Group III-nitride lasers: a materials perspective

ISSN:1369 7021 © Elsevier Ltd 2011SEPTEMBER 2011 | VOLUME 14 | NUMBER 9408

GaN based laser diodes (LDs) have entered widespread use as the

laser source for high-density optical data storage (e.g., Blu-ray

discs), as well as new applications in laser based projectors and

TVs. Data density for optical storage technology is limited by the

wavelength of the light source used to read and write data, with

shorter wavelengths leading to higher data density. GaN along

with InN and AlN form the only short wavelength direct bandgap

materials system that has shown both p- and n-type doping and

sufficiently long device lifetime for successful commercialization.

The need for LDs capable of faster read/write times and 3D

storage continues to drive research related to Blu-ray technology1.

Due to their high spectral purity and efficiency, lasers also make

excellent candidates for light sources in large televisions and

projectors, as well as small, portable pico-projectors (pictured in

Fig. 1). They enable over 90 % of the color gamut to be rendered

with red, green, and blue LDs; over twice the color range available

from competing technologies3.

HDTVs with screen sizes of up to 75” have been commercialized by

Mitsubishi Electric and laser projectors are commercially available from

several manufacturers. Most systems have used GaAs based red LDs

An overview of III-Nitride based laser diodes (LDs) is presented focusing on the materials challenges in each phase of device development. We discuss early breakthroughs leading to the first commercial GaN LDs, covering crystal growth, p-type doping, and defect reduction. Additional device issues, such as polarization effects, strain, and index dispersion are addressed as they apply to the development of blue and green LDs for pico-projector applications. State of the art device results are highlighted. Devices grown on non-polar and semi-polar GaN substrates address many polarization related problems present in c-plane GaN growth. Device results, advantages, and limitations of various non-polar and semi-polar systems are discussed in terms of polarization properties, Indium incorporation, extended defect formation, and critical thickness. A brief description of challenges and progress in UV LDs is also presented.

Matthew T. Hardya*, Daniel F. Feezella, Steven P. DenBaarsa,b, and Shuji Nakamuraa,b

a University of California, Materials Department, Santa Barbara, CA, 93106-5050, USAb University of California, Department of Electrical and Computer Engineering, Santa Barbara, CA, 93106-5960, USA

* Email: [email protected]

Group III-nitride lasers: a materials perspective

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mailto:[email protected]

Group III-nitride lasers: a materials perspective REVIEW

SEPTEMBER 2011 | VOLUME 14 | NUMBER 9 409

and frequency doubled green and blue lasers, although arrays of high

power blue GaN LDs are entering use. Pico-projectors likely represent

the most significant new market for GaN LDs. The pico-projector

market is expected to hit $1.1 billion by 2014 and the green LD market

(both frequency doubled and semiconductor LD) is expected to reach

$500 million by 20164,5. Current pico-projectors use frequency doubled

1060 nm lasers for the green source, however due to low cost, high

efficiency, and low form factor requirements, direct emission GaN LDs

are the most likely choice for blue and green components of RGB light

sources in next generation pico-projectors. Such devices may be in

production as early as 2012.

BackgroundLDs emit light that is both highly coherent and highly directional,

allowing high fiber coupling efficiencies, coherent optics, and high

optical power densities when combined with simple optics. In general,

it is desirable to have LDs with a low threshold current density (Jth).

Device operation below Jth leads only to spontaneous emission, which

is effectively wasted. Jth can be reduced by increasing the modal

gain and reducing the cavity losses. Modal gain is the product of the

fundamental material gain (a metric controlled by the material’s

optical transition efficiency, band structure, and injection level)

and the optical confinement factor, which describes the overlap of

the optical mode with the injected carriers. Internal loss is purely

parasitic and comes from photon scattering out of the optical mode,

dopant-related free-carrier absorption, and metal contact adsorption.

Mirror loss is optical power lost out of a facet and translates directly

into useful LD output power. Mirror loss can be controlled with the

application of high-reflection (HR) or anti-reflection (AR) coatings

on the facets. HR coatings result in a decrease in Jth (due to reduced

losses) but a reduction in output power. Slope efficiency (SE) is the

slope of the optical output power versus current curve above threshold.

