Toward ultimate efficiency: progress and prospects …...Toward ultimate efficiency: progress and...

Toward ultimate efficiency:progress and prospects onplanar and 3D nanostructurednonpolar and semipolar InGaNlight-emitting diodesYUJI ZHAO,1,* HOUQIANG FU,1 GEORGE T. WANG,2 AND

SHUJI NAKAMURA3

1School of Electrical, Computer, and Energy Engineering, Arizona State University, Tempe,Arizona 85287, USA2Sandia National Laboratories, Albuquerque, New Mexico 87185, USA3Materials Department, University of California, Santa Barbara, California 93106, USA*Corresponding author: [email protected]

Received October 3, 2017; revised January 2, 2018; accepted January 2, 2018; publishedFebruary 16, 2018 (Doc. ID 308260)

Nonpolar and semipolar III-nitride-based blue and green light-emitting diodes (LEDs)have been extensively investigated as potential replacements for current polar c-planeLEDs. High-power and low-efficiency-droop blue LEDs have been demonstrated onnonpolar and semipolar planes III-nitride due to the advantages of eliminated or reducedpolarization-related electric field and homoepitaxial growth. Semipolar �202̄1� and�202̄ 1̄� LEDs have contributed to bridging “green gap” (low efficiency in green spectralregion) by incorporating high indium compositions, reducing polarization effects, andsuppressing defects. Other properties, such as low thermal droop, narrow spectral line-width, small wavelength shift, and polarized emission, have also been reported for non-polar and semipolar LEDs. In this paper we review the theoretical background, deviceperformance, material properties, and physical mechanisms for nonpolar and semipolarIII-nitride semiconductors and associated blue and green LEDs. The latest progress ontopics including efficiency droop, thermal droop, green-gap, and three-dimensionalnanostructures is detailed. Future challenges, potential solutions, and applications willalso be covered. © 2018 Optical Society of America

OCIS codes: (250.0250) Optoelectronics; (230.3670) Light-emitting diodes;(120.2040) Displays; (130.5990) Semiconductors; (160.4670) Optical materialshttps://doi.org/10.1364/AOP.10.000246

1. Introduction. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2482. Nonpolar and Semipolar GaN . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 249

2.1. Crystal Orientation and Polarization . . . . . . . . . . . . . . . . . . . . . . . 2502.2. Energy Band Structure . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 251

246 Vol. 10, No. 1 / March 2018 / Advances in Optics and Photonics Review

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https://doi.org/10.1364/AOP.10.000246

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2.3. Substrates for Nonpolar and Semipolar GaN . . . . . . . . . . . . . . . . . 2522.4. Epitaxial Growth of Nonpolar and Semipolar GaN by MOCVD and

MBE . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2553. Blue Nonpolar and Semipolar InGaN LEDs . . . . . . . . . . . . . . . . . . . . . 260

3.1. InGaN Blue LEDs Based on m-Plane and Low-Inclination-AngleSemipolar Planes . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2613.1a. m-Plane InGaN Blue LEDs . . . . . . . . . . . . . . . . . . . . . . . . . 2613.1b. �112̄2� InGaN Blue LEDs . . . . . . . . . . . . . . . . . . . . . . . . . 2623.1c. �101̄ 1̄� InGaN Blue LEDs . . . . . . . . . . . . . . . . . . . . . . . . . 263

3.2. InGaN Blue LEDs Based on High-Inclination-Angle Semipolar Planes 2653.2a. �202̄ 1̄� InGaN Blue LEDs . . . . . . . . . . . . . . . . . . . . . . . . . 2653.2b. �303̄ 1̄� InGaN Blue LEDs . . . . . . . . . . . . . . . . . . . . . . . . . 266

3.3. Thermal Droop . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2673.4. Physical Mechanisms of Efficiency Droop . . . . . . . . . . . . . . . . . . . 2703.5. Auger Recombination. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 270

3.5a. Carrier Leakage . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2723.5b. More Mechanisms for Efficiency Droop . . . . . . . . . . . . . . . . 273

3.6. Modified ABC Model . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2744. Green Semipolar InGaN LEDs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 276

4.1. Material Properties of Semipolar �202̄1� and �202̄ 1̄� Plane for GreenEmission . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2774.1a. High Indium Incorporation . . . . . . . . . . . . . . . . . . . . . . . . . 2784.1b. Reduced Polarization-Related Electric Field . . . . . . . . . . . . . . 2784.1c. Low Active-Region Growth Rate . . . . . . . . . . . . . . . . . . . . . 2804.1d. Surface Morphology . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 282

4.2. High-Performance Green Semipolar InGaN LEDs. . . . . . . . . . . . . . 2824.2a. �202̄1� Green LEDs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2834.2b. �202̄ 1̄� Green LEDs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2854.2c. Device Performance Comparison . . . . . . . . . . . . . . . . . . . . . 285

5. Three-Dimensional Nonpolar and Semipolar Nanostructuresfor InGaN LEDs . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2865.1. Nonpolar Core–Shell Nanostructures. . . . . . . . . . . . . . . . . . . . . . . 2865.2. Semipolar Core–Shell Nanostructures . . . . . . . . . . . . . . . . . . . . . . 290

6. Summary and Outlook . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 292Funding . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 293Acknowledgment . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 293References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 293

Review Vol. 10, No. 1 / March 2018 / Advances in Optics and Photonics 247

Toward ultimate efficiency:progress and prospects onplanar and 3D nanostructurednonpolar and semipolar InGaNlight-emitting diodesYUJI ZHAO, HOUQIANG FU, GEORGE T. WANG, AND

SHUJI NAKAMURA

1. INTRODUCTION

III-nitride semiconductors, i.e., wurtzite (In, Ga, Al)N, are the cornerstones of blueand green light-emitting diodes (LEDs) and laser diodes (LDs), which lay thefoundation for modern solid-state lighting, full-color displays, traffic signals, visiblelight communication (VLC), and so on [1]. The successful demonstration of thefirst bright blue InGaN LEDs on sapphire substrates in the early 1990s [2] sparkedintense research on III-nitride-based materials and optoelectronic devices, the successof which led to the 2014 Nobel Prize in Physics. Currently, commercial blue LEDsare grown on polar c-plane sapphires. Due to the resulting c-plane GaN crystalorientation, the c-plane LEDs suffer from strong polarization-related electricfields inside the active regions, which suppress radiative recombination, reduce theLED efficiency, and lead to the quantum confined Stark effect (QCSE) [3].

Waltereit et al. [4] first proposed that growing devices along nonpolar orientations ofGaN could improve device performance by eliminating polarization and the associatedelectric fields. In the early 2000s, the growth of (101̄0) m-plane GaN quantum wells(QWs) on foreign substrates was demonstrated. However, these m-plane QWs werefound to have high densities of threading dislocations (TDs) and basal plane stackingfaults (BPSFs) [5,6]. Later, by using m-plane bulk GaN substrates, efficient violet andblue LEDs were demonstrated with external quantum efficiency (EQE) of 38.9% and16.8%, respectively [7,8]. Despite this encouraging progress, material challenges, such asanisotropic cracking, pyramidal hillocks, and BPSFs, still remain on m-plane LEDs [9].

After 2006, research interest in the III-nitride community soon shifted to semipolarplanes due to the availability of thick (0001) bulk GaN crystals, from which semipolarbulk GaN substrates can be sliced [10,11]. This allows for high-quality homoepitaxialsemipolar GaN epilayers that exhibit reduced polarization-related effects. The firstsemipolar LEDs were reported on �112̄2� bulk GaN substrates with blue, green,and amber emission [12]. Zhao et al. [13] first demonstrated semipolar �101̄1̄� blueLEDs with over 50% EQE. More importantly, it was also revealed that semipolar blueLEDs exhibited significantly less EQE rollover [14–17] at high current densities (“ef-ficiency droop” [18–22]) and operation temperatures (“thermal droop” [23]). In ad-dition, higher indium incorporation in certain semipolar planes was seen as a potentialroute to bridging the “green gap” [24]. Our group recently reported the �202̄1� greenLEDs with a record high EQE of 28.3% at 20 mA. Other promising properties, such aslow wavelength shift, narrow linewidth [25], and polarized emission [26], were alsoreported for green semipolar LEDs.

In this paper, we review recent progress on the development of high-performance blueand green nonpolar and semipolar InGaN LEDs, primarily focusing on performance


gains made on bulk GaN substrates, while also summarizing recent research efforts onthree-dimensional (3D) nanostructures with nonpolar and semipolar facets on sap-phire for LEDs. Following this introduction section, Section 2 discusses the basicphysics and theoretical backgrounds for nonpolar and semipolar QW structure, wherepolarization and energy band structures are calculated. Section 3 presents the deviceperformance of high-efficiency low-droop nonpolar and semipolar blue LEDs, andefficiency droop and thermal droop are included. The physical mechanisms behindthe efficiency droop and thermal droop are also discussed, including Auger recombi-nation, carrier leakage, modified ABC model, and so on. Section 4 investigates thepotential of semipolar planes in bridging the “green gap.” Topics such as high indiumincorporation, low growth rate, low wavelength shift, and narrow linewidth arediscussed. Section 5 overviews the development of three-dimensional nonpolarand semipolar nanostructures. Finally, Section 6 concludes the paper and discussesnecessary work on these topics in the future.

2. NONPOLAR AND SEMIPOLAR GAN

The nonpolar and semipolar planes of wurtzite GaN crystals have gained considerableattention due to the potential for high-efficiency, low-droop LEDs as a result ofreduced or eliminated detrimental polarization-related effects. A rather flat QW profilecan be obtained on these nonpolar and semipolar planes, which increases the radiativerecombination rate and mitigates the QCSE. In addition, studies have also shown thatLEDs grown on nonpolar and semipolar planes have other beneficial properties, suchas large critical thickness, high indium incorporation, polarization emission, lowgrowth rate, and narrow linewidth. These properties indicate very different materialand optical properties compared with c-plane devices. Figure 1 shows the polarc-plane, nonpolar m-plane, and several widely used semipolar planes of the III-nitridewurtzite crystal structure.

Figure 1

Schematics of polar (c-plane), [semipolar �112̄2�, �101̄ 1̄�, �202̄1�, and �202̄ 1̄�], andnonpolar plane (m-plane) of III-nitride wurtzite crystal structure. The degrees indicatethe inclination angles of the nonpolar and semipolar planes from the c-plane.


2.1. Crystal Orientation and PolarizationDue to the lack of inversion symmetry, wurtzite III-nitride materials have strong spon-taneous polarization along the [0001] c-axis [27]. When an InGaN layer is coherentlygrown on a GaN template, the lattice-mismatch-induced anisotropic strain will resultin piezoelectric polarization. The combination of the spontaneous polarization andpiezoelectric polarization differences at the InGaN/GaN interface induce strongelectric fields inside the InGaN QWs, which result in significant energy band tiltingand QCSE [3,4,27]. This tilted InGaN QW profile reduces the electron–hole wave-function overlap and thus decreases the radiative recombination rate and the LEDefficiency. The polarizations of the c-plane and nonpolar/semipolar planes can be cal-culated based on anisotropic linear elasticity and some simple algebra [27]. Figure 2shows the two coordination systems x–y–z and x0–y0–z0. The unprimed z-axis repre-sents the c-axis of the wurtzite structure and the x- and y- axes are in the basal c-plane.The primed x0-axis and y0-axis are in the substrate surface plane and the z0-axis is alongthe growth direction of a certain plane. The angle θ between the z-axis and z0-axis isdefined as the inclination angle of a certain plane from the c-plane.

When a certain plane is inclined from the c-plane by an angle of θ, the totalpolarization difference ΔPtot along z0 at the InGaN/GaN interface can be expressedas [26,27]

ΔPtot � PInGaNpz � �PInGaN

sp − PGaNsp � cos θ; (1)

Figure 2

Two sets of coordination systems x–y–z and x0–y0–z0 for calculation of strain-inducedpolarization in (a) single InGaN layer and (b) QWs on a GaN template. In x–y–z,[0001] is defined as the z-axis and [12̄10] is chosen as the x-axis for simplicity.In x0–y0–z0, the z0-axis is perpendicular to the template’s surface and the x0-axis servesas the rotation axis for the [0001] inclination. Reprinted with permission from [26].Copyright 2014 The Japan Society of Applied Physics. Reprinted with permissionfrom Romanov et al., J. Appl. Phys. 100, 023522 (2006) [27]. Copyright 2006AIP Publishing LLC.


where PInGaNsp and PGaN

sp are the spontaneous polarization of the InGaN layer and GaNtemplate, respectively. PInGaN

pz is the strain-induced piezoelectric polarization in theInGaN layer, which be expressed as [26,27]

PInGaNpz � e31 cos θεx0x0 �

�e31 cos

3 θεx0x0 �e33 − e15

2sin θ sin 2θ

�εy0y0

��e33 � e15

2sin θ sin 2θ� e33 cos

3 θ

�εz0z0

� ��e31 − e33� cos θ sin 2θ� e15 sin θ cos 2θ�εy0z0 ; (2)

where elements εk 0m0 are the strain tensor components and elements eij are the com-ponents of piezoelectric tensor in Voigt notation. Figure 3 presents the PInGaN

pz andΔPtot of InGaN/GaN with different indium compositions as a function of θ. Sincethe PInGaN

pz is dominant for the InGaN/GaN heterostructure, ΔPtot shows the minimalchange with the addition of spontaneous polarization. ΔPtot becomes zero at θ � 45°(semipolar plane) and θ � 90° (nonpolar plane). The two crossovers are almostnot impacted by the indium composition. The polarization (absolute value) ofseveral common GaN crystal planes are compared as follows: c-plane�0001� > �101̄ 1 �̄ > �202̄1� � �202̄ 1 �̄ > �112̄2�.

2.2. Energy Band StructureThe band structure and electron/hole wave function of QWs on different planes can becalculated using the commercial software SiLENSe developed by the STR Group[28–30] where a one-dimensional Schrödinger-Poisson equation is solved self-consistently with a drift-diffusion model included. The software calculates spontane-ous and piezoelectric polarizations for arbitrary crystal orientations of III-nitridematerials, which are critical for device performance of InGaN LEDs. Figure 4 presentsthe energy band structures of In0.2Ga0.8N�3 nm�∕GaN�15 nm� single quantum well(SQW) LEDs at 100 A∕cm2. For the c-plane LED, Ebi and Epz are antiparallel and Epz

is much larger than Ebi, which results in the QW profile tilted downward along thegrowth direction. Epz of semipolar �101̄1̄� and �202̄1̄� LEDs are also in the oppositiondirection of Ebi but smaller, which results in less QW tilting compared with c-plane

Figure 3

Calculated (a) piezoelectric polarization PInGaNpz and (b) total polarization difference

ΔPtot as a function of semipolar plane orientation θ for InGaN/GaN heterostructurewith In composition of 10%, 20%, 30%, and 40%. (a) Reprinted with permission from[26]. Copyright 2014 The Japan Society of Applied Physics. (b) Reprinted with per-mission from Romanov et al., J. Appl. Phys. 100, 023522 (2006) [27]. Copyright2006 AIP Publishing LLC.


QWs. Semipolar �112̄2� and �202̄1� LEDs have the Epz in the same direction of Ebi,leading to QW tilting upward. For the m-plane LED, Epz is zero and only Ebi existsinside the InGaN QW.With positive bias to achieve current density of 100 A∕cm2, theQW profile is almost flat. Due to the tilted QW profile, the energy difference betweenelectron and hole wave function for the c-plane is smaller than that of nonpolar andsemipolar LEDs. With increasing current density, c-plane LEDs have decreased QWtilting due to the Coulombic screening effect and show a blueshift in the emissionwavelength. However, nonpolar and semipolar LEDs with reduced polarization areless affected by increasing injection current. In terms of electron and hole wave-function overlap, the wave function of the c-plane LEDs is more spatially separatedthan nonpolar and semipolar LEDs, leading to a smaller radiative recombination rateand possibly larger efficiency droop.

2.3. Substrates for Nonpolar and Semipolar GaNCurrently, most high-performance nonpolar and semipolar LEDs are grown on bulkGaN substrates due to much lower dislocation densities compared to heteroepitaxial

Figure 4

Simulated band diagram for In0.2Ga0.8N�3 nm�∕GaN�15 nm� QW grown on (a) polarc-plane; semipolar (b) �112̄2�, (c) �101̄ 1̄�, (d) �202̄1�, (e) �202̄ 1̄�; and (f) nonpolarm-plane at current density of 100 A∕cm2. From left to right, it is n-GaN/i-InGaN/p-GaN. The directions of junction built-in electric field Ebi and piezoelectric electricfield Epz are indicated by arrows whose thickness represents the relative magnitudeof Ebi and Epz.


