)–29(f)]. Other grating patterns, e.g., hexagonal hole pattern and square pattern gratings, were used by Kim et al.¹⁸⁶ and Leem et al.,¹⁸⁷ respectively, as seen in Figs. 29(a)–29(c).

Fig. 29

(a) Schematic of an LED structure with a patterned ITO contact layer. (b) 45-deg tilt SEM image of the hole patterned p-GaN layer. Reproduced with permission from Ref. 186. (c) SEM image of 2-D patterned photoresist, the inset shows enlarged cross-section image. Reproduced with permission from Ref. 187. (d) Drawing of nanostructured MSM PD. Reproduced with permission from Ref. 188. (e) AFM image of 1-D patterned ITO. (f) The L–I characteristics of LEDs fabricated with unpatterned CIO/ITO, 1-D patterned CIO/ITO, and conventional Ni/Au contacts. Reproduced with permission from Ref. 189. (g) Schematic of the metal-coated GaN grating structure with defect cavity. (h–i) Simulated electric-field intensity in the metal-coated GaN grating (h) without and (i) with the defect cavity. Reproduced with permission from Ref. 185.

Large-area soft UV-curing NIL fabrication of NRs for PL enhancement up to 2.5 times was demonstrated in Ref. 191 that can benefit optoelectronic applications.¹⁹² Yet other unmentioned devices are UV PDs as demonstrated in Ref. 188 [Fig. 29(d)], where an Al metal grating with a slit width of 100 nm was employed to increase the photocurrent by 8 times at 322-nm illumination and at the same time achieving faster response and decay time of the PD.

The nanograting structures in this section are summarized in Table 10. The table compares typical device structure, working wavelength, and a general enhancement factor.

Table 10

Summary of III-nitrides UV-photonic devices with NGs approaches.

4.6.1.

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 material [Figs. 30(e)–30(g)]. Microdisks can also be realized in full contact with the substrate [Figs. 30(c) and 30(d)]. Here, we present and discuss the experimental demonstrations and historical progress toward UV microlasers based on microdisk structures.

Fig. 30

The development of UV-emitting microdisks from 1997 to 2018. (a) PL spectra of GaN/AlGaN material before (bottom) and after (top) microdisk fabrication. Reproduced with permission from Ref. 195. (b) Schematic of radial and WG resonant modes in a microdisk. (c) SEM image of the electrically driven directional UV-A microdisk laser. Reproduced with permission from Ref. 196. (d) Schematic of a GaN microdisk in full contact with a sapphire substrate. (e) Lateral and (f) perspective view of freestanding InGaN/GaN-based microdisks on Si substrate. Reproduced with permission from Ref. 197. (g) GaN microdisk on Si post with detail (h) of edge NGs. Reproduced with permission from Ref. 198. (i) FDTD-simulated electric field intensity of a $\lambda =301\mathrm{nm}$ WG mode as well as the mode matching with the corresponding PL emission (j). Reproduced with permission from Ref. 199. (k) PL emission of a $3\text{-}\mu \mathrm{m}$ microdisk at different pulse-injection intensities showing lasing at $\lambda \sim 275\mathrm{nm}$. Reproduced with permission from Ref. 200.

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.

The development of the UV microdisk lasers increased in complexity, including studies on the optical pumping geometry,²⁰³ directional lasing emissions,^196,204,205 electrical injections,¹⁹⁶ foreign substrates,¹⁹⁷198.199.^–200,206207.208.209.210.^–211 different crystal planes,²⁰⁸209.^–210 quantum structures,^200,206207.208.209.210.^–211 and on new material systems, such as GaN/AlN and AlN/AlGaN.^199,200,208209.210.^–211 During this technological progress, smaller lasing thresholds and single-mode lasing were achieved, room temperature operation became the standard, and a UV-C lasing emission [center wavelength $({\lambda }_c)=275\mathrm{nm}$] became a reality.^200,211 Nonetheless, this progress comes with trade-off between different device parameters i.e., the first UV-C microdisk laser required a high pumping energy of $\mathrm{17,000}\mathrm{kW}/{\mathrm{cm}}^2$, whereas the most recent UV-A single mode microdisk laser emitting at 379.25-nm wavelength needs only $225\mathrm{kW}/{\mathrm{cm}}^2$. In Table 11, we illustrate comprehensive historical progress of these III-nitrides-based microdisks UV emitters.