Typically, HR (AR) coatings can be used to decrease (increase) Jth

at the cost (benefit) of SE. SE also depends directly on injection

efficiency, which describes the portion of current above threshold

that goes to stimulated emission, and inversely on internal loss. Thus

low power LDs focus primarily on Jth reduction, while high power

device designs are focused on SE. The most commonly used power

consumption metric is wall plug efficiency (WPE), which is given by

WPE = Pout/(IopVop), where Iop and Vop are the LD operating current and

voltage. WPE is primarily dependant on SE, but achieving low Jth is also

useful. For a full treatment of LD physics, see reference 6.

Early GaN researchWhile the 1970s saw the growth of the first GaN single crystals and

observation of the first optically pumped stimulated emission7, little

progress was made on electrically injected devices until almost three

decades later. The achievement of high crystal quality GaN films

was complicated by the absence of a native substrate. The early

breakthroughs came from using AlN buffer layers grown on sapphire

substrates, and later GaN buffer layers8,9. In addition, modification of

traditional III-V metalorganic chemical vapour deposition (MOCVD)

reactor design and/or low pressure growth was required to prevent

parasitic pre-reactions between the group III metalorganic precursors

and NH3, and convective effects resulting from high GaN growth

temperatures over 1000 °C (as opposed to < 700 °C for conventional

III-Vs)10-12. In spite of drastic improvements in crystal quality,

threading dislocation (TD) densities for GaN on sapphire remained

high; between 7·108 and 1010 cm-2 13-15.

The absence of demonstrations of p-type conductivity limited

early development of III-nitride semiconductors. It wasn’t until 1989

that Amano et al. demonstrated the first p-type Mg doped GaN

using an ex situ low energy electron beam (LEEBI) treatment18. The

mechanism was later clarified by Nakamura et al., who showed H

radicals from NH3 or H2 present in the growth environment passivate

the Mg dopant. LEEBI or a simple thermal annealing in an NH3 free

environment was shown to facilitate out-diffusion of the passivating

H radicals17,18.

Fig. 1 (a)Concept for a RGB laser based projection system embedded in a mobile device. Image provided courtesy of MicroVision, Inc., www.microvision.com. (b) Far field pattern of a green LD under testing. Reprinted with permission from2. © 2010 Japan Society of Applied Physics.

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REVIEW Group III-nitride lasers: a materials perspective

SEPTEMBER 2011 | VOLUME 14 | NUMBER 9410

In addition to crystal quality and doping issues, GaN opto-

electronic heterostructure devices must also contend with extremely

strong polarization-induced electric fields due to the lack of inversion

symmetry in the wurtzite crystal structure. These electric fields cause

a spatial separation of the electron and hole wave functions, which

leads to reductions in radiative recombination rates and radiative

efficiency19. This phenomenon is known as the quantum confined Stark

effect (QCSE). Ultimately this limits the maximum width of quantum

wells (QWs) in device active regions to prevent excessive separation

of electron and hole wave functions. Additionally, polarization-induced

electric fields cause a red shift of emission wavelength at low carrier

density due to bending of the QW band edges, and a blue shift with

increasing carrier density due to the screening of these polarization

fields19,20.

The presence of Mg doping in GaN is a major source of optical

loss, potentially as high as 65 cm-1 21. Device design must consider

the tradeoff with respect to the proximity of Mg doping to the active

region: too far away and hole injection suffers, too close and Mg based

optical loss is high21,22.

Developments in crystal growth and doping led to the first

demonstration of GaN p-n LEDs in 1991, and with further epitaxial

growth improvements, the first AlInGaN based LDs in 199623-25.

These initial LDs were grown on sapphire substrates and employed

n-/p-Al0.15Ga0.85N cladding layers, GaN waveguiding layers and a

26 period In0.2Ga0.8N/In0.05Ga0.95N MQW structure as well as an

InGaN compliance layer beneath the n-AlGaN cladding to prevent

cracking in the AlGaN. Facets were formed by reactive ion etching

due to the difficulty of cleaving GaN on sapphire25. Continued

advancements in epitaxial growth such as the implementation of short-

period strained-layer superlattices (SLSs) for the n-/p-cladding and the

implementation of standard LD technology such as ridge waveguides

and cleaved facets led to device improvements and lifetimes of up to

30 hours26-32.