GaN films grown on foreign substrates. Typically, bulk c-plane GaN are usuallygrown by hydride vapor phase epitaxy (HVPE) [10] and ammonothermal growth[31] methods. The thick GaN crystal is then sliced at a specific angle to exposethe desired nonpolar or semipolar orientations. However, the dimensions of thesenonpolar/semipolar substrates are limited by the thickness of GaN boules, usually∼1 cm. As their small sizes and high costs limit both research development and com-mercialization efforts in universities and industries, methods to grow larger and lessexpensive nonpolar and semipolar GaN are greatly sought after.

Nonpolar and semipolar GaN have also been grown using epitaxial lateral overgrowth(ELOG) on planar sapphire substrates. In ELOG, a SiO2 or SiNx mask is patternedinto a stripe geometry with window regions. The GaN grows out of the window regionand coalesces laterally over the mask region, which blocks vertically propagating TDs[32]. Ni et al. [33] reported the ELOG growth of (112̄2) semipolar GaN on m-planesapphire substrates. Figure 5 shows the scanning electron microscopy (SEM) imagesof two ELOG samples with different SiO2 strip orientations [33]. It is clear that maskorientations play a crucial role in determining the growth dynamics and surface mor-phologies. Figure 5(b) shows the formation of c-plane and a-plane GaN facets, whileFig. 5(d) exhibits two {101̄1} GaN facets. An enhanced version of ELOG involves thetransfer of strip pattern into the underlying GaN buffer layer via drying etching beforethe epitaxial lateral overgrowth, termed as sidewall lateral epitaxial overgrowth(SLEO). Figure 6 presents the sample processing and growth stages for the SLEO[34]. During the initial growth stages of the SLEO, growth conditions were optimizedto ensure a high lateral to vertical growth rate ratio. GaN grows laterally from theexposed vertical sidewall, blocking the incorporation of vertical dislocations origi-nated from the GaN buffer layer. Then it coalesces with the opposite sidewall,and then grows out of the trench region. The rest of the growth stages are similarto those of the conventional ELOG. Though these methods can improve the epilayerquality of GaN above the SiO2 strips, a large density of TDs and BPSFs still remain inGaN epilayers in the window regions and coalescences regions. These result inintrinsic nonuniformity issues in material properties and device performance

Figure 5

(a),(b) Plane and cross-section SEM images of ELOG growth for (112̄2) semipolarGaN with SiO2 strips oriented along [12̄10] a-axis of sapphire. (c),(d) Plane and cross-section SEM images of ELOG growth for (112̄2) semipolar GaN with SiO2 stripsoriented along [0001] c-axis of sapphire. Reprinted with permission from Ni et al.,Appl. Phys. Lett. 90, 182109 (2007) [33]. Copyright 2007 AIP Publishing LLC.


[32–34]. In addition, only very few semipolar orientations are commercially available,i.e., (112̄2) and (101̄3) [35,36].

To improve the conventional ELOG on planar sapphire, the selective growth of semi-polar GaN on patterned sapphire substrates (PSSs) has recently been demonstrated,including (101̄0), (101̄1), (112̄0), (112̄2), and �202̄1� GaN [37–42]. In theory it ispossible using this approach to access any orientation [43], i.e., any nonpolar andsemipolar plane, as long as the patterned sapphire sidewall is precisely inclined.Some semipolar GaN orientations on sapphire substrates, such as �202̄1�, can onlybe achieved on PSSs [35,41], while (303̄1) GaN is only available in the form of smallbulk substrates by slicing thick c-plane GaN substrates [16]. The main challenge is toetch a specific grooved pattern having a sidewall with a particular inclination angle, onwhich the semipolar GaN will be grown.

In 2011, Okada et al. [41] realized the growth of �202̄1� GaN on PSSs. Recently, Hanand co-workers [43,44] have reported �202̄1�GaN on 2 in. PSSs (Fig. 7) with low TDsand BPSFs. InGaN/GaN QW structures on these �202̄1� GaN-on-sapphire templateswere also demonstrated [44]. While this method is still in its infancy and nohigh-performance LEDs have been demonstrated on these semipolar GaN-on-sapphire templates to date, it still represents a promising approach toward enabling

Figure 6

(a) Sample processing steps before SLEO. Traditional ELOG sample processing stopsat SiO2 etching step. (b) Schematic view of the growth stages of SLEO. Reprintedwith permission from Imer et al., Appl. Phys. Lett. 88, 061908 (2006) [34]. Copyright2006 AIP Publishing LLC.


commercial semipolar LEDs due to their relatively inexpensive cost and large sizecompared to nonpolar and semipolar bulk GaN substrates. In addition, some otherforeign substrates, such as SiC and Si, have also been explored for the growth ofnonpolar and semipolar GaN. However, besides high TD density, several other dis-advantages render them not suitable for the commercialization of InGaN LEDs[32,35]. For example, SiC is still too expensive; it is extremely difficult to grow non-polar and semipolar GaN in particular on Si due to the lack of epitaxial relation matchand the melt-back etching effect. Finally, in addition to the growth of planar nonpolarand semipolar GaN, 3D microstructures and nanostructures with nonpolar and semi-polar facets have been heavily studied recently. These efforts are summarized inSection 5. In light of the above, this review will primarily focus on recent advancesusing on nonpolar and semipolar LEDs using bulk GaN substrates and 3Dnanostructures.

2.4. Epitaxial Growth of Nonpolar and Semipolar GaN by MOCVD and MBEThere are two major methods for epitaxially growing III-nitride films and devices:metalorganic chemical vapor deposition (MOCVD) and molecular beam epitaxy(MBE). Planar InGaN LEDs are mainly grown by MOCVD (Sections 3 and 4), whilethe 3D nanostructured InGaN LEDs can be grown by both MOCVD and MBE(Section 5). HVPE is often used to produce bulk GaN substrates [10] and seldomused to grow InGaN LEDs. MOCVD and MBE growth of III-nitrides differ in severalaspects. First, the MOCVD method is the industrial standard for growing III-nitridedevices because of its faster growth rate and lower cost compared with the MBEmethod and the multi-wafer capability, while the MBE method is primarily foundin university research labs. Second, another significant difference is the growth tem-perature and pressure. For example, the GaN growth temperature in MOCVD is usu-ally ∼1000°C–1100°C under atmosphere pressure or low pressure (30–760 Torr) [1],while GaN is grown by MBE at only 700°C–800°C under ultra-high vacuum(10−5–10−11 Torr) [4]. Therefore, MBE-grown GaN usually has worse material qual-ity than MOCVD-grown GaN due to the lower surface mobility caused by the lower

Figure 7

(a) Image of 2 in. �202̄1� GaN on sapphire. (b) Schematic view of the growth process.Growth evolution after (c) 300 nm and (d) 8 μm equivalent thickness of planar growth.Reprinted with permission from Leung et al., Appl. Phys. Lett. 104, 262105 (2014)[44]. Copyright 2014 AIP Publishing LLC.


growth temperature. Third, the MBE method can grow very sharp interfaces betweendifferent layers, while the MOCVD process suffers from the gas flow issues andmemory effects that usually lead to less abrupt interfaces. Other advantages of theMBE process include the availability of in situ monitoring of the device growthand no need to deal with toxic metalorganic precursors. In addition to the differencesbetween MOCVD and MBE GaN growth in general, they also face different materialand growth issues as discussed in the following.

In the MOCVD growth, metalorganic precursors are carried by N2 or H2 or (N2 � H2)into the reactor and react on the heated surfaces of the substrates. Trimethylgallium(TMG) or triethylgallium (TEG), trimethylindium (TMI), trimethylaluminum (TMA)are the sources for Ga, In, and Al, respectively. Ammonia (NH3) is used as the sourcefor nitrogen. Typical V/III ratio for the GaN growth is 1000–3000. The n-type dopantis usually Si from silane (SiH4) and the p-type dopant is usually Mg from bis(cyclo-pentadienyl)magnesium (Cp2Mg). The MOCVD growth of nonpolar and semipolarGaN presents special issues such as surface pyramidal hillocks [45–47] and BPSFs[6]. GaN films grown on nominally on-axis m-plane GaN substrates usually exhibitsurfaces with a high density (∼3 × 103 cm−2) of shallow four-sided pyramidal hill-ocks [Fig. 8(a)], whose dimensions range from 50 to 250 μm [46]. In the pyramidalhillock, two faces are inclined toward the��112̄0� a-directions, one toward the [0001]c-direction, and one toward the [0001̄] c−-direction [Fig. 8(b)]. The panchromatic

Figure 8

(a) Nomarski optical image of m-plane GaN surface. (b) Schematic view of thepyramidal hillocks with four sides. (c) Panchromatic CL image of the pyramidal hill-ock. The inset shows the zoom-in view of the apex of the pyramidal hillock.(d) Optical images for m-plane films with different substrate misorientation anglestoward the [0001̄] c−-direction. (e) RMS roughness of 10 μm × 10 μm AFM imagesversus substrate misorientation angles for m-plane films. (a)–(c) Reprinted with per-mission from Farrell et al., Appl. Phys. Lett. 96, 231907 (2010) [46]. Copyright 2010AIP Publishing LLC. (d),(e) Reprinted from J. Cryst. Growth 313, Farrell et al.,“Effect of carreier gas and substrate misorientation on the structural and optical prop-erties of m-plane InGaN/GaN light-emitting diodes,” 1–7 [47]. Copyright 2010, withpermission from Elsevier.


cathodoluminescence (CL) image [Fig. 8(c)] of the pyramidal hillocks shows almostuniform distribution of TDs (as indicated by the small black dots in the image) on thefour faces, with a TD density of ∼6 × 106 cm−2. In addition, the apex of each pyrami-dal hillock hosts a single TD [the inset of Fig. 8(c)], indicating that dislocations mayfacilitate the nucleation of these pyramidal hillocks. The pyramidal hillocks can besuppressed or eliminated using m-plane substrates with small misorientation anglestoward the [0001̄] c−-direction [Figs. 8(d) and 8(e)]. Smooth surfaces without pyrami-dal hillocks were obtained when the misorientation angles>0.7° [47]. However, whenthe misorientation angle becomes larger than 1.5°, lateral striations parallel to the[112̄0] a-direction occurred due to the step bunching. This trend is also reflected inthe high root-mean-square (RMS) roughness of the atomic force microscopy (AFM)images [Fig. 8(e)]. The density of the pyramidal hillocks formation at differentmisorientation angles relates to the competition between the step-flow growth andthe spiral growth. At a certain misorientation angle, the step-flow growth is favoredover the spiral growth, which results in atomically flat surfaces.

Defect formation in MOCVD also plays a critical role in determining the device per-formance of InGaN LEDs, among which BPSFs in the nonpolar GaN LEDs and misfitdislocations (MDs) in the semipolar GaN LEDs are of particular importance [6,9,48].Figures 9(a)–9(d) clearly show the existence of BPSFs in the m-plane InGaN/GaNQWs from the transmission electron microscopy (TEM) plan-view images [6].These BPSFs were parallel to the [112̄0] direction with a length of ∼100 nm.Two partial dislocations defined the boundaries of each BPSF and no additionalprimary TDs were observed away for the BPSFs. The density of BPSFs was104–105 cm−1 and the density of partial dislocations was ∼109 cm−2. Furthermore,BPSFs generation can also be identified in the QW by high-resolution TEM(HRTEM) images, as shown in Fig. 9(d), where a (0001) half-plane was missing.The formation mechanism of BPSFs is still under investigation, but it is thoughtthat the BPSF formation can help release some compressive stress. And the BPSF

Figure 9

(a)–(c) Plan-view TEM images of m-plane InGaN/GaN QWs. (a) Two-beam bright-field image and (b) multi-beam bright-field image under the g � 01̄10 diffraction con-dition. (c) Dark-field image under the g � 03̄33 diffraction condition. (d) HRTEMimage of m-plane InGaN/GaN QWs with a BPSF generated in a QW. The inset showsthe filtered inversed FFT image that only includes 0002, 0000, and 0002̄. (e) HRTEMimage of a single SF with an I1 type stacking sequence. The arrow shows the missing(0001) half-plane. (f) Schematic view of a SF generated in a QW. Reprinted withpermission from Wu et al., Appl. Phys. Lett. 96, 231912 (2010) [6]. Copyright2010 AIP Publishing LLC.


formation is not related to the existing TDs in the substrates and the plane slip due tothe absence of resolved shear stress for the m-plane GaN [Fig. 10(a)]. Semipolar GaNcan have stress relaxation through the formation of MDs due to resolved shear stress[48]. As shown in Fig. 10(a), the resolved shear stress reaches the maximum at aninclination angle of ∼45°. The resolved shear stress will act upon a pre-existing TDand the TD will glide in the basal c-plane, thus forming a MD [Fig. 10(b)]. Forexample, the formed MD line is along the (112̄2) m-direction in the (11̄00) GaN filmand along the [112̄0] a-direction in the (202̄1) GaN film.

In the MBE growth, high-purity III sources such as Ga, In, and Al are heated andevporated from the effusion cells and the nitrogen source comes in the form of eitherNH3 or nitrogen plasma. The III and V beams will condense on the heated substratesand react to form III-nitride films in a layer-by-layer fashion. Si and Mg are used for n-type and p-type doping, respectively. In contrast to the simple substrate pretreatmentsused in the MOCVD growth (usually solvent and acid cleaning under ultrosonic),proper cleaing of the GaN substrate surface is crucial in the MBE growth. The exposedGaN surfaces can quickly absorb impurities, such as oxygen, from the environment,which can affect the thermal decomposition of GaN and the final surface morphologyof the epi films [49]. Storm et al. [49] proposed complicated cleaning steps (up to13 steps) to prepare the GaN substrate surface before the MBE growth. Furthermore,the different reactivities and chemical sensitivities of Ga-polar and N-polar GaN makeit even harder to obtain high-quality homoepitaxial GaN films. The prevalent TDs inthe MBE c-plane GaN film are the edge-type TDs, as shown in Fig. 11 [49]. PlanarInGaN films can also be grown by the MBE method on different crystal orientations[50]. Generally speaking, with increasing growth temperature, better surface morphol-ogies can be obtained. However, the indium composition decreases with the increas-ing temperature. On the other hand, at the same temperature, crystal orientation playsa critical role in the surface morphology accroding to the AFM images (Fig. 12). Forexample, at a low growth temperature of 575°C, the c-plane InGaN film was ex-tremenly rough with a RMS roughness of 8.91 nm, while the semipolar (202̄1) InGaNshowed a smooth surface with a RMS roughness of only 0.81 nm. Nonpolar m-planeInGaN exhibited unique surface features that are similar to the pyrimidal hillocks inthe MOCVD-grown GaN [45–47], but with a higher density. Among the four samples,

Figure 10

(a) Resolved shear stresses on the basal σyz as a function of inclination angle for(1) InGaN/GaN and (2) AlGaN/GaN. s1-s6 indicate commonly used semipolar planes.(b) Schematic view of the MD-TD configuration due to the slide in the (0001) c-planefor a (112̄2) heterostructure. The in-plane hexagonal basis vectors a1, a2, and a3give possible Burgers vectors for inclination axis [11̄00]. Reprinted with permissionfrom Romanov et al., J. Appl. Phys. 109, 103522 (2011) [48]. Copyright 2011 AIPPublishing LLC.


(202̄1) and m-plane had the higest indium composition at the same temperature [50].In addition, there is another way of obtaining high-quality nonpolar and semipolar III-nitrides through 3D nanostructured templates [51,52]. First, ordered arrays ofIII-nitride nanostructures are grown by selective area growth using metal or dieletricnanohole masks [51]. Then, they continue to grow and merge into a continuous film.This process can filter most of the stacking faults (SFs) and result in high-qualitynonpolar and semipolar III-nitride films [52].

Figure 11

Two-beam bright-field TEM images of c-plane GaN grown on HVPE bulk GaN sub-strates by MBE with diffraction vectors g � �112̄0� and g � �0002�. Reprinted fromJ. Cryst. Growth 456, Storm et al., “Critical issues for homoepitaxial GaN growthby molecular beam epitaxy on hydride vapor-phase epitaxy-grown GaN substrates,”121–132 [49]. Copyright 2016, with the permission from Elsevier.