Table 11

Historical progress in III-nitrides-based microdisk for UV photonics.

In addition to efforts mentioned above, Kneissl et al.¹⁹⁶ demonstrated electrical injection of a spiral-shaped microdisk [Fig. 30(c)] on an InGaAlN laser heterostructure, obtaining a room-temperature directional radial/WG-multimode lasing at 399 nm. So far, this is the only work of an electrically pumped microdisk laser with an emission wavelength in the UV spectral range below 400 nm. Choi et al.¹⁹⁷ demonstrated the first freestanding InGaN/GaN microdisks on Si by selectively dry-etching the substrate [Figs. 30(e) and 30(f)]. This enabled a bottom GaN/air interface instead of GaN/Si interface, creating strain relaxation and improved material properties. Such a microdisk showed WG lasing emission at $\sim 365\mathrm{nm}$ at a 4.3-K temperature and laid the demonstration of III-nitrides on Si platform. Next, Zhang et al.¹⁹⁹ demonstrated a material system (AlN/AlGaN-on-Si) obtaining a single-mode lasing in the UV-B spectral range, being the only report of microdisks in such material system till now. Additionally, they study the structure by FDTD simulation of the WG resonant modes [Figs. 30(i) and 30(j)], obtaining mode matching between the experimental and the simulated results. Mexis et al.²⁰⁷ and Bürger et al.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.

4.6.2.

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 12

III-nitrides-based top-dow-etched nanopillars and nanowires for UV nanophotonics.

A team composed of Li, Xu, Wright et al. utilizing a combination of dry-etching and anisotropic wet etching processes fabricated a set of GaN NWs [Figs. 31(c) and 31(d)].²¹⁵216.^–217,219 These NWs were then isolated and used in several configurations to demonstrate UV-A lasing. First, a single NW ($Ø=0.14\mu \mathrm{m}$, $L=4.7\mu \mathrm{m}$) 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 ${\lambda }_c=370\mathrm{nm}$ [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

Fig. 31

Development of nanopillars and NWs as UV-light lasing structures. (a) GaN NW pair (1, 2, 3, and 4 indicate the facets) formed by NW A and NW B as seen in (b), with their corresponding lasing spectra on (e). Reproduced with permission from Ref. 216. (c) and (d) are dry-etched GaN NWs after 6 and 9 h of additional wet etching. Reproduced with permission from Ref. 215. (f) An array of InGaN/GaN nanopillars before and after (g) Al metal coating. Reproduced with permission from Ref. 214. (h) SEM image of vertical-emitting UV lasing nanopillars with the bottom and upper DBR, and the inset shows a zoom-in SEM image. Reproduced with permission from Ref. 213. (i) and (j) The FDTD-simulated electric field of the WG modes in a micropillar, showing the different radius dependence for light at 365 and 373 nm, respectively. Reproduced with permission from Ref. 214. (k) and (l) The calculated electric field distribution of a nanopillar without and with a metal mirror (Al/SiN) coating. Reproduced with permission from Ref. 218.

Due to the nature of this coupling effect, the lasing threshold was increased to $874\mathrm{kW}/{\mathrm{cm}}^2$; this is about three times higher than the lasing threshold of a single NW. Similarly, a single mode lasing in a discrete GaN NW ($Ø=0.35\mu \mathrm{m}$, $L=5.3\mu \mathrm{m}$) 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 $276\mathrm{kW}/{\mathrm{cm}}^2$. Ultimately, a GaN NW ($Ø=0.2\mu \mathrm{m}$, $L=5.0\mu \mathrm{m}$) 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 $\sim 300\mathrm{kW}/{\mathrm{cm}}^2$. 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 and Hsu et al.²¹⁸ demonstrated the use of Al thin-film coating on nanopillars [Figs. 31(f) and 31(g)], with the goal of providing UV optical confinement in the cavity 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 $420\mathrm{kW}/{\mathrm{cm}}^2$. 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 (${\lambda }_{c1}=365\mathrm{nm}$, ${\lambda }_{c2}=373\mathrm{nm}$). The emission of the peak at 365 nm belongs to the first-order WG, which is exposed to the walls of the nanopillar being prone to losses [Fig. 31(i)]; whereas, the mode corresponding to ${\lambda }_{c2}=373\mathrm{nm}$ is a higher order WG mode and therefore it remains within the nanopillar avoiding the losses encountered by the first-order mode [Fig. 31(j)]. Hsu et al.²¹⁸ studied numerically and experimentally the role of the metal coating in maintaining the WG modes confined within the cavity. As a result of this approach, a two-fold cavity confinement enhancement is achieved as compared with the nanopillars without the metal coating [Figs. 31(k) and 31(l)]. Additionally, they demonstrated that nanopillars with sidewall angles nearer to 90 deg are preferred to improve the cavity quality. Chen et al.²¹³ suggested another approach by utilizing a AlInGaN-based MQW active region existing between a bottom AlGaN-based DBR and a top ${\mathrm{HfO}}_2/{\mathrm{SiO}}_2$-based DBR to create an enhanced vertical-cavity [Fig. 31(h)].