The next major technological breakthrough came from the

implementation of epitaxial laterally overgrown GaN (ELOG) on

sapphire substrates to reduce TD densities from 0.1 – 1 × 1010 cm-2

down to around 5 × 106 cm-2 33-35. As shown in Fig. 2, the

un-patterned window region has high TD density while the wing

region which has overgrown the SiO2 mask shows significantly

reduced TD density. It is remarkable that early GaN devices had any

measurable electroluminescence at all with 1010 cm-2 TD density:

4 orders of magnitude greater than TD densities at which GaAs based

LEDs cease to emit light13. The use of ELOG substrates led to lifetime

improvements of up to an order of magnitude, allowing the first

demonstrations of GaN LDs with lifetimes of 1000s of hours26,36,37.

The success of ELOG led to the development of thick freestanding

GaN substrates grown by hydride vapor phase epitaxy (HVPE). HVPE

is capable of over 100 μm/hr growth rates and can be used to grow

films up to several hundred microns thick. One or more ELO steps

can be used to reduce TD density, resulting in TD densities as low

as 2 × 105 cm-2 38. The substrate is then removed using either laser-

induced lift-off in the case of GaN on sapphire, or a wet etch for

GaN on GaAs38-40. In addition to defect reduction, freestanding GaN

substrates eliminate thermal expansion coefficient mismatch with the

substrate, have excellent thermal conductivity, allow backside contacts

and facilitate high quality cleaved facets. True bulk GaN substrates

Fig. 2 Optical micrographs showing etch pit density of (a) ELOG with substantially reduced defect density in the wing region above the SiO2 mask and (b) GaN/sapphire sample of similar thickness. (c) dependence of mean time to failure (MTTF) on defect density. Figures (a) and (b) reprinted with permission from31 and (c) from21. © 1998 and 2003, IEEE.

(a)

(b)

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with TD densities in the 101 – 104 cm-2 can be grown by high nitrogen

pressure or ammonothermal methods, however such substrates are not

yet readily available41-43.

Violet c-plane LD developmentAll currently available commercial III-nitride LDs are grown by

metalorganic chemical vapor deposition (MOCVD) on the c-plane

of the wurtzite crystal. Violet devices currently account for the

majority of revenue due to their use in high density optical data

storage. A typical device design for violet III-nitride LDs is given in

Fig. 3, and includes n- and p-AlGaN/GaN short-period SLSs (SPSLs)

for low index cladding layers and GaN or InGaN waveguiding layers

to form the separate confinement heterostructure (SCH). The MQW

active region typically has two or three QWs, although there have

been demonstrations of high performance single QW devices44. The

p-waveguiding layer is usually very lightly Mg doped or undoped to

reduce Mg related optical loss, with the p-AlGaN electron blocking

layer (EBL) placed just before or just after the p-waveguiding

layer21,28,45.

In general, there are low power consumption and high optical output

power applications, each requiring different epitaxial growth designs.

Low power consumption design typically focuses on reducing threshold

current density, primarily by increasing the confinement and reducing

optical losses. Conversely, designs for high power involve reducing

the confinement factor to spread out the optical mode and decrease

the optical power density at the facet46. Reduction of optical loss to

improve SE can be achieved by reducing the optical mode overlap with

Mg doped regions. This can be done by reducing Mg doping, moving Mg

doping farther from the active region and adjusting the waveguiding

structure to shift the optical mode towards the n-side of the device21.

Proper design of facet coatings, typically AR on the output facet and HR

coatings on the back facet can maximize SE with only minimal increase

in Jth, and high quality facet coatings can also help protect the output

facet and prevent catastrophic optical damage (COD)47.

In the past 15 years of LD research, continued crystal growth

optimization and processing improvements have resulted in state of

the art continuous wave (CW) output powers and slope efficiencies of

500 mW and 1.6 W/A, respectively, for single lateral mode devices, and

8 W and 2.1 W/A, respectively, for multi-mode devices21,46-51.

The push to c-plane blue and green LDsBlue (440 – 460 nm) semiconductor LDs are an excellent candidate

for the blue component of an RBG light source for high power, plug-in

projectors and small portable projectors known as pico-projectors. In

terms of a high power projector, LDs are attractive due to their highly

directional, high power output and long lifetimes. These qualities apply

equally well to pico-projectors, where the small form factor and high

wall plug efficiency (WPE) of LDs is also advantageous for both blue

and green (515 – 535 nm) light sources52. Laser light sources based on

second harmonic generation from InP/GaAs based LDs have recently

shown much improvement, but GaN based LDs have lower form factor

and have the potential to have higher WPEs, longer lifetimes, and lower

costs53-56.