Figure 12

RMS roughness of 10 μm × 10 μm AFM images for InGaN films with various ori-entations at different growth temperatures. The height scale is 10 nm. Reprinted withpermission from Browne et al., J. Vac. Sci. Technol. B 30, 041513 (2012) [50].Copyright AIP Publishing LLC.


3. BLUE NONPOLAR AND SEMIPOLAR INGAN LEDS

High-brightness InGaN-based blue LEDs were first demonstrated using c-plane GaNon sapphire substrates and are the basis for current commercial blue LEDs. The high-est EQE reported to date for c-plane blue LEDs in the lab is 81.3% at 20 mA fromNichia [53]. However, c-plane LEDs suffer from a reduction of efficiency with in-creasing current density, a phenomenon known as “efficiency droop.” To overcomethis issue, which limits the efficiency of today’s high-power LEDs, nonpolar and semi-polar InGaN LEDs, primarily grown on low dislocation density bulk GaN substrates,have been proposed as a solution. Major results on demonstrated blue/violet nonpolarand semipolar LEDs on bulk GaN substrates are summarized in Table 1. Figure 13compares EQE curves of the state-of-the-art (SOTA) c-plane InGaN blue LEDs onsapphire substrates [53], c-plane InGaN blue LEDs on bulk GaN substrates [54],and semipolar �202̄ 1̄� blue LEDs on bulk GaN substrates [15]. At 100 A∕cm2,

Table 1. Device Structures and Performance of Blue/Violet InGaN LEDs(All the Devices are Grown on Bulk GaN Substrates, Except That [53] and [59]are on Sapphire Substratesa

Plane Structure Method Wavelength (nm) LOP (mW) EQE (%) Droop (%) Reference

(0001) 2.5 nm 6 QWs Pulsed 444 30.4 55.7 48.0 [54](0001) 3 nm 6 QWs Pulsed 460 37.0 80.8 18.3 [53]�101̄0� 8 nm 6 QWs DC 407 23.7 38.9 12.3 [7]�101̄0� 4 nm 5 QWs DC 456 0.24 0.43 9.3 [55]�101̄0� 3 nm 5 QWs DC 435 1.8 3.1 23.1 [56]�101̄0� 2 nm 3 QWs Pulsed 460 5.5 10.5 24.5 [57]�112̄2� 3 nm SQW DC 426 1.8 3.0 17.5 [12]�112̄2� 3 nm 6 QWs Pulsed 480 9.0 17.4 23.4 [58]�112̄2� 5 nm 3 QWs – 441 1.0 1.65 71.4 [59]�101̄ 1̄� 3 nm 6 QWs Pulsed 444 16.2 29.0 16.8 [60]�101̄ 1̄� 2.5 nm 6 QWs DC 411 20.6 33.9 9.1 [61]�101̄ 1̄� 3 nm 6 QWs Pulsed 420 22.8 39.5 4.2 [62]�101̄ 1̄� 3 QWs Pulsed 411 31.1 54.7 3.7 [13]�202̄ 1̄� 3 nm 3 QWs Pulsed 423 30.6 52.2 8.5 [14]�202̄ 1̄� 12 nm SQW Pulsed 446 28.2 52.7 4.7 [15]�303̄ 1̄� 15 nm SQW Pulsed 413 29.8 49.5 9.5 [16]aWavelength is peak electroluminescence (EL) wavelength at 20 mA. LOP and EQE are values at20 mA. Droop is calculated at 100 mA, except that [54] is at 75 mA and [12] at 200 mA. DCindicates direct current.

Figure 13

Comparison of EQE as a function of current density for the SOTA sapphire c-planeblue LED [53], a bulk GaN c-plane [54], and bulk GaN semipolar blue LEDs [15].


the SOTA sapphire c-plane LEDs show efficiency droop of 19% and the efficiency ofbulk GaN c-plane LEDs drops over 50%. The semipolar LEDs show much lowerefficiency droop, especially at high current density: an efficiency droop of only 5%at 100 A∕cm2 and 18% at 300 A∕cm2. In this section, we will review nonpolar andsemipolar blue LEDs developed before and after 2010, focusing on their efficiencydroop and thermal droop. Then a modified ABC model is proposed to explain thereduced-droop characteristics of nonpolar and semipolar blue LEDs.

3.1. InGaN Blue LEDs Based on m-Plane and Low-Inclination-Angle Semipolar PlanesBefore 2010, most of the studies aiming to solve efficiency droop using non-c-planecrystal orientations concentrated on nonpolar planes (primarily m-plane) and semipo-lar planes oriented near 45° from the c-plane, such as the �112̄2� and �101̄ 1̄� planes.

3.1a. m-Plane InGaN Blue LEDs

Figure 14 compares several major reports on m-plane violet/blue LEDs. Chakrabortyet al. [55] first demonstrated anm-plane blue LEDwith an output power of<1 mW andan EQE of <1%. The device structure was comprised of a 2.2 μm Si-doped n-GaNbuffer layer, five periods of In0.17Ga0.83N�4 nm�∕GaN�16 nm�multiple quantumwells(MQWs), 0.3 μm Mg-doped p-GaN, and 40 nm p�-GaN. The low performance waspossibly due to a high BPSF density of 105 cm−1 and TD density of 4 × 109 cm−2.Okamoto et al. [56] obtained a low-dislocation m-plane InGaN (3 nm)/GaN (9 nm)MQW LED on the bulk m-plane GaN substrate with a low misorientation of �0.3°.They used a high V/III ratio of >3000, which was a typical condition for obtaininghigh-quality c-plane GaN, though it has been reported low V/III ratio (<1000) is morebeneficial for the lateral growth of nonpolarGaN films [34]. Despite the highV/III ratio,no TDs or SFs were observed in the scanning transmission electron microscope(STEM) images of the devices and smooth surfaces were also obtained. This LEDshowed an output power of 1.79 mW and an EQE of 3.1% at 20 mA. However, withincreased substrate misorientation, the GaN epitaxial films began to become rough be-cause of the dominant island growthmode. Therefore,m-plane substrate misorientationhas strong impact on the material quality and device performance of m-plane LEDs.

Lin et al. [57] investigated the growth and properties of high indium compositionm-plane In0.26Ga0.74N∕In0.02Ga0.98N MQW LEDs with varying QW thickness andquantum barrier thickness. With increasing QW thickness from 2 nm to 4 nm, the

Figure 14

Comparison of EQE as a function of current for the reported m-plane violet/blueLEDs.


emission wavelength and full width at half-maximum (FWHM) of the LEDs increased[Fig. 15(a)], while the output power and EQE of the LEDs first increased from 2 nmQW to 2.5 QW due to improved carrier injection, and then decreased [Fig. 15(b)]. It isstill not clear what caused the increase in emission wavelength and FWHM and thedecrease in the output power and EQE with increasing QW thickness (>2.5 nm). Boththe emission wavelength and the FWHM of the devices deceased with increasingquantum barrier thickness [Fig. 15(c)]. However, the output power and EQE ofthe devices monotonically increased with quantum barrier thickness from 7.5 nmto 20 nm [Fig. 15(d)]. The physical mechanisms behind these results demand furtherinvestigation. For a blue LED with an electroluminescence (EL) peak wavelength of460 nm, the output power of 5.5 mW and an EQE of 10.5% at 20 mAwere reported.Schmidt et al. [7] demonstrated a highly efficient violet m-plane LED (407 nm) with ahigh output power of 23.7 mWand an EQE of 38.9% at 20 mA. The active region wasthe six-period InGaN (8 nm)/GaN (18 nm) MQWs. The growth conditions were basedon the optimized c-plane GaN film growth: atmospheric pressure, V/III ratio >3000,and temperature ranging from 875°C to 1185°C.

3.1b. �112̄2� InGaN Blue LEDs

Despite these improving results for m-plane blue LEDs, the growth challenges such asBPSFs and alloy inhomogeneity still made it very difficult to fabricate high-qualitym-plane QWs. LEDs grown on semipolar planes oriented near 45° from a c-plane suchas �112̄2� and �101̄1̄� were then extensively studied since there is no polarization at45°. Figure 16 shows the EQE curves of several high-performance �112̄2� blue LEDs.In 2006, Funato et al. [12] demonstrated a blue �112̄2� SQW LED with an outputpower of 1.76 mW and an EQE of 3% at 20 mA, which were comparable to theperformance of the m-plane blue LEDs by Okamoto et al. [56]. The LED structure

Figure 15

Peak photoluminescence (PL) and EL wavelength and EL FWHM versus (a) wellwidth and (c) barrier width at 20 mA. Output power and EQE as a function of (b) wellwidth and (d) barrier width at 20 mA. Reprinted with permission from Lin et al., Appl.Phys. Lett. 94, 261108 (2009) [57]. Copyright 2009 AIP Publishing LLC.


comprised 1.5 μm unintentionally doped GaN, 4.5 μm Si-doped n-GaN and InGaN(3 nm)/GaN (25 nm) SQW, Mg-doped p-AlGaN, and Mg-doped p-GaN. Zhong et al.[58] optimized the device structure and realized an output power of 9 mWand an EQEof 17.4% at 20 mA using a six-period InGaN (3 nm)/InGaN (20 nm) MQW activeregion. For comparison, a �112̄2� blue InGaN (5 nm)/GaN (15 nm) MQWLED grownon the sapphire substrate [59] is also shown in Fig. 16. Kim et al. [59] used a directepitaxial lateral overgrowth method with the two-step growth on the SiO2 patterned(11̄00) sapphire substrates without templates. The SiO2 was first deposited by plasma-enhanced chemical vapor deposition and then patterned in parallel to the [112̄0] sap-phire direction. The InGaN LEDs were grown by MOCVD along the [112̄2] GaNdirection. However, the EQE was less than 2% at 20 mA possibly due to the largedefect densities that resulted from the large lattice mismatch between the GaN andsapphire substrates. This again implies the advantages of using bulk GaN substratesfor efficient semipolar InGaN LEDs. Though the �112̄2� LEDs had comparable per-formance with respect to the m-plane LEDs, they still had inferior performance to thec-plane blue LEDs.

3.1c. �101̄ 1̄� InGaN Blue LEDs

�101̄ 1̄� semipolar blue LEDs have exhibited great promise to further enhance effi-ciency, as shown in Fig. 17. Zhong et al. [60] reported semipolar �101̄ 1̄� LEDson low defect density bulk GaN substrates grown by conventional MOCVD. The de-vice structure comprised 1 μm Si-doped n-GaN, six-period InGaN (3 nm)/GaN(20 nm) MQWs, 10 nm AlGaN, and 200 nm Mg-doped p-GaN. This LED showedan output power of 16.2 mW and an EQE of 29% at 20 mA. Tyagi et al. [61] dem-onstrated a packaged violet �101̄ 1̄� LED (wavelength of 411 nm) with an outputpower of 20.6 mW and an EQE of 33.9% at 20 mA, using a six-period InGaN(2.5 nm)/GaN (23 nm) MQW active region.

Zhao et al. [62] performed systematic optimization of the device structures byMOCVD for �101̄ 1̄� LEDs around 450 nm. The device growth started with 1 μmSi-doped n-GaN, followed by three-period InGaN/GaN MQWs active region endingwith Mg-doped p-GaN last barrier (LB), 16 nm Mg-doped p-AlGaN, 50 nm p-GaN,and a 10 nm p�-GaN contact layer. The optimized device parameters included the QWwidth, quantum barrier thickness, the last barrier thickness, and the Mg doping for thelast barrier and p�-GaN contact layer (Fig. 18). The first three parameters stronglyaffected the EL intensity, while they showed little impact on the emission wavelength.

Figure 16

Comparison of EQE as a function of current for the reported �112̄2� blue LEDs.


According to Figs. 18(a) and 18(b), EL intensity peaked at a QW width of 3 nm and aquantum barrier thickness of 17.5 nm. On the one hand, devices with either a toonarrow QW and/or a too thin quantum barrier would have weak confinement forthe injected carriers within the QWs, therefore decreasing the radiative recombinationrate and the EQE. On the other hand, a thick InGaN/GaN QW is more likely to havepoor structural quality with large defect densities that serve as nonradiative recombi-nation centers. In Fig. 18(c), a too thick LB reduced the EL intensity possibly due topoor hole injection from the p-GaN to the QW. Figure 18(d) showed that a Cp2Mg

Figure 18

EL intensity and emission wavelength as a function of (a) QWwidth, (b) barrier thick-ness, (c) last barrier thickness, and (d) Mg doping for the last barrier and p�-GaNcontact layer for �101̄ 1̄� LEDs. Reprinted with permission from [62]. Copyright2010 The Japan Society of Applied Physics.

Figure 17

Comparison of EQE as a function of current for the reported �101̄ 1̄� blue LEDs.


flow of >0.05 μmol∕min was beneficial to the p�-GaN contact layer, while the ELintensity was nearly constant when varying the Mg doping of the LB. The optimi-zation process led to an efficient blue LED with an output power of 22.75 mWand an EQE of 39.5% at 20 mA. Furthermore, Zhao et al. [13] incorporated backsideroughening and transparent stand packaging techniques into �101̄ 1̄� LEDs. At 20 mAunder pulsed condition, the LED exhibited an output power of 31.1 mW and an EQEof 54.7%. This was the first semipolar LED with over 50% EQE which also exhibitedlow efficiency droop up to 300 mA. Further optimization is likely to lead to evenhigher efficiencies.

3.2. InGaN Blue LEDs Based on High-Inclination-Angle Semipolar PlanesSemipolar planes with high inclination angles such as �202̄ 1̄� (105° from the c-axis)and �303̄ 1̄� (100° from the c-axis) have also recently attracted considerable interest forlow-efficiency-droop blue LEDs, and impressive properties such as large criticalthickness, increased internal quantum efficiency (IQE), low efficiency droop, andlow thermal droop.

3.2a. �202̄ 1̄� InGaN Blue LEDs

Zhao et al. [14] demonstrated �202̄ 1̄� LEDs with an output power of 30.6 mWand anEQE of 52%, which are comparable to the best device performance ever reported foreither nonpolar or semipolar devices. The device structure consisted of 1 μm Si-dopedn-GaN, 10-period InGaN (3 nm)/GaN (13 nm) MQWs, five-period AlGaN/GaNsuperlattice electron blocking layer, and 60 nmMg-doped p-GaN. This device showedlow droop ratios of 0.7% at 35 A∕cm2, 4.3% at 50 A∕cm2, 8.5% at 100 A∕cm2, and14.3% at 200 A∕cm2, while the SOTA c-plane LEDs [53] show nearly 20% efficiencydroop at 100 A∕cm2. The droop ratio is defined as

droop ratio � �EQEMax − EQEJ �∕EQEMax × 100%; (3)

where the EQEMax and EQEJ represent the EQE maximum and the EQE at differentcurrent densities. The reason for such low droop is still not clear and debated and morediscussions on efficiency droop will be presented later. In order to correlate thecarrier dynamics with the low-droop performance of �202̄ 1̄� LEDs, Pan et al. [15]fabricated a small-area (0.1 mm2) LED chip with a SQW GaN�10 nm�∕In0.16Ga0.82N�12 nm�∕GaN �15 nm� active region. Figure 19(a) shows the light output

Figure 19

(a) Light output power and EQE versus current densities for GaN (10 nm)/InGaN(12 nm)/GaN (15 nm) SQW �202̄ 1̄� blue LED under pulsed condition. The insetis the schematic view of the device structure. (b) STEM image of active region ofthe LED. Reprinted with permission from [15]. Copyright 2012 The JapanSociety of Applied Physics.


power (LOP) and EQE of the SQWblue LED. This LED had a LOP of 30mWand EQEof 51% at 20 mA, with low droop ratios of 4.7% at 100 A∕cm2, 14.2% at 200 A∕cm2,and 22.2%at400 A∕cm2. The STEMimage of the active region in Fig. 19(b) indicates auniform QW thickness and a smooth interface. The authors postulated that the combi-nation of reduced electric field (due to the semipolar plane), enhanced carrier uniformity(due to the SQW structure), and smaller potential fluctuation (due to the thick high-quality InGaN layer) increased the effective active region volume, which loweredthe average carrier density. Therefore, the Auger recombination and carrier leakagecould be mitigated and efficiency droop could be reduced [15].