4.6.3.

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 (${\lambda }_c\sim 365\mathrm{nm}$) created from the combination of the ring-shaped 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 ${\mathrm{SiO}}_2$ and 60 nm Al [Fig. 32(b)].²²³ Here, they observed lasing at ${\lambda }_c=370\mathrm{nm}$ under a relatively low optical power density threshold of $0.042\mathrm{kW}/{\mathrm{cm}}^2$, 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.

Fig. 32

Set of unique III-nitrides-based UV-lasing micro and nanostructures. (a) Metal-coated nanoring. Reproduced with permission from Ref. 222. (b) Metal-coated nanostripe. Reproduced with permission from Ref. 223. (c) Hemispherical microcavity on Si post. Reproduced with permission from Ref. 224. (d) Micropyramid on a metal substrate. Reproduced with permission from Ref. 225. (e) Metal-coated nanospiral with calculated electric field distribution (f). Reproduced with permission from Ref. 226.

An innovative approach is presented by Wang et al. demonstrating high $Q$ ($>6000$) WG-mode lasing at room temperature in a GaN micropyramid (${\lambda }_c=367.2\mathrm{nm}$, lasing threshold = 400 to $500\mathrm{kW}/{\mathrm{cm}}^2$).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 ${\lambda }_c=381\mathrm{nm}$. 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 ${\lambda }_c=363.5\mathrm{nm}$ was demonstrated at a low lasing threshold of 0.017 and $0.016\mathrm{kW}/{\mathrm{cm}}^2$, 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.

4.7.1.

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 ($E\parallel c$). Consequently, as shown in Fig. 33(a), only weak light signal of polarization $E\perp c$ can be extracted through the escaping cone. The remaining light signals (entire $E\parallel c$ and part of $E\perp c$ polarization) are captured within the LED and propagate parallel to the Al layer. The $E\parallel c$ 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.

Fig. 33

(a) Schematic of the SP-enhanced AlGaN-based UV-LED. AFM images for the top surfaces of the (b) UV-LED and (c) Al layer. (d) PL spectra of the structures. Reproduced with permission from Ref. 230.

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 $\sim 23\%$, the Ag plasmonic resonance matches the emission wavelength of the GaN QW. In Fig. 34(c), the SP-enhancement 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.

Fig. 34

(a) PL spectrum and carrier lifetime with a schematic of AlGaN/GaN QW shown in the inset. (b) Change of the SP energy of the Ag layer with Al content of the AlGaN barrier layer. (c) Comparing the PL of the metal-free AlGaN/GaN QW and with other metals. Reproduced with permission from Ref. 231.

4.7.2.

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.

Fig. 35

(a) Schematic of the SP-enhanced UV-LED. (b) Top-view SEM image after Al deposition. (c) The output power as a function of the injection current for the LEDs with and without Al layer. Reproduced with permission from Ref. 232.

4.7.3.

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.

4.7.4.

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.

Fig. 36

(a) Schematic of the AlGaN/GaN MQW with Al nanorings. (b) Steps of Al nanorings deposition above the MQW. (c) AFM image for the nanorings above the MQW when using 300-nm silica nanospheres. (d) PL enhancement when measured at 300-K temperature. Reproduced with permission from Ref. 234.

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.

Table 13

Summary of SP-enhanced UV-photonic devices.