The challenges for making longer wavelength III-nitride LDs are

numerous and become more severe with increased lasing wavelength.

To extend the emission wavelength the InN fraction in the active

region must be increased. The 10 % mismatch between the in-plane

lattice constants of GaN and InN leads to high strain in In0.18Ga0.82N

blue and In0.3Ga0.7N green QW active regions, of 1.8 % and 3.0 %,

respectively. For reference, in conventional III-V laser structures

1.5 % is considered very high strain. In addition to directly degrading

material quality, high strain also leads to increased piezoelectric

polarization in the QWs. Larger band offsets for higher In composition

QWs relative to the quantum barriers (QBs) leads to carrier transport

issues in MQW active regions, particularly for holes, which have much

lower mobility than electrons in III-nitrides. Higher In composition

layers suffer from decreased chemical stability leading to reduced

layer homogeneity and broadening of spontaneous emission line

widths in electroluminescence (EL) and photoluminescence (PL) as

well as decreased thermal stability57-59. Formation of non-radiative

recombination centers in the active region causes a drop in the

non-radiative carrier lifetime and a decrease in internal quantum

efficiency60. For example, nitrogen vacancies become increasingly

problematic with high In composition due to the elevated vapor

pressure of nitrogen in InN relative to GaN. The net effect of lower

active region quality is reduced material gain and a broadened

gain spectrum. In addition to the numerous growth challenges, the Fig. 3 Schematic of a typical epitaxial structure for a GaN based ridge waveguide LD.

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dependence of the refractive index on wavelength (index dispersion)

leads to a decreasing index contrast between layers of the optical

waveguide structure, as shown by Fig. 4. For the same cladding and

waveguide structure, the confinement factor decreases with increasing

emission wavelength62,63. The summation of these effects results in

reduced modal gain for green and blue LDs and gain saturation at

lower current densities than for violet LDs64. Careful device design and

processing are required to keep optical losses to a minimum in order

to achieve lasing.

Blue c-plane LDs have been reported with a single mode output

power (slope efficiency) of 120 mW (1 W/A) and a multi-mode power

(slope efficiency) of 1.17 W (1.4 W/A)44,52,65. Green c-plane LDs with

Jth as low as 3.8 kA/cm2 and CW output powers (slope efficiency) of 50

mW (0.33 W/A) have been published64-67. Fig. 5 shows a summary of Jth

as a function of lasing wavelength. For a review of device performance

and progress of green LDs, see reference 72.

Non-polar and semi-polar growth planesEarly theoretical work predicted the absence of polarization related

electric fields for planar thin films grown on substrates orientated

orthogonally to the polar c-plane of GaN, such as the m- and a-planes,

and reduced polarization related fields on semi-polar (SP) growth

planes orientated between the polar and non-polar (NP) planes20,73.

In general, any crystal plane perpendicular to the c-plane, such as m-

and a-planes, is non-polar and any plane inclined between a non-polar

plane and c-plane is a semi-polar plane. Common NP and SP planes

are depicted schematically in Fig. 6, and Fig. 7 shows a common

calculation of the dependence of total polarization on the inclination

angle of the growth plane with respect to the c-plane. While there

is still some ambiguity in the exact polarization magnitude due to

lingering uncertainties in the piezoelectric polarization coefficients

and their variation with alloy composition, most calculations show

a cross-over in the sign of the polarization for semi-polar planes

with an inclination angle relative to c-plane between 40 – 57°74-76.

Polarization of SP InGaN with opposite sign relative to c-plane has

been experimentally confirmed at 58° (112-2) and 62° (101-1-), verifying

the zero crossing occurs at an inclination angle less than 58°77. The

early theoretical work led to a large effort to grow non-polar and

semi-polar GaN on foreign substrates. However, due to high stacking

fault densities in the 103 – 105 cm-1 range, device performance was

severely limited78,79. The stacking fault issue was finally solved by

growing freestanding c-plane GaN substrates by HVPE, and then

slicing and polishing NP and SP orientated substrates with low

dislocation and stacking fault densities40.

Unbalanced bi-axial in-plane strain in NP/SP GaN leads to a

theoretical separation of the top two valence bands. The subsequent

reduction in accessible valence band density of states should lead

to enhanced gain in NP/SP QW based LDs80,81. Early growth studies

showed severely limited In incorporation in m-plane growths

Fig. 5 Summary plot of threshold current density versus lasing wavelength50,51,59,66,68-71,91,92,96,98-100,104,105,109,111,113.