3.2b. �303̄ 1̄� InGaN Blue LEDs

Figure 20(a) shows the Matthews–Blakeslee equilibrium critical thickness fordifferent semipolar orientations for InGaN layers under the assumption of isotropicelasticity [16]. In the Matthews–Blakeslee model, the critical thickness is expressedas [63]

hc ��

b

8πf cos ϕ

��1 − ν cos2 θ

1� v

��ln

αhcb

�; (4)

where hc is the critical thickness, b is the magnitude of the Burgers vector, f is thelattice mismatch, ϕ is the angle between the slip direction and that direction in the filmplane that is perpendicular to the intersection of the slip plane and the interface, ν is thePoisson ratio, θ is the angle between the dislocation line and its Burgers vector, and αis a dimensionless parameter of a GaN crystal estimated by atomic calculations [64].The critical thickness is predominately determined by the resolved shear stress be-cause the resolved shear stress is the driving force for the formation of MDs viathe dislocation slide on the preferable basal c-plane, also referred to as basal plane(BP) slip [48]. From θ � 0° to θ � 45°, the resolved shear stress increases and peaksat θ � 45°; then it decreases with increasing θ [16]. Therefore, as shown in Fig. 20(a),the critical thickness of the semipolar �303̄1̄� plane is larger than both the �101̄1̄� and

Figure 20

(a) Calculated critical thickness based on Matthews–Blakeslee model for �101̄ 1̄�,�202̄ 1̄�, and �303̄ 1̄� semipolar planes. (b)–(d) Panchromatic CL images of �202̄ 1̄�LEDs with 20, 40, and 60 nm SQWactive regions, respectively. (e)–(g) PanchromaticCL images of �303̄ 1̄� LEDs with 20, 40, and 60 nm SQWactive regions, respectively.Reprinted with permission from Becerra et al., Appl. Phys. Lett. 105, 171106 (2014)[16]. Copyright 2014 AIP Publishing LLC.


�202̄ 1̄� planes due to reduced resolved shear stress on the basal plane [16]. Hsu et al.[65] observed the stress relaxation of �303̄ 1̄� InGaN layers near the calculated criticalthickness. Beyond the critical thickness, relaxation occurs along non-basal planes(NBPs), a process also called NBP slip or 2D relaxation. Figures 20(b)–20(g)represent the panchromatic CL images of �202̄ 1̄� and �303̄ 1̄� LEDs with differentSQW thickness. The dark lines parallel to the a-direction corresponded to the BPMDs, which did not have a strong impact on the device performance of InGaN LEDsespecially with thick QWs. Other dark lines that are not parallel to the a-direction wereNBP MDs, which was found to considerably degrade the LED efficiency. With in-creasing QW thickness, the NBP MDs increased dramatically for both semipolarLEDs. However, due to the larger critical thickness, the �303̄ 1̄� LED had significantlyfewer NBP MDs than the �202̄ 1̄� LED at the same QW thickness. Thus, thicker QWscan potentially be grown on the semipolar �303̄ 1̄� GaN substrates, which can help tomitigate efficiency droop due to lower average carrier densities in the QWs.

Figure 21(a) presents the LOP and EQE for a 15 nm thick �303̄ 1̄� SQW blue-violetLED [16]. This device has an output power of more than 1 W at a current density of1 kA∕cm2 with a droop ratio of only 33%, which is promising for high-power ap-plications. In addition, results also showed that increasing QW thickness loweredthe efficiency droop. Figure 21(b) shows the peak EL wavelength and FWHM at dif-ferent current densities. The �303̄ 1̄� LED exhibited only a 7 nm wavelength shift andaround a 15 nm FWHM up to 1 kA∕cm2, possibly due to a much reduced electric fieldand high material quality. These results indicate thick active regions of semipolarLEDs are favorable for reducing efficiency droop.

3.3. Thermal Droopc-plane blue LEDs have not only high efficiency droop but also significant thermaldroop. Thermal droop refers to the reduction of EQE with increasing operating tem-peratures, which is schematically shown in Fig. 22. Though ambient environment hasthe normal room temperature of 20°C–30°C, the junction temperature of InGaN LEDscommonly rises to 80°C–100°C. This high junction temperature can affect the radi-ative and nonradiative recombination processes in the active region and possibly resultin the thermal escaping of hot carriers from the active region. A 20% thermal droop is

Figure 21

(a) Light output power and EQE as a function of current density for GaN (15 nm)/InGaN (15 nm)/GaN (10 nm) SQW �303̄ 1̄� LED under pulsed condition with 1% dutycycle. The inset shows the device under current injection. (b) Peak wavelength andFWHM versus current densities for the device. Reprinted with permission fromBecerra et al., Appl. Phys. Lett. 105, 171106 (2014) [16]. Copyright 2014 AIPPublishing LLC.


observed in c-plane LEDs at 100 A∕cm2 when the temperature is increased fromroom temperature to 100°C [22,23]. Thermal droop is defined as

�EQE�J �20°C − EQE�J �100°C�∕EQE�J �20°C × 100%; (5)

where EQE�J �100°C and EQE�J �20°C is the EQE of LEDs at the junction temperatureof 100°C and 20°C at current density of J , respectively. In addition, the hot/coldfactor can also be used to characterize the thermal performance of a LED, whichis expressed as

Hot∕cold factor � EQE�J �100°C∕EQE�J �20°C: (6)

Semipolar LEDs with a thick QW structure show much improved thermal perfor-mance. Figure 23(a) shows the thermal droop of 12 nm SQW �202̄ 1̄� blue LEDsas a function of temperature [17]. At 100 A∕cm2, the thermal droop is only 9% at100°C with a hot/cold factor of 0.9 while the c-plane blue LEDs show a thermal droop>20% and a hot/cold factor <0.8, meaning that the semipolar LEDs have over 50%reduction in the thermal droop. From 1 to 40 A∕cm2, thermal droop is decreasing withcurrent density; when the current density is further increased to 100 A∕cm2, a slightincrease of thermal droop is observed. This trend can be explained by carrier rateequation model where Shockley-Read-Hall (SRH), radiative, and Auger recombina-tion are considered. SRH recombination dominates at low current densities. Radiativerecombination becomes important with increasing current density, resulting in reduc-tion of thermal droop from 1 to 40 A∕cm2. When the current density becomes evenhigher, Auger recombination becomes dominant and increases the thermal droop.These arguments were further supported by the work from Meyaard et al. [23].

In Fig. 24, Meyaard et al. [23] characterized a c-plane blue InGaN LED with a peakwavelength of 460 nm and decoupled the SRH recombination, Auger recombination,and carrier leakage from the total recombination. Figure 24(a) shows the ratio of theSRH recombination rate to the total recombination rate and Fig. 24(b) was calculatedfrom the Auger recombination rate plus carrier leakage over the total recombinationrate. At a low driving current of 10 mA, the SRH recombination contributes substan-tially to the total recombination: 10.0% at 300 K and 47.3% at 450 K. This means thatthe SRH recombination is responsible for the thermal droop of InGaN LEDs at low

Figure 22

Schematic view of thermal droop in a representative EQE versus temperature curve forInGaN LEDs.


current densities. With increasing current density, the contribution of the SRH recom-bination decreases rapidly. At a high driving current of 2 A, the SRH recombinationaccounts for only 0.6% and 3.1% of the total recombination at 300 K and 450 K, re-spectively. These results indicate that the SRH recombination is not the main contributorto the thermal droop at high current densities. On the other hand, the Auger recombi-nation and carrier leakage are becoming more and more important with increasing cur-rent. At 10 mA, they are only∼10% of the total recombination at all temperatures, whilethey account for 49.3% of the total recombination at 300 K and nearly 60% at 450 K at 2A. Therefore, the Auger recombination and carrier leakage are the dominant lossmechanisms for the thermal droop at high current densities.

Ryu et al. [66] proposed that the efficiency droop of conventional c-plane LEDsresults from the reduced active region volume due to polarization-related electric

Figure 24

Percentage of contribution from (a) SRH recombination and (b) Auger recombinationand carrier leakage to total recombination as a function of temperature at differentcurrent densities. Reprinted with permission from Meyaard et al., Appl. Phys.Lett. 99, 041112 (2011) [23]. Copyright 2011 AIP Publishing LLC.

Figure 23

(a) Thermal droop versus temperature at different current densities for 12 nm SQW�202̄ 1̄� LEDs. Reprinted with permission from [17]. Copyright 2012 The JapanSociety of Applied Physics. (b) EQE as a function of current density with temperature20°C–120°C for 20 nm SQW �303̄ 1̄� LED. This LED has only 10% EQE droop at120°C. Reprinted with permission from Becerra et al., Appl. Phys. Lett. 105, 171106(2014) [16]. Copyright 2014 AIP Publishing LLC.


fields, nonuniform carrier distribution, and indium fluctuations. The current can beexpressed as a function of carrier density as [66]

I � eV eff

�An� Bn2

1� n∕n0� Cn3

1� n∕n0

�; (7)

where e is carrier charge; V eff is the effective active region volume; n is the carrierdensity; A, B, and C are the SRH, radiative, and Auger coefficients, respectively; andn0 is the coefficient for phase-space filling (PSF) effect. This model is called the rateequation and more details will be discussed in the following sections. This reducedactive region volume increases the average carrier density in the QWs and exacerbatesthe Auger recombination and carrier leakage. The thick SQW �202̄ 1̄� blue LEDs havereduced polarization-related electric fields, uniform carrier distribution due to SQW,and thick high-quality homogeneous SQW at high growth temperatures, all contrib-uting to a larger active region volume. This results in a lower average carrier density inthe SQW, which reduces the Auger recombination and carrier leakage and leads to thelow thermal droop performance of the semipolar LEDs. These results indicatethick semipolar QWs can be beneficial in reducing the thermal droop. As mentionedin the previous section, the �303̄ 1̄� plane can achieve even thicker QWs than the�202̄ 1̄� plane. A �303̄ 1̄� LED with 20 nm thick SQW was demonstrated and itstemperature-dependent EQE curves are shown in Fig. 23(b) [16]. This LED exhibitsimproved thermal performance with thermal droop <10% and hot/cold factor>0.9 at 100°C.

3.4. Physical Mechanisms of Efficiency DroopDespite the impressive progress on the reduction of efficiency droop using semipolarorientations for blue InGaN LEDs, the physical origin of the efficiency droop is stillbeing debated. The most popular and widely studied mechanisms are the Auger re-combination [18,19] and carrier leakage [20,21]. The main arguments for the Augerrecombination mechanism are as follows: (1) the Auger recombination plays themajor role at high current/carrier densities as it is proportional to the cube of carrierdensity; (2) the Auger coefficients are larger enough to result in efficiency droop;(3) Auger electrons were experimentally observed in InGaN LEDs and were stronglycorrelated with the droop current. On the other hand, direct observation of carrierleakage was made in InGaN LEDs, and the proposed ABC � f �n� model successfullysimulated the experimental EQE curves. Some other mechanisms will also be dis-cussed, such as reduced effective active region volume [66], carrier delocalization[67], and PSF effect [68–71].

3.5. Auger RecombinationAuger recombination represents a nonradiative pathway for energy dissipation. If theelectron–hole pairs (e-h) recombine, the released energy can be absorbed by anotherelectron (eehprocess) or hole (hhe process) that are subsequently excited to high en-ergy levels, instead of producing photons. Figure 25 schematically shows four pos-sible Auger recombination processes in InGaN LEDs. Because this process involvesthree particles, the Auger recombination rate is proportional to the cube of the carrierdensity. To account for efficiency droop, the Auger recombination coefficient CAuger

must be larger than 10−31 cm6 s−1 [22]. Piprek et al. summarized most of the reportedAuger recombination coefficients based on various device structures [22]. Shen et al.[72] utilized the photoluminescence lifetime measurement and obtained CAuger of∼2.0 × 10−30 cm6 s−1 for c-plane quasi-bulk InGaN samples. Meneghini et al. [73]and Laubsch et al. [74] determined CAuger to be 10−30 cm6 s−1 and 3.5 ×10−31 cm6 s−1 on SQW InGaN LEDs, respectively.


However, the theoretical CAuger based only on the direct intraband Auger recombi-nation process was smaller than the observed experimental results [75,76].Bertazzi et al. [75] obtained a maximum Auger coefficient of ∼10−32 cm6 s−1 basedon direct interband and intraband Auger recombination processes using first-orderperturbation theory. Hader et al. [76] calculated a very small direct band-to-bandCAuger of 3.5 × 10−34 cm6 s−1 using a microscopic many-body model and an 8 × 8

kp band structure model. In order to overcome this discrepancy between theoreticaland experimental CAuger, indirect Auger recombination processes were taken into con-sideration. Kioupakis et al.[71] obtained a larger CAuger with the indirect intrabandAuger recombination process using atomistic first-principle calculations. Delaneyet al. [77] reported a peak Auger coefficient of 2 × 10−30 cm6 s−1 when investigatingthe interband Auger recombination in GaInN using first principle density-functionaland many-body perturbation. In an analytic model, Lin et al. [78] added a drift-induced leakage (CDL) term into the CAuger.

Figure 26(a) shows the direct observation of Auger electrons in InGaN LEDs underbias by electron emission spectroscopy [20,21]. The cesium is deposited on the p-GaNsurface to obtain negative electron affinity so that electrons with energy below theconduction band minimum can be detected with the assistance of energy relaxationin the band-bending region. Electrons with different energies are generated inside theactive region and then emitted from the active region to vacuum, including the en-ergetic Auger electrons. When these Auger electrons reach the surface with enoughenergy, they can be identified by measuring their energies, leading to high-energypeaks in the energy distribution of electrons in vacuum [20,21]. Figure 26(b) comparesthe integrated current of Auger peaks with the droop current. It is obvious that Augercurrent and the droop current are linearly correlated, which suggests that the twophenomena are the same. Since the droop current is also proportional to the cubeof the inject carrier density, other droop mechanisms that do not have the third powerdependence on carrier density are not very likely the major contributors to theefficiency droop [20,21].

Figure 25

Schematic view of four Auger recombination processes: (a) direct intraband, (b) indi-rect intraband, (c) direct interband, and (d) indirect interband. The upward bands areconduction band (CB) and the downward bands are valence band (VB). The nega-tively charaged particles are electrons, and the positively charged particles are holes.The solid arrows are momentum transferred from electron–hole recombination. Thedashed arrows in (b) and (d) represent scattering mechanisms, such as alloy scattter-ing, Coulombic scattering, or electron–phonon coupling [71].


3.5a. Carrier Leakage

Carrier leakage occurs when a portion of electrons escape from the active region ofInGaN LEDs and recombine nonradiatively with holes. It is an inherent process be-cause there are always some electrons with high energy present in the active regionaccording to the Fermi–Dirac distribution. These energetic carriers can fly over thebarrier of the active region and thus contribute to the efficiency droop. Figure 27(a)schematically shows all the nonradiative and radiative recombination processes insidea InGaN LED [19]. The efficiency droop characteristic of LEDs is generally charac-terized by a carrier rate equation model with ABC coefficients (ABC model):

RT � An� Bn2 � Cn3; (8)

Figure 27

(a) Schematics of all the recombination processes in the InGaN LEDs including ra-diative recombination, SRH recombination, Auger recombination, and carrier leak-age. (b) EQE as a function of driving current for two blue c-plane InGaN LEDs.The experimental results are simulated by both the ABC model and ABC � f �n�model. Cho et al., Laser Photon. Rev. 7, 408–421 (2013) [19]. Copyright Wiley-VCH Verlag GmbH & Co. KGaA. Reproduced with permission.

Figure 26

(a) Schematics of band structure of InGaN LEDs under positive bias where differentelectrons are emitted into vacuum and detected. VB is valence band, and EBL is theelectron blocking layer. (b) Integrated high-energy Auger peak versus the droop cur-rent. Weisbuch et al., Phys. Status Solidi A 212, 899–913 (2015) [20]. CopyrightWiley-VCH Verlag GmbH & Co. KGaA. Reproduced with permission. Figures 1(a) and 4(b) reprinted with permission from Iveland et al., Phys. Rev. Lett. 110,177406 (2013) [21]. Copyright 2013 by the American Physical Society.


where RT is the total recombination rate. In order to account for carrier leakage, the“ABC � f �n� model” was proposed [76], which was expressed as

RT � An� Bn2 � Cn3 � f �n�; (9)

where f �n� is the contribution from carrier leakage.

In the high-level injection where the efficiency droop occurs, the drift-induced carrierleakage plays a more important role than the diffusion-induced carrier leakage [79].The drift-induced current JD is proportional to the product of n and total current JT ,JD ∝ nJT . Near the onset of the efficiency droop, radiative recombination dominatesand the JT becomes edBn2, where d is the active region thickness. Therefore, JD canbe rewritten as

JD � edCDn3; (10)

where CD is the third-order recombination coefficient. CD is calculated to be of theorder of 10−29 cm6 s−1 for GaN material [19].

When entering the efficiency droop region with an even high current density, the third-order recombination becomes significant and the JT ∝ n3. The drift-induced currentcan be expressed as

JD � edDDn4; (11)

where DD is the fourth-order recombination coefficient.

Plugging Eqs. (10) and (11) into Eq. (9), the “ABC � f �n� model” can be rewritten as

RT � An� Bn2 � CAugern3 � CDn

3 � DDn4; (12)

where the last components account for the drift-induced leakage. In this model, theonset of the droop current is described by

JONSET � edBA∕CD: (13)

JONSET can be calculated by plugging some typical values of these parametersinto Eq. (10). Using d � 3 nm, B � 10−10 cm3 s−1, A � 107 s−1, and CDL �2.4 × 10−29 cm6 s−1. JONSET of 2 A∕cm2 was obtained, which agrees with the exper-imental onset droop current of 1–10 A∕cm2 [80]. Figure 27(b) shows the fitting ofEQE curves of two blue InGaN LEDs by both the ABC model and the ABC � f �n�model [19,71]. The latter exhibited a better agreement with the experimental data,especially at high current densities. It indicates that Auger recombination and carrierleakage may work together to account for the total efficiency droop.

3.5b. More Mechanisms for Efficiency Droop

Some other mechanisms have also been proposed, providing additional insights intothe efficiency droop issue of InGaN LEDs. As mentioned in Section 3.3, Ryu et al.[66] argued that the effective active region volume should be smaller than the physicalactive region volume because of the poor hole injection, indium fluctuation, andpolarization-induced additional barriers. Using this concept, they obtained reasonablygood fitting of the efficiency droop of a commercial c-plane InGaN LED. Anotherpossible explanation involves carrier delocalization. It is well known that in theInGaN LEDs, the localized exciton recombination is the dominant mechanism ofspontaneous emission [81]. Some carriers can escape from these localized states(i.e., carrier delocalization), and participate in the nonradiative recombination


processes, resulting in efficiency droop [67]. David et al. [68] investigated the PSFeffect as an alternative explanation for efficiency droop. They performed differentialcarrier lifetime measurements on c-plane blue InGaN LEDs and proposed a modifiedABC model incorporating the PSF effect to fit the droop characteristics of the LEDs.More details are discussed in the following section.

3.6. Modified ABC ModelAlthough the IQE curve of c-plane LEDs can be well-fitted with the model, similar A,B, C coefficients were not able to model semipolar LEDs. This indicates a differentABC model has to be used for semipolar LEDs where the different physical propertiesand resulting carrier dynamics must be taken into account [69,70,82]. Although someother models with more recombination components added were also proposed to fitthe experimental data, the ABC model still remains the most popular model used in thecommunity of InGaN LEDs. The popularity of the ABC model is a result of its sim-plicity and flexibility. The A, B, and C coefficients can be extracted by simply fittingexperimental data with the ABC model. The modified ABC model with PSF effect isan advanced version of the conventional ABC model. This model can fit both c-planeand low-droop semipolar LEDs, whereas that latter cannot be fitted by theconventional ABC model.

Figure 28 compares time-resolve photoluminescence (TRPL) spectra and carrier life-times of semipolar and c-plane InGaN LEDs with similar device structures [69,70].The carrier decay in the TRPL can be characterized by Δn � Δn0 exp�−t∕τ�, whereΔn is the excess carrier density, Δn0 is the photogenerated carrier density, and τ is theminority carrier lifetime. In InGaN LEDs, exciton emission is the dominant decayprocess due to the presence of localized states related to the indium fluctuation.Semipolar InGaN LEDs have much smaller carrier lifetime τ compared to that ofc-plane devices, possibly due to the reduced QCSE and material qualities resultingfrom the high growth temperature, lower defect density, and reduced indium fluctua-tions in the QWs. The carrier lifetime is linked to the droop performance of InGaNLEDs by a current density equation and modified ABC model. The current density Jcan be expressed as a function of τ:J � edΔn∕τ, where d is the active region thickness

Figure 28

Representative TRPL spectra of semipolar �202̄ 1̄� LEDs (red line) and c-plane LEDs(blue line) with similar device structures at room temperature. The inset table showsthe carrier lifetime τ of the two devices at different wavelength λ. τ is extracted basedon exponential decay. © 2016 IEEE. Reprinted, with permission, from Fu et al., J.Disp. Technol. 12, 736–741 (2016) [69].


and Δn is the excess carrier density. Therefore, at the same injected current density J,semipolar LEDs will have less excess carrier density Δn than c-plane samples due tothe smaller carrier lifetime τ. This may explain the observed low droop on semipolarLEDs. In order to confirm this hypothesis, a modified ABC model with PSF effect tosimulate the droop characteristics of the semipolar LEDs was used. The modified ABCmodel can be expressed as [68–70]

J � ed

�An� Bn2

1� n∕n0� Cn3

1� n∕n0

�; (14)

IQE � Bn2

1� n∕n0∕�An� Bn2

1� n∕n0� Cn3

1� n∕n0

�: (15)

n is expressed as (Δn� N0) where N 0 is the doping density. Since N0 is much smallerthan Δn, n is roughly equal to Δn. A smaller n0 indicates a stronger PSF effect, andvice versa. The PSF effect comes from the utilization of a Fermi–Dirac distribution athigh carrier densities instead of a Boltzmann distribution. As a result, the radiativerecombination is proportional to n instead of n2 at high carrier density. This effectis accounted for by replacing coefficient B by B∕�1� n∕n0�. Additionally, Hader et al.[83] and David et al. [68]. also assumed the Auger recombination to become subcubic,substituting C by C∕�1� n∕n0�. Figure 29(a) shows IQE curves versus currentdensity for various n0. Changing n0 dramatically modifies the efficiency curves andrelated droop performance [55]. In Fig. 29(b), the droop ratios are calculated for IQEcurves in Fig. 29(a) [69]. It was found that smaller n0 (stronger PSF effect) showslarger efficiency droop.

Based on the discussions above, semipolar LEDs should have a weaker PSF effect dueto the lower carrier density and a larger n0 should be utilized in the simulation. On theother hand, smaller n0 should be applied to c-plane samples because of the strongerPSF effect. Figure 30 presents the simulation results of semipolar [15] and c-planeLEDs [53,54]. The assumption about the light extraction efficiency (LEE) (ηextr)are reasonable with the current technology. The injection efficiency is assumed tobe 100% for all three LEDs. A very good agreement between experimental dataand the theoretical modeling was obtained for a semipolar LED using a weak PSF

Figure 29

(a) Calculated IQE curves as a function of current densities using ABC modelwith different n0 coefficients; (b) calculated droop ratio of different IQE curves in(a). © 2016 IEEE. Reprinted, with permission, from Fu et al., J. Disp. Technol.12, 736–741 (2016) [69].


effect (n0 � 5.0 × 1019 cm−3) and two c-plane LEDs using a strong PSF effect(n0 � 1.0 × 1018 cm−3 or 6.0 × 1018 cm−3). The largest n0 was obtained on semipo-lar LEDs due to the lowest efficiency droop. Furthermore, the chip size can also play arole in the efficiency droop. Large chip size can decrease the average carrier density inthe QWs and thus the efficiency droop. However, from the perspective of cost andfootprint reduction, it is desired to develop high-power small-area LEDs operating athigh current densities. Though the UCSB semipolar LED [15] has a 1 order of mag-nitude smaller chip size, the efficiency droop is much smaller than Nichia c-planeLEDs [53] with comparable IQE at high current densities. These results all indicatethe advantages of developing high-power, low-droop, small-area semipolar blue LEDsfor various lighting applications.

4. GREEN SEMIPOLAR INGAN LEDS

Many studies have shown that the EQE and output power of InGaN and AlGaInPLEDs decrease dramatically when the emission wavelength approaches the greenspectral region, a phenomenon known as the “green gap” [24]. Figure 31 showsthe current status of EQE for green LEDs based on AlGaInP, c-plane InGaN, andsemipolar InGaN [24]. With increasing mole fraction of Al, AlGaInP becomes indirectbandgap and carrier confinement weakens, leading to the emission wavelength notshorter than ∼580 nm. The EQE of c-plane InGaN green LEDs grown on foreignsubstrates decreases dramatically in the green part of the spectrum, possibly dueto high TDs and point defects, increased strain, and stronger polarization at higherindium compositions. One approach to solve the polarization issue is to grow nonpolarm-plane LEDs on bulk GaN substrates. However, these LEDs still had an emissionwavelength of <510 nm, an EQE of <1%, and output power of <1 mW due to poormaterial quality and low indium incorporation [57]. Recently the growth of greenLEDs on the �202̄1� and �202̄ 1̄� semipolar planes has been investigated. Zhao et al.[25] first demonstrated green �202̄ 1̄� LEDs with an emission wavelength of ∼520 nm

and a peak EQE of ∼15%. Even higher EQE of 28.3% was demonstrated on the

Figure 30

Simulated IQE curves as a function of current densities for semipolar �202̄ 1̄� LEDs[15] with weak (solid line) PSF effect, Nicha c-plane LEDs [53] with strong PSF effect(dotted line) and UCSB c-plane LEDs [54] with strong PSF effect (dashed–dottedline). Reported experimental data are also plotted for UCSB semipolar �202̄ 1̄�LED (circle), Nicha c-plane (triangle), and UCSB c-plane LED (diamond).


�202̄1� plane. These results show promise that the semipolar �202̄1� and �202̄ 1̄�planes can potentially bridge the “green gap.”

4.1. Material Properties of Semipolar �202̄1� and �202̄ 1̄� Plane for Green EmissionThe measured EL emission wavelengths under electrical injection for m-plane andsemipolar �202̄1� and �202̄ 1̄� LEDs grown under the same conditions are shownin Fig. 32. The �202̄1� and �202̄ 1̄� LEDs show longer EL wavelengths with largeroutput powers thanm-plane devices. However, it also shows that at longer wavelength,the EL power decreases, which may be due to the degradation of material and strongerpolarization effect at higher indium incorporation. Here, we review the effects of in-dium incorporation and polarization-related electric field on the device performance ofnonpolar and semipolar plane green LEDs.

Figure 31

Current status of EQE for UCSB semipolar, Lumileds c-plane InGaN, and LumiledsAlGaInP LEDs as a function of peak emission wavelength [24]. Lumileds data arenon-thin-film flip-chip devices. All data are collected at 22 A∕cm2 or 35 A∕cm2.Recent UCSB semipolar LEDs are filling the green gap of LEDs. V �λ� is the luminouseye response curve.

Figure 32

Quick-test EL power as a function of wavelength for co-loaded m-plane, �202̄1� and�202̄ 1̄� LEDs. Schematic view of the m-plane, �202̄1� and �202̄ 1̄� planes in thewurzite structure are also shown on the right.


4.1a. High Indium Incorporation

Figure 33(a) shows the EL wavelength of SQW LEDs on m-plane and semipolar�202̄1�, �202̄ 1̄�, �303̄1�, and �303̄ 1̄� orientations at a growth temperature Tg �780°C with varying TMI flows [84]. To ensure similar growth conditions, all the sam-ples were co-loaded into the MOCVD reactor. The device active region is a 3 nmInGaN SQW. The details about the device structure can be found in [84]. As seen,m-plane LEDs have the shortest wavelength. For semipolar LEDs, devices grown onplanes with higher inclination angles showed longer emission wavelengths. X-ray dif-fraction (XRD) analysis shows that at the same growth conditions, the indium in-corporation is decreasing from the �202̄ 1̄� plane (Ga-polar character) to the �202̄1�plane (N-polar character) to the m-plane (nonpolar), which is consistent with theo-retical results [85]. Theoretical calculations indicate that when 60° ≤ θ ≤ 90°, indiumincorporation decreases monotonically with θ due to the increased strain-inducedrepulsive interactions between neighboring indium atoms [85]. Though having thesame θ and strain condition, semipolar �202̄1� and �202̄ 1̄� LEDs showed slightlydifferent indium composition and emission wavelength. Similar differences were alsofound for �303̄1� and �303̄ 1̄� LEDs [86]. These results indicate that surface polaritiesmay also affect the indium incorporation [87]. Figure 33(b) shows the EL peak wave-length versus growth temperature for co-loaded �202̄1� and �202̄ 1̄� SQW LEDs [84].For both samples, the emission wavelength is reduced at high growth temperature dueto strong indium desorption from the growth surface. Furthermore, �202̄ 1̄� SQWLEDs showed longer emission wavelength than �202̄1� LEDs at the same tempera-ture. Therefore, for the same emission wavelength, �202̄ 1̄� LEDs can be grown atrelatively higher temperature (30°C–50°C higher), which can improve the materialquality and reduce the indium fluctuations.

4.1b. Reduced Polarization-Related Electric Field

The polarization-related electric field also plays an important role in determining theemission wavelength since the QW profile can be altered because of it. Figure 34(a)presents the EL intensity versus wavelength for co-loaded m-plane, �202̄1�, and�202̄ 1̄� LEDs. The structure consists of a 3.5 nm thick SQW without indium tin oxideor packaging. The EL spectrum shows that the m-plane LED has the shortest emission

Figure 33

(a) EL peak wavelength versus TMI flow for SQW LEDs grown on �202̄1�, �202̄ 1̄�,�303̄1�, �303̄ 1̄�, and m-planes at growth temperature 780°C. (b) EL peak wavelengthas a function of growth temperature for SQW LEDs grown on semipolar �202̄1� and�202̄ 1̄� planes at TMI flow 160 SCCM. The inset shows semipolar crystal orienta-tions. Reprinted with permission from Zhao et al., Appl. Phys. Lett. 100, 201108(2012) [84]. Copyright 2012 AIP Publishing LLC.


wavelength while the LED grown on the �202̄ 1̄� substrate emits the longest wave-length light. This trend can be a result of combined influences from the indiumincorporation and polarization effects. Figure 34(b) shows the room temperaturephotoluminescence (PL) wavelength versus the EL wavelength of the three LEDs.For �202̄ 1̄� and m-plane LEDs, there is a minimal difference between the PL and ELwavelengths. For �202̄1� LEDs, the PL wavelength is also almost the same as the ELwavelength for the blue-green region. However, in the green region, the EL wave-length is shorter than the PL wavelength, which could be linked to the polarization-related electric field inside the SQW.

In order to further study the polarization effect on the emission wavelength of theInGaN LEDs under bias, the commercial SiLENSe package was used to simulatethe 3 nm In0.33Ga0.67N∕GaN SQW at 20 A∕cm2 for �202̄1�, �202̄ 1̄�, and m-planegreen LEDs. Figure 35(a) shows the band structure and EL spectra of the three LEDs.For the m-plane, Epz is 0 and the QW is nearly flat under 20 A∕cm2. In the case of the�202̄1� plane, Epz is parallel to Ebi and larger than Ebi, which makes the QW tiltedupward compared with them-plane. For the �202̄ 1̄� plane, Epz is antiparallel to Ebi andlarger than Ebi. Therefore, the QW is tilted downward. The consequence of the QW

Figure 35

Simulated (a) band structure and (b) EL spectrum for 3 nm In0.33Ga0.67N∕GaN SQWm-plane, �202̄1� and �202̄ 1̄� green LEDs using SiLENSe.

Figure 34

EL intensity versus wavelength for m-plane, �202̄1� and �202̄ 1̄� LEDs that areco-loaded into the MOCVD; (b) PL wavelength versus the corresponding EL wave-length for all the samples with various indium compositions.


distortion is shown in the EL spectrum in Fig. 35(b). It shows that �202̄1� LEDs havelonger emission wavelength at the same In content than �202̄ 1̄� and m-plane LEDs.

Combining the results from Figs. 33–35, we can study the effect of indium incorpo-ration and polarization-related electric field separately. At the same indium content,the wavelength is �202̄1� > �202̄ 1̄� > m-plane, which is solely determined by thepolarization effects. This means that �202̄1� has the beneficial polarization effects thatredshift the emission wavelength. In the co-load experiment, the wavelength is�202̄ 1̄� > �202̄1� > m-plane, which is influenced by both indium incorporationand polarization effects. As discussed before, the indium incorporation is also�202̄ 1̄� > �202̄1� > m-plane. Despite the advantageous polarization effects, �202̄1�still has a shorter emission wavelength than �202̄ 1̄� in the co-load experiment, be-cause of lesser indium incorporation. On the other hand, �202̄ 1̄� shows the longestemission wavelength due to the largest indium incorporation, regardless of the largeradverse impact from polarization effects. These results indicate that indium incorpo-ration plays a more dominant role in determining the emission wavelength of theInGaN LEDs than polarization effects in the co-load experiment.