Fig. 4 Index contrast Δn of (a) AlGaN and (b) InGaN with respect to GaN for LDs operating at violet (405 nm), blue (450 nm) and green (520 nm) wavelengths. Figure based after61.

(a)

(b)

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co-loaded with c-plane, with compositions up to 2 – 3 times lower82.

Additionally, the surface morphology of m-plane homoepitaxy

exhibited large pyramidal hillocks leading to lateral thickness non-

uniformities in subsequently grown layers83. Further work showed

these pyramids formed due to spiral growth around the termination of

a partially screw-type threading dislocation with the growth surface84.

This led to the use of substrates offcut 1° towards the c-direction to

prevent hillock formation85.

The first demonstrations of violet m-plane LDs in the violet regime

closely followed the availability of free-standing GaN substrates

in 200786,87. Laser stripes oriented along the c-direction showed a

lower Jth relative to those orientated along the a-direction as a result

of anisotropic in-plane gain86,87. Studies of m-plane LEDs showed

very little QW width dependence on internal efficiency due to the

absence of QCSE, in contrast to typical c-plane devices88. This led to

the development of LDs using wide QWs in conjunction with InGaN

waveguiding layers to provide sufficient confinement of the optical mode

without the need for AlGaN cladding layers89,90. These AlGaN-cladding-

free (ACF) structures eliminate cracking issues, poor p-conductivity,

parasitic pre-reactions, and relatively high p-cladding growth

temperatures associated with p-AlGaN growth. m-plane violet LDs with

single mode output power over 500 mW and slope efficiency of 2.0 W/A

have recently been demonstrated by Kaai (now Soraa)91. These device

results are at least equivalent to state of the art c-plane LDs.

Demonstrations of blue m-plane LDs using both ACF and

traditional AlGaN clad structures followed soon after the initial

demonstrations92-96. Again Jth and output power for these initial

devices was competitive with c-plane LDs and showed drastically

reduced blue shift with increasing current density; as little as 10 nm

while c-plane devices often shifted by as much as 30 – 50 nm92,97.

Recently, Soraa demonstrated a world record CW single lateral mode

output power of 750 mW at 447 nm with 23.2 % WPE for LDs grown

on “non-c-plane” substrates98.

Continued effort was made to push m-plane out to longer lasing

wavelengths. Devices emitting in the blue-green regime with a

maximum of 500 nm were reported85,99,100. However, with increasing

wavelength, radiative efficiency, and LED output power dropped off

sharply near 500 nm. TEM investigations showed 104 – 105 cm-1 I1

type stacking faults originating from the QWs in samples emitting

around 500 nm101. These defects are likely responsible for the sudden

drop off in spontaneous emission power. I1 faults can only form due

to a growth error, and as such could be prevented by fully optimized

growth conditions, although further progress on long wavelength

m-plane LDs has not been reported.

There are a large number of low index semi-polar planes that could

potentially be used for long wavelength LDs. Early research focused

on (102-2) and (101-1), due to their relatively high In incorporation,

low piezoelectric polarization, and good LED performance, leading

to early violet LD demonstrations102-105. Stimulated emission was

observed at 514 nm on (112-2), but electrically injected green LDs

have not yet been realized106. TEM investigations discovered misfit

dislocations at each of the AlGaN cladding and InGaN waveguiding

interfaces formed due to relaxation via dislocation glide on the

c-plane107. Relaxation by TD glide is not normally seen in planar

c-plane growth because the most favorable slip system is the c-plane

Fig. 6 Schematic of various GaN crystal planes showing (a) the polar c-plane, non-polar a-plane and semipolar (112

-2) plane and (b) non-polar m-plane

and semi-polar (202-

1) and (101-

1) planes. The inclination angle ϑ with respect to the c-plane is indicated in (a).

Fig. 7 Total polarization discontinuity between (a) InxGa1-xN and GaN and (b) AlyGa1-yN and GaN as a function of inclination angle ϑ with respect to the c-plane. Figure based after74.

(a)

(b)

(a) (b)

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(with a Burger’s vector b→

= 1/3 ⟨112-0⟩), which is parallel to the growth

plane and thus has no resolved shear stress. The same is true for

growth on m- and a-planes, where the growth plane is orthogonal to

the c-plane slip plane. SP planes have non-zero resolved shear stress

on the c-plane and can thus relieve stress by slip on the c-plane. For

the above mentioned (112-2) LD, high densities of misfit dislocations

less than 50 nm away from the active region likely led to high non-

radiative recombination rates, preventing lasing.