4.1c. Low Active-Region Growth Rate

Semipolar planes such as �202̄1� and �202̄ 1̄� show improved performance in the greenspectral region in terms of the EQE and emission wavelength compared with nonpolarm-plane. However, the incorporation of high indium composition results in a largelattice mismatch, which leads to the formation of various defects, such as BPSFsand MDs, for semipolar green InGaN LEDs. For example, dark triangular defects(DTDs) were observed on �202̄1� and �202̄ 1̄� LEDs and many other semipolar planesalso showed the similar defects. Zhao et al. [88] demonstrated that DTDs could beeffectively reduced by using the slow growth rate for the active region, which results inlonger emission wavelength and higher output power as well. Figures 36(a) and 36(b)present the fluorescence microscopy images for semipolar �202̄ 1̄� SQW green LEDs(peak EL wavelength of 510 nm) grown at 0.15 Å/s and 1 Å/s, respectively [88]. LEDsgrown at the slower growth rate have fewer DTDs possibly because of the enhancedadatom diffusion length. TEM characterization in Figs. 37(a) and 37(b) showed thatthe DTDs were associated with the dish-like microstructural voids (50–100 nm large)with clear crystallographic orientations of f0001̄g, f101̄0g, and f101̄1g [88]. Theseorientations were also observed on the pyramidal void defects and V-defects ofc-plane LEDs [89,90]. The surface morphology of the �202̄ 1̄� QWs grown at differentgrowth rates was characterized by AFM, as shown in Figs. 37(c) and 37(d). The sam-ple grown at 0.15 Å/s had a much smoother surface than the sample grown at 1 Å/s. Atthe low growth rate, the adatom diffusion length could be enhanced and these adatomswere able to reach the preferential incorporation sites, resulting in a better surfacemorphology. The observed striated morphologies in both samples stemmed fromthe in-plane diffusion anisotropy of the �202̄ 1̄� plane. Similar morphologies were alsoreported on other planes [91,92].

For �202̄1� plane-based laser structures, similar DTDs have also been observed. Linet al. [93] showed that using AlGaN barriers can largely reduce or eliminate DTDs[Figs. 38(a)–38(c)], as well as increase the emission wavelength in the green spectralregion [Fig. 38(d)]. However, the large difference in growth temperatures betweenhigh-indium QWs and AlGaN barriers limits its applications in green LEDs. Recently,instead of using entire AlGaN barriers, Toshiba work [94,95] and Koleske et al. [96]found using AlGaN interlayer (IL) between the QW and the GaN barrier can alsoimprove the QW quality and increase the In incorporation in c-plane structures.Furthermore, Hardy et al. [97] found that DTDs were also related to post-QWp-GaN annealing. The size of these defects was increased with higher annealing


temperature and longer annealing time. Due to the structural complexity, the physicalmechanism for DTD formation is still under investigation. Figure 37(c) shows thequick-test EL power as a function of peak EL wavelength for devices grown at

Figure 37

(a),(b) Two beam scattering contrast TEM images for �202̄ 1̄� InGaN LEDs. AFMimages of �202̄ 1̄� InGaN QWs grown at (c) 0.15 Å/s and (d) 1 Å/s. Reprinted withpermission from Zhao et al., Appl. Phys. Lett. 102, 091905 (2013) [88]. Copyright2013 AIP Publishing LLC.

Figure 36

Fluorescence microscopy images for semipolar �202̄ 1̄� SQW LEDs grown at(a) 0.15 Å/s and (b) 1 Å/s for peak EL wavelength of 510 nm. (c) Quick-test EL powerversus the peak EL wavelength for green semipolar �202̄ 1̄� SQW LEDs growth at0.15 Å/s, 0.3 Å/s, and 1 Å/s. The inset shows the PL spectrum of semipolar �202̄ 1̄�SQW LEDs at growth rates of 0.15 Å/s and 0.3 Å/s at peak EL wavelength of 515 nm.Reprinted with permission from Zhao et al., Appl. Phys. Lett. 102, 091905 (2013)[88]. Copyright 2013 AIP Publishing LLC.


0.15 Å/s, 0.3 Å/s, and 1 Å/s. With increasing wavelength, the output powerdecreases for all the samples, which shows the same trend as the “green gap.”LEDs with fast growth rates exhibited shorter emission wavelength and lower outputpower. The former is possibly caused by significant indium content in the GaN barrierand associated polarization effects, and the latter is a result of poor material quality. Inaddition, the slow growth rate also contributes to a sharper interface between theInGaN QW and the GaN barrier and more uniform device structure, resulting inimproved device performance [88]. This is related to the effects of growth parameterson the condensation and absorption of indium [98].

4.1d. Surface Morphology

To clarify the different surface characteristics of �202̄1� and �202̄ 1̄� planes, AFM andSEM techniques are utilized. Figures 39(a) and 39(b) present the surface morphologyof the co-loaded LEDs fabricated on �202̄1� and �202̄ 1̄� planes. RMS roughness forthe surface of �202̄1� LEDs is 199 pm while RMS roughness for �202̄ 1̄� LEDs is374 pm. This implies a better surface morphology for �202̄1� LEDs. And the largerRMS roughness of �202̄ 1̄� LEDs may be caused by higher indium incorporation at agiven temperature, which could result in the indium segregation and defect generation.Figures 39(c) and 39(d) show the SEM images of surfaces of the two samples afterphotoelectrochemical (PEC) etching. The morphology of the two samples exhibitsdrastically different features. For N-polar �202̄1� LEDs, PEC etching creates hexago-nal pyramids bound by clear crystallographic facets. However, Ga-polar �202̄ 1̄� LEDsshow only whisker-like structures without showing any crystal orientations. The sim-ilar results have been reported for N-polar and Ga-polar c-plane devices [99,100].Therefore, for �202̄1� LEDs, this roughened hexagonal surface can be used to increasethe light extraction efficiency and thus the external quantum efficiency and outputpower. This is another advantage of �202̄1� LEDs over other planes.

4.2. High-Performance Green Semipolar InGaN LEDsIn 2006, Funato et al. [12] first demonstrated �112̄2� green LEDs with an outputpower of 1.91 mW and an EQE of 4.1% at 20 mA. Later, Sato et al. [101] reporteda better �112̄2� green LED with an output power of 5 mW and an EQE of 10.5% at

Figure 38

Fluorescence microscope images of InGaN MQWs structures with (a) GaN barrier,(b) InGaN barrier, and (c) AlGaN barrier. (d) EL output power as a function of emis-sion wavelength for InGaN MQWs devices with GaN, InGaN, and AlGaN barriers.Reprinted with permission from [93]. Copyright 2010 The Japan Society of AppliedPhysics.


20 mA. Recently developed �202̄1� [102] and �202̄ 1̄� [25] green LEDs have dem-onstrated much more promising device performance in terms of output power, effi-ciency, and spectral response. Table 2 summarizes high-performance semipolar greenLEDs on bulk GaN substrates.

4.2a. �202̄1� Green LEDs

Recently, semipolar �202̄1� planes have received much attention due to the demon-stration of direct emission green LDs [103]. Further studies show that green emittingQWs grown on the �202̄1� plane exhibit high compositional homogeneity comparedto c-plane QWs [104], indicating a possible superior performance for �202̄1� devicesat longer wavelengths. Although green �202̄1� LEDs with relatively high efficiencyhave been reported, the performance of those devices still needs to be improved forhigh-power applications. Our group recently demonstrated a board area (1 mm2) high-power �202̄1� green LED with an output power of 13.3 mWand a peak EQE of 28.2%at 20 mA, which are significantly higher than previous reported semipolar green LEDresults. The LED structures consisted of 1 μm Si-doped n-GaN, followed by an activeregion which consisted of 3.5 nm thick InGaN SQWand GaN barriers. A 20 nm thickMg-doped p-AlGaN electron blocking layer was grown after the active region, fol-lowed by a p-GaN layer and p�-GaN contact layer. To improve the light extractionefficiency, the backside of the LED was roughened [13] and packaged with a verticalstand transparent structure. Figure 40 presents the LOP and EQE as a function ofcurrent for the packaged green LED. At 350 mA, the LED showed an output powerof 120.7 mW and an EQE of 15.5%. However, the efficiency droop is relatively large(24.3% at 100 mA and 47.0% at 350 mA) compared with blue �202̄1� LEDs.Figure 41 shows the EL peak wavelength and FWHM of the LED at different drivecurrents, where the EL spectra of the LED is also plotted in the inset. The wavelengthshifts with injection current are 8.5 nm (1–20 mA), 6.6 nm (20–100 mA), and 3.8 nm(100–350 mA). This amount of blueshift is comparable to that of green LEDs grownon the �112̄2� plane [12,101] and commercial LEDs on the c-plane, but significantlylarger than the values reported on �202̄ 1̄� green LEDs [25]. The observed large

Figure 39

AFM images of the surface morphology for co-loaded (a) �202̄1� plane and(b) �202̄ 1̄� plane LEDs; SEM images of the surface co-loaded (c) �202̄1� planeand (d) �202̄ 1̄� plane LEDs after PEC etching.


blueshift on �202̄1� LEDs is probably due to both the band-filling effect caused bypotential fluctuations and the QCSE. In addition, the spectral boarding with increasingcurrents was observed for the device. The FWHMs are 33.8 nm, 31.1 nm, 32.9 nm,and 40.6 nm at currents of 1 mA, 20 mA, 100 mA, and 350 mA, respectively. Thisspectral boarding effect, possibly due to the compositional fluctuation inside the QWs,is a topic of ongoing research.

Figure 40

Light output power and EQE versus drive currents for a packaged green �202̄1� LEDunder pulsed operation. The inset shows a schematic of the device.

Figure 41

EL peak wavelength and FWHM as a function of currents for the semipolar green�202̄1� LED. The inset shows EL spectra of the green LED from 2 mA to 100 mA.

Table 2. Device Structure and Performance of Green InGaN LEDs Grown on Bulk GaNSubstratea

Plane Structure Method Wavelength (nm) LOP (mW) EQE (%) Droop (%) Reference

�112̄2� 3 nm SQW DC 527 2.3 4.0 31.7 [12]�112̄2� 4 nm 6 QWs DC 516 5.0 10.5 49.8 [101]�112̄2� 4 nm 6 QWs Pulsed 519 9.0 18.9 44.4 [105]�202̄1� 3.5 nm SQW Pulsed 516 9.9 20.4 57.1 [102]�202̄ 1̄� 3 nm SQW Pulsed 518 5.8 11.9 56.5 [25]�202̄1� 3.5 nm SQW Pulsed 529 13.3 28.2 24.3 This workaWavelength is peak EL wavelength at 20 mA. LOP and EQE are the values at 20 mA. All the droopis measured at 100 mA except that [102] and [25] are calculated at 80 mA.


4.2b. �202̄ 1̄� Green LEDs

High indium incorporation observed on the �202̄ 1̄� plane also makes it promising forgreen emission. Zhao et al. [25] demonstrated a �202̄ 1̄� green LED with an outputpower of 5.8 mWand an EQE of 11.9% at 20 mA, which is higher than �112̄2� greenLEDs [12,101] but lower than �202̄1� green LEDs [102]. The device consisted of 1 μmSi-doped n-GaN, 3 nm InGaN SQW active region, 16 nm Mg-doped p-AlGaN elec-tron blocking layer, and 60 nm Mg-doped p-GaN. Moreover, this LED exhibited avery narrow FWHM and an extremely small wavelength shift. Figure 42 compares theEL peak wavelength and FWHM of this �202̄ 1̄� green LED [25] with green LEDsgrown on the c-plane [53], �112̄2� [12], and �202̄1� [102]. The �202̄ 1̄� green LEDshowed almost no wavelength shift compared with other green LEDs with a 10–15 nmwavelength shift. According to Fig. 4(e), a �202̄ 1̄� InGaN QW has reducedpolarization-related effects and nearly flat QW profile because: (1) Ebi and Epz arein opposition direction with similar magnitude and thus cancel each other, and (2) cur-rent injection results in Coulombic screening of the polarization-induced electric field.Furthermore, the �202̄ 1̄� green LED had the smallest FWHM possibly due to reducedindium fluctuation inside the QW.

4.2c. Device Performance Comparison

The EQE and peak EL wavelength of green LEDs grown on different planes are com-pared in Fig. 43. The combined effects of the defective material quality, the low andinhomogeneous indium compositions make m-plane QWs undesirable for greenLEDs. For �112̄2� green LEDs, Sato et al. [105] reported the highest EQE of 18.9%at 520 nm. However, �112̄2� LEDs are prone to formMDs within the QWs even at lowindium compositions [106,107], which severely limits the access to longer emissionwavelength and degrades the device performance. �202̄1� green LEDs were demon-strated with a record high peak EQE of 28.3% with an emission wavelength of530 nm, which are superior to other nonpolar and semipolar green LEDs reported.This is due to higher indium incorporation, better material quality, and advanced chipprocessing technology, such as backside roughening and vertical stand transparentstructure. �202̄ 1̄� LEDs [25] show an extremely low wavelength shift and a narrowspectra linewidth with a decent EQE and emission wavelength. This is related to thecombined effects of Coulombic screening and electric field cancellation between Ebi

and Epz. However, all these green LEDs showed larger efficiency droop comparedwith their counterparts in the blue spectral region. One reason is the inevitable defects

Figure 42

(a) EL peak wavelength and (b) FWHM for green semipolar �112̄2� [12], �202̄1�[102], and �202̄ 1̄� [25], and polar c-plane LEDs [53] at various current densities.Reprinted with permission from [25]. Copyright 2013 The Japan Society ofApplied Physics.


due to high indium composition. Based on the observations and discussions onlow-droop blue LEDs, several methods can be investigated to reduce efficiency droop,such as using thicker QWs and larger chip sizes.

5. THREE-DIMENSIONAL NONPOLAR AND SEMIPOLARNANOSTRUCTURES FOR INGAN LEDS

Over the past several years, nanostructures and microstructures, such as wires/rods,pyramids, and walls/stripes, have been heavily explored as next-generation, “three-dimensional” LED architectures to reduce the efficiency droop and overcome the“green-yellow-red gap” for III-nitrides. Due to their 3D morphology, such structurescan interestingly provide access to nonpolar and semipolar planes without the need forcostly, bulk nonpolar and semipolar GaN substrates. Additionally, due to their smallcross-sectional dimensions and high free surface areas, such structures can accommo-date greater lattice mismatch strain for enhanced In incorporation, be grown on avariety of non-lattice-matched substrates, as well as provide a larger effective activeregion volume compared to standard planar architectures. Here we briefly reviewrecent efforts on nonpolar and semipolar InGaN/GaN nanostructures for LEDsand provide outlooks for future directions.

5.1. Nonpolar Core–Shell NanostructuresPerhaps the most intensive recent efforts for 3D nano-LEDs have focused on radialcore–shell architectures based on c-axis oriented, hexagonal nanowires with nonpolar{101̄0} m-plane sidewalls, and are currently being pursued by academia and industry.A key attractive feature of this architecture is the ability to grow by selective-areaMOCVD n-type GaN nanowires on [108–111] or etch nanowires from [112–117]standard c-plane GaN-on-sapphire epilayers as highly ordered and size-controlled ar-rays. Selective area growth (SAG) of GaN nanorods (nanowires) through holes pat-terned in a dielectric layer (e.g., SiO2) by MOCVD is often achieved by pulsed growthto enhance vertical growth along the h0001i directions [Fig. 44(a)] [118], and it hasalso been found that the carrier gas ratio of H2 to N2 along with Si doping can alsoaffect the vertical growth [110]. Cross-sectional TEM studies showed that threadingdislocations that propagate into the nanowire from the underlying GaN layer tend to

Figure 43

Comparison of the EQE and EL peak wavelength of high-performance green semi-polar �112̄2�, �202̄1�, and �202̄ 1̄�, and nonpolar m-plane LEDs at a driving current of20 mA.


bend toward and terminate at the nanowire sidewalls, effectively reducing thedislocation density in the grown nanowires compared to GaN template [119]. [0001]oriented nanowires can also be realized through top-down etching of (0001) GaN.Anisotropic “dry” Cl2-based plasma etching has been used as a top-down approachto define pillars from InGaN/GaN LED planar heterostructures [120–123]. However,the plasma etch results in tapered pillars with etch-damaged sidewalls.

More recently, Wang and co-workers developed a two-step top-down approach whichcombines an initial inductively coupled plasma etch with a subsequent, crystallo-graphically selective and highly anisotropic KOH-based wet etch [114,115]. This ap-proach removes the dry-etch damaged sidewalls and results in nanowires with highlystraight nonpolar sidewalls [Fig. 44(b)], making them attractive for nanolasers[115,116,117,124,125] and other potential nano-based applications. Additionally, thisapproach offers the ability to realize nanowires with alternative cross sections, such asannular and rectangular, which are difficult to achieve with bottom-up growth(Fig. 45) [126,127]. This cross-sectional shape control offers intriguing possibilitiesfor controlling the optical and other properties.

The growth and properties of m-plane nonpolar InGaN/GaN quantum wells on thebottom-up grown or top-down etched nanowire sidewalls to form core–shell structureshave been studied in detail [110,111,122,128–135]. In addition to the nonpolarInGaN/GaN QWs, six-fold semipolar {101̄1} QWs are observed at the nanowiretip, and depending on growth conditions, a top c-plane facet and correspondingc-plane QWs can also be present [Figs. 46(a) and 46(b)]. Different In content andemission wavelengths are observed on the nonpolar, semipolar, and polar facets,as determined by atom probe tomography, STEM-energy dispersive spectroscopy,and spatially resolved CL. The area and hence the contribution of the semipolarand polar QWs at the nanowire tip can be minimized by higher aspect ratio structuresin which the nonpolar sidewall area dominates. The c-plane facet of the GaN nanowirecan also be grown to extinction prior to InGaN QW growth [122]. For the nonpolarsidewall QWs, both thickness and In composition variation are typically observedalong the vertical nanowire growth direction, resulting in longer wavelength emissionat the top versus the bottom of the nanowires (Fig. 47) [129,130,132,136,137].

Figure 44

(a) c-axis oriented nanowires with {11̄00} nonpolar sidewalls grown by selective areagrowth by MOCVD (scale bar 500 nm). Reprinted with permission from Yeh et al.,Nano Lett. 12, 3257–3262 (2012) [118]. Copyright 2012 American Chemical Society.(b) c-axis oriented nanowires with {11̄00} nonpolar sidewalls etched by a two-stepdry plus wet-etch top-down approach (scale bar 2 μm). Reprinted with permissionfrom Liu et al., Nanoscale 7, 9581–9588 (2015) [117]. Copyright 2015 The RoyalSociety of Chemistry.


A diffusion-based theoretical model was proposed which shows good agreement withexperimental results showing higher In incorporation and thicker sidewall quantumwells going toward the nanowire tips [138]. Recent work has demonstrated improve-ments in the sidewall uniformity for single QWs over a ∼1.2 μm m-plane length byoptimization of growth conditions and a reduction of nanowire fill factor [139].

Vertically integrated, electrically injected core–shell nanowire LEDs have also beendemonstrated [140–145]. It is generally observed that a pronounced blueshift in emis-sion wavelength is observed with increasing bias voltage and injection current, fromred, yellow, or green at lower currents, to blue at higher currents. This is explained by

Figure 46

Nonpolar core–shell nanowire. (a) Schematic and (b) cross-sectional STEM imageshowing structure of nonpolar (11̄00) InGaN/GaN MQWs grown on c-axis orientedGaN nanowires on sapphire. (a) Reprinted with permission from Riley et al., NanoLett. 13, 4317–4325 (2013) [128]. Copyright 2013 American Chemical Society.(b) Reprinted from Wierer et al., Nanotechnology 23, 194007 (2012) [129]. © 2012IOP Publishing. Reproduced with permission. All rights reserved.

Figure 45

Cross-sectional shape controlled GaN nanowires with nonpolar sidewalls fabricatedby two-step top-down approach: (a) c-axis oriented GaN nanotube (scale bar 1 μm)[126]. Reprinted with permission from Li et al., ACS Photon. 2, 1025–1029 (2015)[126]. Copyright 2015 American Chemical Society. (b) Rectangular GaN nanowire(scale bar 250 nm). Reprinted with permission from Li et al., Nanoscale 8, 5682–5687(2016) [127]. Copyright 2016 The Royal Society of Chemistry.


current spreading models of preferential injection of the top In-rich c-plane QWs ortop region of the nonpolar m-plane sidewall QWs at lower injection currents[142,143,145,146]. These results highlight some of the challenges with respect tothe complex 3D architecture in terms of growth and device integration. White-light-emitting, electrically injected core–shell microrod LEDs have also been demonstratedby insertion of yellow- and red-emitting phosphors between the microrods [145].Additionally, optically pumped lasing has been demonstrated for both arrays ofcore–shell nanorods [147], through a random lasing mechanism, and for singlenonpolar core–shell p-i-n nanowires [134].

In addition to nanowires and nanorods, other 3D nonpolar structures being investi-gated are GaN-based nanosheets, nanowalls, or fins. Compared to nanowires andnanorods, such nanowalls can have an even larger effective active region volume[148], while sharing the benefits of reduced dislocation density and lack of polariza-tion fields. High-aspect-ratio GaN fins with a-plane nonpolar sidewalls have beendemonstrated by SAG via MOCVD [148,149]. SAG growth of m-plane nanowalls(Fig. 48) with subsequent growth of InGaN/GaN QWs emitting at ∼436 nm [150]and ∼480 nm [151] have also been demonstrated. As with core–shell nonpolarQW nanowires, the presence of semipolar {101̄1} and (0001) facets presents chal-lenges that require further work. For both nanowalls and nanowires, it is seen thatlonger wavelength emission is observed on the top c-plane polar facet comparedto the m-plane sidewall facets, due to lower In incorporation and the lack ofQCSE [152]. Thus, 3D nonpolar, as with planar nonpolar, structures, do not currentlypresent a compelling solution to the green-yellow gap. Rather, their greater promiselies as ultra-efficient, wavelength-stable blue LEDs driven at the peak of the efficiencycurve before the onset of droop, with their larger effective active region volumecompensating for the lower drive current.

Figure 47

(a) CL intensity image and (b) peak wavelength image for a InGaN SQW grown on ac-axis oriented GaN nanowire taken at T � 16 K [130]. Reprinted with permissionfromMüller et al., Nano Lett. 16, 5340–5346 (2016) [130]. Copyright 2016 AmericanChemical Society. Room temperature micro-PL (c) intensity and (d) average wave-length of an InGaN/GaN MQW core–shell microrod. Both structures show a redshiftin emission going up along the sidewall. Reprinted with permission fromMounir et al.,J. Appl. Phys. 121, 025701 (2017) [137]. Copyright 2017 AIP Publishing LLC.


5.2. Semipolar Core–Shell NanostructuresDue to the higher indium incorporation on certain semipolar planes, 3D semipolarstructures, namely nanowires, pyramids, and stripes, have been explored for greenand longer wavelength LEDs. GaN nanowires grown by metal catalyst assisted(typically Ni), vapor-liquid-solid (VLS) MOCVD typically exhibit growth alongthe h112̄0i a-axis and isosceles triangular cross sections consisting of two semipolar{101̄1} facets and a (0001̄) N-polar c-plane facet [153,154], which is analogous to thegeometry of triangular GaN stripes selectively grown out of h112̄0i oriented stripeopenings [155]. Li and Wang demonstrated the growth of thick InGaN shell layerson the semipolar nanowire sidewalls with up to 40% In concentration and strong op-tical emission out to 700 nm, which is enabled by the 3D compliance of the nanowirestructure [156]. Triangular semipolar core–shell InGaN/GaN nanowire LEDs withelectroluminescence out to 577 nm [157] and optically lasing emission from 365–494 nm, based on growth conditions/In content, have also been demonstrated [158].While these results show promise for long wavelength emitters, the VLS growthmethod leads to a lack of ordering and alignment as well as uniformity both in termsof nanowire size and composition, making vertical device integration difficult.Relatively low growth temperatures required for high-density axial VLS growthcan also lead to increased point defect densities [159–163]. As such, most recent nano-wire LED efforts have focused on the catalyst-free, c-axis oriented nonpolar nanowire/nanorod architecture discussed above.

GaN-based nanopyramids with six-fold semipolar {101̄1} or {112̄2} facets have alsobeen explored for LEDs [164–174]. Similar to c-axis oriented GaN nanorods grownby SAG, nanopyramids are similarly grown from circular openings in a dielectricmask on c-plane GaN under conditions that result in stable semipolar planes withoutthe presence of straight nonpolar sidewalls, as shown in Fig. 49(a). Growth of InGaNlayers and QWs on the semipolar sidewalls generally results in observed higherIn content and redshifts in emission going from the base of the pyramid to the tipas observed by spatially resolved cathodoluminescence [Fig. 49(b)], similar to thecase for InGaN QWs grown on m-plane nonpolar nanowire sidewall facets[164,172,175,176]. The opposite trend with blueshift from base to tip has also beenobserved for closely spaced pyramids [168]. This can be explained by a longer relativeIn surface facet diffusion length for In versus Ga due to the weaker In-N bond strength[176]. Substantially different emission wavelengths can also result from thefacets, corners, and quantum-dot-like tip regions within single pyramids [164].Consequently, electrically injected, nanopyramid-based LEDs tend to exhibit fairlybroad electroluminescence [164,165,167,169–172,174]. Controlled growth and

Figure 48

(a) Schematic and (b) SEM image of a-axis oriented nanostripes with nonpolar (11̄00)sidewalls grown by pulsed-mode MOCVD (scale bar 500 nm). Reprinted with per-mission from Yeh et al., Appl. Phys. Lett. 100, 033119 (2012) [150]. Copyright 2012AIP Publishing LLC.


electrical injection of such structures with different color-emitting regions couldpotentially form the basis for phosphor-free, white LEDs [177].

GaN selectively grown out of stripe openings patterned along the h112̄0i direction isknown to result in triangular cross sections with stable {101̄1} facets over a widerange of growth conditions [155,178,179]. Compared to nanopyramids with similar{101̄1} facets, such stripes have fewer edges where facets meet and where In incor-poration can vary. However, as with other 3D structures, for InGaN layers grown onthe {101̄1} microstripe sidewalls, In content and QW thickness has been observed tochange (generally increasing) going from the base to the apex [178,180,181]. Fenget al. [182] observed relatively constant emission wavelength (∼500–510 nm) on thesidewalls but with significantly longer (550 nm) emission along the apex ridge.Smaller, h112̄0i oriented triangular nanostripes grown through submicrometerSiNx mask openings have also been recently studied. Leute et al. [183] demonstratedelectrically injected blue as well as true green (535 nm) LEDs with {101̄1} QWs,albeit with significantly increased FWHM caused by local inhomogeneities of In con-tent and QW thickness. Rishinaramangalam et al. [175] also realized triangular stripeLEDs with {101̄1} facets and stripe widths around 1.4 μm before the MQW growth[Fig. 49(c)], which exhibited broad electroluminescence spectra and blueshift withincreasing current density. These characteristics were correlated with nonuniformitiesin the QW thickness and In composition across the stripes, which are exacerbated by

Figure 49

Semipolar nanostructures. (a) SEM image GaN nanopyramid array with ∼90 nm thickInGaN layer on {112̄2} semipolar facets. (b) Variation of In content along InGaNlayer on nanopyramid showing increase toward the tip. Reprinted with permissionfrom Ko et al., ACS Photon. 2, 515–520 (2015) [166]. Copyright 2015 AmericanChemical Society. (c) SEM image of a triangular [112̄0] oriented nanostripeInGaN/GaN LED with semipolar {112̄0} sidewalls. Inset: cross-sectional SEM imageof a single stripe. Reprinted with permission from [175]. Copyright 2016 The JapanSociety of Applied Physics.


the poor current spreading in these 3D structures. However, Nakajima et al. [149]reported fairly uniform emission across the sidewalls centered around 575 nm forthree {101̄1} InGaN/GaN QWs grown on triangular nanostripes with widths around250 nm, and PL widths comparably to c-plane planar LEDs in a similar green-yellowwavelength range. It is suggested that at submicrometer scales below the diffusionlength of the precursors, variation of In due to mass transport effects is minimizedcompared to micrometer and larger sized stripes. However, more detailed follow-on TEM analysis of similar structures shows thickness variations from tip to baseas well as rough morphologies and the presence of stacking faults at the base, causedby enhanced surface diffusion of precursors and bending strain from bonding of theovergrown region to the mask [184]. In summary, semipolar nanostructures are prom-ising for green-yellow LEDs, but challenges remain in overcoming nonuniformities inQW thickness and In content due to their 3D morphology, and defects near the bottomof such structures may occur due to interactions with the mask material used inbottom-up SAG approaches.

6. SUMMARY AND OUTLOOK

In summary, we reviewed the recent progress on the device performance and materialproperties of blue and green LEDs grown on nonpolar and semipolar bulk GaN sub-strates, and on the semipolar and nonpolar facets of 3D nanostructures. For blueLEDs, semipolar devices show significantly reduced efficiency droop and thermaldroop with high LOP and EQE. Low efficiency droop is attributed to larger effectiveregion volume and smaller carrier density on these semipolar LEDs (thus less Augerrecombination and carrier leakage), enabled by reduced polarization-related electricfields, large QW thickness, good material quality, and uniform indium distribution.Moreover, both efficiency droop and thermal droop can be reduced by thick SQWdesign of semipolar LEDs due to the reduced QCSE and large critical thickness.Over 50% EQE with a minimal efficiency droop was demonstrated on semipolarLEDs, which makes semipolar blue LED technology comparable with c-plane blueLED technology, especially at high current densities, for high-power applications.Optimizations in light extraction and chip package can further improve the deviceperformances of semipolar LEDs.

For green LEDs, semipolar devices, particularly �202̄1� and �202̄ 1̄� orientations, haveunmatched advantages, such as high efficiency, high indium incorporation, fewer de-fects, low wavelength shift, and narrow spectra linewidth. Semipolar �202̄1� greenLEDs showed a record high peak EQE of 28.2% with an output power of13.3 mW. An outstanding issue for semipolar green LEDs is the relatively large effi-ciency droop, which could be further improved through optimizing QW growth andstructure, and chip processing technology. In addition, the physical origin and forma-tion mechanisms of defects in semipolar green LEDs with high indium compositionare still unclear and need more investigations. The use of AlGaN barriers or ILs mayalso be helpful in reducing these unique defects in semipolar green LEDs, as well asincrease the indium incorporation for green emission. In addition, cubic GaN-basedLEDs provide another route to growing nonpolar green LEDs due to polarization-freecubic crystal structure [185–187].

A continuing concern for nonpolar and semipolar LEDs is the high price of bulk GaNsubstrates compared with the commercially used sapphire substrates for c-planeLEDs, although research and development is continuing to drive down the costand improve the quality of these bulk GaN substrates. For example, Han and co-workers [43,44] have recently demonstrated semipolar GaN on 2 in. sapphire sub-strates with quality comparable to that of commercial c-plane GaN on sapphiresubstrates. This development is expected to facilitate the further development and


commercialization of semipolar LEDs. Alternatively, 3D nanostructures provideanother route toward nonpolar and semipolar LEDs using standard c-plane sapphiresubstrates. Morphologies under investigation include nanowires, nanopyramids, nano-sheets, and nanostripes. While such structures can exhibit long wavelength emission,challenges remain particularly with respect to improving the uniformity of the activeregion growth due to their complex three-dimensional morphologies.

Besides serving as efficient light sources [1,188–190], nonpolar and semipolar LEDsalso exhibits great potential for VLC, known as LiFi, complementary to the RF-basedWiFi. LED-based VLC is a “green” communication technology by leveraging existingsolid-stating lighting systems, with additional advantages such as virtually unlimitedbandwidth, license-free operation, high spatial reuse, and security [191–194]. Due tothe huge InGaN-based lighting infrastructure, the impact of combining lighting andcommunication together will be tremendous. Though LDs show larger modulationbandwidth, LEDs are more cost-effective, energy-efficient, and compatible with cur-rent lighting systems. However, the efficiency droop and low bandwidth of c-planeLEDs hinder their performance in the VLC. −3 dB frequency (f −3 dB) of c-planeLEDs is often limited to a few tens of megahertz due to the presence of electric fields[195] and larger carrier lifetime though micro-LEDs showed improvements of f −3 dBby reducing parasitic resistances and capacitances [196]. f −3 dB of 524 MHz,807 MHz, and 1.03 GHz were demonstrated on m-plane LEDs [195], (202̄ 1̄) super-luminescent diode [197], and (112̄2) LEDs [198], respectively. Further improvementsare expected by device designs, polarization engineering, and substrate orientations.In addition, the InGaN LED-based VLC can also benefit from the recent developmentof other III-nitride electronic [199–202], optoelectronic [203], and photonic devices[204], to form more advanced integrated systems with better performance and morefunctionalities. For example, the monolithic integration of InGaN LEDs with the GaNhigh electron mobility transistors [205–207] can considerably reduce the parasiticeffects and enhance the VLC’s f −3 dB.