In July of 2009, Sumitomo Electric shocked the GaN LD world by

announcing lasing at 531 nm on a novel semi-polar (202-1) plane59. At

that point, the longest reported lasing wavelength was 515 nm on a

c-plane68 device. Sumitomo used quaternary AlInGaN cladding layers

latticed matched to GaN with GaN waveguiding layers to provide

index contrast without exceeding the critical thickness for relaxation

by glide on the c-plane. The critical thickness of (202-1) is 2 – 3 times

the critical thickness of (112-2) due to the difference in the SP plane

orientation relative to the c-plane, providing a larger window for

coherent growth108. A trade off emerges amongst NP/SP planes, where

m-plane has the lowest In incorporation for atmospheric pressure

MOCVD but highest critical thickness, (112-2) has the highest In

incorporation but lowest critical thickness and (202-1) falls in between

making a good compromise for long wavelength LDs with coherent

waveguiding/cladding structures.

Since the initial demonstration, Sumitomo reported CW lasing

at 526 nm with output power up to 36 mW and a reasonably low

Jth of 4.3 kA/cm2 on (202-1)109. In addition, Soraa has published CW

lasing at 521 nm with 60 mW of output power, a Jth of 7.2 kA/cm2,

a slope efficiency of 0.3 W/A, and a 10 000 hour lifetime for their

“non-c-plane” LDs98.

Ultraviolet LDsWhile there has long been interest in InGaN based LDs, interest

in AlGaN based UV LDs has only arisen in the last few years. The

GaN-AlN system spans almost the entire UV spectrum and has the

potential to replace many conventional gas and solid state UV lasers,

offering advantages in terms of cost, reliability, and portability.

AlGaN LDs as a high intensity light source have applications in

chemical analysis, bio-agent detection, photolithography, and

sterilization110,111. Many unique challenges face the development of

AlGaN based LDs. AlGaN QWs do not seem to have the robustness of

InGaN in the face of high dislocation densities, requiring sophisticated

overgrowth techniques to reduce TD density112,113. The poor

surface mobility of Al tends to require high temperature growth

(1100 – 1500 °C) or low growth rates to allow the migration of Al

ad-atoms to a step edge and preserve layer-by-layer growth, however

pre-reactions involving Al containing precursors and NH3 become a

larger problem with higher growth temperature and increasing TMA

gas phase composition112.

In addition to structural quality, doping presents a severe challenge

to electrically injected devices with increased Al content. Both Si and

Mg activation energies increase almost linearly with increasing Al

composition. The problem is particularly severe for Mg doping, which

has an activation energy around 250 meV in GaN and up to 500 meV

in AlN. AlGaN/GaN SPSLs, AlGaN/GaN polarization engineered

heterojunctions and δ-doping have shown some enhancement in hole

concentrations for films with Al compositions up to the 0.4 range.

Additionally, interstitial fluorine may act as a relatively shallow donor,

but must be diffused or implanted ex situ as epitaxially doped F will be

substitutional114. Significant work will be required to achieve good hole

injection through thick, higher Al composition cladding layers.

In spite of these challenges, there has been rapid progress towards

achieving deep UV LDs. Recently, Yoshida et al. demonstrated lasing

at 342 and 336 nm with Jth of 8 and 17 kA/cm2, respectively, claiming

drastically improved device performance due to reduced extended

defect densities in the template111,113. Work on deep UV LEDs has led

to IQE values as high as 69 % for active regions emitting at 247 nm

through the optimization of growth conditions115.

ConclusionViolet and blue c-plane laser diodes are nearing maturity, with very

high output powers and long lifetimes, leading to commercialization

for Blu-ray Discs, high power projectors, and in the near future,

pico-projectors. In spite of the large number of challenges faced,

c-plane green LDs have experienced significant progress, although

current output power, lifetime, and WPE may delay commercialization.

Semi-polar and non-polar LDs have the potential to mitigate

polarization related difficulties and have shown substantial progress,

especially in the blue and green regimes where device performance

now meets or exceeds that of state of the art c-plane LDs. Interest

is rapidly expanding into UV LDs, which may provide high power

UV light sources if structural quality and doping challenges can be

overcome.

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