FUNDING

U.S. Department of Energy (DOE); Sandia National Laboratories; ScienceFoundation Arizona (SFAZ) (Bisgrove Scholar Program); Solid State Lighting andEnergy Electronics Center, University of California Santa Barbara (SSLEEC).

ACKNOWLEDGMENT

Y. Z. and H. F acknowledge the support of the Bisgrove Scholar program from theSFAZ. G. T. W. acknowledges funding from the U.S. DOE, Office of Science, Officeof Basic Energy Sciences, Materials Science and Engineering Division, and Sandia’sLaboratory Directed Research and Development (LDRD) program. Sandia NationalLaboratories is a multimission laboratory managed and operated by NationalTechnology and Engineering Solutions of Sandia, LLC., a wholly owned subsidiaryof Honeywell International, Inc., for the U.S. Department of Energy’s NationalNuclear Security Administration. S. N. acknowledges the SSLEEC at UCSB.

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175. A. K. Rishinaramangalam, M. Nami, M. N. Fairchild, D. M. Shima, G.Balakrishnan, S. Brueck, and D. F. Feezell, “Semipolar InGaN/GaN nanostructurelight-emitting diodes on c-plane sapphire,” Appl. Phys. Express 9, 032101 (2016).

176. S. Zhang, X. Xiu, H. Wang, Q. Xu, Z. Wu, X. Hua, P. Chen, Z. Xie, B. Liu, Y.Zhou, P. Han, R. Zhang, and Y. Zheng, “Epitaxy and optical properties of InGaN/GaN multiple quantum wells on GaN hexagonal pyramids template,” Mater.Lett. 180, 298–301 (2016).

177. S. H. Lim, Y. H. Ko, C. Rodriguez, S. H. Gong, and Y. H. Cho, “Electricallydriven, phosphor-free, white light-emitting diodes using gallium nitride-based


double concentric truncated pyramid structures,” Light Sci. Appl. 5, e16030(2016).

178. F. Scholz, T. Wunderer, M. Feneberg, K. Thonke, A. Chuvilin, U. Kaiser, S.Metzner, F. Bertram, and J. Christen, “GaInN-based LED structures on selec-tively grown semipolar crystal facets,” Phys. Status Solidi A 207, 1407–1413(2010).

179. T. Wunderer, M. Feneberg, F. Lipski, J. Wang, R. A. R. Leute, S. Schwaiger, K.Thonke, A. Chuvilin, U. Kaiser, S. Metzner, F. Bertram, J. Christen, G. J. Beirne,M. Jetter, P. Michler, L. Schade, C. Vierheilig, U. T. Schwarz, A. D. Dräger, A.Hangleiter, and F. Scholz, “Three-dimensional GaN for semipolar light emitters,”Phys. Status Solidi B 248, 549–560 (2011).

180. P. L. Bonanno, S. M. O’Malley, A. A. Sirenko, A. Kazimirov, Z. H. Cai, T.Wunderer, P. Brückner, and F. Scholz, “Intrafacet migration effects in InGaN/GaN structures grown on triangular GaN ridges studied by submicron beamx-ray diffraction,” Appl. Phys. Lett. 92, 123106 (2008).

181. H. Fang, Z. J. Yang, Y. Wang, T. Dai, L. W. Sang, L. B. Zhao, T. J. Yu, and G. Y.Zhang, “Analysis of mass transport mechanism in InGaN epitaxy on ridgeshaped selective area growth GaN by metal organic chemical vapor deposition,”J. Appl. Phys. 103, 014908 (2008).

182. W. Feng, V. V. Kuryatkov, A. Chandolu, D. Y. Song, M. Pandikunta, S. A.Nikishin, and M. Holtz, “Green light emission from InGaN multiple quantumwells grown on GaN pyramidal stripes using selective area epitaxy,” J. Appl.Phys. 104, 103530 (2008).

183. R. A. R. Leute, D. Heinz, J. Wang, T. Meisch, M. Müller, G. Schmidt, S.Metzner, P. Veit, F. Bertram, J. Christen, M. Martens, T. Wernicke, M.Kenissl, S. Jenisch, S. Strehle, O. Rettig, K. Thonke, and F. Scholz, “EmbeddedGaN nanostripes on c-sapphire for DFB lasers with semipolar quantum wells,”Phys. Status Solidi B 253, 180–185 (2016).

184. Y. Nakajima and P. D. Dapkus, “The role of surface diffusion and wing tilt inthe formation of localized stacking faults in high In-content InGaN MQWnanostructures,” Appl. Phys. Lett. 109, 083101 (2016).

185. M. T. Durniak, A. S. Bross, D. Elsaesser, A. Chaudhuri, M. L. Smith, A. A.Allerman, S. C. Lee, S. R. J. Brueck, and C. Wetzel, “Green emitting cubicGaInN/GaN quantum well stripes on micropatterned Si(001) and their strainanalysis,” Adv. Electron. Mater. 2, 1500327 (2016).

186. S. C. Lee, N. Youngblood, Y. B. Jiang, E. J. Peterson, C. J. M. Stark, T.Detchprohm, C. Wetzel, and S. R. J. Brueck, “Incorporation of indium on cubicGaN epitaxially induced on a nanofaceted Si(001) substrate by phase transition,”Appl. Phys. Lett. 107, 231905 (2015).

187. C. J. M. Stark, T. Detchprohm, S. C. Lee, Y. B. Jiang, S. R. J. Brueck, andC. Wetzel, “Green cubic GaInN/GaN light-emitting diode on microstructuredsilicon (100),” Appl. Phys. Lett. 103, 232107 (2013).

188. C. C. Pan, Q. Yan, H. Fu, Y. Zhao, Y. R. Wu, C. G. Van de Walle, S. Nakamura,and S. P. DenBaars, “High optical power and low efficiency droop blue light-emitting diodes using compositionally step-graded InGaN barrier,” Electron.Lett. 51, 1187–1189 (2015).

189. H. Chen, H. Fu, Z. Lu, X. Huang, and Y. Zhao, “Optical properties of highlypolarized InGaN light-emitting diodes modified by plasmonic metallic grating,”Opt. Express 24, A856–A867 (2016).

190. H. Chen, H. Fu, X. Huang, Z. Lu, X. Zhang, J. Montes, and Y. Zhao, “Opticalcavity effects in InGaN core-shell light-emitting diodes with metallic coating,”IEEE Photon. J. 9, 8200828 (2017).


191. S. Rajbhandari, J. J. D. Mckendry, J. Herrnsdorf, H. Chun, G. Faulkner, H. Haas,I. M. Watson, D. O’Brien, andM. D. Dawson, “A review of gallium nitride LEDsfor multigigabit-per-second visible light data communications,” Semicond. Sci.Technol. 32, 023001 (2017).

192. J. Y. Tsao, M. H. Crawford, M. E. Coltrin, A. J. Ficher, D. D. Kolesker, G. S.Subramania, G. T. Wang, J. J. Wierer, and R. F. Karlicek, Jr., “Towardsmart and ultra-efficient solid-state lighting,” Adv. Opt. Mater. 2, 809–836(2014).

193. Z. Lu, P. Tian, H. Chen, I. Baranowski, H. Fu, X. Huang, J. Montes, Y. Fan, H.Wang, X. Liu, R. Liu, and Y. Zhao, “Active tracking system for visible lightcommunication using a GaN-based micro-LED and NRZ-OOK,” Opt.Express 25, 17971–17981 (2017).

194. H. Chen, X. Huang, H. Fu, Z. Lu, X. Zhang, J. A. Montes, and Y. Zhao,“Characterizations of nonlinear optical properties on GaN crystals in polar,nonpolar, and semipolar orientations,” Appl. Phys. Lett. 110, 181110 (2017).

195. A. Rashidi, M. Monavarian, A. Aragon, S. Okur, M. Nami, A.Rishinaramangalam, S. Mishkat-Ul-Masabih, and D. Feezell, “High speed non-polar InGaN/GaN LEDs for visible-light communication,” IEEE Photon.Technol. Lett. 29, 381–384 (2017).

196. R. X. G. Ferreira, E. Xie, J. D. Mckendry, S. Rajbhandari, H. Chun, G. Faulkner,S. Watson, A. E. Kelly, E. Gu, R. V. Penty, I. H. White, D. C. O’Brien, and M. D.Dawson, “High bandwidth GaN-based micro-LEDs for multi-Gb/s visible lightcommunications,” IEEE Photon. Technol. Lett. 28, 2023–2026 (2016).

197. C. Shen, C. Lee, T. K. Ng, S. Nakamura, J. S. Speck, S. P. DenBaars, A. Y.Alyamani, M. M. EI-Desouki, and B. S. Ooi, “High-speed 405-nm superlumi-nescent diode (SLD) with 807-MHz modulation bandwidth,” Opt. Express 24,20281–20286 (2016).

198. D. V. Dinh, Z. Quan, B. Roycroft, P. J. Parbrook, and B. Corbett, “GHzbandwidth semipolar (112) InGaN/GaN light-emitting diodes,” Opt. Lett. 41,5752–5755 (2016).

199. H. Fu, X. Huang, H. Chen, Z. Lu, I. Baranowski, and Y. Zhao, “Ultra-lowturn-on voltage and on-resistance vertical GaN-on-GaN Schottky power diodeswith high mobility double drift layers,” Appl. Phys. Lett. 111, 152102(2017).

200. D. S. Lee, X. Gao, S. Guo, D. Kopp, P. Fay, and T. Palacios, “300-GHz InAlN/GaN HEMTs with InGaN back barrier,” IEEE Electron Device Lett. 32,1525–1527 (2011).

201. H. Fu, X. Zhang, X. Huang, I. Baranowski, H. Chen, Z. Lu, J. Montes, and Y.Zhao, “Demonstration of AlN Schottky barrier diodes with blocking voltageover 1 kV,” IEEE Electron Device Lett. 38, 1286–1289 (2017).

202. H. Fu, X. Huang, H. Chen, Z. Lu, X. Zhang, and Y. Zhao, “Effect of buffer layerdesign on vertical GaN-on-GaN p-n and Schottky power diodes,” IEEE ElectronDevice Lett. 38, 763–766 (2017).

203. X. Huang, H. Fu, H. Chen, X. Zhang, Z. Lu, J. Montes, M. Iza, S. P. DenBaars,S. Nakamura, and Y. Zhao, “Nonpolar and semipolar InGaN/GaNmultiple-quantum-well solar cells with improved carrier collection efficiency,”Appl. Phys. Lett. 110, 161105 (2017).

204. H. Chen, H. Fu, X. Huang, X. Zhang, T. H. Yang, J. A. Montes, I. Baranowski,and Y. Zhao, “Low loss GaN waveguides at the visible spectral wavelengthsfor integrated photonics applications,” Opt. Express 25, 31758–31773 (2017).

205. Z. Li, J. Waldron, T. Detchprohm, C. Wetzel, R. F. Karlicek, Jr., and T. P. Chow,“Monolithic integration of light-emitting diodes and power metal-oxide-semiconductor channel high-electron-mobility transistors for light-emitting


power integrated circuits in GaN on sapphire substrate,” Appl. Phys. Lett. 102,192107 (2013).

206. Y. J. Lee, Z. P. Yang, P. G. Chen, Y. A. Hsieh, Y. C. Yao, M. H. Liao, M. H. Lee,M. T. Wand, and J. M. Hwang, “Monolithic integration of GaN-based light-emitting diodes and metal-oxide-semiconductor field-effect transistors,” Opt.Express 22, A1589–A1595 (2014).

207. Z. Liu, J. Ma, T. Huang, C. Liu, and K. M. Lau, “Selective epitaxial growth ofmonolithically integrated GaN-based light emitting diodes with AlGaN/GaNdriving transistors,” Appl. Phys. Lett. 104, 091103 (2014).

Yuji Zhao received a B.S. degree in microelectronics from FudanUniversity, China in 2008, and a Ph.D. degree in electrical andcomputer engineering from University of California, SantaBarbara (UCSB) in 2012 under the supervision of Nobel LaureateProfessor Shuji Nakamura. He is currently an assistant professor ofelectrical engineering at Arizona State University (ASU), wherehis research interests focus on the physics, materials, and devicesapplications of GaN wide bandgap semiconductors for power, en-

ergy, computing, and photonics applications. He has made important contributions innonpolar and semipolar InGaN light-emitting diodes (LEDs), including the demon-stration of the first “low droop” semipolar InGaN blue LEDs (reported by Science,May 2012) and elucidating the basic growth modes and defect generation for “greengap” InGaN LEDs (reported by Nature Photonics, July 2013). He has authored/co-authored more than 90 journal and conference publications, 2 book chapters, and over10 patents. He is the recipient of 2017 ASU Fulton Outstanding Assistant ProfessorAward, 2016 DoD DTRA Young Investigator Award, 2015 NASA Early CareerFaculty Award, 2015 Bisgrove Scholar Career Faculty Award, and 2010–2013UCSB SSLEC Outstanding Research Award.

Houqiang Fu received a B.S. degree in material physicsfrom Wuhan University, China in 2014, where he was awardedthe National Scholarship of China, and the NationalEncouragement Scholarship of China. He is currently pursuinga Ph.D. degree in electrical engineering at Arizona StateUniversity. His research interests include MOCVD growth anddevice applications of III-nitride wide bandgap semiconductorsincluding LEDs, solar cells, and power devices. He has

authored/co-authored over 40 journal and conference publications, 1 book chapter,and 1 patent.

George T.Wang is a principal member of the technical staff in theAdvanced Materials Sciences Department at Sandia NationalLaboratories. He received his B.S. degree in chemical engineering,with highest honors, from the University of Texas at Austin in1997, and M.S. and Ph.D. degrees in chemical engineering fromStanford University in 1999 and 2002, respectively, where he wasawarded a National Science Foundation Graduate Fellowship andthe David Sen-Lin Lee Fellowship. His thesis work at Stanford

focused on the functionalization of semiconductor surfaces using novel organicreactions under ultra-high vacuum. At Sandia, his primary research efforts have


focused on the synthesis, characterization, and applications of III-nitride nanowiresgrown by metal-organic chemical vapor deposition (MOCVD). In addition, he has ledefforts at Sandia in the 3D growth of high-quality GaN using novel nanostructuringand selective growth techniques. He has also investigated parasitic chemistry and Mgdoping issues in AlGaInN MOCVD using a combination of in situ experimental tech-niques and density functional theory calculations. He has (co)authored more than70 peer-reviewed publications and holds 11 patents.

Shuji Nakamura received B.E., M.S., and Ph.D. degrees in elec-trical engineering from the University of Tokushima, Japan, in1977, 1979, and 1994, respectively. He joined Nichia ChemicalIndustries Ltd., in 1979. In 1989, he started the research of blueLEDs using III-nitride materials. In 1993 and 1995, he developedthe first III-nitride high-brightness blue/green LEDs. He also de-veloped the first III-nitride violet laser diodes (LDs) in 1995. Since2000, he has been a professor of materials at the University of

California, Santa Barbara (UCSB), where he is currently the co-director of theSolid State Lighting and Energy Electronics Center (SSLEEC) and the Cree Chairin Solid State Lighting and Display. He has received over 30 major internationalawards, including the Nishina Memorial Award (1996), the Materials ResearchSociety (MRS) Medal Award (1997), the Institute of Electrical and ElectronicsEngineers (IEEE) Jack A. Morton Award (1998), the British Rank Prize (1998),the Benjamin Franklin Medal Award (2002), the Millennium Technology Prize(2006), the Czochralski Award (2007), the Prince of Asturias Award for TechnicalScientific Research (2008), the Harvey Award (2009), and the Technology &Engineering Emmy Award (2012). He received the Nobel Prize in Physics in2014 “for the invention of efficient blue light-emitting diodes which has enabledbright and energy-saving white light sources.” He is a member of the U.S. NationalAcademy of Engineering. He has published over 550 papers and holds more than 200U.S. patents and over 300 Japanese patents.


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